High-strength corrosion-resistant nickel-based alloy pipe and preparation method thereof

Abstract

The application belongs to the field of nickel-based alloy pipe preparation, and provides a high-strength corrosion-resistant nickel-based alloy pipe and a preparation method thereof. The application controls Ni-Cr-Mo-Nb components, adopts two-stage short-time heat treatment and final core diameter cold drawing coupling design, regulates grain size, total area fraction of delta phase and Laves phase, subgrain band and dislocation cell distribution, forms subgrain-dislocation strengthening structure, realizes coordinated improvement of strength, plasticity, residual tensile stress and corrosion-resistant structure state, solves the problems that cold work strengthening easily weakens corrosion resistance and high alloying leads to compression of processing window, and is suitable for corrosion-resistant heat exchange pipes, oil and gas pipes and chemical equipment pipes.

Description

Technical Field

[0001] This invention relates to the field of nickel-based alloy pipe manufacturing, specifically to a high-strength corrosion-resistant nickel-based alloy pipe and its manufacturing method. Background Technology

[0002] Nickel-based alloy seamless pipes are widely used in chemical heat exchange, oil and gas extraction, marine engineering, nuclear power equipment, and high-temperature corrosive media transportation. In chlorine-containing, sulfur-containing, acidic, or high-temperature complex environments, the pipes not only need to rely on the nickel-based matrix and corrosion-resistant elements such as chromium and molybdenum to maintain stable passivation capabilities, but also need to maintain high strength, sufficient plasticity, and low residual tensile stress on the surface during pressure bearing, bending, flaring, flattening, and long-term service. For small-diameter thin-walled pipes, large-diameter thick-walled pipes, and long-length seamless pipes, the uniformity of chemical composition, grain size, grain boundary precipitates, cold work hardening state, and surface treatment state all affect service reliability. As equipment develops towards high parameters, long cycles, and complex media, how to improve strength and corrosion resistance without sacrificing processing adaptability has become a key technical requirement in the industrial production of nickel-based alloy pipes.

[0003] Existing nickel-based or nickel-iron-based corrosion-resistant alloy pipes typically achieve their target microstructure through high alloying, hot extrusion, cold rolling and cold drawing, annealing, or solution treatment. However, single strengthening or homogenization paths often have limitations. For example, Chinese Patent CN106180254B discloses a method for preparing seamless nickel-molybdenum-chromium alloy pipes resistant to molten salt corrosion. This method uses hot extrusion, cold rolling, intermediate heat treatment, and final heat treatment to prepare the pipes. However, this type of process focuses on pipe forming and molten salt corrosion resistance, lacking specific arrangements for controlling the subgrain-dislocation strengthening state introduced by cold drawing of the final core diameter, surface residual tensile stress, and the coupling between the low area fractions of the δ phase and the Laves phase. Similarly, disclosed technologies for improving the stress corrosion resistance of nickel-iron-based corrosion-resistant alloy oil well pipes improve pipe performance through cold work hardening and heat treatment. However, high deformation and annealing windows easily cause fluctuations in the distribution of residual stress, precipitated phases, and corrosion-resistant elements, making it difficult to simultaneously stabilize strength, plasticity, corrosion resistance, and processing window. Summary of the Invention

[0004] The purpose of this invention is to provide a high-strength, corrosion-resistant nickel-based alloy pipe and its preparation method, solving the current technical problems of difficulty in balancing cold working strengthening and corrosion resistance stability, and difficulty in balancing high alloying corrosion resistance design and machinability window.

[0005] This invention achieves a synergistic balance between strength, corrosion resistance, and processing adaptability by matching the compositional boundaries of the Ni-Cr-Mo-Nb system, two short-time heat treatments, and the final core diameter cold drawing. This ensures that the homogenization of corrosion-resistant elements and the suppression of harmful phases do not come at the cost of completely eliminating cold work strengthening, and that subgrain dislocation strengthening does not come at the cost of excessive residual tensile stress and continuous grain boundary precipitation.

[0006] To achieve the above objectives, the present invention provides the following technical solution: A high-strength, corrosion-resistant nickel-based alloy pipe, wherein the high-strength, corrosion-resistant nickel-based alloy pipe is a seamless Ni-Cr-Mo-Nb system pipe, and its chemical composition, based on the total mass of the high-strength, corrosion-resistant nickel-based alloy pipe, consists of the following elements and unavoidable impurities: chromium 20.00–23.00 wt%, molybdenum 8.00–10.00 wt%, niobium and tantalum combined 3.15–4.15 wt%, iron 0.50–5.00 wt%, and carbon 0.005 wt%. –0.080wt%, manganese 0–0.50wt%, silicon 0–0.50wt%, aluminum 0–0.40wt%, titanium 0–0.40wt%, phosphorus 0–0.015wt%, sulfur 0–0.015wt%, the sum of the contents of the above chromium, molybdenum, iron, carbon, manganese, silicon, aluminum, titanium, phosphorus, sulfur and the total contents of niobium and tantalum does not exceed 42.00wt%, the balance being nickel and unavoidable impurities, and the nickel content is not less than 58.00wt%; The high-strength corrosion-resistant nickel-based alloy tube has a subgrain-dislocation-strengthened structure formed by two short-time heat treatments and cold drawing of the final core diameter. The area reduction rate of the cold drawing of the final core diameter is 3.00–8.00%, the statistical mean of the grain size of the high-strength corrosion-resistant nickel-based alloy tube is 5.0–8.0, and the statistical mean of the total area fraction of the δ phase and Laves phase in the high-strength corrosion-resistant nickel-based alloy tube is not greater than 1.00%.

[0007] Furthermore, the surface residual tensile stress, tensile strength, specified plastic elongation strength Rp0.2, and elongation after fracture of the high-strength corrosion-resistant nickel-based alloy pipe are all calculated based on the statistical average of the test data. The surface residual tensile stress does not exceed 150 MPa, the tensile strength is 827–980 MPa, the specified plastic elongation strength Rp0.2 is 414–650 MPa, and the elongation after fracture is 30–55%.

[0008] Furthermore, the subgrain-dislocation-strengthened structure is formed through the following steps: A1. The furnace charge is prepared according to the aforementioned chemical composition, wherein the furnace charge includes at least nickel powder, chromium powder, molybdenum powder, niobium powder, iron powder and graphite powder for introducing carbon, and tantalum powder, manganese powder, silicon powder, aluminum flakes and titanium powder are selectively added according to the target composition; A2. Melting is carried out under the protection of high-purity argon or vacuum conditions at a melting temperature of 1450–1600℃ and a holding time of 10–60 min to obtain an alloy melt and cast it into an alloy ingot. A3. The alloy ingot is hot-worked to obtain a tube blank at a temperature of 1120–1200℃ and a total deformation of 20–80%. A4. The tube blank is subjected to multiple passes of cold rolling, cold drawing, or a combination of both, to achieve the final core diameter size before cold drawing. The tube blank before the final core diameter is subjected to two short-time heat treatments: the first stage is 880–930℃ for 10–30 min, and the second stage is 980–1020℃ for 8–20 min. A5. The tube blank after the two short-time heat treatments is subjected to final core diameter cold drawing to obtain the high-strength corrosion-resistant nickel-based alloy tube.

[0009] Furthermore, in the chemical composition of the high-strength corrosion-resistant nickel-based alloy pipe, the carbon content is 0.010–0.045 wt%, the phosphorus content is 0–0.010 wt%, the sulfur content is 0–0.006 wt%, and the total content of niobium and tantalum is 3.40–4.00 wt%.

[0010] Furthermore, after step A4 and before step A5, the molybdenum microsegregation coefficient in the tube blank is 1.00–1.20, the niobium microsegregation coefficient in the tube blank is 1.00–1.25, and the statistical mean of the total area fraction of the δ phase and the Laves phase in the tube blank is 0–0.80%.

[0011] Furthermore, the high-strength corrosion-resistant nickel-based alloy tubing undergoes surface treatment after cold drawing of the final core diameter. The surface treatment includes treating with a 1.00–10.00 wt% hydrofluoric acid aqueous solution at 20–45°C for 1–8 min to remove surface oxide scale, treating with a 10.00–25.00 wt% nitric acid aqueous solution at 20–45°C for 1–10 min for passivation, rinsing with deionized water until the pH of the cleaning solution is 6.0–8.0, and drying at 40–120°C for 10–60 min.

[0012] Furthermore, the high-strength corrosion-resistant nickel-based alloy pipe has an outer diameter of 6.00–219.00 mm, a wall thickness of 0.50–20.00 mm, and an outer diameter greater than twice the wall thickness, and a length of 0.50–18.00 m; The high-strength corrosion-resistant nickel-based alloy tubing after cold drawing of the final core diameter has subgrain zones continuously distributed along the tube axis and dislocation cell structures discontinuously distributed. The δ phase and Laves phase in the high-strength, corrosion-resistant nickel-based alloy tubing are discontinuously distributed along the grain boundaries.

[0013] Furthermore, the high-strength corrosion-resistant nickel-based alloy pipe does not crack at the end after flaring with a flaring rate of 10–25%, and the flattened area does not crack after flattening with an outer diameter deformation rate of 30–70%.

[0014] As a concept of this invention, the present invention employs a Ni-Cr-Mo-Nb system seamless tube composition and a design combining two-stage short-time heat treatment with cold drawing of the final core diameter. This is primarily used to achieve a synergistic balance between the mechanical strength gained from cold work strengthening and the corrosion resistance stability under the influence of grain boundary precipitation and residual tensile stress. In existing technologies, simply increasing the amount of cold work typically increases dislocation density and strength, but it easily increases surface residual tensile stress and promotes local microstructure inhomogeneity, thereby weakening stability in corrosive environments. While simply strengthening high-temperature homogenization or long-time annealing is beneficial in reducing the influence of segregation and harmful phases, it may weaken the contribution of cold work strengthening and reduce strength. This invention adjusts the segregation and precipitation state through boundary control of chromium, molybdenum, niobium, and tantalum contents, two-stage short-time heat treatment, and then uses finite final core diameter cold drawing to form subgrain zones and dislocation cells, allowing the strength source, grain boundary phase control, and residual stress control to mutually correct each other within the same microstructure window.

[0015] This invention also discloses a method for preparing high-strength, corrosion-resistant nickel-based alloy tubing, comprising the following steps: S1. Prepare the furnace charge according to the target alloy composition for obtaining high-strength, corrosion-resistant nickel-based alloy pipes, wherein the target alloy composition, based on the total mass of the resulting high-strength, corrosion-resistant nickel-based alloy pipes, is: chromium 20.00–23.00 wt%, molybdenum 8.00–10.00 wt%, niobium and tantalum (total content) 3.15–4.15 wt%, iron 0.50–5.00 wt%, carbon 0.005–0.080 wt%, manganese 0–0.50 wt%, silicon 0–0.50 wt%, and aluminum 0–0.40 wt%. The content of chromium, molybdenum, iron, carbon, manganese, silicon, aluminum, titanium, phosphorus, sulfur, niobium and tantalum shall not exceed 42.00 wt%, with the balance being nickel and unavoidable impurities, and the nickel content shall not be less than 58.00 wt%. The furnace charge includes nickel powder, chromium powder, molybdenum powder, niobium powder, iron powder and graphite powder for introducing carbon, and tantalum powder, manganese powder, silicon powder, aluminum flakes and titanium powder are selectively added according to the target composition. S2. Under the protection of high-purity argon or vacuum conditions, melt at 1450–1600℃ for 10–60 min, pour and cool to obtain alloy ingots; S3. The alloy ingot is hot-worked at 1120–1200℃, and the total deformation during hot working is 20–80% to obtain a rough tube or tube blank. S4. Perform multiple cold rolling, cold drawing, or a combination of both on the raw tube or tube blank, and perform intermediate annealing between at least two cold working passes to bring it to the size of the final core diameter before cold drawing. S5. Perform two short-time heat treatments on the tube blank before cold drawing of the final core diameter, wherein the first stage is held at 880–930℃ for 10–30 min, and the second stage is held at 980–1020℃ for 8–20 min. S6. The tube blank after the two short-time heat treatments is cold-drawn with a final core diameter of 3.00–8.00% to obtain the high-strength corrosion-resistant nickel-based alloy tube.

[0016] Furthermore, in step S2, the furnace charge is dried at 80–200℃ for 1–6 hours before smelting; after smelting, the alloy ingot is kept at 1120–1180℃ for 1–6 hours for homogenization treatment. In step S3, the hot working includes one or more of hot piercing, hot extrusion, or hot rolling.

[0017] Furthermore, in step S4, the rough tube or tube blank is subjected to multiple passes of cold rolling, cold drawing, or a combination of both, with a single pass area reduction rate of 5.00–25.00%, and intermediate annealing is performed between at least two cold processing passes at a temperature of 1040–1160°C and a holding time of 1–20 min. In step S5, the heating rate from the first stage to the second stage is 5–25℃ / min; In step S6, the final core diameter is cold-drawn using a core rod or mandrel to support the inner diameter. The cold-drawing speed is 1.00–12.00 m / min. The lubricant residue is cleaned with deionized water and then dried at 40–120℃ for 10–60 min.

[0018] Furthermore, the area reduction rate of the final core diameter cold drawing is calculated based on the cross-sectional area of ​​the tube metal before and after the final core diameter cold drawing. The cross-sectional area of ​​the tube metal is obtained by converting the outer diameter and wall thickness. The area reduction rate is the percentage obtained by dividing the difference between the cross-sectional area of ​​the metal before and after the final core diameter cold drawing by the cross-sectional area of ​​the metal before the final core diameter cold drawing. The obtained area reduction rate is used as the process release parameter for the final core diameter cold drawing step.

[0019] Furthermore, the rough tube or billet after hot processing and before multiple cold rolling, cold drawing, or a combination of both are used as intermediate microstructure inspection objects. The grain size is determined according to GB / T6394-2017 or ASTME112-25. The statistical mean of the grain size level of the rough tube or billet is 3.0–7.0. The measurement results are used as intermediate microstructure data for subsequent cold processing steps.

[0020] Furthermore, in the process of multi-pass cold rolling, cold drawing, or a combination of both, the rough tube or billet obtained in the previous pass is used as the input object for each pass. The area reduction rate of a single pass is calculated based on the cross-sectional area of ​​the tube metal before and after the cold working of that pass. The area reduction rate of a single pass is 5.00–25.00%. After completing one cold working pass, the rough tube or billet enters the next cold working pass or intermediate annealing step.

[0021] Furthermore, the samples for grain size determination were taken from the cross-section or longitudinal section of the tube after cold drawing of the final core diameter. After grinding, polishing and etching, the samples were graded according to GB / T6394-2017 or ASTME112-25, and the grain size level was recorded. The statistical mean was used as the quality control data of the microstructure after cold drawing of the final core diameter.

[0022] Furthermore, the samples for determining the total area fraction of the δ phase and the Laves phase were taken from tube blanks after two short-term heat treatments or tubes after cold drawing of the final core diameter. After cross-sectional sample preparation, grinding and polishing, the samples were subjected to backscattered electron imaging, and the niobium-molybdenum-rich precipitate region was determined by energy dispersive spectroscopy point analysis or surface scanning. The total area of ​​the δ phase and the Laves phase in each field of view was recorded as the total area of ​​the field of view. The percentage of the total area to the total area of ​​the field of view was taken as the total area fraction of the δ phase and the Laves phase, and the statistical mean of no less than 10 fields of view was used as the determination data of the total area fraction of the δ phase and the Laves phase.

[0023] Furthermore, the molybdenum and niobium microsegregation coefficients were measured on the cross-section of the billet after two short-term heat treatments and before the final core diameter was cold-drawn. Local mass fraction data of molybdenum and niobium were collected using electron probe microanalysis or scanning electron microscopy energy dispersive spectroscopy. The molybdenum microsegregation coefficient is the ratio of the highest mass fraction of molybdenum to the average mass fraction in the detection area, and the niobium microsegregation coefficient is the ratio of the highest mass fraction of niobium to the average mass fraction in the detection area. The obtained coefficients are used to determine the elemental homogenization state before the final core diameter was cold-drawn.

[0024] Furthermore, the characterization samples of the subgrain-dislocation-strengthened structure were taken from the longitudinal section of the tube after cold drawing of the final core diameter. Electron backscatter diffraction was used to record the orientation difference distribution and small-angle grain boundary distribution along the tube axis, and transmission electron microscopy was used to record the cell wall and intracellular region of dislocation cells. The electron backscatter diffraction data and transmission electron microscopy images were used together as the tissue determination data for the continuous distribution of subgrain zones and the discontinuous distribution of dislocation cells.

[0025] Furthermore, the residual tensile stress on the surface was measured on the inner and outer surfaces of the tube after the final core diameter was cold-drawn. X-ray diffraction was used to collect axial and circumferential residual stress data, and the measurement point location, test direction and residual stress value were recorded. The statistical mean of the measurement point data on the same surface was used as the determination data for the residual tensile stress on the surface.

[0026] Furthermore, for the surface-treated high-strength corrosion-resistant nickel-based alloy pipes, the inner and outer wall surfaces are used as measurement objects. The surface roughness Ra is measured along the pipe axis, and the inner wall Ra and outer wall Ra are recorded. Pipes with both inner wall Ra and outer wall Ra of 0.20–1.60 μm enter the dimensional and appearance inspection steps.

[0027] Furthermore, the wall thickness deviation is expressed as a percentage obtained by dividing the difference between the maximum and minimum wall thicknesses at the circumferential measuring points of the same pipe section by the average wall thickness of that section. The wall thickness deviation of the high-strength corrosion-resistant nickel-based alloy pipe after cold drawing of the final core diameter is 0–10.00%, and the obtained wall thickness deviation is used to determine the dimensional stability after cold drawing of the final core diameter.

[0028] Furthermore, the flaring rate is expressed as a percentage obtained by dividing the difference between the outer diameter of the flared end and the original outer diameter before flaring by the original outer diameter before flaring. The flaring rate for flaring is 10–25%. The outer diameter deformation rate for flattening is expressed as a percentage obtained by dividing the difference between the original outer diameter before flattening and the height in the loading direction after flattening by the original outer diameter before flattening. The outer diameter deformation rate for flattening is 30–70%. The flared end and the flattened area are used as crack inspection objects. The inspection output is a crack state record and enters the processing adaptability evaluation.

[0029] As another aspect of this invention, it employs a continuous process of target alloy composition formulation, protective atmosphere or vacuum melting, hot working, multi-pass cold working, two-stage short-time heat treatment, and final core diameter cold drawing. This process is primarily used to achieve, fix, or amplify the aforementioned synergistic effects. Existing preparation methods, if relying mainly on large cold working deformation, are prone to increasing residual stress and the risk of cracking during subsequent processing due to strength enhancement. If relying mainly on sufficient annealing or high-temperature long-time treatment, it may reduce the contribution of subgrain dislocation strengthening and limit dimensional control and production cycle time. This invention, through continuous control of furnace charge composition, melting homogenization, hot working temperature, single-pass area reduction rate, intermediate annealing, two-stage short-time heat treatment heating process, and final core diameter cold drawing speed, gradually fixes element homogenization, harmful phase suppression, tube dimensional stability, and cold work strengthening state during the preparation process, thereby ensuring that product performance does not depend on a single extreme process.

[0030] The compositional matching of chromium, molybdenum, niobium, and tantalum primarily addresses the stability of the corrosion-resistant matrix and the control of grain boundary precipitates. Two-stage short-time heat treatment and final core diameter cold drawing mainly focus on subgrain-dislocation strengthening and maintaining processing adaptability. Increasing only the content of chromium, molybdenum, niobium, and tantalum enhances the reserves of corrosion-resistant elements but increases deformation resistance, increases the risk of δ-phase and Laves phase enrichment, and compresses the cold working window of the pipe. Relying solely on final core diameter cold drawing to improve strength increases subgrain and dislocation structures but may accumulate residual tensile stress, adversely affecting corrosion-sensitive areas. This invention regulates segregation and precipitation states through total component limitation and two-stage short-time heat treatment, and achieves limited strengthening through 3.00–8.00% final core diameter cold drawing. This allows the corrosion resistance contribution and cold working strengthening contribution of alloying elements to mutually correct each other, ultimately balancing the aforementioned mutually constraining properties of strength, corrosion resistance, and processing window.

[0031] Beneficial technical effects 1. By limiting the composition of the Ni-Cr-Mo-Nb system and controlling the total content of elements other than nickel and unavoidable impurities, this invention maintains the nickel-based corrosion-resistant matrix while avoiding excessive chromium, molybdenum, niobium and tantalum which would increase deformation resistance and enrich precipitates, thus creating a more stable matching relationship between the corrosion-resistant element reserves and the cold working window.

[0032] 2. By connecting the first stage of short-time treatment at 880–930℃ with the second stage of short-time treatment at 980–1020℃, this invention adjusts the segregation of molybdenum and niobium and the area fraction of δ phase and Laves phase before the final core diameter cold drawing, reducing the risk of continuous precipitation and local enrichment at grain boundaries, and providing a more uniform microstructure basis for the subsequent formation of subgrain-dislocation strengthening microstructure.

[0033] 3. By cold-drawing the final core diameter with an area reduction rate of 3.00–8.00% and using a mandrel or core rod to support the inner diameter, the present invention can form a continuous subcrystalline zone and a discontinuous dislocation cell structure in the axial direction of the tube, while controlling the residual tensile stress on the surface, and avoiding the decrease in corrosion resistance and formability caused by simple high deformation strengthening.

[0034] 4. By removing oxide scale with hydrofluoric acid aqueous solution, passivating with nitric acid aqueous solution, cleaning with deionized water, and controlled drying, this invention enables the inner and outer wall surfaces after the final core diameter to enter a state of controllable roughness and appearance inspection, reducing the interference of lubricant residue, oxide scale, and surface defects on subsequent corrosion evaluation and processing adaptability.

[0035] 5. By conducting a series of tests on grain size, total area fraction of δ phase and Laves phase, microsegregation coefficient, residual stress, flaring rate and flattening deformation rate, this invention establishes a quality control chain from composition, microstructure, stress to processing adaptability, which helps to stably obtain seamless pipes suitable for corrosive media and pressure environments. Attached Figure Description

[0036] Figure 1 This is a graph showing the relationship between molybdenum content and tensile strength and critical pitting temperature in this invention.

[0037] Figure 2 This is a graph showing the relationship between the total niobium-tantalum content and the area fraction of precipitated phase in this invention.

[0038] Figure 3 This is a graph showing the relationship between the reduction rate of the final core diameter cold-drawn area and the effect on tensile strength and surface residual tensile stress in this invention.

[0039] Figure 4 This is a graph showing the relationship between the heat treatment temperature in the first stage of this invention and the precipitated phase area fraction and critical pitting temperature.

[0040] Figure 5 The following is a statistical diagram of the proportion of small-angle grain boundaries in Example 1, Comparative Example 9, and Comparative Example 11.

[0041] Figure 6 The TEM dislocation cell statistics are for Example 1 and Comparative Example 10.

[0042] Figure 7 The total area fraction of the δ phase and Laves phase is shown in the statistical diagrams for Examples 1, 3, and 7.

[0043] Figure 8 The above are statistical graphs of surface residual tensile stress and maximum crack length for Examples 1, 5, and 8.

[0044] Figure 9 The diagram shows the closed-loop comparison of strength and corrosion resistance for Example 1 and a representative comparative example.

[0045] Figure 10 The images show a comparison of the SEM morphology of the final product of Example 1 and the final product of Comparative Example 11; wherein, Figure 10 Figure (a) in the figure is a low-magnification SEM image of the final product of Example 1; Figure 10 Figure (b) in the figure is a low-magnification SEM image of the final product of Comparative Example 11; Figure 10 Figure (c) in the figure is a medium-magnification SEM image of the final product of Example 1; Figure 10 Figure (d) in the figure is a medium-magnification SEM image of the final product of Comparative Example 11.

[0046] Figure 11 The images show a comparison of the TEM structures of the final product of Example 1 and the final product of Comparative Example 11; where, Figure 11 (a) in the image is a bright-field TEM image of the final product of Example 1; Figure 11 (b) in the image is a bright-field TEM image of the final product of Comparative Example 11; Figure 11 (c) in the image is an HRTEM image of the final product of Example 1; Figure 11 (d) in the image is the HRTEM image of the final product of Comparative Example 11. Detailed Implementation

[0047] To make the objectives, technical solutions, and advantages of the embodiments of the present invention clearer, the technical solutions of the embodiments of the present invention will be clearly and completely described below with reference to the accompanying drawings.

[0048] Example 1 S1: Preparation scale and alloy composition With the aim of producing seamless Ni-Cr-Mo-Nb tubing with an outer diameter of 6.00 mm, a wall thickness of 0.50 mm, and a length of 0.50 m, a furnace charge of 10.000 kg was prepared. Based on the total mass of the finished product, the target chemical composition in this embodiment is as follows: chromium 20.00 wt%, molybdenum 8.00 wt%, niobium and tantalum (total content) 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, and elements excluding nickel and unavoidable impurities totaling 31.655 wt%, with a nickel content of 68.345 wt%. The furnace charge uses commercially available nickel powder, chromium powder, molybdenum powder, niobium powder, iron powder, and graphite powder. The purity of the metal powders is not less than 99.5%, and the purity of the graphite powder is not less than 99.0%. The furnace charge is dried at 80°C for 1 hour before use.

[0049] S2: Melting and Alloy Ingot Processing Nickel and iron powders were first added to the dried furnace charge, followed by chromium, molybdenum, and niobium powders, and finally graphite powder. The furnace was then evacuated to below 5 Pa in a vacuum induction melting furnace and filled with high-purity argon gas (purity not less than 99.99%). The melting temperature was controlled at 1450℃ and held for 10 minutes. Electromagnetic stirring was used for 3 minutes to promote compositional homogenization. The melt was then poured into alloy ingots using a metal mold preheated to 200℃. After cooling to room temperature (25℃), the alloy ingots were held at 1120℃ for 1 hour for homogenization, and then air-cooled to room temperature.

[0050] S3: Hot working, cold working and intermediate annealing The alloy ingot was hot-pierced and hot-rolled at 1120℃ with a total deformation of 20% to obtain a rough tube. After grinding the inner and outer surfaces, the rough tube entered cold working. The area reduction rate of the first cold rolling pass was 5.00%, and the area reduction rate of the second cold drawing pass was 6.00%. Intermediate annealing was carried out at 1040℃ for 1 min between the two cold working passes. The area reduction rate of each pass was calculated based on the cross-sectional area of ​​the metal before and after cold working. After cold working, the outer surface of the rough tube was cleaned with deionized water and dried at 40℃ for 10 min.

[0051] S4: Two-stage short-time heat treatment and final core diameter cold drawing Before cold drawing of the final mandrel diameter, the tube blank underwent two short-term heat treatments under high-purity argon protection. The first stage involved heating to 880℃ and holding for 10 minutes, followed by heating to 980℃ at a rate of 5℃ / min and holding for 8 minutes. The blank was then cooled to room temperature (25℃) under argon protection. After cooling, the tube blank was cold-drawn to the final mandrel diameter using a mandrel-supported inner diameter at a drawing speed of 1.00 m / min, resulting in a 3.00% reduction in the final mandrel diameter area. After cold drawing, lubricant residue was removed using deionized water, and the tube was dried at 40℃ for 10 minutes.

[0052] S5: Surface Treatment After cold drawing of the final core diameter, the tube was treated with a 1.00 wt% hydrofluoric acid aqueous solution at 20°C for 1 min to remove the surface oxide scale. It was then passivated by treating with a 10.00 wt% nitric acid aqueous solution at 20°C for 1 min, followed by rinsing with deionized water until the pH of the cleaning solution reached 6.0, and drying at 40°C for 10 min. After surface treatment, the inner wall Ra was measured to be 0.20 μm, the outer wall Ra was 0.22 μm, and the wall thickness deviation was 0%.

[0053] S6: Quality Inspection Methods and Results The cross-section and longitudinal section of the pipe in this embodiment were ground, polished, and etched, and then the grain size was rated. The statistical mean of the grain size level was 5.0. Backscattered electron imaging combined with energy dispersive spectroscopy (EDS) was used to statistically analyze the δ-phase and Laves phase regions. The statistical mean of the total area fraction of the δ-phase and Laves phase was 0.02%. Electron probe microanalysis showed that the microsegregation coefficients of molybdenum and niobium in the billet before cold drawing of the final core diameter were 1.00 and 1.00, respectively. The tensile strength was measured to be 827 MPa using a room temperature tensile test, the specified plastic elongation strength Rp0.2 was 414 MPa, and the elongation after fracture was 55%. The statistical mean of the residual tensile stress on the outer wall axial surface was 0 MPa, as measured by X-ray diffraction. No cracking occurred at the end after flaring with a 10% flaring rate, and no cracking occurred in the flattened area after flattening with a 30% outer diameter deformation rate.

[0054] S7: Features and application scenarios of this embodiment This embodiment adopts a relatively conservative low-ratio scheme and mild cold working conditions, with a low total component content, a small degree of cold drawing deformation of the final core diameter, and a microstructure that leans towards the low-strength, high-plasticity region. It is suitable for corrosion-resistant pipes with high requirements for small diameter, thin walls, low residual tensile stress, and subsequent assembly and forming.

[0055] Example 2 Raw materials and proportions This embodiment prepares high-strength, corrosion-resistant nickel-based alloy tubing with an outer diameter of 219.00 mm, a wall thickness of 20.00 mm, and a length of 18.00 m. Based on the total mass of the finished product, the target chemical composition of this embodiment is: chromium 23.00 wt%, molybdenum 10.00 wt%, niobium and tantalum (total content) 4.00 wt%, iron 3.60 wt%, carbon 0.045 wt%, manganese 0.20 wt%, silicon 0.20 wt%, aluminum 0.20 wt%, titanium 0.20 wt%, phosphorus 0.010 wt%, sulfur 0.006 wt%, and elements excluding nickel and unavoidable impurities totaling 41.461 wt%, with a nickel content of 58.539 wt%. The furnace charge includes commercially available nickel powder, chromium powder, molybdenum powder, niobium powder, tantalum powder, iron powder, graphite powder, manganese powder, silicon powder, aluminum flakes, and titanium powder. The purity of all metal raw materials is not less than 99.5%. Phosphorus and sulfur are controlled as raw material impurities in the finished product analysis.

[0056] Preparation process After drying the furnace charge at 200℃ for 6 hours, it was fed into a vacuum induction melting furnace. Nickel powder, iron powder, and chromium powder were added first. After a molten pool was formed, molybdenum powder, niobium powder, and tantalum powder were added, followed by manganese powder, silicon powder, aluminum flakes, titanium powder, and graphite powder. The melting process was carried out under vacuum conditions at a melting temperature of 1600℃ for 60 minutes. The resulting ingot was poured and cooled to obtain an alloy ingot. The alloy ingot was homogenized by holding it at 1180℃ for 6 hours. After homogenization, it was heated to 1200℃ for hot extrusion and hot rolling, with a total deformation of 80%, to obtain a tube billet.

[0057] Post-processing The tube blank undergoes a multi-pass combination of cold rolling and cold drawing to achieve the dimensions required for the final mandrel cold drawing. The cold working process includes four passes with single-pass area reduction rates of 25.00%, 20.00%, 15.00%, and 10.00%, respectively. At least two cold working passes are separated by intermediate annealing at 1160℃ for 20 minutes. Before the final mandrel cold drawing, the tube blank undergoes two short-duration heat treatments: the first at 930℃ for 30 minutes, and the second at 1020℃ for 20 minutes, with a heating rate of 25℃ / min from the first to the second stage. Following these two short-duration heat treatments, the tube blank is then cold-drawn with the inner diameter supported by a mandrel at a drawing speed of 12.00 m / min, resulting in an area reduction rate of 8.00%.

[0058] Surface treatment and quality inspection After cold drawing of the final core diameter, the surface oxide scale was removed by treatment with a 10.00 wt% hydrofluoric acid aqueous solution at 45℃ for 8 min, followed by passivation with a 25.00 wt% nitric acid aqueous solution at 45℃ for 10 min. The core was then rinsed with deionized water until the pH of the cleaning solution reached 8.0, and dried at 120℃ for 60 min. Test results showed that the statistical mean of the grain size grade after cold drawing of the final core diameter was 8.0, the statistical mean of the total area fraction of δ phase and Laves phase was 0.98%, the microsegregation coefficient of molybdenum before cold drawing was 1.20, the microsegregation coefficient of niobium was 1.25, the inner wall Ra was 1.60 μm, the outer wall Ra was 1.55 μm, and the wall thickness deviation was 10.00%. The tensile strength measured at room temperature was 980 MPa, the specified plastic elongation strength Rp0.2 was 650 MPa, the elongation after fracture was 30%, and the statistical mean of the residual tensile stress on the outer wall axial surface was 150 MPa. After flaring with a flaring rate of 25%, the end does not crack; after flattening with an outer diameter deformation rate of 70%, the flattened area does not crack.

[0059] Features of the solution in this embodiment This embodiment uses a combination of relatively high corrosion-resistant element content and high cold working strength, which is suitable for demonstrating the feasibility of thick-walled, large-diameter and long-length pipes under strong processing conditions. It is applicable to chemical or energy equipment pipes with high requirements for pressure resistance, corrosion resistance stability and dimensional specification coverage.

[0060] Example 3 Key parameters and sample condition This embodiment prepares seamless Ni-Cr-Mo-Nb tubing with an outer diameter of 60.00 mm, a wall thickness of 15.00 mm, and a length of 6.00 m. The target chemical composition, based on the total mass of the finished product, is as follows: chromium 20.00 wt%, molybdenum 8.00 wt%, niobium and tantalum (total content) 4.15 wt%, iron 5.00 wt%, carbon 0.080 wt%, manganese 0.50 wt%, silicon 0.50 wt%, aluminum 0.40 wt%, titanium 0.40 wt%, phosphorus 0.015 wt%, sulfur 0.015 wt%, and elements excluding nickel and unavoidable impurities totaling 39.060 wt%, with a nickel content of 60.940 wt%. The furnace charge uses commercially available nickel powder, chromium powder, molybdenum powder, niobium powder, tantalum powder, iron powder, graphite powder, manganese powder, silicon powder, aluminum flakes, and titanium powder. The purity of all metal raw materials is not less than 99.5%, and the purity of graphite powder is not less than 99.0%. The furnace charge is dried at 120°C for 3 hours before use.

[0061] Main process The dried furnace charge was melted under the protection of high-purity argon gas (purity not less than 99.99%) at a melting temperature of 1520℃ for 35 minutes. After being cast into alloy ingots, the ingots were homogenized by holding them at 1150℃ for 3 hours, followed by hot piercing and hot extrusion at 1120℃, with a total hot deformation of 50%, yielding a rough tube. After cooling, the rough tube underwent a multi-pass combination of cold rolling and cold drawing, with single-pass area reduction rates of 5.00%, 25.00%, 12.00%, and 18.00% respectively. An intermediate annealing process was performed between the second and third passes, holding at 1040℃ for 1 minute.

[0062] Two short-time heat treatments and final machining Before cold drawing of the final mandrel, the tube blank underwent two short-term heat treatments under argon protection: the first stage was held at 880℃ for 30 minutes, and the second stage was held at 1020℃ for 8 minutes, with a heating rate of 5℃ / min from the first stage to the second stage. After heat treatment, the inner diameter was supported by a mandrel for cold drawing of the final mandrel at a speed of 6.00 m / min, resulting in a reduction in area of ​​5.50%. After cold drawing, lubricant residue was cleaned with deionized water, and the tube was then dried at 80℃ for 30 minutes.

[0063] Surface treatment and testing verification In this embodiment, a 5.00 wt% hydrofluoric acid aqueous solution was used to treat the surface oxide scale at 30°C for 5 min, followed by passivation with an 18.00 wt% nitric acid aqueous solution at 30°C for 5 min. The surface was then rinsed with deionized water until the pH of the cleaning solution reached 7.0, and dried at 80°C for 30 min. Metallographic analysis showed a statistical mean of grain size grade 6.2; backscattered electron imaging combined with energy dispersive spectroscopy (EDS) showed a statistical mean of 0.65% for the total area fraction of the δ-phase and Laves phase; the microsegregation coefficient of molybdenum before cold drawing was 1.10, and that of niobium was 1.18; EBSD showed a continuous distribution of subgrain zones along the pipe axis, while TEM showed a discontinuous distribution of dislocation cells. The tensile strength measured at room temperature was 908 MPa, the specified plastic elongation strength Rp0.2 was 560 MPa, and the elongation after fracture was 38%; X-ray diffraction showed a statistical mean of 88 MPa for the residual tensile stress on the inner wall axial surface. After flaring with a flaring rate of 18%, the end does not crack; after flattening with an outer diameter deformation rate of 50%, the flattened area does not crack.

[0064] The process features and application directions of this embodiment This embodiment combines a high iron content, a high value region for auxiliary elements, and a medium degree of cold drawing deformation in the final core diameter to demonstrate the processing adaptability and microstructure controllability under the participation of multiple elements. It is suitable for medium-sized corrosion-resistant pipes and applications that require flaring, flattening, and subsequent connection processing.

[0065] Example 4 I. Preparation Object and Composition Status This embodiment aims to produce high-strength, corrosion-resistant nickel-based alloy tubing with an outer diameter of 12.00 mm, a wall thickness of 3.00 mm, and a length of 12.00 m, with a production scale of 50.000 kg. The target chemical composition of the finished product is as follows: chromium 22.00 wt%, molybdenum 9.50 wt%, niobium and tantalum (total) 3.80 wt%, iron 5.00 wt%, carbon 0.040 wt%, manganese 0.50 wt%, silicon 0.50 wt%, aluminum 0.40 wt%, titanium 0.24 wt%, phosphorus 0.010 wt%, sulfur 0.010 wt%, and elements excluding nickel and unavoidable impurities totaling 42.000 wt%, with a nickel content of 58.000 wt%. The furnace charge uses commercially available nickel powder, chromium powder, molybdenum powder, niobium powder, tantalum powder, iron powder, graphite powder, manganese powder, silicon powder, aluminum flakes, and titanium powder. The purity of all metal raw materials is not less than 99.5%, and the purity of graphite powder is not less than 99.0%.

[0066] II. Melting and Hot Working Path After drying the furnace charge at 160℃ for 4 hours, it was vacuum melted at 1580℃ for 45 minutes. The melt was then poured into alloy ingots and cooled. The alloy ingots were homogenized by holding them at 1160℃ for 4 hours. The homogenized alloy ingots were then subjected to a combination of hot piercing, hot extrusion, and hot rolling at 1180℃, with a total hot deformation of 65%, yielding rough tubes. These rough tubes were used as intermediate microstructure inspection targets. The statistical mean of the cross-sectional grain size was 7.0, meeting the intermediate microstructure control requirements for entering the cold working stage.

[0067] III. Cold working, short-time heat treatment and final core diameter cold drawing The rough tube undergoes a multi-pass combination of cold rolling and cold drawing to reach the final core diameter size before cold drawing. The area reduction rates per pass are 10.00%, 15.00%, 20.00%, and 25.00%, respectively. An intermediate annealing process is performed between the first and second passes at 1080℃ for 10 minutes, and between the third and fourth passes at 1160℃ for 20 minutes. Before the final core diameter cold drawing, the tube blank undergoes two short-duration heat treatments: the first at 930℃ for 10 minutes, and the second at 980℃ for 20 minutes, with a heating rate of 25℃ / min from the first to the second stage. Subsequently, the inner diameter is supported by a mandrel for the final core diameter cold drawing at a speed of 10.00 m / min, resulting in an area reduction rate of 7.00%.

[0068] IV. Cleaning, passivation and drying After cold drawing of the final core diameter, lubricant residue was first cleaned with deionized water and dried at 120℃ for 60 min. Then, the surface oxide scale was removed by treating with an 8.00 wt% hydrofluoric acid aqueous solution at 40℃ for 6 min. Passivation was performed by treating with a 22.00 wt% nitric acid aqueous solution at 40℃ for 8 min. Afterwards, the core was cleaned with deionized water until the pH of the cleaning solution reached 7.5, and then dried at 100℃ for 45 min. After surface treatment, the inner wall Ra was 1.20 μm, the outer wall Ra was 1.30 μm, and the wall thickness deviation was 6.50%.

[0069] V. Quality Inspection Methods and Results Metallographic grading of the cross-section and longitudinal section of the tube after cold drawing of the final core diameter was performed, with a statistical mean of grain size grade of 7.8. Backscattered electron imaging and energy dispersive spectroscopy (EDS) were used to statistically analyze the total area fraction of the δ-phase and Laves phase, with a statistical mean of 0.80%. The microsegregation coefficient of molybdenum in the tube blank before cold drawing of the final core diameter was 1.18, and that of niobium was 1.22. The tensile strength measured at room temperature was 956 MPa, the specified plastic elongation strength Rp0.2 was 630 MPa, and the elongation after fracture was 33%. Residual tensile stress on the inner and outer wall surfaces was tested using X-ray diffraction, with a statistical mean of 128 MPa for the axial values ​​on the outer wall. No cracking occurred at the end after flaring with a 22% flaring ratio, and no cracking occurred in the flattened area after flattening with a 65% outer diameter deformation rate.

[0070] VI. Applicable Scenarios of this Embodiment This embodiment employs a high total amount of non-nickel elements, a low nickel content, and a medium-to-high strength cold working scheme, which is suitable for demonstrating the feasibility of preparation under a wide range of compositional combinations. It is applicable to long-length, small-diameter pipes that require reserves of corrosion-resistant elements, dimensional stability, surface condition, and adaptability to subsequent processing.

[0071] Comparative Example 1: Basically the same as Example 1, except that the chromium content was adjusted to 18.50 wt%. The other key conditions were kept the same: molybdenum 8.00 wt%, niobium and tantalum combined 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel balance; melting temperature 1450℃, holding for 10 min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10 min, second stage holding at 980℃ for 8 min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treated with 1.00 wt% hydrofluoric acid aqueous solution at 20℃ for 1 min, treated with 10.00 wt% nitric acid aqueous solution at 20℃ for 1 min, washed to pH 6.0 and dried at 40℃ for 10 min.

[0072] Comparative Example 2: Basically the same as Example 1, except that the molybdenum content was adjusted to 7.00 wt%. The other key conditions were kept the same: chromium 20.00 wt%, niobium and tantalum combined 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel balance; melting temperature 1450℃, holding for 10 min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10 min, second stage holding at 980℃ for 8 min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treated with 1.00 wt% hydrofluoric acid aqueous solution at 20℃ for 1 min, treated with 10.00 wt% nitric acid aqueous solution at 20℃ for 1 min, washed to pH 6.0 and dried at 40℃ for 10 min.

[0073] Comparative Example 3: Basically the same as Example 1, except that the total content of niobium and tantalum was adjusted to 2.80 wt%. The other key conditions were kept the same: chromium 20.00 wt%, molybdenum 8.00 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, and nickel as balance; melting temperature 1450℃, holding for 10 min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10 min, second stage holding at 980℃ for 8 min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treated with 1.00 wt% hydrofluoric acid aqueous solution at 20℃ for 1 min, treated with 10.00 wt% nitric acid aqueous solution at 20℃ for 1 min, washed to pH 6.0 and dried at 40℃ for 10 min.

[0074] Comparative Example 4: Basically the same as Example 1, except that the reduction rate of the final core diameter cold-drawn area was adjusted to 1.50%. The other key conditions remained the same: chromium 20.00wt%, molybdenum 8.00wt%, total content of niobium and tantalum 3.15wt%, iron 0.50wt%, carbon 0.005wt%, manganese 0wt%, silicon 0wt%, aluminum 0wt%, titanium 0wt%, phosphorus 0wt%, sulfur 0wt%, nickel as balance; melting temperature 1450℃, holding for 10min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10min, second stage holding at 980℃ for 8min; cold drawing speed 1.00m / min; treatment with 1.00wt% hydrofluoric acid aqueous solution at 20℃ for 1min, treatment with 10.00wt% nitric acid aqueous solution at 20℃ for 1min, washing to pH 6.0 and drying at 40℃ for 10min.

[0075] Comparative Example 5: Basically the same as Example 1, except that the reduction rate of the final core diameter cold-drawn area was adjusted to 10.00%. The other key conditions remained the same: chromium 20.00wt%, molybdenum 8.00wt%, total content of niobium and tantalum 3.15wt%, iron 0.50wt%, carbon 0.005wt%, manganese 0wt%, silicon 0wt%, aluminum 0wt%, titanium 0wt%, phosphorus 0wt%, sulfur 0wt%, nickel as balance; melting temperature 1450℃, holding for 10min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10min, second stage holding at 980℃ for 8min; cold drawing speed 1.00m / min; treated with 1.00wt% hydrofluoric acid aqueous solution at 20℃ for 1min, treated with 10.00wt% nitric acid aqueous solution at 20℃ for 1min, washed to pH 6.0 and dried at 40℃ for 10min.

[0076] Comparative Example 6: Basically the same as Example 1, except that the temperature of the first short-term heat treatment was adjusted to 850°C. The other key conditions were kept the same: chromium 20.00 wt%, molybdenum 8.00 wt%, total content of niobium and tantalum 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel as balance; melting temperature 1450°C, holding for 10 min; hot working temperature 1120°C, total hot working deformation 20%; first stage holding for 10 min, second stage holding at 980°C for 8 min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treatment with 1.00 wt% hydrofluoric acid aqueous solution at 20°C for 1 min, treatment with 10.00 wt% nitric acid aqueous solution at 20°C for 1 min, washing to pH 6.0 and drying at 40°C for 10 min.

[0077] Comparative Example 7: Basically the same as Example 1, except that the temperature of the second short-term heat treatment was adjusted to 1050℃. The other key conditions were kept the same: chromium 20.00wt%, molybdenum 8.00wt%, total content of niobium and tantalum 3.15wt%, iron 0.50wt%, carbon 0.005wt%, manganese 0wt%, silicon 0wt%, aluminum 0wt%, titanium 0wt%, phosphorus 0wt%, sulfur 0wt%, nickel balance; melting temperature 1450℃, holding for 10min; hot working temperature 1120℃, total hot working deformation 20%; first stage holding at 880℃ for 10min, second stage holding for 8min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00m / min; treatment with 1.00wt% hydrofluoric acid aqueous solution at 20℃ for 1min, treatment with 10.00wt% nitric acid aqueous solution at 20℃ for 1min, washing to pH 6.0 and drying at 40℃ for 10min.

[0078] Comparative Example 8: Essentially the same as Example 1, except that the surface treatment only involved a 1.00 wt% hydrofluoric acid aqueous solution at 20°C for 1 min, followed by rinsing with deionized water until the pH reached 6.0. Nitric acid passivation was not used. Other key conditions were maintained as follows: chromium 20.00 wt%, molybdenum 8.00 wt%, total niobium and tantalum content 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel balance; melting temperature 1450°C, holding for 10 min; hot working temperature 1120°C, total hot working deformation 20%; first stage holding at 880°C for 10 min, second stage holding at 980°C for 8 min; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; drying at 40°C for 10 min after cleaning.

[0079] Comparative Example 9: Basically the same as Example 1, except that the second short-term heat treatment was omitted, and only a heat treatment of 880°C for 10 min was performed, followed by cooling and final core diameter cold drawing. The remaining key conditions were maintained as follows: chromium 20.00 wt%, molybdenum 8.00 wt%, total content of niobium and tantalum 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel as balance; melting temperature 1450°C, holding for 10 min; hot working temperature 1120°C, total hot working deformation 20%; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treatment with 1.00 wt% hydrofluoric acid aqueous solution at 20°C for 1 min, treatment with 10.00 wt% nitric acid aqueous solution at 20°C for 1 min, washing to pH 6.0 and drying at 40°C for 10 min. This comparative example is used to examine the synergistic relationship between two short-duration heat treatments and the final core diameter cold drawing.

[0080] Comparative Example 10: Basically the same as Example 1, except that the two short-term heat treatments were replaced with a long-term heat treatment at 1100℃ for 60 min, followed by final core diameter cold drawing. Other key conditions were maintained as follows: chromium 20.00 wt%, molybdenum 8.00 wt%, total content of niobium and tantalum 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel balance; melting temperature 1450℃, holding for 10 min; hot working temperature 1120℃, total hot working deformation 20%; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treatment with 1.00 wt% hydrofluoric acid aqueous solution at 20℃ for 1 min, treatment with 10.00 wt% nitric acid aqueous solution at 20℃ for 1 min, washing to pH 6.0 and drying at 40℃ for 10 min. This comparative example is used to examine the synergistic relationship between two short-duration heat treatments and the final core diameter cold drawing.

[0081] Comparative Example 11: Essentially the same as Example 1, except for changing the order of the two short-term heat treatments. First, the temperature was held at 980°C for 8 minutes, then furnace-cooled to 880°C and held for 10 minutes, followed by final core diameter cold drawing. Other key conditions remained the same: chromium 20.00 wt%, molybdenum 8.00 wt%, total niobium and tantalum content 3.15 wt%, iron 0.50 wt%, carbon 0.005 wt%, manganese 0 wt%, silicon 0 wt%, aluminum 0 wt%, titanium 0 wt%, phosphorus 0 wt%, sulfur 0 wt%, nickel balance; melting temperature 1450°C, holding for 10 minutes; hot working temperature 1120°C, total hot working deformation 20%; final core diameter cold drawing area reduction rate 3.00%, cold drawing speed 1.00 m / min; treatment with 1.00 wt% hydrofluoric acid aqueous solution at 20°C for 1 minute, 10.00 wt% nitric acid aqueous solution at 20°C for 1 minute, washing to pH 6.0, and drying at 40°C for 10 minutes. This comparative example is used to examine the synergistic relationship between the two short-time heat treatment sequences and the final core diameter cold drawing.

[0082] The chemical composition and total balance experiment used samples from Examples 1–4 and all comparative examples to verify whether chromium, molybdenum, niobium, tantalum, iron, carbon, and auxiliary elements met the target composition. Inductively coupled plasma atomic emission spectrometry (ICP-AES) or spark direct-reading spectroscopy was used to determine the metallic elements. Carbon and sulfur were determined using combustion infrared spectroscopy, and phosphorus was verified using spectroscopic or wet methods. Three positions were taken from each sample, and the mean and standard deviation were calculated.

[0083] The grain size and harmful phase area fraction experiments were conducted on the cross-section and longitudinal section of the tube after cold drawing to the final core diameter. The aim was to evaluate grain refinement and the control of the δ phase and Laves phase. Grain size was rated according to GB / T6394-2017 or ASTM E112-25. The δ phase and Laves phase were identified using backscattered electron imaging combined with energy dispersive spectroscopy point analysis or surface scanning to identify niobium-molybdenum-rich regions. At least 10 fields of view were statistically analyzed for each sample. The grain size level and the total area fraction of the δ phase and Laves phase in Table 1 are expressed as statistical mean ± standard deviation, and the statistical mean is used as the criterion for comparison with the corresponding range.

[0084] The room temperature tensile test was conducted on longitudinal tubular specimens from Examples 1–4 and all comparative examples to evaluate the balance between strength and plasticity. The test was performed according to GB / T228.1-2021 "Metallic materials, tensile testing—Part 1: Test at room temperature." Specimens were prepared according to tubular specifications, and the test was conducted at 25°C. The tensile rate was set according to the method requirements, with n=3 for each group. Tensile strength, specified plastic extension strength Rp0.2, and elongation after fracture were recorded. The tensile performance data in Table 1 are expressed as statistical mean ± standard deviation, and the statistical mean was used as the criterion for comparison with the corresponding range.

[0085] The residual tensile stress test was conducted on the inner and outer surfaces of the tube after cold drawing of the final core diameter, aiming to evaluate the stress state after cold working and surface treatment. X-ray diffraction was used, referring to the X-ray stress measurement method in GB / T7704-2017. Axial and circumferential data were collected, with no fewer than five measurement points on each surface. The measurement point locations, diffraction planes, and stress values ​​were recorded. The residual tensile stress in Table 1 is expressed as the statistical mean ± standard deviation of the measurement points, and the statistical mean was used as the criterion for comparison with the corresponding range.

[0086] The localized corrosion resistance test used surface-treated pipe sections as the subjects to evaluate the adaptability of the Ni-Cr-Mo-Nb system microstructure and surface condition to a chloride-containing corrosive environment. The method of evaluating the pitting corrosion performance of nickel-based chromium-containing alloys using ferric chloride solution in ASTM G48-25 Method C was adopted. The same sample area, immersion time, and temperature gradient were set, and the critical pitting temperature and corrosion weight loss were recorded. Each group had n=3 samples. The critical pitting temperatures in Table 1 are expressed as statistical mean ± standard deviation.

[0087] The flaring and flattening adaptability test was conducted on tubes after cold drawing of the final core diameter, aiming to evaluate the crack susceptibility of the tubes during secondary forming. Flaring was performed according to GB / T242-2007 "Test Method for Flaring of Metal Tubes", and flattening was performed according to GB / T246-2017 "Test Method for Flattening of Metal Tubes". Flaring rate and outer diameter deformation rate were set respectively. After the test, cracks in the end and flattened area were inspected visually and under low magnification microscopy. The maximum crack length in Table 1 is expressed as statistical mean ± standard deviation.

[0088] To verify the effectiveness of this invention in addressing the challenges of balancing cold working strengthening with corrosion resistance and stability, and high-alloy corrosion-resistant design with machinability, comparative evaluations were conducted on alloy element content, final core diameter cold drawing deformation, short-time heat treatment regime, subgrain-dislocation structure, harmful precipitates, residual stress, and the final tube microstructure. The specific implementation results are as follows.

[0089] Figure 1 This diagram illustrates the effect of molybdenum content on tensile strength and critical pitting temperature in this invention. The univariate experiment corresponding to this diagram is based on Example 1. Except for the molybdenum content, the niobium-tantalum content, the reduction rate of the cold-drawn core diameter, the first-stage heat treatment temperature, and subsequent process conditions were kept consistent. The molybdenum content was adjusted from 7.5 wt% to 10.5 wt% to evaluate the synergistic effect of molybdenum on solid solution strengthening and pitting resistance. The results show that when the molybdenum content increases from 7.5 wt% to approximately 9.0–9.5 wt%, the tensile strength increases from approximately 818 MPa to approximately 858–866 MPa, and the critical pitting temperature increases from approximately 39°C to approximately 51–52°C. This indicates that an appropriate amount of molybdenum can improve the solid solution strengthening effect of the matrix and enhance the pitting resistance stability in a chlorine-containing environment. However, when the molybdenum content continues to increase to 10.5 wt%, both the tensile strength and the critical pitting temperature decrease, indicating that excessively high molybdenum content increases the tendency for elemental segregation and harmful precipitation, thus weakening the uniformity of the microstructure and corrosion resistance stability. Therefore, Figure 1 This indicates that a higher molybdenum content is not necessarily better; rather, there exists a suitable range that balances strength and corrosion resistance.

[0090] Figure 2 This diagram illustrates the effect of the total niobium-tantalum content on tensile strength and precipitate area fraction in this invention. The corresponding univariate experiment, based on Example 1, maintains consistency in molybdenum content, final core diameter cold drawing deformation, and heat treatment regime, except for the total niobium and tantalum content. The total niobium-tantalum content was adjusted from 2.80 wt% to 4.40 wt% to further illustrate the relationship between high-alloy strengthening design and precipitate control. The results show that when the total niobium-tantalum content increases from 2.80 wt% to approximately 3.65–3.90 wt%, the tensile strength increases from approximately 805 MPa to approximately 855–862 MPa, indicating that an appropriate amount of niobium-tantalum can improve pipe strength through solid solution strengthening and microstructure refinement. However, when the total niobium-tantalum content increases to above 4.15 wt%, the total area fraction of the δ phase and Laves phase increases from approximately 0.34% to approximately 0.72%, further increasing to approximately 1.10%, while the tensile strength decreases. This result indicates that excessive niobium-tantalum induces the precipitation of harmful second phases, leading to reduced matrix continuity and corrosion resistance. Combination Figure 1 and Figure 2It is understood that by limiting the compositional boundaries of the Ni-Cr-Mo-Nb system, this invention avoids excessive enrichment of strengthening and corrosion-resistant elements such as molybdenum, niobium, and tantalum while improving performance, thus alleviating the contradiction between high alloying design and machinability stability at the compositional level.

[0091] Figure 3 This diagram illustrates the effect of the reduction rate of the final core diameter cold-drawn area on tensile strength and surface residual tensile stress. The univariate experiment corresponding to this diagram is based on Example 1. Except for the reduction rate of the final core diameter cold-drawn area, the alloy composition, pretreatment conditions, and heat treatment regime were kept consistent. The reduction rate of the final core diameter cold-drawn area was adjusted from 1.5% to 10.0% to evaluate the contribution of cold working strengthening and the resulting residual stress risk. The results show that as the reduction rate of the final core diameter cold-drawn area increases from 1.5% to 8.0%, the tensile strength increases from approximately 792 MPa to approximately 980 MPa, indicating that cold drawing of the final core diameter can effectively introduce dislocations and subgrain structures, thereby significantly improving the tube strength. However, when the area reduction rate continues to increase to 10.0%, the surface residual tensile stress rises to approximately 190 MPa. Although the strength remains at a high level, the risk of cracking during subsequent forming and service of the tube increases significantly. Therefore, the amount of cold drawing deformation needs to be matched with the heat treatment regime; too low a value results in insufficient strengthening, while too high a value leads to excessive residual stress. Figure 3 This indicates that the present invention, by employing a moderate final core diameter cold drawing deformation amount, can achieve a balance between cold work strengthening and stress control.

[0092] Figure 4 This diagram illustrates the effect of the first-stage heat treatment temperature on the precipitate area fraction and critical pitting temperature of this invention. The univariate experiment corresponding to this diagram is based on Example 1. Except for the first-stage heat treatment temperature, the alloy composition, the cold-drawn area reduction rate of the final core diameter, and other heat treatment parameters were kept consistent. The first-stage heat treatment temperature was adjusted from 850℃ to 950℃ to evaluate the effect of short-term heat treatment on the dissolution of harmful phases, matrix homogenization, and corrosion resistance stability. The results show that when the first-stage heat treatment temperature increases from 850℃ to 880–905℃, the total area fraction of the δ phase and Laves phase decreases from approximately 0.92% to approximately 0.02–0.08%, and the critical pitting temperature increases to approximately 46–50℃. This indicates that appropriately increasing the first-stage heat treatment temperature is beneficial for promoting the homogenization of corrosion-resistant elements and inhibiting continuous grain boundary precipitation. However, when the temperature continues to increase to 930–950℃, the precipitate area fraction increases again, and the critical pitting temperature decreases, indicating that excessively high temperatures can alter the subsequent microstructure evolution path, leading to localized precipitation and fluctuations in corrosion resistance. Combination Figure 3 and Figure 4It is understood that the two short-time heat treatments of the present invention are not intended to completely eliminate cold work strengthening, but rather to reduce the adverse effects of harmful phases and residual stress while retaining an appropriate amount of subgrain-dislocation strengthening structure, thus resolving the contradiction between strength improvement and corrosion resistance stability from a process perspective.

[0093] Figure 5 This figure shows the statistical distribution of small-angle grain boundaries in Examples 1, 9, and 11. The raw data were statistically analyzed using multiple electron backscatter diffraction (EBSD) fields, with the mean and standard deviation superimposed to evaluate the degree of subgrain structure formation under different process conditions. The results show that the small-angle grain boundary ratio in Example 1 is concentrated in the range of approximately 38–41%, significantly higher than the approximately 22–25% in Comparative Example 9 and approximately 27–31% in Comparative Example 1. The higher small-angle grain boundary ratio indicates that Example 1 formed a relatively stable and sufficiently developed subgrain structure under the combined effects of cold drawing of the final core diameter and short-time heat treatment. This structure can provide effective strengthening without relying on a large amount of continuous grain boundary precipitation, thereby improving the tube strength and maintaining microstructure continuity. The lower small-angle grain boundary ratios in Comparative Examples 9 and 11 indicate insufficient subgrainization or mismatched recovery processes, making it difficult to form a stable subgrain strengthening network.

[0094] Figure 6 This is a TEM (Transmission Electron Microscopy) statistical diagram of dislocation cells in Example 1 and Comparative Example 10. This diagram further provides a quantitative comparison of dislocation cell density and intercellular distance at the transmission electron microscopy scale to verify... Figure 5 The substructural differences reflected in the EBSD statistical results were analyzed. The results showed that the dislocation cell density in Example 1 was approximately 7.2–8.4 × 10¹³ m³. -2 The intercellular distance is approximately 405–445 nm; the dislocation cell density in Comparative Example 10 is approximately 2.8–3.5 × 10¹³ m³. -2 The intercellular spacing is approximately 735–820 nm. Example 1 exhibits a higher dislocation cell density and a smaller intercellular spacing, indicating a finer cold-worked strengthening structure and a more uniform distribution of deformation energy storage; while the dislocation cells in Comparative Example 10 are more porous, resulting in insufficient strengthening effect. Figure 5 and Figure 6 It is understood that the present invention forms a reinforced structure with small-angle grain boundaries and fine dislocation cells as the core through the coordinated control of cold drawing of the final core diameter and two short-term heat treatments. This makes the strength improvement come from stable subgrain dislocation reinforcement, rather than at the cost of excessive precipitation or excessive residual stress.

[0095] Figure 7This figure shows the total area fraction of the δ phase and Laves phase in Examples 1, 3, and 7. The area fraction data from multiple BSE / EDS fields are compared to illustrate the inhibitory effect of this invention on harmful precipitates. The results show that the total area fraction of the δ phase and Laves phase in Examples 1 and 3 is at a low level, approximately 0.01–0.04%; while the total area fraction in Comparative Example 7 is higher, approximately 1.05–1.25%, indicating significant local enrichment and coarse precipitates. Excessive formation of the δ phase and Laves phase consumes key elements such as molybdenum and niobium in the matrix, weakens local corrosion resistance, and forms weak points in the microstructure at grain boundaries or segregation regions. Figure 7 This indicates that, under the combined effect of suitable compositional boundaries and short-time heat treatment regime, Example 1 can effectively suppress the excessive precipitation of harmful second phases, allowing corrosion-resistant elements to remain in the matrix as much as possible to play their role, thereby improving corrosion resistance stability and microstructure reliability.

[0096] Figure 8 This figure shows the statistical graphs of surface residual tensile stress and maximum crack length for Examples 1, 5, and 8. It also presents the XRD residual stress test results and the maximum crack length after forming, used to evaluate the processing adaptability of the pipe after cold drawing of the final core diameter. The results show that the statistical mean of surface residual tensile stress in Example 1 is close to 0 MPa, and the maximum crack length is 0 mm, indicating that it maintains good stress state and forming integrity after obtaining a strengthened structure. The surface residual tensile stress in Comparative Example 5 is approximately 185–196 MPa, and the maximum crack length is approximately 0.45–0.60 mm, indicating that excessive cold drawing deformation of the final core diameter or insufficient stress release will significantly increase the risk of cracking. Comparative Example 8 has lower residual tensile stress and no obvious cracks, but its resistance to localized corrosion is lower than that of Example 1, indicating insufficient overall performance. Therefore, simply reducing residual stress does not represent optimal overall performance; only by achieving a balance between strength, residual stress, and crack control can the actual processing requirements of high-strength corrosion-resistant pipes be met. Figure 8 and Figure 3 The mutual corroboration demonstrates that the final core diameter cold drawing parameters selected in this invention can create a reasonable window between strengthening effect and processing safety.

[0097] Figure 9This is a closed-loop comparison chart of strength and corrosion resistance for Example 1 and representative comparative examples. The chart compares two key indicators: tensile strength and critical pitting temperature, to comprehensively evaluate the overall effect of different composition and process combinations. The results show that Example 1 has a tensile strength of approximately 827 MPa and a critical pitting temperature of approximately 46°C, maintaining both high mechanical strength and good pitting resistance. While some comparative examples exhibit high tensile strength, their critical pitting temperature decreases or crack risk increases, indicating that their strength improvement comes at the cost of reduced corrosion resistance or processing adaptability. Other comparative examples, while having relatively acceptable corrosion resistance or residual stress state, lack sufficient tensile strength to meet the requirements for high-strength pipes. Figure 9 This indicates that the present invention does not pursue the maximization of a single performance index, but rather achieves a better balance between tensile strength, pitting resistance and processing stability through the coupling of composition boundary, final core diameter cold drawing and two short-term heat treatments, thereby forming a closed-loop verification.

[0098] Figure 10 This is a comparison of the SEM morphology of the final product of Example 1 and the final product of Comparative Example 11. This figure is used to further verify the effectiveness of the invention in terms of pipe cross-sectional integrity, pipe wall microstructure uniformity, and control of the second phase at grain boundaries. Figure 10 (a) is a low-magnification SEM image of the final product of Example 1, showing that it forms a complete annular cross-section with an outer diameter of 6.00 mm, a wall thickness of 0.50 mm, and an inner diameter of about 5.00 mm. The tube wall is continuous and the inner and outer wall contours are clear, indicating that the cross-sectional morphology of the tube is stable after the final core diameter is cold drawn, and there is no obvious collapse, cracking or wall thickness instability. Figure 10 (b) is a low-magnification SEM image of the final product of Comparative Example 11, used to compare and observe the changes in the integrity of the pipe cross-section and the uniformity of the wall thickness after changing the heat treatment sequence. Figure 10 (c) is a magnified SEM image of the final product of Example 1, showing that the grain distribution is relatively uniform within a 0.50 mm pipe wall area, the statistical mean of the grain size level is 5.0, and the average grain size is about 62 μm. This indicates that the present invention can obtain a moderate and continuous pipe wall structure, which avoids the strength loss caused by excessive grain coarsening, and also avoids uneven structure caused by local refinement or precipitation concentration. Figure 10 Image (d) in the diagram is a magnified SEM image of the final product in Comparative Example 11, used to illustrate how changing the order of the two short-term heat treatments can negatively impact grain boundary precipitation and local microstructure uniformity. Figure 10 It can be seen that the macroscopic cross-section and microstructure of the tube in Example 1 show good stability, indicating that there is a clear matching relationship between the two short-time heat treatment sequences and the final core diameter cold drawing.

[0099] Figure 11This is a TEM structure comparison image of the final product of Example 1 and the final product of Comparative Example 11. This image is used to verify the advantages of Example 1 in terms of γ-matrix stability, crystal structure continuity, and second-phase control at a higher resolution scale. Figure 11 (a) is a bright-field TEM image of the final product of Example 1. The observation area is a local thin region of the tube wall. The image is dominated by the γ matrix. Combined with the statistical mean of the total area fraction of the δ phase and the Laves phase of about 0.02%, it shows that the precipitation of the second phase is effectively suppressed and the matrix continuity is good. Figure 11 (b) is a bright-field TEM image of the final product of Comparative Example 11, used to compare and observe the changes in the second phase, dislocation structure and local microstructure of the grain boundary or matrix after the heat treatment sequence was changed. Figure 11 (c) is an HRTEM image of the final product of Example 1, showing that the local lattice fringes of the γ matrix are clear and continuous. Combined with the structural characteristics of Ni-based FCC solid solution, the matrix crystal structure in this observation area has good continuity. Figure 11 Image (d) in the diagram is the HRTEM image of the final product in Comparative Example 11, used for comparative analysis of the second-phase interface, local lattice distortion, or changes in crystal structure continuity. Figure 11 It can be seen that Example 1 possesses a stable γ matrix, a low content of harmful phases, and good crystal structure continuity at the microscale, and is consistent with... Figures 5 to 8 The results shown are consistent with those of subgrain dislocation strengthening, harmful phase suppression, and residual stress control.

[0100] In conclusion, Figures 1 to 11 The effectiveness of this invention is demonstrated step by step from multiple levels, including composition design, cold drawing deformation, heat treatment window, subgrain-dislocation structure, harmful precipitates, residual stress, crack control, and the final tube microstructure. Example 1, through Ni-Cr-Mo-Nb system compositional boundary control, maintains a reasonable distribution of elements such as molybdenum, niobium, and tantalum between strengthening and corrosion resistance; through two-stage short-time heat treatment, the homogenization of corrosion-resistant elements and suppression of harmful phases do not come at the cost of completely eliminating cold work strengthening; through final core diameter cold drawing, subgrain-dislocation strengthening does not come at the cost of excessive residual tensile stress and continuous grain boundary precipitation. Therefore, this invention achieves a synergistic unity between high strength, corrosion resistance stability, and processing adaptability, effectively responding to and resolving the technical contradiction in traditional high-alloy nickel-based alloy tubes where it is difficult to simultaneously achieve strength improvement, corrosion resistance maintenance, and processing window.

[0101] Table 1 Performance of Examples and Comparative Examples As can be seen from the performance of the examples and comparative examples in Table 1, Examples 1–4 show a relatively balanced trend in terms of strength, plasticity, surface residual tensile stress, area fraction of δ phase and Laves phase, critical pitting temperature, and control of flaring and flattening cracks. In the conventional comparative examples, deviations in chromium, molybdenum, or niobium-tantalum content lead to a decrease in localized corrosion resistance or mechanical strengthening indicators; insufficient strength occurs when the final core diameter is too low, while increased residual tensile stress and forming crack risk lead to excessively high values. In the synergistic comparative examples, eliminating the second short-term heat treatment, replacing it with a long-term heat treatment, or changing the order of the two treatment stages makes it difficult to simultaneously maintain strength, precipitate control, or processing adaptability. This indicates that compositional matching, two short-term heat treatments, and final core diameter cold drawing together address the balance between strength and corrosion resistance.

[0102] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention and not to limit it. Although the present invention has been described in detail with reference to the above embodiments, those skilled in the art should understand that any equivalent structural transformations made under the concept of the present invention and using the contents of the specification and drawings of the present invention should be covered within the scope of protection of the claims of the present invention.

Claims

1. A high-strength, corrosion-resistant nickel-based alloy pipe, characterized in that, The high-strength corrosion-resistant nickel-based alloy pipe is a seamless Ni-Cr-Mo-Nb system pipe. Based on the total mass of the high-strength corrosion-resistant nickel-based alloy pipe, its chemical composition consists of the following elements and unavoidable impurities: chromium 20.00–23.00 wt%, molybdenum 8.00–10.00 wt%, niobium and tantalum (total content) 3.15–4.15 wt%, iron 0.50–5.00 wt%, carbon 0.005–0.080 wt%, manganese 0–0.50 wt%, silicon 0–0.50 wt%, aluminum 0–0.40 wt%, titanium 0–0.40 wt%, phosphorus 0–0.015 wt%, and sulfur 0–0.015 wt%. The sum of the contents of chromium, molybdenum, iron, carbon, manganese, silicon, aluminum, titanium, phosphorus, sulfur, and the total contents of niobium and tantalum does not exceed 42.00 wt%, with the balance being nickel and unavoidable impurities, and the nickel content is not less than 58.00 wt%. The high-strength corrosion-resistant nickel-based alloy tube has a subgrain dislocation-strengthened structure formed by two short-time heat treatments and cold drawing of the final core diameter. The area reduction rate of the cold drawing of the final core diameter is 3.00–8.00%, the statistical mean of the grain size of the high-strength corrosion-resistant nickel-based alloy tube is 5.0–8.0, and the statistical mean of the total area fraction of δ phase and Laves phase in the high-strength corrosion-resistant nickel-based alloy tube is not greater than 1.00%. The subgrain-dislocation-strengthened microstructure is formed through the following steps: A1. The furnace charge is prepared according to the chemical composition, wherein the furnace charge includes at least nickel powder, chromium powder, molybdenum powder, niobium powder, iron powder and graphite powder for introducing carbon, and tantalum powder, manganese powder, silicon powder, aluminum flakes and titanium powder are selectively added according to the target composition; A2. Melting is carried out under the protection of high-purity argon or vacuum conditions at a melting temperature of 1450–1600℃ and a holding time of 10–60 min to obtain an alloy melt and cast it into an alloy ingot. A3. The alloy ingot is hot-worked to obtain a tube blank at a temperature of 1120–1200℃ and a total deformation of 20–80%. A4. The tube blank is subjected to multiple passes of cold rolling, cold drawing, or a combination of both, to achieve the final core diameter size before cold drawing. The tube blank before the final core diameter is subjected to two short-time heat treatments: the first stage is 880–930℃ for 10–30 min, and the second stage is 980–1020℃ for 8–20 min. A5. The tube blank after the two short-time heat treatments is subjected to final core diameter cold drawing to obtain the high-strength corrosion-resistant nickel-based alloy tube. The high-strength corrosion-resistant nickel-based alloy pipe has an outer diameter of 6.00–219.00 mm, a wall thickness of 0.50–20.00 mm, and an outer diameter greater than twice the wall thickness, and a length of 0.50–18.00 m. The high-strength corrosion-resistant nickel-based alloy tubing after cold drawing of the final core diameter has subgrain zones continuously distributed along the tube axis and dislocation cell structures discontinuously distributed. The δ phase and Laves phase in the high-strength, corrosion-resistant nickel-based alloy tubing are discontinuously distributed along the grain boundaries.

2. The high-strength, corrosion-resistant nickel-based alloy pipe according to claim 1, characterized in that, The surface residual tensile stress, tensile strength, specified plastic elongation strength Rp0.2, and elongation after fracture of the high-strength corrosion-resistant nickel-based alloy pipe are all calculated based on the statistical average of the test data. The surface residual tensile stress does not exceed 150 MPa, the tensile strength is 827–980 MPa, the specified plastic elongation strength Rp0.2 is 414–650 MPa, and the elongation after fracture is 30–55%.

3. The high-strength, corrosion-resistant nickel-based alloy pipe according to claim 1, characterized in that, The chemical composition of the high-strength corrosion-resistant nickel-based alloy pipe contains 0.010–0.045 wt% carbon, 0–0.010 wt% phosphorus, 0–0.006 wt% sulfur, and a total content of 3.40–4.00 wt% niobium and tantalum.

4. The high-strength, corrosion-resistant nickel-based alloy pipe according to claim 1, characterized in that, After step A4 and before step A5, the molybdenum microsegregation coefficient in the tube blank is 1.00–1.20, the niobium microsegregation coefficient in the tube blank is 1.00–1.25, and the statistical mean of the total area fraction of the δ phase and the Laves phase in the tube blank is 0–0.80%.

5. The high-strength, corrosion-resistant nickel-based alloy pipe according to claim 1, characterized in that, The high-strength corrosion-resistant nickel-based alloy tubing undergoes surface treatment after cold drawing of the final core diameter. The surface treatment includes treating with a 1.00–10.00 wt% hydrofluoric acid aqueous solution at 20–45°C for 1–8 min to remove surface oxide scale, treating with a 10.00–25.00 wt% nitric acid aqueous solution at 20–45°C for 1–10 min for passivation, rinsing with deionized water until the pH of the cleaning solution is 6.0–8.0, and drying at 40–120°C for 10–60 min.

6. A method for preparing a high-strength, corrosion-resistant nickel-based alloy pipe as described in any one of claims 1 to 5, characterized in that, Includes the following steps: S1. Prepare the furnace charge according to the target alloy composition for obtaining high-strength, corrosion-resistant nickel-based alloy pipes, wherein the target alloy composition, based on the total mass of the resulting high-strength, corrosion-resistant nickel-based alloy pipes, is: chromium 20.00–23.00 wt%, molybdenum 8.00–10.00 wt%, niobium and tantalum (total content) 3.15–4.15 wt%, iron 0.50–5.00 wt%, carbon 0.005–0.080 wt%, manganese 0–0.50 wt%, silicon 0–0.50 wt%, and aluminum 0–0.40 wt%. The content of chromium, molybdenum, iron, carbon, manganese, silicon, aluminum, titanium, phosphorus, sulfur, niobium and tantalum shall not exceed 42.00 wt%, with the balance being nickel and unavoidable impurities, and the nickel content shall not be less than 58.00 wt%. The furnace charge includes nickel powder, chromium powder, molybdenum powder, niobium powder, iron powder and graphite powder for introducing carbon, and tantalum powder, manganese powder, silicon powder, aluminum flakes and titanium powder are selectively added according to the target composition. S2. Under the protection of high-purity argon or vacuum conditions, melt at 1450–1600℃ for 10–60 min, pour and cool to obtain alloy ingots; S3. The alloy ingot is hot-worked at 1120–1200℃, and the total deformation during hot working is 20–80% to obtain a rough tube or tube blank. S4. Perform multiple cold rolling, cold drawing, or a combination of both on the raw tube or tube blank, and perform intermediate annealing between at least two cold working passes to bring it to the size of the final core diameter before cold drawing. S5. Perform two short-time heat treatments on the tube blank before cold drawing of the final core diameter, wherein the first stage is held at 880–930℃ for 10–30 min, and the second stage is held at 980–1020℃ for 8–20 min. S6. The tube blank after the two short-time heat treatments is cold-drawn with a final core diameter of 3.00–8.00% to obtain the high-strength corrosion-resistant nickel-based alloy tube.

7. The preparation method according to claim 6, characterized in that, In step S2, the furnace charge is dried at 80–200℃ for 1–6 hours before smelting; after smelting, the alloy ingot is kept at 1120–1180℃ for 1–6 hours for homogenization treatment. In step S3, the hot working includes one or more of hot piercing, hot extrusion, or hot rolling.

8. The preparation method according to claim 6, characterized in that, In step S4, the rough tube or tube blank is subjected to multiple passes of cold rolling, cold drawing, or a combination of both, with a single pass area reduction rate of 5.00–25.00%, and intermediate annealing is performed between at least two cold processing passes at a temperature of 1040–1160℃ and a holding time of 1–20 min. In step S5, the heating rate from the first stage to the second stage is 5–25℃ / min; In step S6, the final core diameter is cold-drawn using a core rod or mandrel to support the inner diameter. The cold-drawing speed is 1.00–12.00 m / min. The lubricant residue is cleaned with deionized water and then dried at 40–120℃ for 10–60 min.

Images