High-elasticity high-conductivity alloy material for connector terminals

Abstract

The application relates to the technical field of alloy materials, and discloses a high-elasticity and high-conductivity alloy material for a connector terminal, which is made of the following raw materials in parts by weight: cathode copper 895.75-930.5 parts, electrolytic nickel 18-22 parts, electrolytic cobalt 4-6 parts, copper-silicon intermediate alloy 40-60 parts, copper-magnesium intermediate alloy 5-10 parts and copper-phosphorus intermediate alloy 2.5-6.25 parts. The preparation method comprises the following steps: single solid solution base strip is prepared through smelting, casting, heat soaking, hot rolling, quenching and annealing; the strip is subjected to multi-pass cold rolling and low-temperature pre-cluster treatment; the strip is subjected to micro-plastic extension deformation under the action of tensile stress under heating; finally, in-situ nonlinear heat treatment is carried out in a hydrogen atmosphere, slow cooling in the furnace is carried out after high-temperature heat preservation, and rapid cooling is carried out to below room temperature when the temperature is reduced to a specified temperature. Through the cooperation of specific alloying and deformation heat treatment, the yield strength, conductivity and stress relaxation resistance of the material are improved.

Description

Technical Field

[0001] This invention relates to the field of alloy materials technology, specifically to a high-elasticity, high-conductivity alloy material for connector terminals. Background Technology

[0002] With the rapid development of fields such as electronic communications, automotive electrical systems, and smart terminals, connector products are gradually evolving towards miniaturization and high density. The copper alloy materials used to manufacture connector terminals not only need to have high conductivity to ensure low-loss transmission of current and signals, but also must have sufficient room temperature yield strength and elastic properties to maintain stable mechanical contact pressure between the connecting components.

[0003] Currently, precipitation-strengthened copper alloys with added elements such as nickel and silicon are commonly used in industry as connector terminal materials. These alloys undergo deformation heat treatment to induce the precipitation of nanoscale second phases from the matrix, thereby increasing strength by hindering dislocation movement. However, in actual processing and long-term service, existing alloy materials still exhibit significant limitations. When terminal components are stamped and bent, existing materials are prone to localized strain concentration in areas subjected to complex bending deformations. Due to the lack of effective microscopic regulation and pinning mechanisms for internal dislocation slippage, rough wrinkles or even microcracks easily appear on the outer surface of the bent area, directly limiting the processing yield of the terminal structure.

[0004] Furthermore, connector terminals are typically exposed to high ambient temperatures and continuously subjected to elastic bending stress during actual operation. Under prolonged thermal stress coupling, residual processing stress within existing precipitation-strengthened copper alloys gradually releases, triggering dislocation recovery and rearrangement, resulting in significant stress relaxation. This microscopic stress relaxation manifests macroscopically as an irreversible decrease in the positive contact force of the terminals, easily leading to increased contact resistance and heat generation, severely weakening the reliability of the connector. Although the industry often attempts to improve matrix performance through multi-element composite additions, improper control of trace element ratios and heat treatment solution removal can significantly increase the electron scattering cross-section due to excessive solid solution atoms, resulting in a substantial decrease in material conductivity. Therefore, effectively suppressing bending strain concentration and improving the high-temperature stress relaxation resistance of materials from a microstructural design perspective, while ensuring high strength and high conductivity, is a key technical challenge that needs to be addressed in the field of high-elasticity and high-conductivity copper alloys.

[0005] Therefore, this invention proposes a high-elasticity, high-conductivity alloy material for connector terminals to address the shortcomings of existing technologies. Summary of the Invention

[0006] To address the shortcomings of existing technologies, this invention provides a high-elasticity and high-conductivity alloy material for connector terminals, which solves the problems of local strain concentration during bending and forming of existing copper alloys for connector terminals, and residual stress relaxation leading to poor contact under long-term high-temperature service conditions.

[0007] In a first aspect, the present invention provides a high-elasticity, high-conductivity alloy material for connector terminals, said alloy material being made from raw materials comprising the following parts by weight: Cathode copper: 895.75–930.5 parts; 18-22 parts of electrolytic nickel; 4-6 parts of electrolytic cobalt; 40-60 parts of copper-silicon master alloy; 5-10 parts of copper-magnesium master alloy; 2.5 to 6.25 parts of copper-phosphorus master alloy.

[0008] By employing a composite alloying process using cathode copper, electrolytic nickel, and electrolytic cobalt in specific mass ratios, combined with an intermediate alloy containing silicon, magnesium, and phosphorus, a microstructure that synergistically combines precipitation strengthening and dislocation pinning can be formed within the material. This contributes to achieving a combination of high conductivity, room temperature yield strength, and resistance to stress relaxation.

[0009] In terms of its microstructure evolution, nickel, cobalt, and silicon in the system tend to desolvate from the supersaturated solid solution during heat treatment, thereby precipitating silicide nanoparticles. During this process, cobalt can participate in the lattice construction of the precipitated phase, entering the lattice by directly replacing some nickel atoms.

[0010] Lattice substitution alters the intrinsic lattice constant of the precipitate, increasing the lattice mismatch between the precipitate and the copper matrix. This often helps to enhance the precipitate's resistance to dislocation slip.

[0011] Trace amounts of magnesium and phosphorus in the system tend to accumulate in lattice defect regions such as dislocation lines and stacking faults. Considering the size difference between magnesium and phosphorus atoms and the matrix copper atoms, they will aggregate in the dislocation core region driven by the reduction of the system's elastic strain energy, forming a Mg-P precluster structure.

[0012] Some of the remaining cobalt atoms dissolved in the matrix undergo short-range drift and combine with the already formed Mg-P preclusters to construct a composite gas around the dislocation core. This gas structure, enveloping the vicinity of the dislocation line, can, to some extent, compensate for and counteract the lattice distortion energy caused by dislocations, forming a thermodynamic pinning effect on the dislocation network. Therefore, when the material is subjected to external bending loads, the aforementioned composite gas can restrict the activation of dislocation sources and long-range dislocation movement, reduce local strain concentration, and thus improve the overall load-bearing capacity and stress relaxation resistance of the material at high temperatures.

[0013] Preferably, the copper-silicon master alloy is a binary alloy ingot containing alloyed elemental silicon with a silicon mass fraction of 10 wt%, the copper-magnesium master alloy is a binary alloy ingot containing alloyed elemental magnesium with a magnesium mass fraction of 10 wt%, and the copper-phosphorus master alloy is a binary alloy ingot containing alloyed elemental phosphorus with a phosphorus mass fraction of 8 wt%.

[0014] Replacing pure silicon and pure magnesium with binary alloy ingots of a specific mass fraction not only reduces the oxidation loss rate of active elements during high-temperature smelting but also improves the dissolution kinetics of silicon and magnesium in copper melt to some extent. This helps reduce macroscopic component segregation and the formation of non-metallic inclusions, thus providing a basis for the uniform distribution of nano-precipitates and pre-clusters in the matrix in subsequent processes. Adding binary alloy ingots with a specific phosphorus mass fraction can, on the one hand, purify the copper alloy melt using the self-deoxidation effect of phosphorus, and on the other hand, introduce a relatively appropriate amount of free phosphorus atoms into the matrix.

[0015] In a second aspect, the present invention provides a method for preparing a high-elasticity and high-conductivity alloy material for connector terminals, comprising the following steps: S1. The raw materials in the specified proportion are added to the melting furnace and melted, and cast into alloy ingots. The alloy ingots are homogenized and hot rolled into strips in multiple passes. After online water quenching and surface descaling, they are sent to a continuous annealing furnace for heat preservation and annealing. After exiting the furnace, they are forced to quench and cool to room temperature to obtain a basic strip with a single supersaturated solid solution phase structure. S2. The base strip is subjected to multi-pass cold rolling deformation, followed by low-temperature pre-clustering treatment under a protective atmosphere; S3. Apply pure tensile stress to the treated strip under heating conditions to induce micro-plastic elongation deformation. S4. The strip roll that has undergone micro-deformation is loaded into an annealing furnace with hydrogen as the working atmosphere and continuously subjected to in-situ nonlinear heat treatment. First, it is heated to a constant high temperature and held at a heating rate. Then, it is slowly cooled with the furnace at a constant cooling rate. When it drops to the specified temperature, it is rapidly cooled to room temperature by high-pressure gas forced convection purging.

[0016] The main process for preparing the basic strip aims to effectively dissolve the added alloying elements into the copper matrix to obtain a single supersaturated solid solution, which often provides the necessary solute atoms for subsequent phase transformation. After multiple cold rolling passes, a dislocation network of a certain density is introduced into the matrix, serving as a channel for short-range atomic diffusion. In the subsequent low-temperature heat treatment, driven by thermodynamics, trace amounts of specific atoms such as magnesium and phosphorus tend to agglomerate towards the dislocation core region, gradually constructing a pre-cluster structure. During the micro-deformation stage, the superposition of temperature and tensile stress induces a small elastic expansion of the lattice. At this point, due to the reduced diffusion barrier of atoms, some cobalt atoms in the solid solution state are more likely to drift directionally towards and combine with the pre-cluster region. The final nonlinear heat treatment, with its high-temperature holding stage promoting the desolvation of the silicon-containing precipitates, followed by slow furnace cooling, helps release accumulated processing stress. When the temperature drops to a specific point, the applied forced convection quenching largely cuts off the long-range atomic diffusion pathways, thus preserving the nanoscale phase size and dislocation cluster structure.

[0017] Preferably, in step S1, the specific smelting process is as follows: cathode copper, electrolytic nickel, and electrolytic cobalt are added to a high-vacuum induction furnace, and smelted at 5.0 × 10⁻⁶ ℃. -2 Heated to 1150–1200°C under a vacuum of Pa to completely melt it; Argon gas was introduced to atmospheric pressure, followed by the sequential addition of copper-silicon master alloy, copper-magnesium master alloy, and copper-phosphorus master alloy. After electromagnetic stirring was turned on for 10–15 minutes, the alloy ingot was formed by vertical semi-continuous casting at a casting temperature of 1120–1160°C.

[0018] The alternation of high-vacuum induction melting and argon purging, from a metallurgical perspective, helps reduce the absorption of ambient gases by the high-temperature copper molten metal and decreases the loss of easily oxidized alloying elements. Furthermore, the addition of an intermediate alloy followed by electromagnetic stirring generally promotes the homogenization of the temperature and concentration fields within the melt during the early solidification stage, thus suppressing macroscopic compositional segregation in the ingot to a certain extent.

[0019] Preferably, in step S1, the specific process for heat treatment of the base strip is as follows: The alloy ingot is placed in a heating furnace and heated to 900-950°C for 2-4 hours. After being hot rolled in multiple passes using a hot rolling mill, it is immediately water-quenched online. After milling the surface of the hot-rolled strip to remove the oxide scale, it is sent into a continuous annealing furnace with a protective atmosphere and held at 900-920℃ for 1-1.5 hours. After exiting the furnace, it is subjected to single-layer forced quenching and cooling by high-pressure water drainage.

[0020] Prolonged high-temperature homogenization can often promote the diffusion and homogenization of segregated elements in the as-cast microstructure. The hydraulic forced quenching following hot rolling deformation and continuous annealing provides a high cooling rate, which can effectively suppress the unplanned precipitation of non-coherent phases during the strip cooling process, thus helping to maintain a relatively pure phase structure inside the solid solution.

[0021] Preferably, in step S2, the total thickness reduction of the multi-pass cold rolling is controlled to be 60% to 65%, and the rolling oil temperature is controlled to be ≤40℃ during the rolling process; The process of low-temperature pre-clustering treatment is as follows: the cold-rolled strip is placed in a bell-shaped annealing furnace under nitrogen protection, the furnace temperature is set to 250-280℃, and the temperature is kept constant for 2-2.5 hours. Then it is taken out of the furnace and air-cooled to room temperature.

[0022] Controlling the total thickness reduction within a specific range is primarily used to accumulate suitable deformation energy storage for the matrix; simultaneously, limiting the rolling oil temperature can reduce the tendency for premature precipitation of solute atoms caused by deformation heat generation. Combined with a relatively low annealing furnace temperature, it is sufficient to activate the short-range diffusion of atoms such as magnesium and phosphorus, without inducing a large number of nucleations of the main precipitate phase, which allows the pre-clustering process to proceed relatively independently to a certain extent.

[0023] Preferably, the specific implementation method of step S3 is as follows: The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 150-180°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a micro-plastic elongation deformation of 1.5%-2.5%.

[0024] By applying pure tensile stress at a specific plate temperature using a tension straightener, in addition to improving the flatness of the strip, the resulting controlled microplastic elongation is accompanied by minute slippage of internal grain boundaries and local dislocations. Considering that this process does not induce significant recovery softening, this microscopic evolution provides a certain kinetic energy and spatial basis for the rearrangement and distribution of specific trace amounts of solute atoms within the strip.

[0025] Preferably, in step S4, the purity of the hydrogen gas is ≥99.999 vol%. The heating and holding process for the in-situ nonlinear heat treatment is as follows: Turn on the strong convection circulating fan and start the heating power to raise the temperature of the coil center to 480℃ at a rate of 12-20℃ / min; keep it at 480℃ for 1-2 hours.

[0026] Hydrogen of a specified purity is introduced as the working atmosphere. Its excellent thermal conductivity and reducing properties accelerate heat transfer within the strip coil and suppress surface oxidation. Setting a relatively fast heating rate helps the strip cross the precipitation-sensitive temperature range, which to some extent reduces the probability of the formation of coarse equilibrium phases and guides the internal precipitates to desolvate in a fine and dispersed form.

[0027] Preferably, in step S4, the in-situ nonlinear heat treatment cooling and quenching process is as follows: After turning off the main heating source, adjust the auxiliary heating power and circulating air volume to make the coil temperature cool down linearly and slowly with the furnace at a constant cooling rate of 10-15℃ / h. When the thermocouple reading shows that the center of the roll has dropped to 350°C, immediately start the water-cooled heat exchanger and maintain the high-pressure hydrogen circulation in the furnace. Use the high-pressure hydrogen for forced convection purging to rapidly cool the roll to below 100°C within 20-30 minutes.

[0028] Slow in-furnace cooling after heat preservation and solvent removal helps to gradually release residual thermal stress inside the material and stabilize the precipitated phase structure. When the strip is cooled to a specific point, forced convection quenching is switched on. The engineering purpose of this is to suppress further long-range diffusion growth of atoms by rapid cooling after the precipitated phase size development and internal stress have been reduced to a certain level, thereby preserving the current microstructure at a lower temperature.

[0029] This invention provides a high-elasticity, high-conductivity alloy material for connector terminals. It has the following advantages: 1. This invention employs a specific ratio of cathode copper, electrolytic nickel, and electrolytic cobalt combined with an intermediate alloy containing silicon, magnesium, and phosphorus. Cobalt atoms partially replace nickel atoms in the precipitated phase to increase lattice mismatch, and trace amounts of magnesium and phosphorus are used to construct pre-clusters in the dislocation core region, thereby guiding cobalt atoms to combine and form composite gas clusters. This composition-designed precipitated phase strengthening and dislocation thermodynamic pinning mechanism restricts the long-range movement of dislocations within the lattice, improving the room-temperature yield strength and conductivity of the alloy material while reducing interference with the electron transport channels of the matrix.

[0030] 2. In the smelting stage, this invention uses binary master alloys of copper-silicon, copper-magnesium, and copper-phosphorus with specific mass fractions to replace elemental pure metal raw materials, thereby reducing the oxidation loss rate of active alloying elements at high temperatures. By leveraging the self-deoxidizing effect of the phosphorus-containing master alloy, the melt is purified, and the dissolution kinetics of solute atoms in the copper melt are improved. This helps to reduce macroscopic component segregation and the formation of non-metallic inclusions, providing a compositional basis for the uniform distribution of subsequent nano-precipitates and pre-clusters in the matrix.

[0031] 3. The preparation process of this invention combines low-temperature pre-clustering treatment, micro-deformation with tensile stress, and in-situ nonlinear heat treatment in a series. It utilizes the coupling effect of temperature and micro-plastic extension to reduce the diffusion barrier of solute atoms, driving atoms to directionally agglomerate towards dislocation regions. Combined with slow furnace cooling during the heat treatment stage and forced airflow quenching at specific temperature nodes, it orderly releases residual stress from material processing while cutting off further atomic diffusion channels and solidifying the microstructure, thereby improving the material's stress relaxation resistance under high-temperature service conditions. Attached Figure Description

[0032] Figure 1 The X-ray diffraction test spectra of each sample of the present invention are shown; wherein, (a) is a comparison diagram of local slow scanning diffraction peaks of the crystal plane, and (b) is a comparison diagram of characteristic diffraction peaks of the micro-area. Figure 2 The temperature-dependent internal friction spectra of each strip sample of the present invention are shown. Figure 3 This is a size distribution diagram of the nano-precipitated phase inside each strip sample of the present invention; Figure 4 This is a two-dimensional distribution diagram of yield strength and electrical conductivity of the present invention; Figure 5 The stress relaxation curves of each strip sample of the present invention under a high temperature environment of 150°C are shown. Figure 6 This is a comparison chart of the minimum R / t value test for a 90-degree bend according to the present invention. Detailed Implementation

[0033] The technical solutions in the embodiments of the present invention will be clearly and completely described below. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0034] The main raw materials and reagents used in the following examples and comparative examples have the following sources and specifications. Reagents not specifically mentioned are all commercially available analytical grade or higher grade products.

[0035] Cathode copper (CAS No.: 7440-50-8) is a commercially available high-purity industrial metal raw material with elemental copper purity ≥ 99.99 wt%.

[0036] Electrolytic nickel (CAS No.: 7440-02-0) is a commercially available high-purity industrial metal raw material with a purity of elemental nickel ≥99.99wt%.

[0037] Electrolytic cobalt (CAS No.: 7440-48-4) is a commercially available high-purity industrial metal raw material with elemental cobalt purity ≥ 99.98 wt%.

[0038] The copper-silicon master alloy is a commercially available standard binary alloy ingot containing alloyed elemental silicon (CAS No.: 7440-21-3) and with a silicon mass fraction of 10wt%.

[0039] The copper-magnesium master alloy is a commercially available standard binary alloy ingot containing alloyed elemental magnesium (CAS No.: 7439-95-4) with a magnesium mass fraction of 10 wt%.

[0040] The copper-phosphorus master alloy is a commercially available standard binary alloy ingot containing alloyed elemental phosphorus (CAS No.: 7723-14-0) with a phosphorus mass fraction of 8wt%.

[0041] Preparation Example 1: This preparation example provides a method for preparing a basic strip, including the following steps: 930.5 parts by weight of cathode copper, 18 parts by weight of electrolytic nickel, and 4 parts by weight of electrolytic cobalt were added to a high-vacuum induction furnace and heated at 5.0 × 10⁻⁶ mm. -2 Heated to 1150℃ under a vacuum of Pa to completely melt it. Argon gas was introduced to atmospheric pressure, followed by the sequential addition of 40 parts by weight of copper-silicon master alloy, 5 parts by weight of copper-magnesium master alloy, and 2.5 parts by weight of copper-phosphorus master alloy. After electromagnetic stirring for 10 minutes, the alloy ingot was drawn into a vertical semi-continuous casting method at a casting temperature of 1120℃. The alloy ingot was placed in a heating furnace and heated to 900℃ for 2 hours for homogenization. It was then hot-rolled in multiple passes using a hot rolling mill to a strip with a thickness of 10 mm, followed by online water quenching. After milling the surface of the hot-rolled strip to remove the oxide scale, it was sent to a continuous annealing furnace with a protective atmosphere and held at 900℃ for 1 hour. After exiting the furnace, it was subjected to single-layer forced quenching under high pressure to room temperature, obtaining a basic strip with a single supersaturated solid solution phase structure.

[0042] Preparation Example 2: This preparation example provides a method for preparing a basic strip, including the following steps: 913.25 parts by weight of cathode copper, 20 parts by weight of electrolytic nickel, and 5 parts by weight of electrolytic cobalt were added to a high-vacuum induction furnace and heated at 5.0 × 10⁻⁶ mm. -2Heated to 1180℃ under a vacuum of Pa to completely melt it. Argon gas was introduced to atmospheric pressure, followed by the sequential addition of 50 parts by weight of copper-silicon master alloy, 8 parts by weight of copper-magnesium master alloy, and 3.75 parts by weight of copper-phosphorus master alloy. After electromagnetic stirring for 12 minutes, the alloy ingot was drawn into a vertical semi-continuous casting method at a casting temperature of 1140℃. The alloy ingot was placed in a heating furnace and heated to 920℃ for 3 hours for homogenization. It was then hot-rolled in multiple passes using a hot rolling mill to a strip with a thickness of 12 mm, followed by online water quenching. After milling the surface of the hot-rolled strip to remove the oxide scale, it was sent to a continuous annealing furnace with a protective atmosphere and held at 910℃ for 1.2 hours. After exiting the furnace, it was subjected to single-layer forced quenching under high pressure to room temperature to obtain a basic strip with a single supersaturated solid solution phase structure.

[0043] Preparation Example 3: This preparation example provides a method for preparing a basic strip, including the following steps: 895.75 parts by weight of cathode copper, 22 parts by weight of electrolytic nickel, and 6 parts by weight of electrolytic cobalt were added to a high-vacuum induction furnace and heated at 5.0 × 10⁻⁶ mm. -2 Heated to 1200℃ under a vacuum of Pa to completely melt it. Argon gas was introduced to atmospheric pressure, followed by the sequential addition of 60 parts by weight of copper-silicon master alloy, 10 parts by weight of copper-magnesium master alloy, and 6.25 parts by weight of copper-phosphorus master alloy. After electromagnetic stirring for 15 minutes, the alloy ingot was drawn into a vertical semi-continuous casting method at a casting temperature of 1160℃. The alloy ingot was heated to 950℃ in a heating furnace for 4 hours and then hot-rolled in multiple passes using a hot rolling mill to a strip with a thickness of 15 mm, followed by online water quenching. After milling the surface of the hot-rolled strip to remove oxide scale, it was sent to a continuous annealing furnace with a protective atmosphere and held at 920℃ for 1.5 hours. After exiting the furnace, it was subjected to single-layer forced quenching under high pressure to room temperature to obtain a basic strip with a single supersaturated solid solution phase structure.

[0044] Example 1: This embodiment provides a method for preparing a high-elasticity, high-conductivity alloy material for connector terminals, including the following steps: The base strip obtained in Preparation Example 2 was subjected to multi-pass cold rolling using a 20-roll reversible cold rolling mill, with the total thickness reduction controlled at 62%. The rolling oil temperature was controlled to be ≤40℃ during the rolling process. The cold-rolled strip was then coiled and placed in a bell-type annealing furnace for low-temperature pre-clustering treatment under nitrogen protection. The furnace temperature was set at 265℃ and held at that temperature for 2.2 hours. The strip was then removed from the furnace and air-cooled to room temperature.

[0045] The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 165°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a 2.0% micro-plastic elongation deformation.

[0046] The slightly deformed strip coil is loaded into a full-hydrogen high-convection annealing furnace. The working atmosphere is hydrogen with a purity ≥99.999 vol%. In-situ nonlinear heat treatment is performed continuously without compromising the furnace's airtightness. First, the high-convection circulating fan at the bottom of the furnace platform is turned on and the heating power is started, causing the center temperature of the strip coil to rise to 480℃ at a rate of 15℃ / min. After holding at 480℃ for 1.5 hours, the main heating source is turned off, and the auxiliary heating power and circulating air volume are adjusted to allow the strip coil temperature to cool linearly and slowly to 350℃ at a constant cooling rate of 12℃ / h.

[0047] When the temperature measuring thermocouple shows that the center of the coil has dropped to 350°C, the water-cooled heat exchanger outside the furnace is immediately started and the high-pressure hydrogen is kept circulating inside the furnace. The high-pressure hydrogen is used for forced convection purging to rapidly cool the coil to below 100°C within 25 minutes. Then, nitrogen is used for replacement and the furnace is opened to discharge the material. After conventional surface cleaning and slitting, a high-elasticity and high-conductivity alloy material is obtained.

[0048] Example 2: This embodiment provides a method for preparing a high-elasticity, high-conductivity alloy material for connector terminals, including the following steps: The base strip obtained in Preparation Example 1 was subjected to multiple cold rolling passes using a 20-roll reversible cold rolling mill, with the total thickness reduction controlled at 60%. During the rolling process, the rolling oil temperature was controlled to be ≤40℃. The cold-rolled strip was then coiled and placed in a bell-type annealing furnace for low-temperature pre-clustering treatment under nitrogen protection. The furnace temperature was set at 250℃ and held at that temperature for 2 hours. Subsequently, the strip was removed from the furnace and air-cooled to room temperature.

[0049] The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 150°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a 1.5% micro-plastic elongation deformation.

[0050] The slightly deformed strip coil is loaded into a full-hydrogen high-convection annealing furnace. The working atmosphere is hydrogen with a purity ≥99.999 vol%. In-situ nonlinear heat treatment is performed continuously without compromising the furnace's airtightness. First, the high-convection circulating fan at the bottom of the furnace platform is turned on and the heating power is started, causing the center temperature of the strip coil to rise to 480℃ at a rate of 12℃ / min. After holding at 480℃ for 1 hour, the main heating source is turned off, and the auxiliary heating power and circulating air volume are adjusted to allow the strip coil temperature to be linearly and slowly cooled to 350℃ at a constant cooling rate of 10℃ / h along with the furnace.

[0051] When the temperature measuring thermocouple shows that the center of the coil has dropped to 350°C, the water-cooled heat exchanger outside the furnace is immediately started and the high-pressure hydrogen is kept circulating inside the furnace. The high-pressure hydrogen is used for forced convection purging to rapidly cool the coil to below 100°C within 20 minutes. Then, nitrogen is used for replacement and the furnace is opened to discharge the material. After conventional surface cleaning and slitting, a high-elasticity and high-conductivity alloy material is obtained.

[0052] Example 3: This embodiment provides a method for preparing a high-elasticity, high-conductivity alloy material for connector terminals, including the following steps: The base strip obtained in Preparation Example 3 was subjected to multiple cold rolling passes using a 20-roll reversible cold rolling mill, with the total thickness reduction controlled at 65%, and the rolling oil temperature controlled at ≤40℃ during the rolling process. The cold-rolled strip was then coiled and placed in a bell-type annealing furnace for low-temperature pre-clustering treatment under nitrogen protection. The furnace temperature was set at 280℃ and held at that temperature for 2.5 hours, after which it was removed from the furnace and air-cooled to room temperature.

[0053] The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 180°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a 2.5% micro-plastic elongation deformation.

[0054] The slightly deformed strip coil is loaded into a full-hydrogen high-convection annealing furnace. The working atmosphere is hydrogen with a purity ≥99.999 vol%. In-situ nonlinear heat treatment is performed continuously without compromising the furnace's airtightness. First, the high-convection circulating fan at the bottom of the furnace platform is turned on and the heating power is started, causing the center temperature of the strip coil to rise to 480℃ at a rate of 20℃ / min. After holding at 480℃ for 2 hours, the main heating source is turned off, and the auxiliary heating power and circulating air volume are adjusted to allow the strip coil temperature to be linearly and slowly cooled to 350℃ at a constant cooling rate of 15℃ / h.

[0055] When the temperature measuring thermocouple shows that the center of the coil has dropped to 350°C, the water-cooled heat exchanger outside the furnace is immediately started and the high-pressure hydrogen is kept circulating inside the furnace. The high-pressure hydrogen is used for forced convection purging to rapidly cool the coil to below 100°C within 30 minutes. Then, nitrogen is used for replacement and the furnace is opened to discharge the material. After conventional surface cleaning and slitting, a high-elasticity and high-conductivity alloy material is obtained.

[0056] Example 4: This embodiment provides a method for preparing a high-elasticity, high-conductivity alloy material for connector terminals, including the following steps: The base strip obtained in Preparation Example 2 was subjected to multiple cold rolling passes using a 20-roll reversible cold rolling mill, with the total thickness reduction controlled at 65%. During the rolling process, the rolling oil temperature was controlled to be ≤40℃. The cold-rolled strip was then coiled and placed in a bell-type annealing furnace for low-temperature pre-clustering treatment under nitrogen protection. The furnace temperature was set at 250℃ and held at that temperature for 2.5 hours, after which the strip was removed from the furnace and air-cooled to room temperature.

[0057] The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 180°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a 1.5% micro-plastic elongation deformation.

[0058] The slightly deformed strip coil is loaded into a full-hydrogen high-convection annealing furnace. The working atmosphere is hydrogen with a purity ≥99.999 vol%. In-situ nonlinear heat treatment is performed continuously without compromising the furnace's airtightness. First, the high-convection circulating fan at the bottom of the furnace platform is turned on and the heating power is started, causing the center temperature of the strip coil to rise to 480℃ at a rate of 18℃ / min. After holding at 480℃ for 1.2 hours, the main heating source is turned off, and the auxiliary heating power and circulating air volume are adjusted to allow the strip coil temperature to cool linearly and slowly to 350℃ at a constant cooling rate of 10℃ / h.

[0059] When the temperature measuring thermocouple shows that the center of the coil has dropped to 350°C, the water-cooled heat exchanger outside the furnace is immediately started and the high-pressure hydrogen is kept circulating inside the furnace. The high-pressure hydrogen is used for forced convection purging to rapidly cool the coil to below 100°C within 20 minutes. Then, nitrogen is used for replacement and the furnace is opened to discharge the material. After conventional surface cleaning and slitting, a high-elasticity and high-conductivity alloy material is obtained.

[0060] Comparative Example 1: Compared with Example 1, the difference is that no electrolytic cobalt and copper-magnesium master alloy were added when preparing the basic strip, so that the melt does not contain cobalt and magnesium elements. The total amount of copper elements introduced by the cathode copper and master alloy in the system was adjusted to 974.7 parts by weight, and the rest were the same.

[0061] Comparative Example 2: Compared with Example 1, the difference is that the low-temperature pre-clustering treatment step is skipped, and the cold-rolled strip is directly passed through a strip bending straightener with an online electromagnetic induction heating section for medium-temperature stress coupling micro-deformation. All other aspects are the same.

[0062] Comparative Example 3: Compared with Example 1, the difference is that the medium-temperature stress-coupled micro-deformation is replaced with room-temperature conventional cold rolling treatment, that is, the strip that has completed the low-temperature pre-clustering treatment is subjected to a 2.0% thickness reduction deformation by a cold rolling mill at room temperature, instead of applying pure tensile stress under the plate temperature of 165°C. All other aspects are the same.

[0063] Comparative Example 4: Compared with Example 1, the difference lies in the change of the heating rate parameter in the all-hydrogen strong convection annealing furnace, which reduces the heating rate of the coil center temperature from 15°C / min to the conventional 2°C / min, while the rest are the same.

[0064] Comparative Example 5: Compared with Example 1, the difference lies in changing the dynamic cooling and desolvation step in the in-situ nonlinear heat treatment. Instead of the slow step cooling process of 12℃ / h, the strip is kept at a constant temperature of 480℃ for 1.5h and then kept at a constant temperature of 480℃ for 3h. The subsequent rapid cooling operation is then carried out directly. All other aspects are the same.

[0065] Comparative Example 6: Compared with Example 1, the difference is that the in-situ gas quenching and locking step when the temperature drops to 350°C is cancelled. That is, when the temperature measuring thermocouple shows that the center of the coil has dropped to 350°C, the water-cooled heat exchanger and high-pressure hydrogen forced convection purging are not started. Instead, the coil is allowed to cool slowly to room temperature naturally with the furnace. Everything else is the same.

[0066] Test Example 1: Pure copper samples were selected as the reference standard, and alloy materials prepared in Example 1, Comparative Example 1, and Comparative Example 5 were used as test objects. Block samples with dimensions of 15mm × 15mm were cut along the strip rolling direction.

[0067] The surfaces of each sample were mechanically ground with wet sandpaper to remove surface oxides and initial macroscopic scratches. A 30% (v / v) phosphoric acid aqueous solution was prepared as a polishing solution, and electrolytic polishing was performed at room temperature with a polishing voltage of 5V and a polishing time of 60s to remove the residual stress layer generated during mechanical grinding. The samples were then ultrasonically cleaned with anhydrous ethanol and dried with cold air.

[0068] X-ray diffraction data were acquired from the treated sample using an X-ray diffractometer. The test conditions were set as follows: Cu-Kα radiation source (λ=1.54056Å), tube voltage 40kV, and tube current 40mA.

[0069] To determine the matrix lattice constant, a slow-step scanning was performed on the Cu(311) diffraction peak in the high diffraction angle region. The scanning range was set to 2θ = 88.5°–91.0°, the step size was set to 0.01°, and the dwell time per step was 3 s. After data acquisition, the actual lattice constant of the copper matrix in each sample was calculated using the Nelson-Riley extrapolation method.

[0070] For characterizing the precipitated phase structure, an ultra-slow scan was performed in the range of 2θ = 43° to 46.5° with a step size of 0.02° and a dwell time of 8 s per step to improve the signal-to-noise ratio of weak diffraction peaks. After obtaining the diffraction pattern, considering the tailing interference of the strong diffraction peak of Cu(111) in the copper matrix (around 43.3°), the background was subtracted, and peak fitting was used to finely separate overlapping peaks. The positions of the main diffraction peaks of the precipitated phase were extracted and compared with the standard PDF card for calibration.

[0071] The experimental data are shown in Table 1: Table 1: Calculated and measured X-ray diffraction peak positions and lattice constants of each sample pure copper reference sample 89.927 3.61494 - Example 1 89.941 3.61452 44.783 Comparative Example 1 89.986 3.61317 45.105 Comparative Example 5 89.968 3.61369 44.821 in conclusion: According to Table 1 and Figure 1 Compared with the pure copper reference sample, the peak position of the copper matrix (311) in each alloy sample shifted to a certain extent, indicating that the alloying elements and subsequent heat treatment process affected the lattice state of the copper matrix. Specifically, the matrix lattice constant of Example 1 was 3.61452 Å, which is close to the 3.61494 Å of the pure copper reference sample and higher than that of Comparative Examples 1 and 5. This result indicates that under the heat treatment conditions used in Example 1, the residual solute atoms in the matrix have a relatively small impact on the lattice constant, suggesting that these process conditions may promote the precipitation of some solute atoms from the matrix, thereby reducing the degree of matrix lattice distortion.

[0072] The matrix lattice constant of Comparative Example 1 is 3.61317 Å, which is lower than that of Example 1 and Comparative Example 5, indicating that even without the addition of cobalt and under appropriate heat treatment conditions, a relatively large number of solid solution atoms causing lattice distortion may still remain in the copper matrix. The matrix lattice constant of Comparative Example 5 is 3.61369 Å, which is between that of Example 1 and Comparative Example 1, indicating that under the same or similar alloy systems, different heat treatment conditions have a certain influence on the degree of solute desolvation.

[0073] Regarding the diffraction peaks of the precipitated phase, the characteristic peak position of the main precipitated phase in Comparative Example 1 is 45.105°. Based on the peak separation fitting and phase analysis results, its main precipitated phase can be attributed to the Ni₂Si phase. The characteristic peak positions of the main precipitated phases in Example 1 and Comparative Example 5 are 44.783° and 44.821°, respectively, both shifted towards lower angles relative to Comparative Example 1. Based on the precipitation behavior and theoretical inference, its main precipitated phase can be attributed to the (Ni,Co)₂Si type precipitated phase. This phenomenon indicates that the addition of cobalt participates in the formation of the precipitated phase and may enter the precipitated phase lattice by partially replacing nickel atoms, thereby causing changes in the precipitated phase lattice parameters and diffraction peak positions.

[0074] Compared to Comparative Examples 1 and 5, Example 1 exhibits a matrix lattice constant closer to that of the pure copper reference sample, while simultaneously forming a cobalt-containing (Ni,Co)₂Si type precipitate. These results demonstrate that the alloy composition and nonlinear dynamic cooling and solvent extraction process employed in Example 1 are beneficial for adjusting the distribution of solute atoms in the matrix and the structure of the precipitated phase, thus providing experimental evidence for reducing matrix lattice distortion and improving the matching relationship between the precipitated phase and the copper matrix.

[0075] Test Example 2: Strip materials at specific manufacturing stages were selected as test subjects. Example 1 was prepared as a slightly deformed intermediate strip after undergoing a strip bending and straightening process, and Example 1 was prepared as a final strip after the entire process. To isolate a single process variable, Comparative Example 2's intermediate strip (without pre-clustering, only undergoing bending and straightening) and Comparative Example 3's intermediate strip (after pre-clustering, room temperature cold rolling was used instead of bending and straightening) were prepared simultaneously.

[0076] The strips in the above four states were cut along the rolling direction and processed into long strips of standard size 30mm × 5mm. Mechanical shear burrs on the edges were removed using fine sandpaper, and the strips were ultrasonically cleaned in anhydrous ethanol for 10 minutes before drying for later use.

[0077] The specimens were tested using a dynamic thermomechanical analyzer (DMA). The test fixture was selected in tensile mode to accommodate the geometric characteristics of the thin strip material. The dynamic alternating load frequency was set to 1 Hz, the static pre-tension force was set to 0.5 N, and the dynamic strain amplitude was controlled at 15 μm to ensure that the specimen was in the purely elastic deformation range and to avoid secondary plastic deformation during the test.

[0078] A variable-temperature scanning program was performed under a protective atmosphere of high-purity nitrogen (flow rate of 50 mL / min). The initial test temperature was set at 50℃, and the temperature was increased to 350℃ at a constant heating rate of 3℃ / min. The system acquired the storage modulus and loss modulus in real time and calculated and output the internal loss value.

[0079] Extract characteristic data from the internal friction curve during the heating process. Based on the solid internal friction theory, identify and record the peak temperature (corresponding temperature position) of the Köster peak caused by dislocation dragging solute atoms and the peak intensity of the internal friction peak.

[0080] The experimental data are shown in Table 2: Table 2: Characteristic data of temperature-dependent internal friction test (Köster peak) for each sample Comparative Example 3 Intermediate state (pre-clustered + room temperature cold rolling) 193.5 9.41 Comparative Example 2 Intermediate state (no pre-clustering + medium-temperature stretching) 224.8 6.78 Example 1 Intermediate state (pre-clustered + medium-temperature stretching) 276.4 4.12 Example 1 Final state (end of process completed) 281.2 (Residual weak peak) 0.85 in conclusion: According to Table 2 and Figure 2 The data shows that the Köster peak of the intermediate strip (cold-rolled at room temperature) in Comparative Example 3 appears at a lower temperature (193.5℃), and the peak internal friction reaches 9.41 × 10⁻⁶. -3 In solid-state physics, the Köster peak characterizes the process by which dislocations drag solute atoms under applied alternating stress. The lower peak temperature in Comparative Example 3 indicates that the internal dislocations have lower slip activation energies, are more free, and have weaker binding forces from the solute atoms.

[0081] In Example 1, the Köster peak temperature of the intermediate strip shifted significantly to the high-temperature region to 276.4℃, while the peak internal friction decreased to 4.12 × 10⁻⁶. -3 The significant increase in peak temperature indicates that dislocation movement has encountered strong resistance, requiring a higher external thermal energy to overcome the local pinning effect. This change confirms that under the driving force of tensile tension at medium temperature (165℃), the micro-elastic expansion of the matrix lattice causes short-range drift of Co atoms with larger atomic radii, which then combine with previously formed Mg-P preclusters to construct a stable complex gas in the dislocation core region. This complex gas creates a high-intensity pinning effect on the dislocation lines.

[0082] Further comparison with the data from the intermediate-state tape in Example 2 revealed the absence of the low-temperature pre-clustering step, resulting in a Köster peak temperature drop to 224.8℃. This reflects that, without Mg-P clusters as early-stage segregation nuclei, while intermediate-temperature stretching can induce some solute atom segregation, it cannot form a sufficiently stable gas cluster structure, leading to a significant decrease in pinning energy. This demonstrates a sequential synergistic dependence between the two-step segregation mechanism of pre-clustering and stress coupling.

[0083] In Example 1, the internal friction curve of the final state strip tends to flatten out, the original Köster peak almost disappears, and the internal friction value drops to 0.85 × 10⁻⁶. -3This reflects the significant reduction in free dislocations and solid-solution atoms within the material after the all-hydrogen annealing process. The dislocation network, previously strongly pinned by gas clusters, did not statically recover and disappear during the rapid heating phase; instead, it directly became a favorable nucleation site for the newly generated precipitates. The dislocation lines were completely fixed by the nanoprecipitate entity, thus losing their ability to undergo microscopic slip and generate internal friction. This verifies the mechanism by which high-density defects are thermodynamically frozen and transformed into an effective strengthening phase.

[0084] Test Example 3: The final state alloy strips obtained in Examples 1, 4, and 6 were selected as test objects. Samples were cut from the strips and thinned along the thickness direction. Precision grinding and polishing equipment was used to thin both sides of the samples to a thickness of approximately 30 μm to facilitate X-ray penetration and obtain effective scattered signals. The thinned samples were cleaned in anhydrous ethanol, dried, and then used for testing.

[0085] The above samples were tested using a small-angle X-ray scattering instrument. The test source was Cu-Kα rays with a wavelength of 0.154 nm; the X-ray tube operating voltage was 40 kV and the operating current was 40 mA. During the test, the distance between the sample and the two-dimensional array detector was adjusted to 1500 mm to cover the required small-angle scattering vector range.

[0086] The sample was fixed on the sample stage, and the test chamber was kept under vacuum to reduce the influence of air scattering on the test results. Each sample underwent cumulative exposure acquisition, with a single exposure time of 1200s, to obtain a two-dimensional small-angle scattering spectrum.

[0087] The acquired two-dimensional scattering spectra were imported into analysis software for processing. Using omnidirectional angular integrals, the two-dimensional scattering spectra were converted into one-dimensional scattering intensity curves; subsequently, system background scattering and matrix background interference were subtracted. Based on the particle size distribution model, the distance distribution function was calculated using the indirect Fourier transform method, and the curve of precipitate volume distribution as a function of particle size was further inverted to obtain the curve. Characteristic parameters such as peak center particle size, full width at half maximum (FWHM), and main volume distribution intervals were extracted based on the inversion results.

[0088] The experimental data are shown in Table 3: Table 3: Test results of characteristic parameters of nano-precipitate size distribution inside each sample Example 1 7.35 3.12 5.21-9.48 Comparative Example 4 13.82 7.95 9.04-21.65 Comparative Example 6 17.18 9.84 11.63-28.72 in conclusion: According to Table 3 and Figure 3According to the data, the peak center particle size of the nano-precipitated phase in Example 1 was 7.35 nm, the main volume distribution range was 5.21 nm to 9.48 nm, and the full width at half maximum (FWHM) was 3.12 nm. This result indicates that the precipitated phase size in Example 1 is mainly distributed within a small particle size range, and the distribution width is relatively narrow. Therefore, it can be seen that under the heat treatment conditions used in Example 1, the growth of the precipitated phase is somewhat limited, and the size distribution of the nano-precipitated phase is relatively concentrated.

[0089] The peak center particle size of Comparative Example 4 was 13.82 nm, the main volume distribution range was 9.04 nm to 21.65 nm, and the full width at half maximum (FWHM) was 7.95 nm. Compared with Example 1, the peak particle size of the precipitated phase in Comparative Example 4 was increased, and the particle size distribution range was widened. This result indicates that under the heating regime used in Comparative Example 4, the precipitated phase may have undergone a longer nucleation, growth, or coarsening process, resulting in a larger final precipitated phase size and a higher degree of dispersion.

[0090] The peak center particle size of Comparative Example 6 was 17.18 nm, the main volume distribution range was 11.63 nm to 28.72 nm, and the full width at half maximum (FWHM) was 9.84 nm. Compared with Example 1 and Comparative Example 4, the average size of the precipitated phase in Comparative Example 6 was further increased, and the distribution range was further broadened. This result indicates that without rapid quenching and locking treatment, the material may undergo a relatively long diffusion process during subsequent cooling, causing some precipitated phases to continue to grow or coarsen, thereby forming precipitated phase particles with larger particle sizes.

[0091] In summary, the results show that Example 1 exhibits a smaller peak center particle size and a narrower particle size distribution range compared to Comparative Examples 4 and 6. This indicates that the nonlinear heat treatment process employed in Example 1, particularly the rapid cooling step at a specific temperature node, effectively limits the continued growth of the precipitated phase during subsequent thermal processes, keeping the precipitated phase within a smaller size range. These test results provide experimental evidence for the control of the nano-precipitated phase size distribution by the process of this invention.

[0092] Test Example 4: The final-state alloy strips prepared in Examples 1 to 4 and Comparative Examples 1 to 6 were selected as test objects. Samples were taken along the rolling direction of the strips and processed into rectangular cross-section metal sheet tensile specimens according to the requirements of GB / T228.1. At the same time, long strips with dimensions of approximately 1000 mm × 10 mm were cut as conductivity test specimens. Before testing, oil, oxide film and other adhering substances on the surface of the specimens were removed.

[0093] Tensile specimens were mounted on a universal testing machine, and an extensometer was used to measure strain. The test temperature was 25℃, and the tensile rate was 1 mm / min. Stress-strain curves were recorded during the test, and the stress value corresponding to 0.2% non-proportional elongation of the specimen was taken as the yield strength. Five samples were tested in parallel under each process condition, and the arithmetic mean was taken as the test result.

[0094] The electrical conductivity of the strip was tested using an eddy current conductivity meter. Before testing, the instrument was calibrated using a standard test block with known conductivity. During testing, the probe was placed vertically against the strip surface, and the test area was kept flat and clean. Ten measurements were taken at different locations for each sample, and the conductivity value expressed in terms of the International Standard for Annealed Copper (%IACS) was recorded. After removing obviously abnormal data, the arithmetic mean was taken as the conductivity test result for that sample.

[0095] The experimental data are shown in Table 4: Table 4: Test results of combined yield strength and electrical conductivity of each specimen

[0096] in conclusion: According to Table 4 and Figure 4 The data shows that the yield strength of Examples 1 to 4 ranges from 762.1 MPa to 808.6 MPa, and the electrical conductivity ranges from 45.8% IACS to 48.2% IACS, indicating that each example maintains good electrical conductivity at a high strength level.

[0097] The yield strength and electrical conductivity of Comparative Example 1 were lower than those of the embodiments, indicating that the strengthening effect and electrical conductivity of the material were affected to some extent in the absence of corresponding trace element adjustment. The electrical conductivity of Comparative Examples 2 and 3 were close to those of some embodiments, but the yield strength was significantly reduced, indicating that the combination of low-temperature pre-clustering treatment and medium-temperature stress treatment has a certain effect on improving strength.

[0098] Comparative Example 4 exhibits lower yield strength and electrical conductivity than the embodiments, indicating that the heating regime affects the final overall performance, considering its heat treatment conditions. Comparative Example 5 shows a yield strength close to some embodiments, but its electrical conductivity is significantly lower, suggesting that conventional isothermal treatment cannot balance strength and electrical conductivity. Comparative Example 6 has higher electrical conductivity, but a decreased yield strength, indicating that eliminating the rapid cooling step results in relatively insufficient strength retention.

[0099] In summary, Examples 1 to 4 exhibit a good match between yield strength and electrical conductivity. This result demonstrates that the combination of alloy composition and heat treatment process used in this invention is beneficial for obtaining copper alloy strips with both high strength and high electrical conductivity.

[0100] Test Example 5: The final alloy strips prepared in Examples 1 to 4, Comparative Example 1, and Comparative Example 6 were selected as test objects. Long strip samples were cut parallel to the rolling direction, and the edges of the samples were mechanically ground to reduce the influence of edge burrs or microcracks on the test results.

[0101] The stress relaxation test system was equipped with an ambient heating chamber. Based on the room temperature yield strength Rp0.2 of each specimen obtained from the tensile test, and combined with the formula for calculating the deflection of an elastic cantilever beam, the initial loading displacement required to make the maximum surface bending stress of the specimen reach 80% of its yield strength was determined.

[0102] During testing, one end of the sample was fixed in a rigid fixture, and a predetermined displacement was applied to the free end of the sample using a loading probe to bring the sample to a set initial stress state. The testing device was then placed in a constant temperature aging chamber preheated to 150°C for heat preservation testing, with a total testing time of 1000 hours.

[0103] During the test, the change of the reaction force at the free end of the specimen over time was recorded using a load sensor. Load data were extracted at 100h, 500h, and 1000h of heat preservation, and the load values ​​at each time point were compared with the initial load values ​​to calculate the corresponding percentage of residual stress.

[0104] The experimental data are shown in Table 5: Table 5: Test results of high-temperature stress relaxation performance of each sample at 150℃

[0105] in conclusion: According to Table 5 and Figure 5 According to the data, after being kept at 150°C for 1000 hours, the residual stress of Examples 1 to 4 remained between 90.72% and 92.14%, indicating that each example has good stress retention capability under long-term high-temperature loading conditions.

[0106] Comparative Example 1 exhibited a residual stress of 68.21% after 1000 hours, lower than that of the other examples. Considering its compositional design, this result indicates that the lack of appropriate trace element adjustments may affect the dislocation stability, precipitate strengthening effect, or matrix structure stability of the material, making it more susceptible to stress decay under high-temperature loading conditions.

[0107] Comparative Example 6 exhibited a residual stress of 72.84% after 1000 hours, which was also lower than that of the other examples. Considering the aforementioned precipitate size distribution test results, this difference may be related to changes in the size and distribution of the precipitates after the rapid cooling step at specific temperature nodes was removed. Larger precipitate sizes or wider grain size distributions may reduce their confinement effect on dislocation movement, thereby affecting the long-term stress retention performance of the material.

[0108] In summary, Examples 1 to 4 exhibited high residual stress retention rates at 150°C for 1000 hours. This demonstrates that the combination of alloy composition and heat treatment process employed in this invention is beneficial for improving the stress relaxation resistance of copper alloy strips under high-temperature service conditions.

[0109] Test Example 6: The final alloy strips obtained in Examples 1 to 4, Comparative Examples 2, 4, and 6 were selected as test objects. Long strip-shaped bending specimens were cut along both the strip rolling direction and perpendicular to the strip rolling direction. Specimens sampled parallel to the rolling direction were designated as longitudinal, and those sampled perpendicular to the rolling direction were designated as transverse. The edges of the specimens were finely chamfered and polished with sandpaper to remove burrs formed during mechanical shearing and reduce the influence of edge defects on the bending test results.

[0110] The test was conducted using a precision CNC bending testing machine. The test followed GB / T232, the standard procedure for bending of metallic materials, with the punch pressing rate set to 2 mm / min. During the test, a set of bending punches with different fillet radii were used to induce a 90° bending deformation in the specimen. The bending radius was adjusted progressively from large to small until visible cracks appeared on the tensile side of the specimen.

[0111] After each 90° bend, the specimen is unloaded and removed, and its outer bent surface is observed using a 10x to 20x stereomicroscope. If no continuous cracks are observed on the tensile side of the specimen, the specimen is deemed acceptable at that bending radius; if continuous cracks or obvious cracking are observed, the specimen is deemed unacceptable at that bending radius. The minimum acceptable bending radius R for each specimen under the condition of no continuous cracks is recorded, and this value is divided by the actual thickness t of the specimen to obtain the corresponding R / t value. The smaller the R / t value, the better the bending performance of the specimen, indicating that it can bend at a smaller bending radius without producing visible cracks.

[0112] The experimental data are shown in Table 6: Table 6: Test Results of Minimum R / t Value of Strip Material at 90 Degree Bending for Each Sample Example 1 0.1 0.3 Example 2 0.2 0.4 Example 3 0.1 0.3 Example 4 0.2 0.5 Comparative Example 2 0.6 1.3 Comparative Example 4 1.3 2.6 Comparative Example 6 0.8 1.5 in conclusion: According to Table 6 and Figure 6The data show that the minimum longitudinal R / t values ​​of Examples 1 to 4 are 0.1 to 0.2, and the minimum transverse R / t values ​​are 0.3 to 0.5, both lower than those of Comparative Examples 2, 4, and 6. This result indicates that Examples 1 to 4 exhibit good resistance to bending cracking under 90° bending conditions, and can maintain the bending surface without cracking at a relatively small bending radius.

[0113] The minimum longitudinal R / t value of Comparative Example 2 was 0.6, and the minimum transverse R / t value was 1.3; the minimum longitudinal R / t value of Comparative Example 6 was 0.8, and the minimum transverse R / t value was 1.5, both higher than those of the Examples. This result indicates that the bending forming performance of the strip decreases when the low-temperature pre-clustering treatment is omitted or the rapid cooling step at specific temperature nodes is eliminated. Combined with the aforementioned microstructure and performance test results, this difference may be related to changes in the size and distribution of precipitated phases, the state of residual dislocations, the microstructure stability near grain boundaries, and changes in local strain concentration behavior.

[0114] Comparative Example 4 showed the highest minimum longitudinal R / t value (1.3) and minimum transverse R / t value (2.6) among the tested samples. This result indicates that the strip exhibits relatively poor bending formability under heat treatment conditions with lower heating rates. This may be due to the enhanced recovery process resulting from the lower heating rate during heat treatment, the premature nucleation and coarsening of precipitated phases, and the decreased uniformity of the microstructure near grain boundaries, which makes local strain concentration easier during bending deformation.

[0115] In summary, Examples 1 to 4 all exhibited low minimum R / t values ​​under both longitudinal and transverse bending conditions. This indicates that the alloy composition and heat treatment process combination employed in this invention is beneficial for improving the bending and forming properties of the strip and enhancing its crack resistance during subsequent bending processes.

[0116] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A high-elasticity, high-conductivity alloy material for connector terminals, characterized in that, The alloy material is made from raw materials comprising the following parts by weight: Cathode copper: 895.75–930.5 parts; 18-22 parts of electrolytic nickel; 4-6 parts of electrolytic cobalt; 40-60 parts of copper-silicon master alloy; 5-10 parts of copper-magnesium master alloy; 2.5 to 6.25 parts of copper-phosphorus master alloy.

2. The high-elasticity, high-conductivity alloy material for connector terminals according to claim 1, characterized in that, The copper-silicon master alloy is a binary alloy ingot containing alloyed elemental silicon with a silicon mass fraction of 10 wt%, and the copper-magnesium master alloy is a binary alloy ingot containing alloyed elemental magnesium with a magnesium mass fraction of 10 wt%.

3. The high-elasticity, high-conductivity alloy material for connector terminals according to claim 1, characterized in that, The copper-phosphorus master alloy is a binary alloy ingot containing alloyed elemental phosphorus with a phosphorus mass fraction of 8 wt%.

4. A method for preparing a high-elasticity, high-conductivity alloy material for connector terminals according to any one of claims 1-3, characterized in that, Includes the following steps: S1. The raw materials in the specified proportion are added to the melting furnace and melted, and cast into alloy ingots. The alloy ingots are homogenized and hot rolled into strips in multiple passes. After online water quenching and surface descaling, they are sent to a continuous annealing furnace for heat preservation and annealing. After exiting the furnace, they are forced to quench and cool to room temperature to obtain a basic strip with a single supersaturated solid solution phase structure. S2. The base strip is subjected to multi-pass cold rolling deformation, followed by low-temperature pre-clustering treatment under a protective atmosphere; S3. Apply pure tensile stress to the treated strip under heating conditions to induce micro-plastic elongation deformation. S4. The strip roll that has undergone micro-deformation is loaded into an annealing furnace with hydrogen as the working atmosphere and continuously subjected to in-situ nonlinear heat treatment. First, it is heated to a constant temperature at a heating rate and held at that temperature. Then, it is slowly cooled with the furnace at a constant cooling rate. When it drops to the specified temperature, it is rapidly cooled to room temperature by forced convection of gas.

5. The method for preparing the high-elasticity and high-conductivity alloy material for connector terminals according to claim 4, characterized in that, In step S1, the specific smelting process is as follows: cathode copper, electrolytic nickel, and electrolytic cobalt are added to a vacuum induction furnace, and smelted at 5.0 × 10⁻⁶ ℃. -2 Heated to 1150–1200°C under a vacuum of Pa to completely melt it; Argon gas was introduced to atmospheric pressure, followed by the sequential addition of copper-silicon master alloy, copper-magnesium master alloy, and copper-phosphorus master alloy. After electromagnetic stirring was turned on for 10–15 minutes, the alloy ingot was formed by vertical semi-continuous casting at a casting temperature of 1120–1160°C.

6. The method for preparing the high-elasticity and high-conductivity alloy material for connector terminals according to claim 4, characterized in that, In step S1, the specific heat treatment process for the base strip is as follows: The alloy ingot is placed in a heating furnace and heated to 900-950°C for 2-4 hours. After being hot rolled in multiple passes using a hot rolling mill, it is immediately water-quenched online. After milling the surface of the hot-rolled strip to remove the oxide scale, it is sent into a continuous annealing furnace with a protective atmosphere and held at 900-920℃ for 1-1.5 hours. After exiting the furnace, it is subjected to single-layer forced quenching and cooling through a water drain.

7. The method for preparing the high-elasticity, high-conductivity alloy material for connector terminals according to claim 4, characterized in that, In step S2, the total thickness reduction of the multi-pass cold rolling is controlled to be 60% to 65%, and the rolling oil temperature is controlled to be ≤40℃ during the rolling process. The process of low-temperature pre-clustering treatment is as follows: the cold-rolled strip is placed in a bell-shaped annealing furnace under nitrogen protection, the furnace temperature is set to 250-280℃, and the temperature is kept constant for 2-2.5 hours. Then it is taken out of the furnace and air-cooled to room temperature.

8. The method for preparing the high-elasticity and high-conductivity alloy material for connector terminals according to claim 4, characterized in that, The specific implementation method of step S3 is as follows: The pretreated strip is passed through a strip bending straightener with an online electromagnetic induction heating section. The actual temperature of the strip in the straightening area is controlled at 150-180°C. Pure tensile stress is applied to the strip through the front and rear tension rollers and bending roller group of the bending straightener, causing the strip to produce a micro-plastic elongation deformation of 1.5%-2.5%.

9. The method for preparing the high-elasticity, high-conductivity alloy material for connector terminals according to claim 4, characterized in that, In step S4, the purity of the hydrogen gas is ≥99.999 vol%. The heating and holding process for the in-situ nonlinear heat treatment is as follows: Turn on the strong convection circulating fan and start the heating power to raise the temperature of the coil center to 480℃ at a rate of 12-20℃ / min; keep it at 480℃ for 1-2 hours.

10. The method for preparing the high-elasticity, high-conductivity alloy material for connector terminals according to claim 4, characterized in that, In step S4, the cooling and rapid quenching process of the in-situ nonlinear heat treatment is as follows: After turning off the main heating source, adjust the auxiliary heating power and circulating air volume to make the coil temperature cool down linearly and slowly with the furnace at a constant cooling rate of 10-15℃ / h. When the thermocouple reading shows that the center of the roll has dropped to 350°C, immediately start the water-cooled heat exchanger and keep the hydrogen circulating in the furnace. Use the hydrogen for forced convection purging to rapidly cool the roll to below 100°C within 20 to 30 minutes.

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