Oxide dispersion strengthened copper alloy stranded wire and preparation method thereof

Abstract

The invention discloses an oxide dispersion strengthened copper alloy stranded wire and a preparation method thereof, and belongs to the technical field of copper-based composite materials and powder metallurgy, the alloy components comprise, by mass, 0.15%-0.55% of Al, 0.03%-0.15% of Y, 0.01%-0.06% of Ti and the balance Cu, Al2O3 and Y3Al5O12 composite dispersed phases are dispersed and distributed in a matrix, and the composite dispersed phases are dispersed and distributed in the matrix. The preparation method comprises the steps of gas atomization powder preparation, gradient oxygen partial pressure internal oxidation, hydrogen reduction, cold isostatic pressing, hot extrusion, cold drawing, stranding and stable annealing, the tensile strength of the stranded wire is not lower than 500 MPa, the electric conductivity is not lower than 82% IACS, and the strength retention rate is not lower than 90% after heat preservation at 800 DEG C for 100 h.

Description

Technical Field

[0001] This invention belongs to the field of copper-based composite materials and powder metallurgy technology, specifically relating to oxide dispersion reinforced copper alloy stranded wire and its preparation method. Background Technology

[0002] High-strength, high-conductivity copper alloys are irreplaceable key conductor materials in aerospace, nuclear industry, high-speed rail transportation, and other fields. While traditional precipitation-strengthened copper alloys, such as the Cu-Cr-Zr system, exhibit high strength and conductivity at room temperature, the Cr precipitates, which strengthen the alloy, undergo significant coarsening at high temperatures, leading to a sharp decrease in strength above 500°C. This inherent defect severely limits the application of precipitation-strengthened copper alloys in high-temperature, long-term service conditions.

[0003] Oxide dispersion strengthened copper alloys achieve strengthening by introducing thermodynamically stable nanoscale oxide particles into a copper matrix. Because the melting point of the oxide dispersion phase is much higher than that of the copper matrix and no solid solution reaction occurs between them, it exhibits excellent high-temperature structural stability. Currently, commercially available oxide dispersion strengthened copper products are represented by the GLIDCOP series, which uses an internal oxidation method to generate Al2O3 nanoparticles in situ within Cu-Al alloy powder, followed by hot extrusion densification to obtain rod or sheet products.

[0004] However, existing oxide dispersion strengthening technology for copper alloys still has the following shortcomings. First, the dispersion phase system is singular, relying solely on Al2O3 particles for strengthening. Under extreme high-temperature environments, Al2O3 particles still exhibit a certain degree of Ostwald ripening tendency, and after long-term service, the particle size increases, leading to a weakening of the pinning effect. Second, traditional internal oxidation processes use Cu2O powder as an oxidant and mechanically mix it with alloy powder, resulting in uneven oxygen source distribution. This leads to a wide distribution of Al2O3 particle size within the alloy and the appearance of coarse oxide clusters in some areas. Third, existing densification processes mostly employ high-speed pressing or simple sintering, resulting in a low density of the obtained billet. Residual porosity becomes a weak point in electrical conductivity and mechanical properties. Fourth, systematic preparation processes for stranded wire products have not been reported. After multiple strands of monofilament are stranded, there is a lack of microstructure stabilization treatment, and residual stress release during service may cause the stranded wire to loosen or experience performance fluctuations.

[0005] Therefore, there is an urgent need to develop an oxide dispersion-reinforced copper alloy stranded wire and its system preparation method that combines higher temperature stability, more uniform dispersed phase distribution, and better overall performance. Summary of the Invention

[0006] To address the technical problems in existing technologies, such as the single dispersed phase system, insufficient uniformity of dispersed phase distribution, limited densification degree, and lack of systematic preparation processes for stranded wire products, this invention provides oxide dispersion strengthened copper alloy stranded wire and its preparation method.

[0007] The oxide dispersion-stretched copper alloy stranded wire of the present invention is composed of multiple oxide dispersion-stretched copper alloy monofilaments. The chemical composition of the alloy, by mass percentage, includes: Al 0.15%~0.55%, Y 0.03%~0.15%, Ti 0.01%~0.06%, with the balance being Cu and unavoidable impurities. Al₂O₃ particles and Y₃Al₅O₃ particles are dispersedly distributed in the alloy matrix. 12 The composite oxide dispersion phase consists of particles, with a total volume fraction of 0.3%–2.0%. The average particle size of Al₂O₃ particles is 3–15 nm, and Y₃Al₅O₃ particles are also present. 12 The average particle size is 5~25 nm, and the two oxide particles are uniformly distributed within the grain and grain boundary regions of the copper matrix. The tensile strength of the stranded wire is not less than 500 MPa, the conductivity is not less than 82% IACS, the strength retention rate is not less than 90% after holding at 800°C for 100 h, and the elongation after fracture is not less than 5%.

[0008] This invention also provides a method for preparing the above-mentioned oxide dispersion reinforced copper alloy stranded wire, comprising the following steps: preparing Cu-Al-Y-Ti alloy powder by gas atomization; subjecting the alloy powder to internal oxidation treatment in a gradient oxygen partial pressure atmosphere to in-situ oxidize Al and Y elements to generate nano-sized Al2O3 and Y3Al5O3. 12 Composite dispersed phase; the inner oxide powder is reduced to remove the surface copper oxide; the reduced powder is cold isostatically pressed into a blank; the blank is hot-extruded and densified; the extruded blank is cold-drawn into monofilaments through multiple passes; multiple monofilaments are twisted into a shape; the stranded wire is stabilized and annealed in a protective atmosphere.

[0009] This invention introduces Y and Ti microalloying elements into the Cu-Al binary system, utilizing the competitive oxidation and recombination reactions between Y and Al during the internal oxidation process to generate Al2O3 and Y3Al5O3 in situ within the copper matrix. 12 Two nano-dispersed phases. Y3Al5O 12Yttrium aluminum garnet (YAG) exhibits extremely high thermodynamic stability and excellent interfacial compatibility with the copper matrix, effectively suppressing the high-temperature coarsening of Al2O3 particles and achieving the technical effects of synergistic strengthening and high-temperature stabilization of the two dispersed phases. The addition of trace amounts of Ti promotes the formation of a transitional interface layer between the dispersed phase and the copper matrix, enhancing the interfacial bonding strength. The gradient oxygen partial pressure internal oxidation process achieves uniform and fine precipitation of dispersed phase particles through precise staged control of the oxidation atmosphere. A two-step densification route combining cold isostatic pressing and hot extrusion ensures near-full densification of the billet. Stable annealing after stranding eliminates residual processing stress and further optimizes the interfacial structure between the dispersed phase and the matrix. Attached Figure Description

[0010] Figure 1 This is a process flow diagram for preparing the oxide dispersion reinforced copper alloy stranded wire of the present invention.

[0011] Figure 2 This is a dark-field image of the alloy prepared in Example 1, showing Al2O3 and Y3Al5O. 12 Distribution morphology of dual-dispersed phase particles in a copper matrix.

[0012] Figure 3 This is a high-resolution transmission electron microscope image of the interface region between the dispersed phase particles and the copper matrix in Example 1, showing the atomic arrangement structure of the Al2TiO5 transition interface layer.

[0013] Figure 4 The Al2O3 particles and Y3Al5O in the alloy prepared in Example 1 12 Histogram of particle size distribution.

[0014] Figure 5 This is a scanning electron microscope image of the room temperature tensile fracture surface of the stranded wire in Example 1.

[0015] Figure 6 A bar chart showing the tensile strength retention rate of each embodiment and comparative example after heat treatment at 800°C for 100 h. Detailed Implementation

[0016] The technical solution of the present invention will be further described in detail below with reference to the accompanying drawings and specific embodiments. The following embodiments are used to illustrate the present invention, but are not intended to limit the scope of the present invention.

[0017] like Figure 1 As shown, the preparation method of oxide dispersion strengthened copper alloy stranded wire of the present invention has the following overall process route: alloy smelting and gas atomization powdering → gradient oxygen partial pressure internal oxidation → hydrogen reduction → cold isostatic pressing into billet → hot extrusion densification → multi-pass cold drawing → stranding and forming → stabilization annealing. The technical points and process parameter ranges of each process are described below.

[0018] The alloy powder was prepared using a vacuum induction melting combined with nitrogen atomization. First, the raw materials were fed into the vacuum induction melting furnace according to the designed composition. The selection of raw materials is crucial to the final product quality. Electrolytic copper plates with a purity of no less than 99.95% were used as the copper raw material, and high-purity aluminum ingots with a purity of no less than 99.9% were used. Yttrium was introduced through a Cu-10%Y master alloy, and titanium was introduced through a Cu-5%Ti master alloy. The use of master alloys lowers the melting temperature of high-melting-point alloying elements and promotes their uniform dissolution in the copper melt. The feeding sequence was as follows: first, the electrolytic copper plates were placed in the crucible; after the vacuum level in the furnace was evacuated to below 10 Pa, high-purity argon gas (purity no less than 99.999%) was introduced to a slightly positive pressure protection to prevent oxidation and gas absorption by the copper melt. The induction heating power was gradually increased. After the copper was completely melted, the aluminum ingots were added and electromagnetically stirred for 5-10 minutes to ensure complete dissolution. Subsequently, the Cu-Y master alloy and Cu-Ti master alloy were added sequentially, and stirring was continued for 10-15 minutes to ensure uniform composition. The melting temperature is controlled at 1250~1350°C, and the superheat (i.e., the difference between the temperature of the molten alloy and its liquidus temperature) is controlled within the range of 100~200°C. Superheat is a key parameter affecting the particle size and morphology of the atomized powder. Too low a superheat leads to poor flowability of the molten alloy, low atomization efficiency, and a high proportion of coarse powder; too high a superheat increases energy consumption and may cause volatilization loss of the light element aluminum. After the temperature and composition of the molten alloy stabilize, the bottom-injection guiding system is activated to introduce the molten alloy into the atomization chamber. The atomization medium is high-purity nitrogen (purity not less than 99.99%), and the atomization pressure is 2.5~4.0 MPa. The high-pressure gas flow breaks the molten alloy into fine droplets, which rapidly cool and solidify during flight to form powder particles. The cooling rate can reach 10... 3 ~10 5 The rapid solidification process, characterized by a K / s ratio, allows Al, Y, and Ti elements to remain in a supersaturated solid solution state within the copper matrix, laying the microstructural foundation for the uniform precipitation of dispersed phases during subsequent internal oxidation. The collected alloy powder was sieved using a vibrating sieve to obtain fine powder with a particle size less than 75 μm for later use. The chemical composition of the powder was determined by inductively coupled plasma optical emission spectrometry (ICP-OES), and the oxygen content was measured using an oxygen-nitrogen analyzer, with the oxygen content required to be controlled below 0.02%. The loose packing density of the powder was 3.5–4.5 g / cm³. 3 The tap density is 4.5~5.5 g / cm³. 3 The powder flowability, measured by a Hall effect flowmeter, is no greater than 30 s / 50g. By adjusting the atomization pressure and superheat, powder particles with good sphericity and a satellite powder ratio of less than 5% can be obtained, which is crucial for uniform oxygen penetration during subsequent internal oxidation. Nitrogen was chosen instead of argon as the atomization medium because nitrogen has extremely low solubility in copper and does not react with copper. Furthermore, nitrogen has a slightly higher specific heat capacity and thermal conductivity than argon, which is beneficial for increasing the droplet cooling rate and obtaining a finer powder particle size distribution.

[0019] Internal oxidation is one of the core processes of this invention. Unlike the traditional method of using Cu2O powder mixture to provide an oxygen source, this invention uses a gradient oxygen partial pressure atmosphere for internal oxidation, achieving selective in-situ oxidation of Al and Y elements by precisely controlling the oxygen partial pressure in the furnace atmosphere. In the traditional Cu2O powder mixing method, Cu2O powder decomposes and releases oxygen atoms during heating. However, due to the differences in particle size and density between Cu2O powder and alloy powder, even with thorough mechanical mixing, it is difficult to achieve a completely uniform distribution of the oxygen source at the microscale. This invention uses a gas-phase oxygen supply method, controlling the oxygen partial pressure by introducing a precisely proportioned nitrogen-oxygen mixture into the furnace tube. The volume fraction of oxygen in the mixture is calculated based on the target oxygen partial pressure and the total pressure inside the furnace. The flow rates of nitrogen and oxygen are controlled separately by a high-precision mass flow controller, achieving a mixing accuracy of ±2%. Internal oxidation is carried out in a tubular atmosphere furnace with high-purity alumina material and a tube diameter of not less than 60 mm. Alloy powder is spread evenly in a corundum boat, with a powder thickness not exceeding 15 mm to ensure sufficient oxygen penetration. The thickness of the powder layer is a crucial process parameter affecting the uniformity of internal oxidation. An excessively thick powder layer can lead to inconsistent oxidation levels between upper and lower layers, resulting in a gradient structure with over-oxidation on the surface and insufficient oxidation in the core. The internal oxidation process consists of two stages: the first stage is maintained at 400–450°C for 2–4 hours, with the oxygen partial pressure controlled at 50–200 Pa. This stage primarily achieves preferential oxidation of Al to generate γ-Al₂O₃ nanoparticles. Within this temperature and oxygen partial pressure range, the thermodynamic driving force for Al oxidation is much greater than that for Y. According to Wagner's internal oxidation theory, the diffusion coefficient of Al in copper and the oxide formation free energy determine that Al₂O₃ nucleates preferentially over Y₂O₃. The lower temperature in the first stage is beneficial for increasing the ratio of Al₂O₃ nucleation rate to growth rate, thereby obtaining a large number of extremely fine Al₂O₃ nuclei. In the second stage, the temperature is raised to 480–550°C and maintained for 2–6 hours. During this stage, the oxygen partial pressure is increased to 200–800 Pa, which, while continuing the oxidation of Al, promotes the solid-state reaction between Y and the already formed Al₂O₃. The solid-state reaction between Y₂O₃ and Al₂O₃ follows the sequence 3Y₂O₃ + 5Al₂O₃ → 2Y₃Al₅O₃. 12 The stoichiometric relationship yields Y3Al5O with a garnet structure. 12 (YAG) phase. The formation temperature of the YAG phase is usually above 450°C, so the second stage heating to the range of 480~550°C just meets the kinetic conditions for the formation of the YAG phase. Some Y elements are also directly oxidized to form Y2O3, which then undergoes a solid-state interdiffusion reaction with the surrounding Al2O3, eventually transforming into Y3Al5O. 12The total oxidation time within the two-stage gradient was 4–10 h. Controlling the temperature and oxygen partial pressure in the first stage facilitated the formation of numerous fine Al2O3 nuclei in the copper matrix, avoiding coarse particles and uneven distribution caused by rapid oxidation. In the second stage, moderately increasing the temperature and oxygen partial pressure promoted the formation of Y3Al5O3 nuclei. 12 The composite phase was fully formed. Through this gradient control strategy, Al2O3 with an average particle size of 3~15 nm and Y3Al5O were finally obtained. 12 A bidispersed phase structure with an average particle size of 5–25 nm. Y3Al5O 12 The particle size is slightly larger than that of Al2O3 because the formation of the YAG phase involves solid-state interdiffusion reactions between the two oxides, requiring a greater atomic diffusion distance. During the internal oxidation process, trace amounts of Ti are partially oxidized to form TiO2, which subsequently undergoes an interfacial reaction with Al2O3 to form a nanoscale Al2TiO5 transition layer. This transition layer improves the interfacial bonding between the dispersed phase and the copper matrix, reducing interfacial thermal mismatch stress. Since the oxidation potential of Ti is between that of Al and Cu, only a portion of Ti is oxidized. The unoxidized Ti remains in the copper matrix in solid solution form, providing a weak solid solution strengthening effect without significantly reducing conductivity.

[0020] After internal oxidation, hydrogen reduction is performed to remove the copper oxide (CuO and Cu2O) layer formed on the surface of the powder particles during oxidation. The presence of this copper oxide layer severely affects the metallurgical bonding between particles during subsequent pressing and sintering, reducing the density of the billet and the conductivity of the final product. Reduction is carried out in a tube furnace under a high-purity hydrogen atmosphere (purity not less than 99.999%) at a flow rate of 0.5–2.0 L / min, a reduction temperature of 350–450°C, and a reduction time of 1–3 h. The chemical reactions during reduction are CuO + H2 → Cu + H2O and Cu2O + H2 → 2Cu + H2O. The water vapor generated is carried out of the furnace by the hydrogen flow. Excessively high reduction temperatures can cause undesirable growth of the dispersed phase particles already formed in the copper matrix. In particular, Al2O3 particles may coarsen at temperatures exceeding 500°C through the Ostwald ripening mechanism, weakening the dispersion strengthening effect. If the reduction temperature is too low, the copper oxide will not be reduced sufficiently, and the residual CuO or Cu2O will decompose and release oxygen during subsequent hot extrusion, forming porosity defects inside the billet. The oxygen content of the reduced powder is precisely measured using an oxygen-nitrogen analyzer, and is required to be controlled within 1.05 to 1.15 times the theoretical oxygen content of the dispersed phase, indicating that the free copper oxide on the surface has been basically removed while the oxygen in the dispersed phase has not been reduced. The reduced powder has a metallic copper luster and is stored under dry nitrogen protection in a glove box to prevent re-oxidation.

[0021] Cold isostatic pressing (COP) involves loading reduced powder into a cylindrical rubber mold for isostatic pressing. The mold wall thickness is 5–8 mm, and the material is natural rubber or polyurethane elastomer. Before loading the powder, it is slightly compacted to increase the packing density. The powder-filled rubber mold is then sealed and placed in the high-pressure chamber of a cold isostatic press, using hydraulic oil as the pressurizing medium. The applied pressure is 200–350 MPa, with a pressurization rate of 5–10 MPa / min to prevent mold breakage due to sudden pressure changes. The holding time is 3–8 min. Cold isostatic pressing uses uniform hydrostatic pressure to apply load to the powder, resulting in a more uniform density distribution in all directions of the compact, which is the fundamental difference from uniaxial molding. After demolding, a cylindrical compact with a relative density of 80%–88% is obtained. The compact has sufficient green strength (greater than 5 MPa) to support subsequent handling and assembly operations. The diameter of the preform is designed according to the specifications of the subsequent hot extrusion die and the extrusion ratio, and is generally 60~120 mm, with the length-to-diameter ratio controlled within the range of 2~4. Cold isostatic pressing can achieve uniform densification of the powder, avoid the density gradient problem commonly found in uniaxial pressing, and provide a preform with a uniform structure for subsequent hot extrusion.

[0022] Hot extrusion densification is carried out on a horizontal hydraulic extrusion press with an extrusion pressure of no less than 500 tons. The cold isostatically pressed billet is loaded into a copper cladding made of industrial pure copper (T2 copper) with a wall thickness of 2-3 mm. The copper cladding isolates the billet from direct contact with the extrusion cylinder, preventing the powder billet from sticking to the die material at high temperatures, and also provides lubrication and sealing. The two ends of the cladding are welded together under argon protection to prevent oxidation of the billet. The sealed clad billet is heated to 800-900°C in a resistance heating furnace and held for 30-60 minutes. The purpose of holding is to ensure uniform temperature distribution in the billet; the holding time depends on the billet dimensions. The heating atmosphere is nitrogen protection. After holding, the billet is quickly (within 10 seconds) transferred to the extrusion cylinder for extrusion to minimize heat transfer loss and subsequent temperature drop. The extrusion ratio is controlled at 12:1-25:1, and the extrusion speed is 5-15 mm / s. During extrusion, the hydraulic system monitors the changes in extrusion pressure. A normal extrusion pressure curve shows an initial peak followed by a steady-state condition, with fluctuations in steady-state extrusion pressure not exceeding ±10%. The relative density of the extruded billet reaches over 99.5%, essentially achieving full densification. The copper cladding is removed in subsequent processing via chemical etching with dilute nitric acid or machining. The high temperature and large deformation during hot extrusion result in sufficient metallurgical bonding between powder particles. Simultaneously, the extrusion shear deformation promotes a certain degree of streamline distribution of dispersed phase particles along the extrusion direction, which is beneficial for the formation of fiber texture during subsequent drawing processes and has a positive effect on improving the mechanical properties along the wire axis. X-ray diffraction analysis shows that the extruded billet exhibits significant... <111> and <200> The fibrous texture exhibits increased strength with increasing extrusion ratio. The selection of extrusion temperature requires balancing two factors: excessively low temperatures result in high deformation resistance, insufficient compaction of the billet, and a higher likelihood of extrusion cracks; excessively high temperatures may cause slight coarsening of the dispersed phase particles and grain growth in the copper matrix. After systematic optimization, 850°C was determined to be the optimal extrusion temperature balancing compaction and microstructural stability. At this temperature, dynamic recrystallization of copper proceeds fully while the coarsening of the dispersed phase particles is extremely limited.

[0023] After surface finishing, the extruded billet undergoes multiple cold drawing passes. Surface finishing includes removing the copper cladding layer and surface oxide scale, completed using a centerless lathe or chemical pickling. Drawing is performed on a chain-type wire drawing machine using cemented carbide (YG6) dies with an entry semi-cone angle of 6-8° and a sizing zone length of 30%-50% of the drawing diameter. The drawing lubricant is a copper-specific wire drawing oil, which has excellent lubricity and cooling properties, reducing frictional heat and die wear during the drawing process. Before each drawing pass, the wire surface is cleaned to remove residual oil film and copper shavings from the previous pass. The single-pass reduction in cross-section is controlled at 10%-20%. Excessive deformation in a single pass can lead to cracks or wire breakage on the wire surface, while insufficient deformation results in low drawing efficiency and may cause uneven work hardening. The total reduction in cross-section reaches 85%-95%. Due to the strong pinning effect of dispersed phase particles on dislocation movement, oxide dispersion-strengthened copper alloys exhibit a significantly higher work hardening rate than pure copper and precipitation-strengthened copper alloys. Therefore, a well-planned intermediate annealing process is necessary to restore the material's plastic deformation capacity. Intermediate annealing is performed when the cumulative reduction in cross-section reaches 60%–70%, at a temperature of 350–500°C for 30–90 minutes, using a nitrogen or nitrogen-hydrogen mixture (hydrogen component 5%–15%). The purpose of intermediate annealing is to eliminate work hardening, restore the material's plastic deformation capacity, and maintain the strengthening effect of the dispersed phases. Because of the pinning effect of the oxide dispersed phases, the copper matrix does not undergo complete recrystallization after intermediate annealing; only recovery and partial recrystallization occur, thus retaining a high strength level after annealing. The selection of the intermediate annealing temperature needs to balance the degree of plasticity recovery and strength preservation. If the temperature is too low (below 350°C), dislocation recovery will be insufficient and plasticity recovery will be limited; if the temperature is too high (above 500°C), significant recrystallization and softening may occur. Metallographic observation shows that after annealing at 400°C for 60 min, approximately 40%–60% of the deformed structure underwent recovery and local recrystallization, while the remaining area retained the fibrous structure characteristics of the drawn state. After intermediate annealing, drawing continued to the target monofilament diameter, which ranged from 0.3 to 1.5 mm. Throughout the drawing process, the dimensional accuracy of the wire was monitored using an online diameter gauge, requiring a diameter tolerance not exceeding ±0.01 mm. After drawing, the monofilaments underwent surface quality inspection, requiring a smooth surface free of cracks, scratches, and oxidation discoloration.

[0024] Multiple monofilaments are twisted together on a stranding machine. The stranding employs a concentric layer stranding structure with a pitch ratio of 10 to 16 times the outer diameter. The pitch ratio is a key parameter affecting the flexibility and mechanical properties of the stranded wire. A pitch ratio that is too small will lead to excessive twisting of the strand and greater bending stress on the monofilaments, while a pitch ratio that is too large will result in a loose stranded structure and decreased overall mechanical properties. The stranded wire structure is selected according to application requirements; typical structures include 1+6 type seven-core stranded wire and 1+6+12 type nineteen-core stranded wire. The 1+6 type structure uses one monofilament as the center, with six monofilaments arranged at equal intervals on the outer layer, typically twisted in a right-hand (S-direction). The 1+6+12 type structure adds twelve more monofilaments to the seven-core structure, with the inner and outer layers twisted in opposite directions to improve the structural stability of the stranded wire. During the stranding process, the tension control accuracy is no less than ±5%, and the tension is set according to the monofilament diameter and material strength, generally 2% to 5% of the monofilament's breaking force. Tension control is achieved through the servo tension control system of the stranding machine. The tension of each feed reel is independently adjustable to ensure that each filament is subjected to uniform force and that the stranded wire surface is smooth without any loose strands. The stranding speed is set according to the equipment capacity and product quality requirements, and is generally 5~30 m / min.

[0025] After stranding, a stabilization annealing treatment is performed, which is a crucial step to ensure the long-term performance stability of the stranded wire. Annealing is carried out in a continuous annealing furnace or a tubular annealing furnace, with a high-purity nitrogen atmosphere or a nitrogen-hydrogen mixture (hydrogen component 5%–15%), an annealing temperature of 300–500°C, and an annealing time of 30–120 min. The advantage of using a nitrogen-hydrogen mixture atmosphere instead of pure nitrogen is that trace amounts of hydrogen can reduce the extremely thin oxide film that may form on the wire surface during stranding, maintaining the metallic luster and good electrical contact performance of the stranded wire. The heating rate of the annealing furnace is controlled at 3–8°C / min to avoid temperature overshoot and localized overheating of the stranded wire caused by rapid heating. The main functions of stabilizing annealing are threefold: First, it eliminates bending stress and contact stress between adjacent filaments introduced during stranding, ensuring the stranded wire is in a low residual stress state, which is crucial for its long-term service in vibrating environments. Second, it promotes short-range atomic diffusion and interface relaxation at the interface between dispersed phase particles and the copper matrix, optimizing the interface bonding state and releasing stress concentration at the interface, thus forming a more stable semi-coherent or coherent interface relationship. Third, it moderately reduces dislocation density without causing significant recrystallization softening, forming ordered dislocation cells or subgrain boundary structures through dislocation rearrangement, achieving an optimal balance between strength and plasticity. After annealing, the Vickers hardness of the stranded wire decreases by approximately 8% to 15% compared to before annealing, but the elongation increases by approximately 50% to 80%, resulting in a significant improvement in overall mechanical properties. The annealing temperature should not exceed 500°C, otherwise the copper matrix grains will grow significantly (from the submicron level in the drawn state to the several micron level), resulting in a substantial decrease in strength. If the annealing temperature is below 300°C, the stress relief effect will be insufficient (residual stress will only be reduced by 20%~30%), and the interface optimization effect will be limited. After annealing, the stranded wire is allowed to cool naturally to ambient temperature at room temperature before final testing and packaging.

[0026] Example 1

[0027] The alloy's chemical composition, by mass percentage, is: Al 0.35%, Y 0.08%, Ti 0.03%, with the balance being Cu. The atomization pressure for gas atomization powder production is 3.0 MPa, the superheat is 150°C, and powder with a particle size less than 75 μm is used. The gradient oxidation process involves the first stage at 420°C and an oxygen partial pressure of 100 Pa for 3 h, and the second stage at 500°C and an oxygen partial pressure of 400 Pa for 4 h. Hydrogen reduction is performed at 400°C for 2 h. Cold isostatic pressing is carried out at 280 MPa for 5 min. Hot extrusion is performed at 850°C for 45 min, with an extrusion ratio of 18:1 and an extrusion speed of 10 mm / s. The extruded billet is first drawn to a reduction in cross-section of 65%, then intermediately annealed at 400°C for 60 min, and subsequently drawn further to a single wire diameter of 0.5 mm, resulting in a total reduction in cross-section of approximately 92%. Seven monofilaments were twisted together in a 1+6 configuration with a pitch ratio of 12 times the outer diameter. The stranded wire was annealed at 400°C for 60 min in a nitrogen-hydrogen mixed atmosphere (hydrogen gas integral 10%).

[0028] like Figure 2 As shown, the microstructure of the alloy prepared in Example 1 was observed using a transmission electron microscope. Two types of nano-dispersed phase particles with different morphologies were uniformly distributed within the copper matrix. One type consisted of nearly spherical γ-Al₂O₃ particles with an average particle size of approximately 8 nm and a number density of approximately 3.5 × 10⁻⁶. 23 m -3 ,like Figure 4 As shown, the particle size distribution ranges from 3 to 15 nm, with a standard deviation of only 2.8 nm, indicating that the gradient oxygen partial pressure internal oxidation process achieves an extremely narrow particle size distribution. The γ-Al₂O₃ particles exist in a spinel structure, with approximately 4.2% lattice mismatch between their (111) crystal planes and the (111) crystal planes of the copper matrix, belonging to the category of semi-coherent interfaces. Another type is the polyhedral Y₃Al₅O₃. 12 The particles have an average diameter of approximately 12 nm and a number density of approximately 8.0 × 10⁻⁶. 22 m -3 The particle size distribution ranges from 5 to 25 nm. Y3Al5O 12 The particles exhibit typical garnet structure characteristics, and their cubic crystal system provides them with extremely high thermodynamic stability at high temperatures. Two dispersed phases are uniformly distributed within the grain and at grain boundaries. Intragranular particles primarily hinder dislocation movement through the Orowan mechanism, while grain boundary particles suppress grain boundary migration and grain growth through the Zener pinning effect. The spacing between the dispersed phase particles is statistically estimated to be approximately 30–80 nm, a range that precisely matches the optimal strengthening spacing for the Orowan bypass mechanism. Selected area electron diffraction analysis confirmed the presence of γ-Al₂O₃ and Y₃Al₅O₃. 12The presence of two phases, the diffraction ring indices and the standard PDF card (γ-Al2O3: PDF #10-0425, Y3Al5O) 12 (PDF #33-0040) is a complete match. Figure 3 As shown, at the interface between the dispersed phase particles and the copper matrix, a Ti-containing transition layer with a thickness of approximately 1–2 nm was observed by high-resolution transmission electron microscopy. Energy dispersive spectroscopy analysis indicated that this transition layer was rich in Ti and O elements, and was inferred to be the Al₂TiO₅ phase. This transition layer effectively reduced the lattice mismatch between the dispersed phase and the copper matrix, transforming the interface from an incoherent to a semi-coherent state, enhancing the interfacial bonding force, and facilitating the effective transfer of load between the dispersed phase and the matrix. Figure 5 As shown, the fracture morphology of the alloy was observed under a scanning electron microscope. The room temperature tensile fracture surface exhibited typical ductile fracture characteristics. The fracture surface was covered with uniform and fine dimples with an average diameter of about 0.5~2 μm. Nano-dispersed phase particles or imprints left by particle debonding were visible at the bottom of the dimples, indicating that dispersed phase particles are the preferred sites for micropore initiation. However, due to the Ti-induced interfacial transition layer enhancing interfacial bonding, the initiation of micropores requires greater stress concentration. Therefore, the alloy exhibits high fracture toughness.

[0029] The mechanical and electrical properties of the stranded wire in Example 1 were tested as follows: the room temperature tensile strength was 536 MPa, the yield strength was 498 MPa, the elongation after fracture was 7.2%, and the conductivity was 84.5% IACS. After the stranded wire sample was kept at 800°C in a high-purity argon atmosphere for 100 h, the tensile strength was 497 MPa, with a strength retention rate of 92.7%. In comparison, the strength retention rate of Cu-0.35%Al binary alloy (excluding Y and Ti) stranded wire prepared using the same process was only 83.6% under the same high-temperature holding conditions, indicating that Y3Al5O 12 The introduction of composite dispersed phases significantly improves the stability of high-temperature structures.

[0030] Example 2

[0031] The alloy's chemical composition, by mass percentage, is: Al 0.18%, Y 0.04%, Ti 0.015%, with the balance being Cu. The atomization pressure for gas atomization powder production is 3.5 MPa, and the superheat is 180°C. The gradient internal oxidation process is performed in two stages: the first stage at 410°C and an oxygen partial pressure of 80 Pa for 2.5 h, and the second stage at 490°C and an oxygen partial pressure of 300 Pa for 3 h. Hydrogen reduction is performed at 380°C for 2.5 h. Cold isostatic pressing is performed at 250 MPa for 5 min. Hot extrusion is performed at 870°C for 40 min, with an extrusion ratio of 15:1 and an extrusion speed of 12 mm / s. The monofilament is drawn to a diameter of 0.8 mm, with a total section reduction of approximately 88%, and the intermediate annealing temperature is 380°C. The seven-core stranded wire is annealed in a nitrogen atmosphere at 380°C for 45 min.

[0032] Example 2 shows that the room temperature tensile strength of the stranded wire is 502 MPa, the yield strength is 461 MPa, the elongation after fracture is 8.5%, and the conductivity is 87.3% IACS. After holding at 800°C for 100 h, the tensile strength is 456 MPa, with a strength retention rate of 90.8%. This example has a lower dispersed phase content and higher conductivity, but its strength and high-temperature retention rate are slightly lower than those of Example 1.

[0033] Example 3

[0034] The alloy chemical composition, by mass percentage, is: Al 0.50%, Y 0.13%, Ti 0.05%, with the balance being Cu. The atomization pressure for gas atomization powder production is 2.8 MPa, and the superheat is 130°C. The gradient internal oxidation process is carried out in two stages: the first stage at 440°C and an oxygen partial pressure of 150 Pa for 3.5 h, and the second stage at 530°C and an oxygen partial pressure of 600 Pa for 5 h. The hydrogen reduction temperature is 420°C, and the reduction time is 1.5 h. The cold isostatic pressing pressure is 320 MPa, and the holding time is 6 min. The hot extrusion temperature is 830°C, the holding time is 50 min, the extrusion ratio is 20:1, and the extrusion speed is 8 mm / s. The monofilament is drawn to a diameter of 0.5 mm, with a total section reduction of approximately 93%, and the intermediate annealing temperature is 420°C. The 19-core stranded wire is annealed at 420°C for 90 min in a nitrogen-hydrogen mixed atmosphere (hydrogen component 8%).

[0035] Example 3 shows that the room temperature tensile strength of the stranded wire is 568 MPa, the yield strength is 531 MPa, the elongation after fracture is 5.3%, and the conductivity is 82.1% IACS. After holding at 800°C for 100 h, the tensile strength is 528 MPa, with a strength retention rate of 93.0%. This example has the highest dispersed phase content, the best strength and high-temperature stability, but the conductivity and elongation are slightly reduced.

[0036] Comparative Example 1 is a Cu-Al binary alloy scheme without Y and Ti. The alloy chemical composition, by mass percentage, is: Al 0.35%, with the balance being Cu. The preparation process is exactly the same as in Example 1. The room temperature tensile strength of the stranded wire is 478 MPa, and the conductivity is 86.2% IACS. After holding at 800°C for 100 h, the tensile strength is 399 MPa, with a strength retention rate of only 83.6%. Compared with Example 1, the high-temperature strength retention rate of Comparative Example 1 is about 9 percentage points lower, which is due to the lack of Y3Al5O. 12 Due to the synergistic pinning effect of the dispersed phase, individual Al2O3 particles underwent a certain degree of Ostwald ripening during long-term high-temperature holding. Transmission electron microscopy confirmed that the average particle size of Al2O3 particles in Comparative Example 1 increased from 9 nm to 22 nm after holding at 800°C for 100 h, while the Al2O3 particles in Example 1 only increased from 8 nm to 11 nm.

[0037] Comparative Example 2 uses a traditional Cu2O powder internal oxidation process instead of the gradient oxygen partial pressure internal oxidation process of this invention. The alloy composition is the same as in Example 1, but the internal oxidation step is changed to: the alloy powder and Cu2O powder are mixed uniformly in a molar ratio and internally oxidized at 450°C in a nitrogen atmosphere for 6 h. The remaining process steps are the same as in Example 1. The room temperature tensile strength of the stranded wire is 508 MPa, and the conductivity is 83.8% IACS. After holding at 800°C for 100 h, the strength retention rate is 88.1%, which is lower than 92.7% in Example 1. Microscopic observation revealed that the size distribution range of the dispersed phase particles in Comparative Example 2 is 3~45 nm, which is significantly wider than the 3~25 nm range in Example 1, and local clusters of coarse Al2O3 particles with a diameter exceeding 40 nm appeared. This is because the distribution of Cu2O powder in the alloy powder is locally uneven, resulting in inconsistent precipitation of the dispersed phase due to fluctuations in the oxygen source concentration.

[0038] Comparative Example 3 investigated the effect of not performing stabilizing annealing on the properties of the stranded wire. The alloy composition and all process parameters were the same as in Example 1, but no stabilizing annealing was performed after stranding. The room temperature tensile strength of the stranded wire was 549 MPa, slightly higher than the 536 MPa of Example 1, because the residual stress introduced by stranding had not yet been eliminated. However, the elongation after fracture of this stranded wire was only 4.1%, lower than the 7.2% of Example 1, and it exhibited obvious early wire breakage in the bending fatigue test. Creep tests at 300°C for 500 h showed that the creep strain of the stranded wire in Comparative Example 3 was 2.3 times that of Example 1, indicating that the presence of residual stress accelerated high-temperature creep.

[0039] Comparative Example 4 is a Ti-free Cu-Al-Y ternary alloy scheme. The alloy chemical composition by mass percentage is: Al 0.35%, Y 0.08%, and the balance Cu. The preparation process is exactly the same as in Example 1. The room temperature tensile strength of the stranded wire is 525 MPa, and the conductivity is 84.8% IACS. After holding at 800°C for 100 h, the strength retention rate is 91.2%. Compared with Example 1, the high-temperature retention rate of Comparative Example 4 is about 1.5 percentage points lower. High-resolution transmission electron microscopy analysis shows that the interface between the dispersed phase particles and the copper matrix in Comparative Example 4 is in a non-coherent state, with obvious moiré fringes at the interface, while the Ti-induced Al2TiO5 transition layer in Example 1 transforms the interface into a semi-coherent state. The non-coherent interface is more likely to become a site for vacancy aggregation and microcrack initiation at high temperatures, which is the microstructural reason for the slightly lower high-temperature performance of Comparative Example 4.

[0040] The performance testing methods for each embodiment and comparative example are as follows. Tensile strength and elongation after fracture were determined according to GB / T228.1-2021 "Metallic materials, tensile testing—Part 1: Tests at room temperature," using an electronic universal testing machine with a maximum load of not less than 10 kN, a test rate of 1 mm / min, and a gauge length of 100 mm. At least three parallel specimens were tested under each test condition, and the average value was taken as the final result. The standard deviation of the three specimens should not exceed 5% of the average value. Conductivity was determined according to GB / T 3048.2-2007 "Electrical properties of wires and cables—Part 2: Tests for resistivity of metallic materials," using the four-terminal method at a constant temperature of 20°C ± 0.5°C, a measuring current of 100 mA, and a measuring length of not less than 1 m. Conductivity was expressed in terms of the International Standard for Annealed Copper (IACS), with the resistivity of pure copper at 20°C being 1.7241 μΩ·cm corresponding to 100% IACS. The high-temperature holding test was conducted in a tubular resistance furnace under high-purity argon gas protection (flow rate 0.5 L / min). The heating rate was 5°C / min, the holding temperature was 800°C, and the holding time was 100 h. After cooling to room temperature with the furnace, tensile tests were performed. The high-temperature holding test examines the microstructural stability of the dispersed phase under extreme temperature conditions and its ability to inhibit matrix grain growth, which is a core indicator for evaluating the high-temperature service reliability of oxide dispersion-strengthened copper alloys. Microstructure observation was performed using a FEI Talos F200X field emission transmission electron microscope with an accelerating voltage of 200 kV and equipped with a Super-X energy dispersive spectroscopy detector. Samples were prepared using a dual-jet electrolytic thinning method with a 30% nitric acid methanol solution (volume fraction) as the electrolyte. The electrolysis temperature was -30°C, and the perforation voltage was 15–20 V. Statistical analysis of the dispersed phase particles was based on dark-field images of at least five different regions, with no fewer than 200 particles counted in each region. The particle size distribution was fitted using a log-normal distribution.

[0041] like Figure 6 As shown, based on the performance data of the above examples and comparative examples, the following pattern can be derived: Within the Al content range of 0.15% to 0.55%, as the Al content increases, the total amount of dispersed phase increases, the alloy strength improves but the conductivity decreases; the two are inversely proportional. Example 2 (Al 0.18%) has a conductivity of 87.3% IACS and a tensile strength of 502 MPa; Example 3 (Al 0.50%) has a conductivity of 82.1% IACS and a tensile strength of 568 MPa; Example 1 (Al 0.35%) falls between the two, representing the optimal balance between strength and conductivity. The physical essence of this trade-off is that while dispersed phase particles provide dislocation pinning strengthening, they also increase the contribution to free electron scattering. Each 0.1% increase in the volume fraction of dispersed phase reduces the conductivity by approximately 1-2% IACS.

[0042] The introduction of Y element causes the simultaneous formation of Al2O3 and Y3Al5O in the copper matrix. 12 Two diffuse phases, Y3Al5O 12 The thermodynamic stability of the particles is higher than that of Al2O3, and the synergistic presence of the two types of particles can effectively inhibit the Ostwald ripening of single Al2O3 particles, thus improving the high-temperature strength retention rate by 5-9 percentage points. A comparison between Comparative Example 1 (without Y, retention rate 83.6%) and Example 1 (containing Y 0.08%, retention rate 92.7%) clearly shows that Y3Al5O 12 The dispersed phase makes a decisive contribution to high-temperature stability. The driving force of Ostwald ripening comes from the surface energy difference between particles of different sizes, with larger particles growing by consuming smaller particles. When two dispersed phases with different chemical compositions coexist in the system, the dissolution-redeposition process between particles of the same type is spatially blocked by particles of different types, effectively reducing the flux of long-range atomic diffusion, thereby significantly slowing down the Ostwald ripening rate.

[0043] Ti microalloying, by forming an Al2TiO5 transition layer at the dispersed phase / matrix interface, transforms the interface from incoherent to semi-coherent, improving interfacial bonding strength and load transfer efficiency. This results in an approximately 10 MPa increase in room temperature strength while also improving high-temperature creep resistance. While the difference between Comparative Example 4 (without Ti, 91.2% retention) and Example 1 (containing 0.03% Ti, 92.7% retention) is not as significant as that of Y, the interfacial modification effect of Ti is more pronounced under long-term creep conditions. The lattice constant of the Al2TiO5 transition layer, between Al2O3 and Cu, acts as a bridge for the gradual transition of lattice parameters, transforming the originally sharp interfacial stress concentration into a gentle stress distribution gradient.

[0044] Compared to the traditional Cu2O powder-mixed internal oxidation process, the gradient oxygen partial pressure internal oxidation process can significantly narrow the size distribution of dispersed phase particles, eliminate coarse particle clusters, and improve the high-temperature retention rate by approximately 4 percentage points. A comparison between Comparative Example 2 (Cu2O powder-mixed method, dispersed phase size distribution 3–45 nm, retention rate 88.1%) and Example 1 (gradient oxygen partial pressure method, dispersed phase size distribution 3–25 nm, retention rate 92.7%) shows that the uniformity of the dispersed phase has an impact on material properties as significant as the total amount of dispersed phase. Coarse oxide clusters not only fail to provide effective dislocation pinning but also become weak points for crack initiation, reducing the material's fracture toughness and fatigue performance.

[0045] Stabilizing annealing significantly improves the plasticity, fatigue life, and creep resistance of stranded wire products. While Comparative Example 3 (unannealed) showed slightly higher room temperature strength than Example 1 (annealed), its elongation decreased by approximately 3 percentage points, and its high-temperature creep strain increased by approximately 1.3 times. The partial recovery and rearrangement of dislocations during annealing transforms the stress state in the stranded wire from high-stress triaxial tension to a low-stress near-equilibrium state, significantly improving the reliability of the stranded wire under complex load conditions.

[0046] The synergistic mechanism of the aforementioned technical features can be analyzed from the following levels. At the dispersion strengthening level, Al2O3 particles, with their high number density, provide resistance to Orowan dislocation bypassing and are the main source of strengthening contribution; Y3Al5O... 12 Although the particle number density is low, its large size and higher elastic modulus make it a strong barrier to dislocation movement, and the two work together to form a multi-scale dispersion-strengthened system. Regarding high-temperature stability, Y3Al5O... 12 The formation of Y3Al5O consumes some of the Al-O coordination environment around the Al2O3 particles, increasing the average spacing between the remaining Al2O3 particles and reducing the driving force of Ostwald ripening; at the same time, Y3Al5O 12 The particles themselves possess a garnet structure with extremely low solubility at high temperatures, enabling them to serve as anchor points for long-term stability in the microstructure. At the interface engineering level, the Ti-induced Al2TiO5 transition layer eliminates stress concentration at the dispersed phase / matrix interface, preventing interface debonding and microcrack initiation during high-temperature service. Furthermore, the semi-coherent interface exhibits higher dislocation pinning efficiency than the incoherent interface. At the processing and annealing level, the high-density dislocations introduced by cold drawing form a dislocation entanglement-particle pinning composite structure with the nano-dispersed phase particles. After annealing, some dislocations recover but do not undergo complete recrystallization, preserving subgrain boundaries and dislocation cell structures conducive to strengthening.

[0047] The above description is merely a preferred embodiment of the present invention and is not intended to limit the invention. Various modifications and variations can be made to the present invention by those skilled in the art. Any modifications, equivalent substitutions, improvements, etc., made within the spirit and principles of the present invention should be included within the scope of the claims of the present invention.

Claims

1. An oxide dispersion-strengthened copper alloy stranded wire, characterized in that, The stranded wire is made of multiple oxide dispersion-strengthened copper alloy monofilaments. The chemical composition of the alloy, by mass percentage, includes: Al 0.15%~0.55%, Y 0.03%~0.15%, Ti 0.01%~0.06%, with the balance being Cu and unavoidable impurities. Al2O3 particles and Y3Al5O3 particles are dispersed in the alloy matrix. 12 The composite oxide dispersion phase consists of particles, the total volume fraction of which is 0.3% to 2.0%, the average particle size of which is 3 to 15 nm, and the Y3Al5O3 particles. 12 The average particle size is 5~25 nm, the tensile strength of the stranded wire is not less than 500 MPa, and the conductivity is not less than 82% IACS.

2. The oxide dispersion reinforced copper alloy stranded wire according to claim 1, characterized in that, The Al2O3 particles and Y3Al5O 12 An Al2TiO5 transition interface layer exists between the particles and the copper matrix, and the thickness of the transition interface layer is 1~3 nm.

3. The oxide dispersion reinforced copper alloy stranded wire according to claim 1, characterized in that, The number density of the Al2O3 particles is 1.0 × 10⁻⁶. 23 ~5.0×10 23 m -3 The Y3Al5O 12 The number density of the particles is 3.0 × 10⁻⁶. 22 ~1.5×10 23 m -3 .

4. The oxide dispersion reinforced copper alloy stranded wire according to claim 1, characterized in that, The tensile strength retention rate of the stranded wire after being kept at 800°C for 100 hours is not less than 90%.

5. The oxide dispersion reinforced copper alloy stranded wire according to claim 1, characterized in that, The diameter of the single filament of the stranded wire is 0.3~1.5 mm, the stranded wire structure is a concentric layer stranded structure, and the stranding pitch ratio is 10~16 times the outer diameter.

6. The method for preparing oxide dispersion reinforced copper alloy stranded wire according to any one of claims 1 to 5, characterized in that, Includes the following steps: Step 1, Alloy Powder Preparation: Cu-Al-Y-Ti alloy powder was prepared by gas atomization, with an atomization pressure of 2.5~4.0MPa, and powder with a particle size of less than 75 μm was selected. Step 2, gradient internal oxidation treatment: The alloy powder is internally oxidized in a gradient oxygen partial pressure atmosphere, wherein the first stage temperature is 400~450°C, the oxygen partial pressure is 50~200 Pa, and the holding time is 2~4 h; the second stage temperature is 480~550°C, the oxygen partial pressure is 200~800 Pa, and the holding time is 2~6 h. Step 3, reduction treatment: reduce in a hydrogen atmosphere at 350~450°C for 1~3 h; Step 4, cold isostatic pressing to form the billet: pressure is 200~350 MPa; Step 5, hot extrusion densification: temperature is 800~900°C, extrusion ratio is 12:1~25:1; Step 6, cold drawing: the total section reduction rate is 85%~95%, and the intermediate annealing temperature is 350~500°C; Step 7: Twist and shape; Step 8, Stabilization Annealing: Anneal at 300~500°C for 30~120 min in a protective atmosphere.

7. The preparation method according to claim 6, characterized in that, In step two, the first stage primarily achieves the preferential oxidation of Al to generate γ-Al₂O₃ nanoparticles, while the second stage promotes the solid-state reaction between Y and Al₂O₃ to form Y₃Al₅O₃. 12 The total internal oxidation time for the two phases is 4–10 h.

8. The preparation method according to claim 6, characterized in that, The holding time for hot extrusion in step five is 30-60 min, the extrusion speed is 5-15 mm / s, and the relative density of the extruded billet is not less than 99.5%.

9. The preparation method according to claim 6, characterized in that, In step six, when the cumulative cross-sectional reduction rate reaches 60%~70%, intermediate annealing is performed. The intermediate annealing time is 30~90 min, and the annealing atmosphere is nitrogen or a nitrogen-hydrogen mixture.

10. The preparation method according to claim 6, characterized in that, The protective atmosphere described in step eight is high-purity nitrogen or a nitrogen-hydrogen mixture with a hydrogen gas fraction of 5% to 15%.

Images