Nitrogen-doped niobium-titanium alloy and preparation method and application thereof

Abstract

The invention belongs to the technical field of superconducting materials, and particularly relates to a nitrogen-doped niobium-titanium alloy and a preparation method and application thereof. The nitrogen-doped niobium-titanium alloy provided by the invention comprises the following components in percentage by mass: 0.08 to 0.19 weight percent of nitrogen, 52 to 54 weight percent of niobium and 45.81 to 47.92 weight percent of titanium, the nitrogen-doped niobium-titanium alloy contains a beta phase, an alpha phase and a nitride precipitated phase. According to the nitrogen-doped niobium-titanium alloy, a certain content of nitrogen is doped in the niobium-titanium alloy, the nitrogen is effectively dissolved in the alloy in a solid mode, and by changing the microstructure and phase composition, the nitrogen-doped niobium-titanium alloy keeps good superconducting performance and meanwhile has high mechanical performance.

Description

TECHNICAL FIELD

[0001] The application belongs to the technical field of superconducting materials, and particularly relates to a nitrogen-doped niobium-titanium alloy and a preparation method and application thereof. BACKGROUND

[0002] The niobium-titanium alloy not only has high specific strength, good mechanical properties, excellent corrosion resistance, but also has good biocompatibility, stable superconducting performance, high ductility and excellent processing performance, which makes the niobium-titanium alloy irreplaceable in the fields of medical treatment, aerospace, transportation, superconducting, electronics, energy and power transmission. The niobium-titanium alloy is the main material of the commercialized superconducting wire at present, which is mainly due to its excellent superconducting performance in a low magnetic field environment and good mechanical processing performance. However, with the continuous progress of science and technology, the application range and scale of the niobium-titanium alloy are continuously extended and expanded, and people have put forward higher requirements for its application; the niobium-titanium alloy with a Ti content of 47wt.% is currently mainly applied, and in order to reduce the loss in the application process, the niobium-titanium alloy is more used as a multi-core wire, and the diameter of the wire is usually kept at the micron level, which has a higher requirement on the mechanical properties of the niobium-titanium alloy raw material, and how to maintain the superconducting performance of the niobium-titanium alloy while improving its mechanical properties is still a research hotspot of the niobium-titanium superconducting alloy. SUMMARY

[0003] Therefore, the application provides a nitrogen-doped niobium-titanium alloy, a preparation method and application thereof. The nitrogen-doped niobium-titanium alloy provided by the application has good mechanical properties and superconducting performance.

[0004] In order to solve the above technical problems, the application provides a nitrogen-doped niobium-titanium alloy, which comprises the following components in mass percentage: nitrogen 0.08-0.19wt%; niobium 52-54wt%; titanium 45.81-47.92wt%; The nitrogen-doped niobium-titanium alloy contains a beta phase, an alpha phase and a nitride precipitated phase.

[0005] Preferably, the volume percentage of the beta phase in the nitrogen-doped niobium-titanium alloy is greater than or equal to 99%, and the total volume percentage of the alpha phase and the nitride precipitated phase is less than or equal to 1%.

[0006] The application also provides a preparation method of the nitrogen-doped niobium-titanium alloy. nitrogenizing treatment is performed on the niobium-titanium alloy powder to obtain a nitrogenized niobium-titanium alloy powder; the temperature of the nitrogenizing treatment is 550-650 DEG C; sintering is performed on the nitrogen-doped niobium-titanium alloy powder to obtain a nitrogenized niobium-titanium alloy; solid solution treatment is performed on the nitrided Nb-Ti alloy to obtain the nitrogen-doped Nb-Ti alloy.

[0007] Preferably, the holding time of the nitriding treatment is 0.5-2h, the nitrogen-containing gas used in the nitriding treatment includes ammonia, and the flow rate of the nitrogen-containing gas is 0.8-1m 3 / h.

[0008] Preferably, the particle size of the Nb-Ti alloy powder is 50-100μm.

[0009] Preferably, the nitriding treatment is performed in a nitriding furnace, the atmosphere in the nitriding furnace is stirred and mixed by a stirring fan, and the rotating speed of the stirring fan is 140-160r / min.

[0010] Preferably, the sintering includes discharge plasma sintering. The discharge plasma sintering has a vacuum degree of ≤1Pa, a temperature rising rate of 100-200℃ / min, a temperature of 1000-1200℃, a holding time of 15-30min, and an axial pressure of 40-50MPa.

[0011] Preferably, before the sintering, the nitrogen-doped Nb-Ti alloy powder is pre-pressed and formed in a graphite mold, the nitrogen-doped Nb-Ti alloy powder and the graphite mold are separated by graphite paper, and the axial pressure of the pre-pressing is 45-55MPa.

[0012] Preferably, the solid solution treatment has a vacuum degree of ≤1Pa, a temperature rising rate of 150-300℃ / min, a holding temperature of 1000-1200℃, and a holding time of 30-60min; after the holding, cooling is performed under a protective atmosphere, the protective atmosphere includes argon or nitrogen, the flow rate of the protective atmosphere is 45-55mL / min, and the temperature after the cooling is below 80℃.

[0013] The application further provides an application of the nitrogen-doped Nb-Ti alloy or the nitrogen-doped Nb-Ti alloy prepared by the preparation method as a superconducting material.

[0014] This invention provides a nitrogen-doped niobium-titanium alloy, comprising the following components by mass percentage: 0.08~0.19 wt% nitrogen, 52~54 wt% niobium, and 45.81~47.92 wt% titanium; the nitrogen-doped niobium-titanium alloy contains β phase, α phase, and nitride precipitates. This invention dops a certain amount of nitrogen into the niobium-titanium alloy and effectively dissolves the nitrogen into the alloy, thereby altering the microstructure and phase composition to give the nitrogen-doped niobium-titanium alloy high mechanical properties and good superconducting properties. Test results from the embodiments of this invention show that the optimal properties of the nitrogen-doped niobium-titanium alloy provided by this invention are: compressive yield strength of 1096 Pa, microhardness of 369.4 HV, self-corrosion potential of -0.2016 V, and self-corrosion current density of 2.6096 × 10⁻⁶. -8 A·cm -2 R ct The value is 872190 Ω·cm -2 The critical superconducting transition temperature is 8.7K. Attached Figure Description

[0015] Figure 1 This is a schematic diagram illustrating the process of preparing nitrogen-doped niobium-titanium alloy in an embodiment of the present invention; Figure 2 SEM image of Nb-47Ti alloy powder; Figure 3 This is a real-world diagram of the spark plasma sintering process in Example 1; Figure 4 The XRD patterns are shown for the nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, and the Nb-47Ti alloy powder in Comparative Example 3. Figure 5 The images show actual photos of the nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, and the Nb-47Ti alloy powder in Comparative Example 3. (a) is an image of powder F; (b) is an image of powder F550; (c) is an image of powder F600; (d) is an image of powder F650; (e) is an image of powder F700; and (f) is an image of powder F750. (F represents the original powder, i.e., unnitrided powder; 550, 600, 650, and 700 represent the nitriding temperature (°C) of the powder, respectively.) Figure 6 The images show SEM images of the nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, and the Nb-47Ti alloy powder in Comparative Example 3, where (a) is the SEM image of F; (b) is the SEM image of F550; (c) is the SEM image of F600; (d) is the SEM image of F650; (e) is the SEM image of F700; and (f) is the SEM image of F750. Figure 7EDS spectra of nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, and Nb-47Ti alloy powder in Comparative Example 3; Figure 8 The graph shows the results of the compression performance test, where (a) is a graph showing the changes in sintering temperature and axial shrinkage rate of different samples with sintering time; and (b) is a graph showing the changes in the vacuum degree inside the cavity of different samples with sintering time during the sintering process. Figure 9 The XRD patterns are of the sintered products prepared in Examples 1-3 and Comparative Examples 1-3. Figure 10 The images show the metallographic microstructures, where (a) is the metallographic microstructure of Raw, (b) is the metallographic microstructure of NT550, (c) is the metallographic microstructure of NT600, (d) is the metallographic microstructure of NT650, (e) is the metallographic microstructure of NT700, and (f) is the metallographic microstructure of NT750. (Where Raw represents the sintered sample without nitriding treatment; NT550, NT600, NT650, and NT700 represent samples sintered after nitriding at different temperatures, and the numbers indicate the nitriding temperature.) Figure 11 The images are SEM images of the sintered samples, where (a) is the SEM image of the Raw sample, (b) is the SEM image of the NT550 sample, (c) is the SEM image of the NT600 sample, (d) is the SEM image of the NT650 sample, (e) is the SEM image of the NT700 sample, and (f) is the SEM image of the NT750 sample. Figure 12 The inverse pole diagrams are for Raw, NT600, and NT650, where (a) is Raw, (b) is NT600, and (c) is NT650. Figure 13 The following are grain size distribution diagrams for Raw, NT600, and NT650, where (a) is Raw, (b) is NT600, and (c) is NT650. Figure 14 The distribution of local mismatched angles for Raw, NT600, and NT650 is shown, where (a) is Raw, (b) is NT600, and (c) is NT650. Figure 15 The selected electron diffraction patterns are shown for two precipitated phases inside the NT650 grain and their corresponding calibrations. (a) and (c) are bright field images of the acicular α phase and the TiN phase at the grain boundary inside the NT650, respectively. (b) and (d) are the diffraction spot calibrations for regions 1 and 2. Figure 16 Images (a) to (c) are the original high-resolution images of the interface between the α phase and β phase inside the NT650 matrix, and magnified images of region 1 and region 2, respectively. Figure 17 EDS energy dispersive spectroscopy elemental distribution diagram of the precipitates inside NT650 grains; Figure 18 The image shows a bar chart of the micro Vickers hardness of different sintered samples. Figure 19 The compressive stress-strain curves of the sintered samples of Examples 1-3 and Comparative Examples 1-3 are shown. Figure 20 The figures show macroscopic views and microstructures of the longitudinal sections of Raw, NT600, and NT650 compression specimens, where (a), (b), and (c) are macroscopic views of the longitudinal sections of Raw, NT600, and NT650 compression specimens, respectively; and (d), (e), and (f) are microstructures of the corresponding longitudinal sections of (a), (b), and (c), respectively. Figure 21 Polarization curves for different sintered samples; Figure 22 Impedance spectra of different sintered samples are shown, where (a) is the Nyquist plot and (b) is the high-frequency magnified plot. Figure 23 The equivalent circuit diagrams of electrochemical impedance spectroscopy are shown, where (a) is the equivalent circuit diagram of Raw, NT550, NT600, NT650, and NT750, and (b) is the equivalent circuit diagram of NT700. Figure 24 (a) shows the magnetic susceptibility curves of Raw, NT600, and NT650 as a function of temperature. Figure 24 (b) is a line graph showing the changes in the superconducting critical transition temperature for Raw, NT600, and NT650. Figure 25 In Figure (a), the hysteresis loops of Raw, NT600, and NT650 vary with magnetic field strength. Figure 25 (b) shows the critical current density of Raw, NT600, and NT650 at 4.2K as a function of the magnetic field. Figure 26 XRD patterns for N1000, N1100, N1200, R1000, R1100, and R1200; Figure 27 These are metallographic micrographs, where a~f are metallographic micrographs of R1000, R1100, R1200, N1000, N1100 and N1200 respectively; Figure 28 These are SEM images, where a~f are SEM images of R1000, R1100, R1200, N1000, N1100, and N1200, respectively; (R represents nitrogen-free samples, and the number represents the solution treatment temperature; N represents nitrogen-containing samples, and the number represents the solution treatment temperature;) Figure 29The images are TEM images of N1200, where (a) and (c) are bright field images of TiN at the grain boundaries and the internal needle-like α phase, respectively; (b) and (d) are the diffraction spot markings of corresponding regions 1 and 2. Figure 30 The following are the EDS energy spectra of N1200, where (a) is the EDS energy spectrum of TiN at the grain boundary and (b) is the EDS energy spectrum of the α-Ti phase inside the matrix. Figure 31 The microhardness histograms are for R1000, R1100, R1200, N1000, N1100, and N1200. Figure 32 The results of room temperature compression performance tests for R1000, R1100, R1200, N1000, N1100 and N1200 specimens are shown. (a) is the stress-strain curve and (b) is the yield strength histogram. Figure 33 The potentiodynamic polarization curves are for R1000, R1100, R1200, N1000, N1100, and N1200. Figure 34 The impedance spectrum test results are for R1000, R1100, R1200, N1000, N1100, and N1200, where (a) is the Nyquist plot and (b) is the high-frequency magnified plot. Figure 35 In Figure (a), the magnetic susceptibility of samples R1000, R1100, R1200, N1000, N1100, and N1200 varies with temperature. Figure 35 (b) in the figure is a line graph showing the changes in the superconducting critical transition temperature of samples R1000, R1100, R1200, N1000, N1100, and N1200. Figure 36 (a) shows the hysteresis loop curves of NT600 alloy at different solution temperatures as a function of magnetic field strength. Figure 36 (b) in the middle is J c Curve showing how the magnetic field changes. Detailed Implementation

[0016] This invention provides a nitrogen-doped niobium-titanium alloy, comprising the following components in weight percentage: Nitrogen 0.08~0.19 wt%; Niobium 52~54wt% Titanium 45.81~47.92 wt%; The nitrogen-doped niobium-titanium alloy contains β phase, α phase and nitride precipitates.

[0017] The nitrogen-doped niobium-titanium alloy provided by the present invention contains 0.08 to 0.19 wt% nitrogen, specifically 0.086 wt%, 0.095 wt%, or 0.13 wt%.

[0018] The nitrogen-doped niobium-titanium alloy provided by the present invention comprises 52-54 wt% niobium, specifically 52.79 wt%, 52.84 wt%, 52.91 wt%, 53 wt%, or 53.5 wt% by mass percentage.

[0019] The nitrogen-doped niobium-titanium alloy provided by the present invention comprises 45.81 to 47.92 wt% titanium, specifically 46 wt%, 46.9 wt%, 47.074 wt%, 47.115 wt%, or 47.5 wt% by mass percentage.

[0020] The nitrogen-doped niobium-titanium alloy provided by this invention contains a β phase, an α phase, and nitride precipitates. The volume percentage of the β phase in the nitrogen-doped niobium-titanium alloy can be ≥99%, and the total volume percentage of the α phase and nitride precipitates in the nitrogen-doped niobium-titanium alloy can be ≤1%. As a specific embodiment of this invention, the volume percentage of the β phase in the nitrogen-doped niobium-titanium alloy can be 99.3% or 99.5%, the volume percentage of the α phase can be 0.4%, 0.5%, or 0.6%, and the volume percentage of the nitride precipitates can be 0.1% or 0.2%.

[0021] This invention also provides a method for preparing the nitrogen-doped niobium-titanium alloy described in the above technical solution, comprising the following steps: Niobium-titanium alloy powder is subjected to nitriding treatment to obtain nitrided niobium-titanium alloy powder; the nitriding treatment temperature is 550~650℃. The nitrogen-doped niobium titanium alloy powder was sintered to obtain a nitrided niobium titanium alloy. The nitrogen-doped niobium-titanium alloy was obtained by solution treatment.

[0022] In this invention, unless otherwise specified, all materials are conventional commercially available products.

[0023] This invention involves nitriding niobium-titanium alloy powder to obtain nitrided niobium-titanium alloy powder. In one specific embodiment, the niobium-titanium alloy powder can be Nb-47Ti alloy powder; the particle size of the niobium-titanium alloy powder can be 43~142μm, or 50~100μm, specifically 53μm, 60μm, 70μm, 80μm, or 90μm; in this invention, the Nb-47Ti alloy powder can be commercially produced Nb-47Ti alloy powder, which can be prepared using a plasma rotating electrode process (PREP). In one specific embodiment of the present invention, the nitriding treatment can be carried out in a nitriding furnace. The nitriding treatment may include the following steps: spreading niobium-titanium alloy powder evenly on the surface of a stainless steel mesh, placing the stainless steel mesh with the niobium-titanium alloy powder layer by layer on a stainless steel fixture, and then placing it in a nitriding furnace. Under a first protective atmosphere, the furnace is heated together. After the temperature reaches the temperature required for nitriding treatment, nitrogen-containing gas is introduced for nitriding treatment. After nitriding treatment, a first cooling is performed under nitrogen-containing gas, and a second cooling is performed under a second protective atmosphere to obtain the nitrided niobium-titanium alloy powder.

[0024] In one specific embodiment of the present invention, the stainless steel mesh has a mesh size of 500 mesh, and the stainless steel mesh can be circular with a diameter of 145-155 mm, specifically 150 mm; the mass of niobium-titanium alloy powder on the surface of each layer of stainless steel mesh can be less than 30 g, specifically 25-30 g; the number of layers of stainless steel mesh on the stainless steel tooling can be 2-3 layers, and the distance between adjacent layers can be 20-28 mm, specifically 25 mm; the inner diameter of the nitriding furnace can be 200 mm, and the depth can be 400-600 mm. In another specific embodiment of the present invention, the first protective atmosphere can be ammonia, the purity of which can be 99.99% or higher, and the flow rate of the first protective atmosphere can be 0.48-0.52 m³ / h. 3 / h, which can be specifically 0.5m 3 / h. In this invention, the nitriding treatment temperature is 550~650℃, specifically 550℃, 600℃, or 650℃; the holding time for the nitriding treatment can be 0.5~2h, or 1~1.5h; the nitrogen-containing gas used in the nitriding treatment process may include ammonia, the purity of which can be above 99.99%, and the flow rate of which can be 0.8~1m³. 3The nitrogen-containing gas can be stirred and mixed using a stirring fan in the nitriding furnace at a speed of 140~160 r / min, specifically 145 r / min, 150 r / min, or 155 r / min; the stirring power can be 0.48~0.52 kW, specifically 0.5 kW. In one specific embodiment of the invention, the flow rate of the nitrogen-containing gas during the first cooling process can be 0.48~0.52 m³ / h. 3 / h, which can be specifically 0.5m 3 / h; the temperature after the first cooling can be below 300℃, or it can be 250~280℃; the second protective atmosphere can include nitrogen or argon, the purity of the second protective atmosphere can be above 99.99%, and the flow rate of the second protective atmosphere can be 0.48~0.52m³ / h. 3 / h, which can be specifically 0.5m 3 / h; the temperature after the second cooling can be below 40℃, or it can be 30~35℃.

[0025] After obtaining the nitrided niobium-titanium alloy powder, the present invention sinters the nitrogen-doped niobium-titanium alloy powder to obtain the nitrided niobium-titanium alloy. As a specific embodiment of the present invention, the process may further include: placing the nitrogen-doped niobium-titanium alloy powder in a graphite mold for pre-pressing, wherein the nitrogen-doped niobium-titanium alloy powder and the graphite mold can be separated by graphite paper, the thickness of which can be 0.48~0.52mm, specifically 0.5mm; the axial pressure of the pre-pressing can be 45~55MPa, specifically 50MPa.

[0026] In one specific embodiment of the present invention, the sintering may include plasma discharge sintering (SPS); the vacuum degree of the plasma discharge sintering may be ≤1 Pa, specifically 0.5 Pa or 0.8 Pa; the heating rate of the plasma discharge sintering may be 100~200℃ / min, specifically 120℃ / min, 150℃ / min or 180℃ / min; the temperature of the plasma discharge sintering may be 1000~1200℃, specifically 1000℃, 1100℃ or 1200℃; the holding time of the plasma discharge sintering may be 15~30 min, or 20~25 min; the axial pressure of the plasma discharge sintering may be 40~50 MPa, or 43~45 MPa.

[0027] As a specific embodiment of the present invention, the sintering process may further include: unloading the axial pressure from the sintered system and cooling it to below 40°C under vacuum conditions to obtain a nitrided niobium-titanium alloy.

[0028] After obtaining the nitrided niobium-titanium alloy, the present invention performs solid solution treatment on the nitrided niobium-titanium alloy to obtain the nitrogen-doped niobium-titanium alloy. In one specific embodiment of the present invention, the vacuum degree of the solution treatment can be ≤1 Pa, specifically 0.5 Pa or 0.8 Pa; the heating rate of the solution treatment can be 150~300℃ / min, specifically 180℃, 200℃, 250℃ or 280℃; the holding temperature of the solution treatment can be 1000~1200℃, specifically 1000℃, 1100℃ or 1200℃; the holding time of the solution treatment can be 30~60 min, specifically 35 min, 40 min, 45 min, 50 min or 55 min; after the holding period, cooling can be performed under a protective atmosphere, which may include argon or nitrogen, and the flow rate of the protective atmosphere can be 45~55 mL / min, specifically 50 mL / min; the temperature after cooling can be below 80℃, or 50~70℃; the present invention has no special requirements for the cooling method, as long as the required temperature can be achieved.

[0029] In one specific embodiment of the present invention, the solution treatment can be performed in a discharge plasma sintering apparatus. The specific process is as follows: the surface of the sintered sample is polished and then placed between two electrodes of the discharge plasma sintering apparatus. The sintered sample and the electrodes are separated by a graphite block, and the graphite block and the sintered sample are separated by 2mm thick graphite paper. The present invention does not have a specific limitation on the polishing method; conventional methods in the art are acceptable. The present invention removes the carbon contamination layer formed on the alloy surface during sintering through polishing. In the present invention, the two end faces of the sintered sample are parallel and flat, without obvious defects. To ensure sufficient contact between the graphite block and the upper and lower electrodes, and between the upper and lower electrodes and the sintered sample during heating, axial pressure needs to be applied. The axial pressure can be 10~15MPa, specifically 12MPa, 13MPa, or 14MPa. The solution treatment in the discharge plasma sintering apparatus of the present invention can be performed under pulsed current conditions, which can accelerate electron diffusion while ensuring a relatively fast heating rate.

[0030] Figure 1 This is a schematic diagram of the process for preparing nitrogen-doped niobium-titanium alloy in an embodiment of the present invention. The present invention combines spark plasma sintering (SPS) with gas powder nitriding treatment to achieve controllable nitrogen doping in the niobium-titanium alloy system. The preparation method provided by the present invention can effectively control the morphology, compositional uniformity and nitrogen content of the alloy powder, and significantly improve the sintering density and subsequent mechanical properties of the alloy.

[0031] The present invention also provides the application of the nitrogen-doped niobium-titanium alloy described in the above technical solution or the nitrogen-doped niobium-titanium alloy prepared by the preparation method described in the above technical solution as a superconducting material.

[0032] To further illustrate the present invention, the technical solutions provided by the present invention will be described in detail below with reference to the embodiments, but they should not be construed as limiting the scope of protection of the present invention.

[0033] The raw materials and chemical reagents used in the examples are shown in Table 1, and the equipment used in the examples is shown in Table 3.

[0034] Table 1. Raw materials and chemical reagents used in the examples.

[0035] The Nb-47Ti alloy powder used in the examples is a commercially produced Nb-47Ti alloy powder with a particle size of 53 μm, prepared by plasma rotating electrode atomization method. Its chemical composition is shown in Table 2. Figure 2 The image shows a SEM image of Nb-47Ti alloy powder, which indicates that the powder has good sphericity.

[0036] Table 2 Chemical composition of Nb-47Ti alloy powder (wt.%)

[0037] Table 3. Equipment used in sample preparation and material testing in the examples.

[0038] Example 1 Weigh 60g of spherical Nb-47Ti alloy powder with a particle size of 53μm, and evenly spread it on the surface of two layers of 500-mesh stainless steel mesh (150mm in diameter, 30g of Nb-47Ti alloy powder per layer). Then, place each layer on a stainless steel fixture (25mm apart) and put it into a nitriding furnace with an inner diameter of 200mm and a depth of 500mm. Follow the 0.5m... 3 Ammonia gas with a purity of 99.99% is introduced into the furnace at a flow rate of / h. Under the protective atmosphere of ammonia, the furnace temperature is raised to 550℃, and the flow rate of ammonia gas is adjusted to 1m³ / h. 3 The furnace is held at a constant temperature for 2 hours (nitriding treatment is performed; during the holding process, the ammonia gas inside the furnace is uniformly stirred by a 0.5kW agitator at a speed of 150r / min); after the holding period, the ammonia gas flow rate is adjusted to 0.5m³ / h. 3 After the furnace temperature is reduced to 300℃, the ammonia gas supply is stopped, and the flow is maintained at 0.5m... 3 Nitrogen gas with a purity of 99.99% was introduced into the furnace at a flow rate of / h and cooled to 30°C. The nitrogen supply was then stopped, and the powder was removed to obtain nitrided niobium-titanium alloy powder, denoted as F550. 20g of nitrided niobium-titanium alloy powder was weighed and placed into a graphite mold with an inner diameter of 20mm and a height of 40mm. The powder was separated from the mold by a 0.5mm thick graphite paper. The powder was pre-pressed into shape in the mold under an axial pressure of 50MPa using a pre-pressing device. The mold was then placed in a spark plasma sintering apparatus and heated to 1200℃ at a heating rate of 100℃ / min under a vacuum of 10Pa. The temperature was then maintained at 1200℃ and 40MPa for 10min. After the axial pressure was released, the mold was cooled to 20℃ in the furnace under vacuum and the sample was removed. The resulting nitrided niobium-titanium alloy with a diameter of 20mm and a height of 12mm was denoted as NT550. After the surface of the nitrided niobium-titanium alloy was polished, it was placed between two electrodes in a discharge plasma sintering apparatus. The nitrided niobium-titanium alloy and the electrodes were separated by a graphite block, and the graphite block and the nitrided niobium-titanium alloy were separated by 2 mm thick graphite paper. An axial pressure of 10 MPa was applied between the graphite block and the upper and lower electrodes, and between the upper and lower electrodes and the nitrided niobium-titanium alloy. The alloy was then held at 1000℃, 1100℃, and 1200℃ for 30 min, respectively. After the holding time, the alloy was cooled to 50℃ in a nitrogen atmosphere with a flow rate of 50 mL / min to obtain a nitrogen-doped niobium-titanium alloy. The composition of the nitrogen-doped niobium-titanium alloy included: 0.086 wt% nitrogen, 52.84 wt% niobium, and 47.074 wt% titanium.

[0039] Example 2 Nitrided niobium-titanium alloy powder was prepared according to the method in Example 1, except that the holding temperature of the nitriding treatment was 600°C, and the nitrided niobium-titanium alloy powder was denoted as F600. Nitrided niobium titanium alloy was prepared according to the method of Example 1, except that F600 was used as raw material and the resulting niobium titanium alloy was denoted as NT600. Nitrogen-doped niobium-titanium alloys were prepared according to the method in Example 1, except that NT600 was used as the raw material. The composition of the nitrogen-doped niobium-titanium alloys prepared included: 0.095 wt% nitrogen, 52.79 wt% niobium, and 47.115 wt% titanium.

[0040] Example 3 Nitrided niobium-titanium alloy powder was prepared according to the method in Example 1, except that the holding temperature of the nitriding treatment was 650°C, and the nitrided niobium-titanium alloy powder was denoted as F650. Nitrided niobium titanium alloy was prepared according to the method of Example 1, except that F650 was used as raw material and the resulting niobium titanium alloy was denoted as NT650. Nitrogen-doped niobium-titanium alloys were prepared according to the method of Example 1, except that NT650 was used as the raw material; the composition of the prepared nitrogen-doped niobium-titanium alloy included: 0.19 wt% nitrogen, 52.91 wt% niobium, and 46.9 wt% titanium.

[0041] Comparative Example 1 Nitrided niobium-titanium alloy powder was prepared according to the method in Example 1, except that the holding temperature of the nitriding treatment was 700°C, and the nitrided niobium-titanium alloy powder was denoted as F700. Nitrided niobium titanium alloy was prepared according to the method of Example 1, except that F700 was used as raw material and the resulting niobium titanium alloy was denoted as NT700. Nitrogen-doped niobium-titanium alloys were prepared according to the method in Example 1, except that NT700 was used as the raw material.

[0042] Comparative Example 2 Nitrided niobium-titanium alloy powder was prepared according to the method in Example 1, except that the holding temperature of the nitriding treatment was 750°C, and the nitrided niobium-titanium alloy powder was denoted as F750. Nitrided niobium titanium alloy was prepared according to the method of Example 1, except that F750 was used as raw material and the resulting niobium titanium alloy was denoted as NT750. Nitrogen-doped niobium-titanium alloys were prepared according to the method in Example 1, except that NT750 was used as the raw material.

[0043] Comparative Example 3 (without nitriding treatment) The Nb-47Ti alloy powder is denoted as F; 20g of Nb-47Ti alloy powder was weighed and placed into a graphite mold with an inner diameter of 20mm and a height of 40mm. The Nb-47Ti alloy powder was separated from the mold by a 0.5mm thick graphite paper. The powder in the mold was pre-pressed into shape under an axial pressure of 50MPa using a pre-pressing device and then placed in a spark plasma sintering equipment. Under a vacuum of 10Pa, the temperature was raised to 1200℃ at a heating rate of 100℃ / min, and then held at 1200℃ and 40MPa for 10min. After the axial pressure was unloaded, the sample was cooled to 20℃ in the furnace under vacuum and then removed. The sintered sample with a diameter of 20mm and a height of 12mm was obtained and denoted as Raw. After the surface of the sintered sample was polished, it was placed between the two electrodes of the discharge plasma sintering equipment. The sintered sample and the electrodes were separated by a graphite block, and the graphite block and the sintered sample were separated by a 2 mm thick graphite paper. An axial pressure of 15 MPa was applied between the graphite block and the upper and lower electrodes, and between the upper and lower electrodes and the sintered sample. The samples were then held at 1000℃, 1100℃, and 1200℃ for 30 min, respectively. After the holding time was completed, the samples were cooled to 50℃ in a nitrogen atmosphere with a flow rate of 50 mL / min to obtain a niobium-titanium alloy.

[0044] Figure 3 This is a real-world diagram of the spark plasma sintering process in Example 1.

[0045] The performance of the products prepared in the examples was tested according to the following methods.

[0046] 1) X-ray diffraction analysis XRD patterns of the nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, and the Nb-47Ti alloy powder in Comparative Example 3, were obtained using a PANalytical Empyrean X-ray diffractometer. Figure 4 As shown; the detection conditions were: Cu-Ka as the target material, working voltage of 40 kV, working current of 40 mA, scanning test step size of 0.02°, and scanning range of 20~90°. Characteristic peaks of the β phase were observed in the XRD pattern. The diffraction peak crystal planes mainly included (110), (200), (211), and (220), which appeared near 38°, 55°, 70°, and 82°, respectively. These are typical diffraction peaks of the β phase. As the nitriding temperature increased, the diffraction intensity corresponding to the four diffraction peak crystal planes of the powder gradually weakened. No nitrides or other peaks were observed in the XRD pattern of the alloy powder. This indicates that after nitriding at different temperatures, nitrogen atoms may not have significantly penetrated into the crystal lattice of the powder or a small phase transformation may have occurred that could not be detected. The nitrided powders at 550℃ and 600℃ did not show any new peaks compared to the original powder. The phase shift or obvious peak position shift indicates that nitrogen diffusion and solid solution are still limited at these two nitriding temperatures. After nitriding at 650℃, a slight peak position shift was observed in the powder. This is because nitrogen atoms begin to enter the niobium-titanium matrix in small amounts at higher temperatures, causing slight deformation of the crystal lattice. Similar changes were observed at nitriding temperatures of 700℃ and 750℃. At this temperature, the diffusion of nitrogen atoms was further intensified, and slight structural changes may have occurred in local areas. Furthermore, under high temperature and long-term conditions, subtle changes in peak intensity and width appeared. This was caused by the more complete diffusion of nitrogen into the matrix material, resulting in a slight stress effect on the crystal structure.

[0047] Figure 5 Images show physical samples of the nitrided niobium-titanium alloy powders from Examples 1-3 and Comparative Examples 1-2, and the Nb-47Ti alloy powder from Comparative Example 3, where (a) is a physical sample of F; (b) is a physical sample of F550; (c) is a physical sample of F600; (d) is a physical sample of F650; (e) is a physical sample of F700; and (f) is a physical sample of F750. Figure 5It can be seen that after nitriding at different temperatures, the macroscopic morphology of the niobium-titanium alloy powder changed differently after cooling. As the nitriding temperature increased, the macroscopic color of the alloy powder changed significantly. The powder changed from gray to golden yellow as the nitriding temperature increased. Combined with the structural analysis of the XRD pattern, it can be concluded that nitrogen may have entered the niobium-titanium alloy lattice after nitriding. This macroscopic change in the powder reflects the feasibility of the nitriding scheme.

[0048] 2) Scanning electron microscopy analysis The nitrided niobium-titanium alloy powders in Examples 1-3 and Comparative Examples 1-2, as well as the Nb-47Ti alloy powder in Comparative Example 3, were analyzed using a GeminiSEM 300 scanning electron microscope (SEM) and an equipped Ultim Max 65 energy dispersive spectrometer (EDS). The SEM images are shown below. Figure 6 As shown, (a) is the SEM image of F; (b) is the SEM image of F550; (c) is the SEM image of F600; (d) is the SEM image of F650; (e) is the SEM image of F700; and (f) is the SEM image of F750. The obtained EDS spectrum is shown below. Figure 7 As shown in Table 4, the surface nitrogen content of different test samples is listed.

[0049] Depend on Figure 6 It can be seen that the Nb-47Ti alloy powder has a smooth surface and good sphericity. Its surface has obvious rapid cooling cellular and dendritic structures, and there are no obvious deformations or powder defects. After undergoing nitriding treatment at different temperature gradients, the powder did not undergo significant changes and still maintained similar morphological characteristics to the original powder.

[0050] Table 4. Surface nitrogen content (wt.%) of Nb-47Ti alloy powder at different nitriding temperatures

[0051] Combine Table 4 and Figure 7 ( Figure 7 The peaks of Nb, Ti and N are visible. As the nitriding temperature increases, the nitrogen content on the surface of the alloy powder continuously increases. The nitrogen content on the powder surface after two cycles of holding at 750℃ for 2 hours is as high as 11.8%, indicating that a large number of nitrogen atoms dissolve into the NbTi crystal during the high-temperature nitriding process, and the powder does not undergo obvious deformation or cracking.

[0052] 3) Compression performance test Compression tests were conducted on sintered specimens using a WDW-200 universal testing machine. The specimens were processed into cylindrical compression specimens with a diameter of 6 mm and a height of 9 mm, strictly in accordance with the requirements of GB / T7314-2005. Before the test, the specimens were polished with 600-grit sandpaper to remove machining marks and oxide films from the surface, in order to avoid the influence of processing stress on the test results and thus avoid errors. The specimens were placed in the center of the machine, ensuring that the upper and lower pressure plates of the universal testing machine were parallel to the upper and lower surfaces of the specimens. The testing machine was started and the test data was recorded when the pressure plates just began to contact the upper and lower surfaces of the specimens. The compression rate was 1.0 mm / min throughout the compression process.

[0053] Figure 8 The graph shows the results of the compression performance test, where (a) shows the changes in sintering temperature and axial shrinkage rate of different samples with sintering time; and (b) shows the changes in the vacuum degree inside the cavity of different samples with sintering time during the sintering process. Figure 8 As can be observed in (a), the sintering process of SPS-prepared niobium-titanium alloy mainly consists of three stages: initial sintering, middle sintering, and late sintering. In the initial stage, when the sintering time is less than 6 minutes, neither the sintering temperature nor the axial shrinkage rate changes significantly. This is because the equipment uses infrared thermometry, and temperatures above 569℃ can be detected. In the early sintering stage, low current and low temperature do not cause significant axial shrinkage; in fact, reverse expansion may occur due to thermal expansion and contraction. At this time, the powder surface is activated, and impurities and water vapor adhering to the surface are expelled. Figure 8In (b), the vacuum level inside the cavity of different samples during sintering also shows a slight decrease with sintering time. As the temperature increases, the diffusion mechanism gradually transitions from surface diffusion to grain boundary diffusion and bulk diffusion. The undoped niobium titanium alloy exhibits a high axial shrinkage rate at lower temperatures, while in the nitrogen-doped niobium titanium alloy, the addition of nitrogen atoms has a multi-faceted effect on the axial shrinkage rate. On the one hand, nitrogen atoms, as interstitial atoms, enter the body-centered cubic (BCC) lattice, causing lattice expansion, thereby increasing the elastic interaction between particles at lower sintering temperatures and inhibiting rapid particle displacement, resulting in a slightly lower axial shrinkage rate at low temperatures compared to the undoped sample. On the other hand, the presence of nitrogen atoms enhances the diffusion driving force at high temperatures by introducing point defects and lattice distortion, thus accelerating densification in the later stages. Therefore, the doped sample requires higher temperatures to achieve a comparable diffusion rate and pore closure, resulting in a lower axial shrinkage rate in the high-temperature region during the later stages of sintering. The significant increase indicates that nitrogen doping alters diffusion kinetics, making the densification process more temperature-sensitive. In addition to the above two points, the vacuum degree versus sintering time graph also shows that nitrogen atoms combine to form nitrogen gas, especially in niobium-titanium alloys nitrided after nitriding at 650℃. This leads to drastic changes in vacuum degree in the later stages of sintering. Furthermore, the chemical reactivity of nitrogen at high temperatures may cause partial precipitation of nitrides (TiN), further altering sintering behavior and shrinkage characteristics. Extending the sintering time further amplifies the diffusion-driven cumulative effect in nitrogen-doped samples. Therefore, it can be seen that the axial shrinkage rate of the material gradually increases with increasing nitriding temperature. For each niobium-titanium alloy at different nitriding temperatures, only the axial displacement changes during sintering, i.e., the axial shrinkage rate changes. This phenomenon also implies that its density gradually increases with increasing nitriding temperature. However, it should be noted that excessively long sintering times may induce abnormal grain growth, leading to uneven microstructure and deteriorating material properties.

[0054] 4) Phase analysis X-ray diffraction analysis was performed on the sintered products prepared in Examples 1-3 and Comparative Examples 1-3, and the obtained XRD patterns are shown below. Figure 9As shown in the figure, the original sintered sample without nitriding is still composed of a single β phase. The bulk alloy phase structure of NT550, NT600 and NT650 after sintering is still the same as that of the unnitrided bulk alloy. However, the intensity and width of the diffraction peaks are increasing and gradually shifting to smaller angles to different degrees. The regular change of the diffraction peak produced by the crystal plane (110) is the most obvious. This indicates that the nitrogen atoms dissolved into the β phase lattice and produced lattice distortion, which increased the interplanar spacing. According to the Bragg equation, once the interplanar spacing increases, the 2θ angle where the diffraction peak is located will decrease accordingly. Therefore, the diffraction peak will shift to the left. Furthermore, no α-phase or other nitride phases were detected. This is likely because there were few precipitates, which were detected by XRD. The intensity of the diffraction peak of NT700 decreased and the shift angle returned to the same as that of the unnitrided sintered bulk alloy. This is because nitrogen atoms dissolved in the lattice exist in other forms (such as precipitates). Subsequently, TiN and α-Ti phases were detected in the bulk alloy sintered after nitriding at 750℃. This indicates that nitrogen atoms escaped from the interstitial spaces during sintering and combined with Ti to form TiN, which promoted the precipitation of the α-Ti phase. In summary, there are three situations in which nitrogen atoms dissolve into the β phase: (1) they exist in the octahedral and tetrahedral interstitial spaces of the β phase as dissolved atoms; (2) they escape from the lattice and combine with Ti to form the TiN phase; (3) the addition of nitrogen promotes the precipitation of the α-Ti phase.

[0055] 5) Microstructural morphology analysis The sintered samples were observed under an OLYMPUS CX51 inverted microscope. First, the samples were hot-mounted using an XQ-1 mounting machine. Then, they were polished with wet / dry diamond sandpaper in the following order: 400 grit, 800 grit, 1000 grit, 1500 grit, and 2000 grit. Next, a rough polish was performed using W2.5 diamond polishing paste on a polishing machine with a woolen cloth, followed by a fine polish with a 50 nm OPS silica suspension. After polishing, the samples underwent metallographic etching. The etching solution was prepared with a volume ratio of HF:HNO3:H2O = 2:1:97. The etching solution was evenly applied to the polished surface using a dropper, with an etching time of approximately 30 seconds. After etching, the samples were immediately rinsed with water, then dehydrated with alcohol and dried. Finally, the samples were observed under a metallographic microscope. The resulting metallographic microstructure is shown in the figure below. Figure 10As shown in the figures, (a) is the metallographic microstructure of Raw, (b) is the metallographic microstructure of NT550, (c) is the metallographic microstructure of NT600, (d) is the metallographic microstructure of NT650, (e) is the metallographic microstructure of NT700, and (f) is the metallographic microstructure of NT750. The figures show that the Raw microstructure consists of uniform equiaxed β grains, and a small amount of needle-like precipitates are visible within the matrix. As the nitriding temperature gradually increases, the grains begin to show a significant refining trend at 650℃, and the number of needle-like precipitates increases. In the NT650 matrix, needle-like precipitates are visible within the matrix, as well as another type of precipitate originating at grain boundaries and distributed in a chain-like pattern within the matrix. Precipitates also form at the grain boundaries.

[0056] The sintered sample was examined using a scanning electron microscope (SEM) to obtain SEM images, such as... Figure 11 As shown, (a) is the SEM image of Raw, (b) is the SEM image of NT550, (c) is the SEM image of NT600, (d) is the SEM image of NT650, (e) is the SEM image of NT700, and (f) is the SEM image of NT750. It can be observed that a few needle-like precipitates with lengths in the micrometer range and widths in the nanometer range exist in Raw, and the grains are coarse at this time, mainly composed of β-phase equiaxed grains. After nitriding, the grains of NT550, NT600, NT650, and NT700 are significantly refined, and with the increase of nitriding temperature, more and more precipitates of the sintered alloy are observed. Initially, only a few dispersed needle-like precipitates can be seen within the grains of Raw and NT550. When the nitriding temperature reaches 600℃, some precipitates originating from the grain boundaries can be observed within the grains. The needle-like precipitates gradually connect into chains, and some precipitates are also present at the grain boundaries. This phenomenon is more obvious when the nitriding temperature reaches above 650℃. It can be clearly observed that many needle-like precipitates are dispersed at the grain boundaries and inside the grains of NT650 and NT700, and the current situation of connecting into chains from the grain boundaries to the grains is more obvious. The microstructure of NT750 has undergone a qualitative change. The internal microstructure of the matrix is ​​composed of β phase and TiN and α-Ti needle-like precipitates inside it, which is consistent with the XRD pattern results of the sintered bulk alloy.

[0057] The nitrogen content of NT600 and NT650, which exhibit the most significant changes in precipitated phases and grain size, was detected using an oxygen, nitrogen, and hydrogen analyzer. The results are listed in Table 5. The testing conditions included: the sample size was 5×5×2mm, the samples were polished to 600 mesh with SiC sandpaper, and cleaned with an acetone ultrasonic cleaner. Three samples were measured for each group of samples, and the average value was taken.

[0058] Table 5. Nitrogen content (wt.%) in sintered samples

[0059] 6) Electron backscattering diffraction analysis Electron backscattered diffraction (EBSD) was used to analyze the sintered samples. Sample preparation: After grinding and rough polishing, 10×10×3mm sample blocks were subjected to electrolytic polishing for 30–50 s. For the niobium-titanium alloy electrolyte, a volume ratio of 1:7 of perchloric acid and ethanol was selected. The polishing voltage was maintained at 30V, and the polishing current was 0.3–0.6A. EBSD data were collected using a FEI Nova NanoSEM 450 with an operating voltage of 30kV and a scan step size of 1μm. The raw data were analyzed using TSL-OIM (TexSEM Laboratories-Orientationimaging microscopy) analysis software, and the obtained data are listed in Table 6. Figure 12 The inverse pole figures (IPF) are for Raw, NT600, and NT650, where (a) is Raw, (b) is NT600, and (c) is NT650. Figure 13 The figures show the grain size distribution of Raw, NT600, and NT650, where (a) is Raw, (b) is NT600, and (c) is NT650.

[0060] Table 6. EBSD test data of Raw, NT600, and NT650 niobium-titanium sintered alloys.

[0061] The phase volume content and grain diameter in the alloy were obtained based on EBSD testing. Table 7 shows the volume percentage of β phase, α phase and nitride precipitates and the average grain size in Raw, NT600 and NT650.

[0062] Table 7. Volume percentage of each phase and grain size of Raw, NT600, and NT650

[0063] from Figure 12 and Figure 13 It can be seen that the grain orientation distribution of the three sintered samples is relatively chaotic, mainly dominated by the three directions of (001), (111) and (101). Combined with the EBSD test results in Table 6, it can be seen that NT650 also has (0001). and The grain orientation distribution was observed, and it was clearly observed that the nitrogen-containing NT600 and NT650 niobium-titanium sintered samples had finer grain sizes than the raw samples, indicating that nitrogen doping can effectively inhibit grain growth. As the nitrogen content gradually increased, the small-angle grain boundaries and subgrains of the niobium-titanium alloy gradually decreased. This is likely because after nitrogen doping, nitrogen atoms segregated at the grain boundaries, hindering further grain growth and inhibiting the formation of small-angle grain boundaries and subgrains. In niobium-titanium alloys, nitrogen, as an interstitial atom, affects dynamic recovery and enhances dynamic recrystallization in the following ways, thereby leading to a reduction in small-angle grain boundaries and subgrains: (1) Solute drag effect of nitrogen atoms Effect), nitrogen atoms accumulate at grain boundaries and dislocations, reducing the ability of dislocation climb and cross-slip, and limiting dynamic recovery. This phenomenon has been reported in many steel materials containing ferrite and austenite structures; (2) the formation of nitrides (such as TiN) or α-Ti phases, and the dispersed particles also hinder dislocation movement and grain boundary migration, further inhibiting grain growth and reducing dynamic recovery efficiency.

[0064] Figure 14 The distribution of local mismatch angles (KAM) in Raw, NT600, and NT650 niobium-titanium sintered alloys is shown, where (a) is Raw, (b) is NT600, and (c) is NT650. Figure 14 It can be seen that the KAM value of the Raw sample is lower, and the color is mainly blue, indicating that the overall strain of the material is smaller, the internal dislocation density is lower, the mismatch angle distribution is more uniform, no obvious strain concentration region is formed, and there are more small-angle grain boundaries. Compared with Raw, the KAM values ​​of NT600 and NT650 are generally higher, with more green and some yellow in local areas. The yellow-green color is mostly concentrated at the grain boundaries, indicating that the stress is mainly concentrated at the grain boundaries, and the stress in the intragranular region is relatively lower. It is very likely that nitrogen doping promotes dislocation accumulation at the grain boundaries, so that the strain energy is stored near the grain boundaries. This is similar to the results of the inverse pole figure and grain orientation distribution analysis. The solute drag effect of nitrogen atoms is very likely to exist. The segregation of nitrogen atoms at grain boundaries hinders dislocation movement (i.e., the pinning effect on dislocations) and grain boundary shift, causing strain to accumulate at the grain boundaries. The stress concentration at the grain boundaries is more intense in the NT650 sample than in the NT600 sample, forming a continuous high-strain zone. This indicates that the local strain of the NT650 sample is greater after the nitrogen content is further increased. The grain boundary strengthening caused by this stress concentration can improve the strength of the material, but at the same time, this local stress concentration can also lead to the risk of grain boundary slip or intergranular fracture. Therefore, when nitrogen-doped niobium-titanium alloys, it is necessary to balance the mechanical properties and avoid excessive nitrogen doping leading to increased brittleness.

[0065] 6) Transmission electron microscopy analysis Samples were taken from the target specimen using an electrical discharge wire cutter, cutting thin slices of 10mm × 10mm × 0.5mm. These slices were then successively ground with 180-mesh, 400-mesh, 800-mesh, 1000-mesh, and 2000-mesh wet sandpaper to a thickness of 60-70μm. Afterward, they were punched to obtain circular slices with a diameter of 3mm. These circular slices were then embedded in the bottom of a wooden stopper and successively ground with 2000-mesh, 2400-mesh, and 3000-mesh wet sandpaper to a thickness of 40μm. The slices were then sectioned using a Leica EM UC7 ultramicrotome, and subsequently ion-thinned using a GatanPIP Model 695 ion thinner. The preparation of the specimens should adhere to three standards: the specimens should be as thin as possible, the specimens should possess sufficient strength, and the experiments should be repeatable. After the sample was prepared, the crystal structure, second phase and dislocations of the alloy were characterized by a Thermo Scientific FEITalos F200X transmission electron microscope (TEM) at a voltage of 20 kV, and the elemental content of the phases in the alloy was quantitatively analyzed.

[0066] Figure 15 The figures show the selected electron diffraction (SEED) patterns of two precipitated phases within the NT650 grains, along with their corresponding calibrated patterns. (a) and (c) are bright-field images of the acicular α phase, matrix β phase, and TiN phase at the grain boundaries within NT650, respectively. (b) and (d) are the diffraction spot calibrations for regions 1 and 2, respectively. The figures reveal a face-centered cubic TiN phase at the grain boundaries of NT650, while a close-packed hexagonal acicular α phase is observed within the grains. Furthermore, the acicular α phase is abundant within the grains, with lengths ranging from nanometers to micrometers and widths at the nanometer level. Figure 15 From the electron diffraction patterns of the α phase and the matrix β phase in (b), it can be seen that the β phase and the α phase are nearly parallel, while from... Figure 15 The selected area electron diffraction pattern in (d) shows that TiN at the grain boundaries has a standard face-centered cubic structure.

[0067] Figure 16 Images (a) to (c) show the original high-resolution images of the α-phase and β-phase interface within the NT650 matrix, and magnified views of regions 1 and 2, respectively. Figure 16 In (a), a transition layer with a thickness of about 3 nm can be seen at the interface between the α phase and the β phase. Figure 16 From (b) and (c), we can see that the α phase... It is nearly parallel to the β phase. as well as Crystalline.

[0068] EDS elemental distribution analysis was performed on the acicular phases at grain boundaries and within the matrix. Figure 17The image shows the elemental distribution of the precipitates inside the NT650 grains using EDS energy dispersive spectroscopy. Figure 17 (a) represents the TiN phase at the grain boundary. Figure 15 As shown in (c), the crystal structure at the grain boundaries is face-centered cubic. Among the Nb, Ti, and N elements, only NbN and TiN exhibit a face-centered cubic structure. According to EDS energy dispersive spectroscopy results, this phase can only be a TiN precipitate. This is because during powder nitriding, N element first accumulates on the powder surface, and the nitrogen content decreases gradually from the powder surface to the powder interior. After sintering, the N content at the grain boundaries becomes even richer, leading to the precipitation of TiN at these nitrogen-rich grain boundaries. The presence of TiN at the grain boundaries increases the stress at the grain boundaries and also inhibits dislocation movement, thus enhancing the mechanical properties of the material. Figure 17 (b) is the EDS energy spectrum elemental distribution diagram of the needle-like α phase inside the matrix. As can be seen from the figure, a small amount of N element is also enriched inside the α phase. These nitrogen elements are dissolved inside the α phase, which reflects that the presence of N element promotes the precipitation of α phase inside the niobium-titanium alloy, and the size of the precipitated phase ranges from nanometer to micrometer.

[0069] 7) Mechanical property analysis The HVS-1000 Vickers microhardness tester was used to test the microhardness of niobium-titanium alloy. First, a rectangular sample with dimensions of 10mm × 10mm × 6mm was cut using an EDM wire cutter for microhardness testing. Before testing, the test surface (10mm × 10mm) needed to be polished to avoid the influence of microparticles and surface impurities on the results. The five-point testing method was used, with a load set at 100gf, a holding time of 15s, and a diamond indenter angle of 136°. A diamond-shaped indentation with widths d1 and d2 was generated on the sample surface. The Vickers hardness was calculated as follows: Where F is the load (N) and S is the indentation surface area (mm²). 2 ), d1d2 are the average lengths of the indentation diagonals (mm), α is the included angle (°) of the diamond square pyramid indenter, the diamond included angle is 136°, therefore the formula can also be written as .

[0070] Figure 18The graph shows the micro Vickers hardness histograms of different sintered samples. To reduce error, ten points were randomly selected on the sample surface for measurement and error bar calculation. The graph shows that the hardness value of the sintered samples gradually increases with increasing nitriding temperature. Although the increase in hardness between NT600 and NT550 is not significant, the microhardness of both NT550 and NT600 is still higher than that of Raw. This is because nitrogen atoms diffuse into the interstitial spaces of the niobium-titanium lattice, resulting in solid solution strengthening. Furthermore, due to the concentration gradient of nitrogen atoms diffused into the powder from the outside to the inside during powder nitriding, the nitrogen concentration at the grain boundaries is relatively high after the formation of the sintering neck during sintering. This leads to stress concentration at the grain boundaries, resulting in grain boundary strengthening. However, with increasing nitrogen content, the microhardness of NT650 has increased by 91.3% compared to Raw. This is because NT650 not only has a higher nitrogen content than Raw, but also... The presence of more nitrogen atoms and the addition of N elements led to grain refinement. Furthermore, numerous new phases precipitated along grain boundaries and within the grains were formed in the β phase. This synergistic effect of precipitation strengthening, grain refinement, grain boundary strengthening, and solid solution strengthening resulted in a continued increase in the microhardness of NT650. Similarly, NT700 and NT750, due to the addition of even more N elements, exhibited an increasing number of precipitates. These strengthening effects further enhanced their microhardness. Notably, NT750, due to excessive N diffusion, exhibited a very high proportion of TiN and α phases, with TiN accounting for 0.4% and α phases for 0.6% by volume. The lower hardness of the β phase compared to TiN and α-Ti phases resulted in uneven hardness distribution.

[0071] Axial compressive properties at room temperature were tested for different sintered samples. Figure 19 Table 8 shows the compressive stress-strain curves of the sintered samples of Examples 1-3 and Comparative Examples 1-3. The table also shows the yield strength, limiting compressive strength, and fracture compressive strain values ​​of different sintered samples.

[0072] Table 8. Compressive performance data of sintered samples from Examples 1-3 and Comparative Examples 1-3

[0073] Because the test was stopped when the sample was compressed to 70% strain, the samples Raw, NT550, NT600 and NT650 had not yet shown obvious fracture, so there is no corresponding ultimate compressive strength and fracture strain in Table 8.

[0074] from Figure 19As shown in Table 8, the yield strength of the sintered samples exhibits a continuous increasing trend. The yield strength of the nitrided samples is significantly improved compared to the untreated samples. The yield strengths of NT550 and NT600 are increased by 113% and 126% respectively compared to Raw. This is because, at nitriding temperatures of 550℃ and 600℃, although the nitrogen content is very low, the uniform diffusion of nitrogen atoms into the matrix plays a role in solid solution strengthening. Furthermore, the yield strengths of NT550 and NT600 are very close. The higher yield strength of NT650 is due to the fact that the nitrogen content of NT650 is a critical point. At the nitriding temperature of 650℃, not only is there a solid solution strengthening effect, but the microstructure also shows that the NT650 samples have more grain boundaries, finer grains, and a large number of acicular phases precipitated in the matrix compared to Raw, NT550, and NT600. Alloys with fine-grained microstructures can disperse stress concentration within more grains during plastic deformation under external forces, thus reducing stress concentration. Simultaneously, the smaller grains provide larger grain boundary areas, making crack propagation more difficult. While NT700 and NT750 exhibit very high yield strength, their excessive nitrogen content leads to excessive stress at grain boundaries and within the matrix. Furthermore, the TiN present in the NT750 microstructure itself is a phase with poor fracture toughness. Therefore, apart from the NT700 and NT750 samples showing obvious fracture strain points, the NT650, NT600, and NT550 samples all produced strain hardening curves similar to the Raw curve. This indicates that at these three nitriding temperatures, nitrogen-containing niobium-titanium alloys not only improve yield strength but also possess relatively good toughness and plasticity, with NT600 exhibiting the best overall compressive strength. Figure 20 These are macroscopic and microscopic views of the longitudinal sections of Raw, NT600, and NT650 compression specimens. (a), (b), and (c) are macroscopic views of the longitudinal sections of the Raw, NT600, and NT650 compression specimens, respectively; (d), (e), and (f) are microscopic views of the corresponding longitudinal sections of (a), (b), and (c), respectively. Figure 20 It can be seen that the strain levels of the three compression specimens are consistent. Except for the difference in grain size between NT600 and Raw, the compression streamlines of the longitudinal sections of the two compression specimens are basically the same. However, cracks were found near the streamlines of the longitudinal section of the NT650 compression specimen, indicating that NT600 has the same good plasticity as Raw.

[0075] 8) Corrosion resistance analysis Electrochemical polarization curve test The potentiodynamic polarization curves of niobium-titanium sintered samples were measured using a Corrtest CS350 electrochemical workstation (Wuhan Corrtest Instruments Co., Ltd.), and their electrochemical characteristics were investigated. The samples were circular discs with a diameter of 20 mm and a thickness of 3 mm. The sample surface was ground and polished with silicon carbide sandpaper at a grit of 400–2000 grit. To ensure good adhesion and contact between the copper wire and the sample, LX30 conductive silver paste was used to firmly bond the copper wire to the untested side of the sample. Electrochemical tests were performed using a 3.5 wt% NaCl solution as the electrolyte in a conventional three-stage battery, with an active electrode exposure area of ​​3.14 cm². 2 The sample, platinum electrode, and saturated calomel electrode (SCE) were used as the working electrode, auxiliary electrode, and reference electrode, respectively. Before polarization curve testing, the sample was exposed to 3.5 wt% NaCl solution for 1 hour to identify the open circuit potential. Then, potentiodynamic polarization scanning was performed. The potentiodynamic polarization curve was tested using a Gamry instrument with a scanning range of -1 to 2.5 V and a scanning speed of 1 mV / s. Finally, the self-corrosion potential Ecorr and self-corrosion current density Icorr of the corresponding sample were calculated from the obtained data using the Tafel linear extrapolation method.

[0076] Electrochemical impedance spectroscopy (EIS) The electrochemical impedance spectroscopy sample preparation and testing methods employed the same equipment as the polarization test, with a sinusoidal voltage amplitude of 5mV and a test frequency range of 10. 5 ~10 -2 The Nyquist plot of the material was obtained by testing. Subsequently, the measured impedance data was fitted using ZView software by fitting an equivalent circuit diagram. The corrosion resistance of the relevant samples was evaluated by combining the solution resistance, charge transfer resistance and other related data obtained after fitting with the corrosion potential obtained from the polarization curve and the self-corrosion current density.

[0077] Figure 21 Polarization curves for different sintered samples. Figure 22 Impedance spectra of different sintered samples are shown, where (a) is the Nyquist plot and (b) is the high-frequency magnified plot. Figure 23 The equivalent circuit diagrams of electrochemical impedance spectroscopy (EIS) are shown below. Rs is the solution resistance, CPE is the constant phase angle element related to double-layer capacitance, Rct is the charge transfer resistance, and W is the Weber impedance. Among them, (a) is the equivalent circuit diagram of Raw, NT550, NT600, NT650, and NT750, and (b) is the equivalent circuit diagram of NT700. Table 9 shows the main parameters of the potentiodynamic polarization curves of different sintered samples, and Table 10 shows the EIS fitting results of different sintered samples.

[0078] Table 9. Main parameters of the potentiodynamic polarization curves of the sintered samples from Examples 1-3 and Comparative Examples 1-3.

[0079] Table 10. EIS fitting results of sintered samples from Examples 1-3 and Comparative Examples 1-3

[0080] Figure 21 The figure shows the potentiodynamic polarization curves of the sintered samples. It can be seen from the figure that all samples exhibited passivation and secondary passivation regions; in some samples, the passivation region was not even completely finished within the testing range. Table 9 shows the self-corrosion potential E obtained after fitting the polarization curves using CView software. corr and corrosion current density I corr The corrosion potential is an electrochemical parameter that mainly represents the thermodynamic corrosion tendency of a material. When characterizing the tendency of a material to corrode, the higher the corrosion potential (the more positive), the lower the tendency of the material to corrode. The corrosion current density reflects the corrosion reaction rate and is a kinetic parameter. The lower its value, the slower the rate of corrosion.

[0081] Figure 21 All samples exhibited self-passivation in the anodic polarization region. During self-passivation, the corrosion current density of the material slowly increased with the corrosion potential, indicating that all samples formed a passivation film with certain corrosion resistance. However, with increasing voltage, the passivation film could not withstand the Cl... - The corrosion was gradually destroyed, and pitting corrosion appeared. Table 9 clearly shows that as the nitriding temperature increased, the nitrided niobium-titanium alloy exhibited both a low corrosion potential and a low self-corrosion current density. This phenomenon can be attributed to several possibilities: First, a dense and stable passivation film formed on the surface of the Nb-47Ti material, inhibiting the corrosion rate. Even with a low corrosion potential, the passivation film can significantly reduce the activity of the metal, thereby reducing the self-corrosion current density. Second, although the material's corrosion potential is low, the charge transfer process is limited due to the blocking effect of Nb and Ti corrosion products (such as Nb2O5, TiO2, etc.), which also reduces the self-corrosion current density.

[0082] Regarding the results of Tafel fitting of the polarization curves, although a higher corrosion potential (more positive) generally indicates greater thermodynamic stability, lower self-corrosion current density, and lower corrosion rate, when these two factors contradict each other, it is sufficient to focus on the corrosion current density exhibited in the potentiodynamic polarization test, as it more directly reflects the kinetic information of the corrosion rate. Table 9 also shows that the corrosion current densities of NT650, NT600, and NT550 are all lower than those of Raw, while the corrosion current densities of NT700 and NT750 are higher. This indicates that trace amounts of nitrogen doped into niobium-titanium alloys can improve their corrosion resistance. This is due to the increased density of the sintered sample caused by nitrogen solid solution. NT600 has the lowest corrosion current density, at 9.44278 × 10⁻⁶. -8 A·cm -2 The results clearly show that its corrosion resistance is the best among the six groups of samples. Conversely, if the nitrogen content is too high, such as NT700 and NT750, too many precipitates will appear, causing its corrosion resistance to decrease. At this time, even if the density is high enough, it cannot change the adverse effect of the precipitates on the corrosion resistance.

[0083] from Figure 22 As can be seen, the Nyquist plot of NT700 consists of two parts: a high-frequency region and a low-frequency region. In the high-frequency region, charge transport and double-layer capacitance effects typically occur at the electrode interface, manifesting as capacitive impedance, with a shape resembling an arc called a capacitive arc. In the low-frequency region, the presence of straight lines generally indicates the kinetics of electrochemical reactions, usually dominated by diffusion. The reaction rate at the electrode surface is relatively slow at this point, and Cl- ions in the NaCl solution need to diffuse to the surface to participate in the reaction, which has a certain impact on the corrosion of the sample. Furthermore, the capacitive arc portion in the high-frequency region of NT700 is not obvious, with a relatively large proportion of straight lines, indicating that its... The corrosion resistance was the worst among all samples: except for NT700, the Nyquist plots of all other samples consisted of a single capacitive arc, and similar results to the potentiodynamic polarization test were observed. The capacitive arcs of NT550, NT600 and NT650 were larger than the Raw arc, while the capacitive arc radius of NT750, which had more nitrogen doping, was smaller than the Raw arc due to more precipitates. This further indicates that the niobium-titanium alloy with trace amounts of nitrogen doping has better corrosion resistance. Among them, NT600 showed the largest trend in capacitive arc radius, that is, the higher the impedance modulus (Z) in the low-frequency range, the lower the overall corrosion rate, indicating that its corrosion resistance was the best among all sintered samples.

[0084] As shown in Table 10, the fitting errors are all below 5%, indicating that the equivalent circuit was selected correctly. Electrochemical impedance spectroscopy can analyze the protective performance of passivation films or corrosion products formed on the material surface. A higher charge transfer resistance (Rct) value indicates stronger corrosion resistance of the material. As can be seen from the table, with the increase of nitrogen content, the Rct value shows a trend of first increasing, then decreasing, and then increasing again. The passivation films produced by the trace nitrogen-doped niobium-titanium alloys NT550, NT600, and NT650 are more stable and have better corrosion resistance. Among them, NT600 has the largest Rct value and the best corrosion resistance. The fitting results are consistent with the trend shown by the polarization curves above.

[0085] 9) Superconducting performance analysis Measurement of critical transition temperature (Tc) The critical superconducting transition temperature (Tc, K) is an important parameter for evaluating superconducting performance. When the temperature drops below a certain critical value, the material transitions from a normal state to a superconducting state; the corresponding critical value is the critical superconducting transition temperature Tc. The superconducting transition temperature Tc is generally measured using magnetic susceptibility. This paper uses the third-generation Magnetic Property Measurement System (MPMS3) from Quantum Design, USA, to measure the trend of the critical transition temperature of the alloy. The sample size was cut into 2mm × 2mm × 3mm pieces, and the surface was polished to remove wire cutting marks and oxide film. Measurements were performed under a 10Oe external magnetic field using zero-field cooling (ZFC), with a temperature range of 2~22K. The MT curve of the sample was obtained. The temperature corresponding to the diamagnetic signal transition point in the MT curve is the superconducting transition temperature Tc, which is the temperature corresponding to the first point where the magnetic moment of the ZFC curve coincides with the zero magnetic moment line. The transition width is T90%~T10%.

[0086] Measurement of critical current density (Jc) The critical superconducting current density Jc is another important parameter for evaluating superconducting performance and a crucial indicator for measuring the quality of commercially available superconducting materials. The testing method for critical current density is essentially the same as that for critical transition temperature, and the testing equipment and sample requirements are identical, differing only in parameters. For critical current density testing, the sample needs to be at a temperature of 4.2K, with a magnetic field range of 0 to 6T, ultimately obtaining a hysteresis loop (MH curve). Finally, the critical current density Jc is calculated using the Bean model based on the obtained hysteresis loop. This invention uses cuboid alloy samples with dimensions of 2×2×3mm. The calculation formula for the bulk samples is as follows: In the formula, Jc Critical temperature, unit: A / cm 2△M (M+-M-) is the difference in magnetic moment per unit volume between the applied and demagnetized fields under the same hysteresis loop, expressed in emu / cm. 3 a and b represent the width and length of the block sample, respectively, and a < b, in cm.

[0087] Figure 24 In the middle (a), the magnetic susceptibility of Raw, NT600 and NT650 as a function of temperature (MT) is shown. Figure 24 (b) is a line graph showing the changes in the superconducting critical transition temperature for Raw, NT600, and NT650. Figure 24 As shown in (a), all three samples exhibit a rapid superconducting transition trend, a characteristic of high-quality superconducting phase formation. The superconducting critical transition temperature (CCT) of the Raw sintered sample is 9.0 K, while the CCTs of NT600 and NT650 decrease by 1.4 K (15.6%) and 2.1 K (23.3%) respectively compared to Raw. This indicates that with increasing nitriding temperature, nitrogen content increases, and some nitrogen atoms enter the niobium-titanium lattice, causing lattice distortion. Simultaneously, the precipitation of the TiN phase reduces the volume fraction of the superconducting phase. These factors combined cause the superconducting transition temperature (Tc) to gradually decrease. The superconducting transition relies on controlling the mutual vibration of the crystal lattice. This suggests that for niobium-titanium alloys, excessive defects in the matrix and the distribution of non-superconducting impurity phases affect the superconducting transition temperature. Whether it's the lattice distortion caused by trace amounts of nitrogen dissolving into the matrix, or the increased TiN precipitate due to higher nitrogen content, both have an adverse effect on the critical transition temperature.

[0088] Nitrogen solid solution induces localized distortion of the niobium-titanium (BCC) lattice, altering the lattice vibrational frequency and thus affecting the phonon-electron coupling strength. This is one of the main microscopic mechanisms underlying the decrease in Tc, especially in NT650. Due to the high nitrogen content, not only is lattice distortion exacerbated, but the precipitation of the TiN phase further reduces the volume fraction of the superconducting phase, resulting in a significant decrease in Tc compared to the raw sample. This indicates that the addition of nitrogen needs to be controlled within a reasonable range to avoid compromising the integrity of the superconducting phase.

[0089] Figure 25(a) shows the variation of hysteresis loops of Raw, NT600, and NT650 with magnetic field strength B(T) measured under magnetic field conditions of 4.2K and -6T to 6T. For superconductors, the hysteresis loop generally reaches its maximum value at zero magnetic field. Subsequently, the magnetization and the width of the hysteresis loop decrease under the influence of increasing magnetic field. Raw, NT600, and NT650 do indeed exhibit the characteristics of a standard superconductor. However, NT600 and NT650 show a second peak effect, also known as the "fishtail effect," as the external magnetic field reaches a moderate intensity. For low-temperature metal or alloy superconductors like NbTi alloys, the flux lattice begins to soften near the upper critical magnetic field. The softened flux lines can more easily find and be pinned to defects with lower potential energy. For NbTi superconducting alloys... α -Ti precipitates are pinned phases, which is why NT600 and NT650 exhibit this second peak effect, resulting in an increased current density and a sharp peak. Furthermore, the NT650... α The Ti precipitate is more abundant than NT600, so the peak of NT650 is wider and higher. However, this second peak effect usually appears near the upper critical field, which indicates that the addition of nitrogen element changes the original upper critical field of Nb-47Ti.

[0090] Figure 25 (b) shows the critical current density of Raw, NT600, and NT650 at 4.2K as a function of the magnetic field, calculated based on the Bean model. J c - B 4.2K is one of the environments in which niobium-titanium superconducting alloys are most commonly used. From Figure 25 (a) It can be seen that at a temperature of 4.2K, the NT600's J c The values ​​are almost identical to those in the Raw data, and even slightly higher in a non-magnetic field environment. This indicates that nitrogen, in the form of trace solid solutions, enters the Nb-47Ti alloy to a certain extent, improving its critical current density at 4.2 K, while that of NT650... J c The value is lower than that of the Raw value. This is primarily due to the presence of a small amount of TiN precipitation in NT650. The transformation temperature of TiN (4.5 K) is lower than that of the Nb-47Ti alloy. Based on the critical current density test results, the presence of TiN is detrimental to the critical current density of the Nb-47Ti alloy. Secondly, for niobium-titanium alloys, while the increased number of grain boundaries increases the effective current-carrying area and grain connectivity, it does not benefit the current-carrying capacity of the niobium-titanium superconducting alloy under an applied magnetic field. This leads to a decrease in the critical current density. Therefore, even though NT650 contains substances that could improve flux pinning ability...α The reason why this phase still cannot increase its critical current density is that, not to mention the NT650... α There are not many phases, and their distribution is uneven; there are only more evenly distributed and fine phases. α The presence of the phase has a beneficial effect on the critical current density of niobium-titanium alloys.

[0091] In summary, the superconducting properties of nitrogen-containing Nb-Ti alloys exhibit a clear trend with changes in nitrogen content, with trace amounts of nitrogen (NT600) slightly improving the properties through the synergistic effect of solid solution and precipitated phases. J c And slightly reduced T c Excess nitrogen (NT650 and above) leads to increased lattice distortion and an increase in non-superconducting phases, resulting in... T c Significant decline J c On the contrary, it is suppressed. This relationship between structure and performance can be summarized as follows: moderate nitrogen doping refines the grains and increases pinning centers, which helps to improve current carrying capacity and performance under magnetic fields; while excessive nitrogen content, non-superconducting precipitates such as TiN block the superconducting channel, and the grain boundary defects caused by grain boundary stress concentration together lead to a decrease in superconducting performance.

[0092] 10) Effects of solution treatment Effect of solution treatment on phase composition Nitrogen-doped niobium-titanium alloys obtained by solution treatment of NT600 at 1000℃, 1100℃, and 1200℃ were designated N1000, N1100, and N1200, respectively; niobium-titanium alloys obtained by solution treatment of raw materials at 1000℃, 1100℃, and 1200℃ were designated R1000, R1100, and R1200, respectively. X-ray diffraction analysis was performed on N1000, N1100, N1200, R1000, R1100, and R1200, and the XRD patterns were obtained as follows: Figure 26 As shown in the figure, the phase composition of samples N1000, N1100, N1200, R1000, R1100, and R1200 is the same as that of the Raw sample, remaining in the β phase, with no other new phase precipitation detected. This indicates that within the selected heat treatment temperature range, the undoped Nb-47Ti alloy maintains β phase stability without significant phase transformation.

[0093] As the solution temperature increases, the intensity of the main diffraction peak gradually increases, indicating that a higher solution temperature is beneficial to the stability of the β phase. It can be seen that, compared with Raw, the diffraction peaks of R1000, R1100 and R1200 do not shift significantly with the increase of solution temperature, indicating that different solution temperatures have little effect on the lattice parameters of Raw and still maintain a relatively stable β phase structure. However, in nitrogen-containing samples (N1000, N1100, N1200), the positions of the diffraction peaks shifted significantly. With increasing solution temperature, the N-series samples shifted to smaller angles first (N1000) and then to larger angles (N1100, N1200) compared to NT600. This phenomenon indicates that at lower temperatures, nitrogen solution leads to lattice expansion, increasing the lattice constant, thus causing the XRD diffraction peaks to shift to smaller angles. As the solution temperature increases, the amount of nitrogen dissolved in the β phase changes due to the increase in precipitates caused by nitrogen, and the lattice distortion gradually decreases, leading to a shift of the diffraction peaks to larger angles. This further illustrates that with increasing solution temperature, nitrogen leads to a reduction in intracrystalline lattice distortion.

[0094] In summary, the XRD results show that: (1) the phase structure remains basically stable after solid solution treatment, whether it is raw or nitrogen-containing NT600; (2) after solid solution treatment, nitrogen-containing NT600 reduces the amount of nitrogen dissolved due to the increase of nitrogen-containing precipitates, which is beneficial to the release of lattice distortion.

[0095] Effects of solution treatment on microstructure R1000, R1100, R1200, N1000, N1100, and N1200 were observed using an optical microscope and a scanning electron microscope to obtain metallographic micrographs and SEM images. Figure 27 These are metallographic micrographs, where a~f are metallographic micrographs of R1000, R1100, R1200, N1000, N1100 and N1200 respectively; Figure 28 The image is a SEM image, where a~f are SEM images of R1000, R1100, R1200, N1000, N1100 and N1200 respectively.

[0096] from Figure 27 As can be seen from the metallographic images, the microstructure and grain size of R1000, R1100 and R1200 obtained after three different temperature solution treatments are basically consistent, and there are few precipitates in the matrix and no excessive changes in microstructure morphology. This indicates that the microstructure of the nitrogen-free Nb-47Ti alloy is relatively stable as the solution treatment temperature increases. Figure 28SEM images (a) to (c) further demonstrate that the solution temperature has little effect on its microstructure. Although the N1000, N1100, and N1200 NT600 obtained after three different solution temperatures did not show significant changes in grain size, the amount of precipitated phases within the matrix gradually increased with increasing solution temperature. This included both the dispersed α phase within the matrix and the TiN at the grain boundaries, as well as the chain-like α phase originating from the grain boundaries, all of which gradually increased in quantity. Figure 28 Images (d) to (f) further confirm the distribution and morphological evolution of the precipitates. As the solution temperature increases, the number of needle-like α-phase precipitates inside the matrix increases, and TiN and chain-like α-phases at the grain boundaries gradually appear. This indicates that the addition of nitrogen promotes the formation of precipitates, which evolve continuously with the increase of solution temperature. In N1000, the precipitates are mainly distributed inside the matrix, with small particles and relatively dispersed morphology. In N1100, the number of precipitates inside the matrix increases, and obvious precipitates are seen at the grain boundaries in N1200. The precipitates inside the matrix gradually connect into a chain-like distribution, indicating that the growth and aggregation effects of the precipitates are enhanced.

[0097] In summary, the microstructure of the undoped alloys (R series) remained basically stable at different temperatures without significant changes. In contrast, the nitrogen-containing alloys (N series) showed an increasing number of precipitates with rising solution temperature, which began to aggregate towards the grain boundaries. This indicates that nitrogen promotes the growth of precipitates and affects the microstructure of the alloy.

[0098] The N1200 was examined using transmission electron microscopy (TEM), and the TEM images were obtained, as shown below. Figure 29 As shown, (a) and (c) are bright-field images of TiN at the grain boundaries and the internal acicular α phase of N1200, respectively; (b) and (d) are the diffraction spot markings for regions 1 and 2, respectively. TiN and the acicular α phase were found in the transmission results of N1200. These two precipitates were not observed in the transmission results of the untreated NT600 sample. However, the two precipitates, TiN and the acicular α phase, which are only present in NT650, were observed in the N1200 sample after heat treatment at 1200℃. Furthermore, the size of the α phase in the N1200 sample was much smaller than that in NT650, reaching the nanometer scale. Figure 29 The selected electron diffraction pattern in (b) shows that TiN at the N1200 grain boundaries has a standard face-centered cubic structure, while Figure 29 In (d), the selection of the needle-like α phase and the matrix β phase in the electron diffraction pattern shows that the α phase... It is nearly parallel to the β phase. as well as The presence of these precipitates on the crystal planes leads to improved mechanical properties, and N1200 also has fewer dislocations and defects than NT600 and NT650.

[0099] EDS elemental distribution analysis was performed on N1200, and the EDS spectrum was obtained, as shown below. Figure 30 As shown, (a) is the EDS energy spectrum of TiN at the grain boundaries, and (b) is the EDS energy spectrum of the α-Ti phase inside the matrix. Although no TiN phase was observed at the grain boundaries in the NT600 sample, heat treatment promoted the precipitation of N elements enriched at the grain boundaries and Ti elements at the grain boundaries. The presence of TiN will enhance the mechanical properties of the material. As can be seen from the figure, the α phase of N1200 is also similar to that of NT650, with a small amount of N elements enriched inside. These nitrogen elements are dissolved inside, and the precipitated phase of N1200 is finer than that of N650, reaching a length of several hundred nanometers.

[0100] Effect of solution treatment on mechanical properties Figure 31 The graph shows the microhardness histograms for R1000, R1100, R1200, N1000, N1100, and N1200. From the graph, it can be seen that after solution heat treatment, the microhardness of Raw (N1000) initially decreases slightly from 179 HV to 172 HV for R1100, then increases again to 177 HV for R1200 as the solution temperature increases. Overall, the microhardness of Raw (N1000) does not change significantly after solution treatment, remaining basically around 175 HV. Furthermore, after solution treatment at 1100℃ and 1200℃, the microhardness becomes more uniform across the matrix. This is likely because after high-temperature solution treatment, the grain boundaries of Raw (N1000) are more tightly connected, and structural defects gradually decrease, resulting in a more uniform overall microstructure. In contrast, the microhardness of NT600 (N1000) decreases slightly with increasing solution temperature after solution treatment, starting at 311.7 HV. The HV increases to 369.4 HV for N1200. This change corresponds to the microstructure. Although the grain size does not change significantly from N1000 to N1200, N1200 has more precipitates than N1000 and N1100. Its second-phase strengthening effect is more significant, which further enhances the strengthening effect of the alloy. This results in N1200 having higher hardness than N1100, N1000 and NT600.

[0101] In summary, the effects of solution treatment on the hardness of R-series and N-series niobium-titanium alloys differ significantly. For the R-series, the main effect is to improve the uniformity of the microstructure, while for the N-series, precipitation and distribution at grain boundaries and within the matrix are the dominant factors in hardness variation. The higher the solution temperature, the more pronounced the precipitation strengthening effect.

[0102] The room temperature compression performance of R1000, R1100, R1200, N1000, N1100, and N1200 samples was tested, and the results are as follows: Figure 32As shown, Figure 32 (a) shows the stress-strain curves of R1000, R1100, R1200, N1000, N1100, and N1200 at room temperature. Figure 32 (b) is a histogram of yield strengths for R1000, R1100, R1200, N1000, N1100, and N1200. Figure 32 It can be seen that compared to the untreated sample (Raw, 443 MPa), the yield strength of R1000, R1100, and R1200 samples showed almost no change. This phenomenon is similar to the results of microhardness testing and is also close to the results of microstructure changes. However, for N1000, N1100, and N1200, compared to the untreated NT600 (1000 MPa), the yield strength continuously increased with increasing solution treatment temperature, from 1067 MPa for N1000 to 1096 MPa for N1200. This indicates that solution treatment effectively improved the strength of NT600. This strengthening trend is attributed to precipitation strengthening. Previous microstructural observations revealed fine and dispersed α-Ti precipitates in the nitrogen-containing samples, which formed a more stable and uniform distribution, especially after solution treatment. These precipitates act as dislocation pinning centers, effectively hindering dislocation slip and improving material strength.

[0103] The three stress-strain curves after NT600 solution treatment also show that the deformation strengthening process in the latter half follows the same trend as that in Raw, indicating that it not only improves strength but also has good plasticity similar to Raw. This is of great significance for materials that require both high strength and good plasticity in practical applications.

[0104] Effect of solution treatment on corrosion resistance Following the electrochemical polarization curve testing method described above, potentiodynamic polarization scanning tests were performed on R1000, R1100, R1200, N1000, N1100, and N1200 in a 3.5 wt.% NaCl solution. The obtained potentiodynamic polarization curves are shown below. Figure 33As shown in the figure, the cathode curves of the N-series and R-series niobium-titanium sintered alloys are relatively standard. All six groups of samples exhibit passivation and secondary passivation regions at the anode, demonstrating different levels of active dissolution. The corrosion current density of all six samples initially increases rapidly and then tends to increase slowly. This phenomenon is due to the formation of oxide films such as Nb₂O₅ and TiO₂ on the sample surface during electrochemical corrosion. Nb₂O₅ acts as the main protective oxide layer, improving corrosion resistance. For the N-series, TiN provides additional stability in localized areas. Furthermore, nitrogen solid solution optimizes the microstructure of the oxide film, improving its density and stability. These changes create an isolation effect between the sample surface and the corrosive liquid, significantly delaying the corrosion process. The passivation film gradually corrodes and ruptures as the corrosion potential increases, allowing the sample to come into contact with the corrosive liquid and undergo chemical corrosion. At this point, the corrosion current density increases sharply with the increase of the corrosion potential, and the chemical reaction between the sample surface and the corrosive liquid proceeds violently.

[0105] Table 11 Main parameters of the potentiodynamic polarization curves of samples with different solution treatments

[0106] Table 11 shows the electrochemical parameters obtained after fitting the polarization curves using Cview, combined with... Figure 33 As shown in Table 11, the corrosion resistance of samples N100, N1100, and N1200 is superior to that of sample NT600 (Ecorr: -0.5013 V, Icorr: 9.44278 × 10⁻⁶). -8 A·cm -2 The same situation is also reflected in the Raw data; R1100 and R1200 have significantly higher values ​​than Raw (Ecorr: -0.38458 V, Icorr: 4.3298 × 10⁻⁶ V). -7 A·cm -2 The corrosion current density was lower for R1000 than for R1200. Solution treatment improved the microstructure interface, and the nitrogen solution enhanced the surface protection of the samples. With increasing solution temperature, the corrosion potential of the three R1000 samples at different solution temperatures gradually increased, while the corrosion current density first decreased and then increased. The corrosion potentials of N1000, N1100, and N1200 also gradually increased, while the corrosion current density first increased and then decreased. Among the six groups of solution-treated samples, based on the fitting results of the corrosion current density, N1000 exhibited the best corrosion resistance, while R1000 showed the worst.

[0107] Figure 34Impedance spectral results of samples R1000, R1100, R1200, N1000, N1100, and N1200 in 3.5 wt.% NaCl solution are shown. (a) is the Nyquist plot, and (b) is a high-frequency magnified plot. The figures show that the impedance of the three nitrogen-undoped groups (R1000, R1100, and R1200) is relatively low, especially in the higher frequency range, where their Nyquist curves exhibit smaller arcs, indicating low polarization impedance and strong electron transport capability. The impedance of N1000, N1100, and N1200 increased significantly, and the Nyquist curve shifted towards the high impedance region with a larger arc. This indicates enhanced interfacial polarization after doping, which is likely due to the doping element altering the interfacial electronic structure, increasing the interfacial barrier, and thus affecting carrier transport behavior. At the same heat treatment temperatures (1000℃, 1100℃, 1200℃), the impedance of nitrogen-doped (N series) materials was much higher than that of undoped (R series) materials. This may be related to the increased defect concentration and grain boundary structure changes caused by doping. As the temperature increased (from 1000℃ to 1200℃), the impedance of both R and N groups showed a trend of first increasing and then slowing down. This may be related to grain growth and the evolution of the internal curves of the material. After trace nitrogen doping, Nb-47Ti showed a significant increase in impedance, indicating that nitrogen doping has a significant impact on the electrochemical behavior of the material.

[0108] Table 12 EIS fitting results for samples with different solution treatments

[0109] Table 12 shows the EIS fitting results for samples R1000, R1100, R1200, N1000, N1100, and N1200. From the table, we can see the R series... s The values ​​remain largely unchanged, indicating that the solution resistance of Raw material does not change significantly at different solution temperatures, and the overall conductivity of the material is relatively stable. However, the N-series R... s The values ​​varied drastically, especially for the N1200 sample, where R... s Up to 235.8 Ω·cm -2 This indicates that nitrogen doping caused significant changes after solution treatment at 1200℃, which corresponds to the metallographic structure. After solution treatment at 1200℃, the precipitated phase in the N1200 sample increased, leading to a decrease in overall conductivity. Charge transfer resistance R ct It is an important parameter affecting the kinetics of electrochemical reactions. A higher value usually means increased charge transport resistance, i.e., enhanced interfacial polarization. Undoped (R series) R ct The value gradually decreases with increasing solution temperature; however, the R value for (N series) ctThe different trends indicate that high-temperature treatment after doping may enhance the grain boundary barrier effect, and the dopant element may accumulate at the grain boundaries, forming a potential barrier. As the solution temperature increases, the enrichment at the grain boundaries decreases, therefore R ct It also gradually decreased.

[0110] Effect of solution treatment on superconducting properties The superconducting transition temperature and critical current density of R1000, R1100, R1200, N1000, N1100, and N1200 samples were tested using the MPMS comprehensive physical property testing system. Figure 35 In Figure (a), the magnetic susceptibility of samples R1000, R1100, R1200, N1000, N1100, and N1200 as a function of temperature (MT) is shown. Figure 35 (b) in the figure is a line graph showing the changes in the superconducting critical transition temperature for samples R1000, R1100, R1200, N1000, N1100, and N1200. Figure 35 As can be seen in (a), NT600 exhibits clear superconducting transition characteristics at different solution temperatures, indicating that all three samples are high-quality superconducting phases. The magnetic susceptibility of the samples decreases rapidly near the superconducting transition temperature, indicating that they are in a superconducting state. Figure 35 As shown in (b), with the increase of solution temperature, the critical superconducting transition temperature of the N series gradually increases, from 8.2K for N1000 to 8.7K for N1200, while the critical transition temperature of NT600 is only 7.6K. This indicates that solution treatment can significantly improve the critical transition temperature of nitrogen-doped NT600. This shows that high-temperature solution treatment has a positive effect on optimizing the superconducting properties of nitrogen-doped niobium-titanium alloys. The reason for this Tc increase is related to the effect of solution treatment on the microstructure of the sample. First, there is the effect of precipitated phases. The microstructure observation above shows that the α-Ti phase in unsolidified NT600 is unevenly distributed and has agglomeration. With the increase of solution temperature, the α-Ti phase distribution in the NT600 sample becomes more uniform, reducing the inhomogeneity of the microstructure and providing good conditions for the continuous formation of superconducting phases. In addition, high-temperature solution treatment promotes the homogenization of nitrogen-doped niobium-titanium alloys and reduces the adverse effects of non-superconducting impurities and material defects, thereby further improving the Tc value.

[0111] Overall, solution treatment has a positive effect on the superconducting transition temperature of NT600, optimizing the size and distribution of its internal α phase and increasing the superconducting critical transition temperature. In particular, the N1200 sample exhibits the highest Tc value (8.7 K), indicating that a higher solution temperature is beneficial for increasing the superconducting transition temperature of NT600.

[0112] Figure 36(a) shows the hysteresis loops of NT600 alloy at different solution treatment temperatures under magnetic field conditions of 4.2K and -6T to 6T, as a function of magnetic field strength. B For superconductors, the hysteresis loop initially reaches its maximum at zero magnetic field, and then gradually decreases in both magnetic susceptibility and hysteresis loop width as the magnetic field strength increases. Figure 36 In (a), we can see that NT100, NT1100, and NT1200 also possess the characteristic of hysteresis loop changing with magnetic field strength, similar to NT600, which is typical of standard superconductors. However, they differ from... Figure 25 In (a) the hysteresis loop of NT600 as a function of magnetic field strength, the three N-series samples after solution treatment did not exhibit a second peak effect similar to that of the NT600 sample. The second peak effect typically signifies the arrival of the upper critical magnetic field for low-temperature metals or alloy superconductors like niobium-titanium superconducting alloys. The disappearance of the second peak indicates that the overall flux pinning ability of the sample tends to stabilize. This is because high-temperature solution treatment leads to a more uniform size and distribution of the α-Ti pinning phase, making the flux pinning effect more even across the entire magnetic field range and avoiding localized "re-pinning" behavior. This indicates that the upper critical magnetic field of the NT600 sample was enhanced after solution treatment.

[0113] Figure 36 (b) in the figure is calculated based on the Bean model. J c The curves show the critical current density changing with magnetic field strength at different solution temperatures, exhibiting a consistent decreasing trend. J c The value decreases with increasing magnetic field, consistent with typical superconducting behavior of low-temperature alloys. Among them, the N1200 sample maintains a relatively high value under low magnetic fields (0~2T). J c This indicates that its optimized structure has a higher current-carrying capacity.

[0114] comprehensive T c and J c The changing trends indicate that solution treatment synergistically enhances the superconducting performance of NT600 from multiple perspectives by controlling lattice distortion, optimizing α-Ti precipitation, reducing TiN discontinuities, and decreasing dislocation density. Particularly in the N1200 sample, the stability, homogeneity, and grain boundary continuity of the microstructure are optimal, exhibiting the highest performance. T c (8.7K) and good J c performance.

[0115] Although the above embodiments have provided a detailed description of the present invention, they are only some embodiments of the present invention, and not all embodiments. People can obtain other embodiments based on these embodiments without creative effort, and these embodiments all fall within the protection scope of the present invention.

Claims

1. A nitrogen-doped niobium-titanium alloy, characterized in that, The components include the following components by mass percentage: Nitrogen 0.08~0.19 wt%; Niobium 52~54wt% Titanium 45.81~47.92 wt%; The nitrogen-doped niobium-titanium alloy contains β phase, α phase and nitride precipitates.

2. The nitrogen-doped niobium-titanium alloy according to claim 1, characterized in that, The nitrogen-doped niobium-titanium alloy has a β phase volume percentage of ≥99% and a total volume percentage of α phase and nitride precipitates of ≤1%.

3. The method for preparing the nitrogen-doped niobium-titanium alloy according to claim 1 or 2, characterized in that, Includes the following steps: Niobium-titanium alloy powder is subjected to nitriding treatment to obtain nitrided niobium-titanium alloy powder; the nitriding treatment temperature is 550~650℃. The nitrogen-doped niobium titanium alloy powder was sintered to obtain a nitrided niobium titanium alloy. The nitrogen-doped niobium-titanium alloy was obtained by solution treatment.

4. The preparation method according to claim 3, characterized in that, The holding time for the nitriding treatment is 0.5~2 hours, and the nitrogen-containing gas used in the nitriding process is ammonia, with a flow rate of 0.8~1 m³ / h. 3 / h.

5. The preparation method according to claim 3, characterized in that, The particle size of the niobium-titanium alloy powder is 50~100μm.

6. The preparation method according to any one of claims 3 to 5, characterized in that, The nitriding process is carried out in a nitriding furnace, and the atmosphere in the nitriding furnace is stirred and mixed using a stirring fan with a rotation speed of 140~160 r / min.

7. The preparation method according to claim 3, characterized in that, The sintering includes discharge plasma sintering; The vacuum degree of the discharge plasma sintering is ≤1Pa, the heating rate is 100~200℃ / min, the temperature is 1000~1200℃, the holding time is 15~30min, and the axial pressure is 40~50MPa.

8. The preparation method according to claim 3 or 7, characterized in that, The process before sintering also includes: placing the nitrogen-doped niobium titanium alloy powder in a graphite mold for pre-pressing, wherein the nitrogen-doped niobium titanium alloy powder and the graphite mold are separated by graphite paper, and the axial pressure of the pre-pressing is 45~55MPa.

9. The preparation method according to claim 3, characterized in that, The vacuum degree of the solution treatment is ≤1 Pa, the heating rate is 150~300℃ / min, the holding temperature is 1000~1200℃, and the holding time is 30~60min. After the holding time is completed, the temperature is lowered under a protective atmosphere, which includes argon or nitrogen, the flow rate of the protective atmosphere is 45~55mL / min, and the temperature after cooling is below 80℃.

10. The application of the nitrogen-doped niobium-titanium alloy according to claim 1 or 2, or the nitrogen-doped niobium-titanium alloy prepared by the preparation method according to any one of claims 3 to 9, as a superconducting material.

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