Cu-15ni-8sn alloy and method for producing the same

Abstract

The application provides a Cu-15Ni-8Sn alloy and a preparation method thereof. The preparation method comprises the following steps: ingredient preparation, smelting and casting according to the alloy component ratio, homogenization treatment, hot rolling, solid solution and water quenching, first-stage cold rolling, first-stage aging, second-stage cold rolling, second-stage aging of the solid solution alloy in sequence, wherein the deformation of the first-stage cold rolling is greater than that of the second-stage cold rolling, and the temperature of the first-stage aging is higher than that of the second-stage aging, and finally, tension annealing is carried out to obtain the Cu-15Ni-8Sn alloy. The application induces high proportion Σ3 n special grain boundaries through multi-stage deformation heat treatment, and then uses grain boundary engineering to inhibit discontinuous precipitation and improve high-temperature microstructure stability, and cooperates with Co-V-Zr-Ce quaternary micro-alloying to further strengthen the matrix, so that a Cu-15Ni-8Sn alloy material with high proportion of special grain boundaries and low stress relaxation rate is obtained.

Description

Technical Field

[0001] This invention relates to the field of alloy materials technology, and more specifically, to a Cu-15Ni-8Sn alloy and its preparation method. Background Technology

[0002] Cu-15Ni-8Sn alloys possess superior properties such as high strength, electrical conductivity, wear resistance, high elasticity, and high corrosion resistance. Their high elasticity makes them stand out in the field of elastic components, and their strength is comparable to beryllium bronze. Furthermore, they exhibit excellent resistance to stress relaxation at high temperatures, ensuring minimal deformation of components during prolonged high-temperature use. Based on their excellent elasticity and comprehensive properties, Cu-15Ni-8Sn alloys are considered ideal elastic materials for connectors in high-temperature electrical connectors. However, Cu-15Ni-8Sn alloy ingots have significant defects. Due to their high Sn content and low melting point, macroscopic segregation and dendritic segregation, resulting in uneven composition distribution, are easily generated during the casting process. This segregation severely affects the performance stability and consistency of the alloy.

[0003] To address these issues, existing technologies primarily employ microalloying for control. Reported microalloying elements include Si, Ti, Cr, Zr, Nb, Al, Mn, Fe, P, Ag, and rare earth elements (La, Y). For example, Chinese patent application CN114086027A discloses a Cu-Ni-Sn alloy resistant to high-temperature softening. This alloy utilizes the addition of elements such as Al, Zn, Si, and Nb, combined with a combined deformation heat treatment process involving continuous casting, primary cold deformation (70-90%), primary aging, secondary cold deformation (50-70%), and secondary aging. The aim is to suppress discontinuous precipitation and improve the softening resistance temperature. For example, Chinese patent application CN115896537A discloses a high-strength, corrosion-resistant Cu-Ni-Sn alloy, which incorporates elements such as Al, Zn, Cr, Nb, and Zr, and employs a multi-stage homogenization treatment and a process route of hot rolling pre-aging, single cold rolling (70-80%), aging, and secondary cold rolling (20-30%) to improve the alloy's strength and corrosion resistance. Furthermore, Chinese patent application CN119351914A discloses a preparation process for a Cu-Ni-Sn alloy, which improves the alloy's comprehensive mechanical properties (strength-ductility product) by introducing recrystallization heat treatment and small-deformation secondary cold rolling (≤45%). Chinese patent application CN121228045A discloses alloys with added Si and Zr; Chinese patent application CN121228046A discloses alloys with added Al and Zr; and Chinese patent application CN114457256B discloses alloys with added Ag, La, and Er. In addition, although Ti, a common strengthening element in the Cu-Ni-Sn system, can form Ni3Ti phase to produce precipitation strengthening, the addition of Ti can easily reduce the conductivity of the alloy and narrow the aging process window, resulting in a decrease in the uniformity of the microstructure.

[0004] Currently, existing technologies, including the aforementioned solutions, still have the following problems: First, insufficient attention is paid to the regulation of grain boundary characteristic distribution; existing technologies generally neglect special grain boundaries (especially Σ3). n The role of twin boundaries in high-temperature stress relaxation resistance has been limited by the failure to actively increase the proportion of special grain boundaries through processing methods, thus restricting further improvement in relaxation resistance. Secondly, there is still room for improvement in stress relaxation resistance. Existing alloys lack microstructural stability under long-term service at higher temperatures (≥150℃), and grain boundaries easily become weakening zones and crack initiation sources, making it difficult to meet the requirements of extreme working conditions. For example, the stress relaxation rate of the traditional Cu-15Ni-8Sn alloy at 150℃ / 50h is as high as 8.51%, and even after microalloying improvements, there is still room for further reduction. Thirdly, the combination of microalloying elements tends to be similar, and the synergistic effect is insufficient. Existing patents are mostly concentrated on elements such as Si, Ti, Cr, Zr, and Al, failing to fully explore the deep-level physical metallurgical synergistic effect among multiple elements. Summary of the Invention

[0005] The main objective of this invention is to provide a Cu-15Ni-8Sn alloy and its preparation method, so as to solve the problem of insufficient stress relaxation resistance of Cu-15Ni-8Sn alloy in the prior art.

[0006] To achieve the above objectives, according to one aspect of the present invention, a method for preparing a Cu-15Ni-8Sn alloy is provided, comprising, by weight percentage: Ni 14.5~15.5%, Sn 7.5~8.5%, Co 0.1~1.0%, V 0.05~0.5%, Zr 0.02~0.2%, Ce The alloy composition is 0.01~0.15%, with the balance being Cu and unavoidable impurities. The preparation method includes the following steps: Step S1, batching, melting and casting according to the alloy composition ratio to obtain an ingot; homogenizing the ingot to obtain a homogenized ingot; Step S2, hot rolling the homogenized ingot to obtain a hot-rolled alloy; solution-cooling and water-quenching the hot-rolled alloy sequentially to obtain a solution-treated alloy; Step S3, cold-rolling the solution-treated alloy sequentially through first-stage cold rolling, first-stage aging, second-stage cold rolling, and second-stage aging to obtain an aging alloy; the deformation amount of the first-stage cold rolling is 30~55%, the temperature of the first-stage aging is 400~450℃, the deformation amount of the second-stage cold rolling is 10~35%, the temperature of the second-stage aging is 350~420℃, and the deformation amount of the first-stage cold rolling is greater than that of the second-stage cold rolling, and the temperature of the first-stage aging is greater than that of the second-stage aging; Step S4, tension annealing the aging alloy to obtain a Cu-15Ni-8Sn alloy.

[0007] Furthermore, in step S3, the holding time for the first-stage aging is 0.5 to 3 hours; and / or the holding time for the second-stage aging is 1 to 4 hours.

[0008] Further, in step S3, the deformation amount of the first-stage cold rolling is 35~45%; and / or the temperature of the first-stage aging is 410~430℃, and the holding time is 0.5~1.5h; and / or the deformation amount of the second-stage cold rolling is 15~25%; and / or the temperature of the second-stage aging is 370~390℃, and the holding time is 1~2h.

[0009] Furthermore, in step S3, the deformation amount of the first-stage cold rolling is greater than that of the second-stage cold rolling, with a difference of 10-30%; and / or the aging temperature of the first stage is greater than that of the aging temperature of the second stage, with a difference of 20-60℃.

[0010] Further, in step S1, the homogenization treatment temperature is 850~890℃, the holding time is 8~10h, and the atmosphere is argon; and / or after homogenization treatment, the obtained material is cooled in the furnace to 450~550℃ and then air-cooled to obtain a homogenized ingot.

[0011] Further, in step S2, the initial rolling temperature of hot rolling is 840~860℃, the holding time is 1~2h, and the final rolling temperature is 750~800℃; and / or the cumulative deformation of hot rolling is ≥85%; and / or the solution temperature is 840~890℃, the holding time is 2~6h; and / or the cooling rate of water quenching is 30~60℃ / s.

[0012] Furthermore, the solution treatment temperature is 850~870℃, and the holding time is 3~5h; and / or the tension annealing temperature is 300~400℃, the holding time is 1~2h, and the applied tension is 5~20MPa.

[0013] Furthermore, by weight percentage, the Cu-15Ni-8Sn alloy comprises Ni 14.5~15.5%, Sn 7.5~8.5%, Co 0.3~0.8%, V 0.1~0.4%, Zr 0.05~0.2%, Ce 0.02~0.1%, with the balance being Cu and unavoidable impurities; and / or 0.3%≤Co+V+Zr+Ce≤1.8%.

[0014] Furthermore, in the Cu-15Ni-8Sn alloy, Σ3 n The proportion of grain boundaries to the total grain boundary length is ≥60%; and / or in Cu-15Ni-8Sn alloys, the volume fraction of discontinuous precipitates is ≤15%.

[0015] According to another aspect of the present invention, a Cu-15Ni-8Sn alloy is provided, which is obtained by the preparation method of the Cu-15Ni-8Sn alloy described above according to the present invention.

[0016] This invention provides a method for preparing a Cu-15Ni-8Sn alloy with high resistance to stress relaxation based on grain boundary engineering, which induces a high proportion of Σ3 through multi-stage deformation heat treatment. n Special grain boundaries, multi-stage processes (first-stage cold rolling, first-stage aging, second-stage cold rolling, second-stage aging) break dislocation entanglement through intermediate aging, promoting dislocation recombination and grain boundary migration, thus creating Σ3 n The nucleation and growth of twin boundaries provide favorable conditions, which in turn allows for the use of grain boundary engineering to suppress discontinuous precipitation and improve the stability of the microstructure at high temperatures. Simultaneously, this invention, in conjunction with Co-V-Zr-Ce quaternary microalloying, further strengthens the matrix, jointly addressing the problems of as-cast segregation, discontinuous precipitation, and insufficient stress relaxation resistance in Cu-15Ni-8Sn alloys. Ultimately, this results in a Cu-15Ni-8Sn alloy material with a high proportion of special grain boundaries and a low stress relaxation rate. Attached Figure Description

[0017] The accompanying drawings, which form part of this application, are used to provide a further understanding of the invention. The illustrative embodiments of the invention and their descriptions are used to explain the invention and do not constitute an undue limitation of the invention. In the drawings:

[0018] Figure 1 The stress relaxation curve according to Embodiment 1 of the present invention is shown;

[0019] Figure 2 The stress relaxation curves of Comparative Example 3 according to the present invention are shown;

[0020] Figure 3 The stress relaxation curves of Comparative Example 4 according to the present invention are shown. Detailed Implementation

[0021] It should be noted that, unless otherwise specified, the embodiments and features described in this application can be combined with each other. The present invention will now be described in detail with reference to the accompanying drawings and embodiments.

[0022] Terminology Explanation:

[0023] Macroscopic segregation: The phenomenon of uneven distribution of alloying elements along the cross section on a macroscopic scale of an ingot, manifested as local enrichment or depletion of components, is usually caused by solute redistribution during solidification.

[0024] Dendritic segregation, also known as microsegregation or intracrystalline segregation, refers to the compositional differences that occur in the dendrite branches and interdendritic regions during solidification due to the different solubilities of solute atoms in the solid and liquid phases.

[0025] Discontinuous precipitation (DP): a cellular precipitation morphology that nucleates at grain boundaries and grows into the grain. It is usually driven by grain boundary migration and leads to a significant decrease in the strength and plasticity of the alloy. It is a common aging failure mechanism of Cu-Ni-Sn alloys.

[0026] Σ3 n Special grain boundaries: refers to grain boundaries whose coincidence density, conforming to the Coincidence Site Lattice (CSL) model, is inversely 3. n The grain boundaries, where n is an integer. Where Σ3 n Grain boundaries typically correspond to twin boundaries, and they have low interfacial energy and high resistance to intergranular failure, playing an important role in improving the high-temperature stability of alloys.

[0027] Twin boundary: The interface between two parts of a crystal that are coherently connected by a twin relationship. It belongs to the low-energy grain boundary and can be induced to form during deformation heat treatment. It helps to hinder dislocation movement and grain boundary sliding.

[0028] Grain Boundary Character Distribution (GBCD): This involves controlling the distribution ratio of different grain boundary types in polycrystalline materials through appropriate deformation and heat treatment processes, particularly enhancing the distribution of low-energy special grain boundaries (such as Σ3). n The method of improving material properties by adjusting the ratio of )

[0029] Stress relaxation rate: The degree to which the internal stress of a material gradually decreases over time under constant strain conditions. It is usually expressed as the percentage of the difference between the initial stress and the residual stress to the initial stress. It is a key indicator for evaluating the high-temperature service stability of elastic elements.

[0030] As described in the background section of this invention, the existing technology suffers from insufficient stress relaxation resistance of Cu-15Ni-8Sn alloys. To address this problem, in a typical embodiment of this invention, a method for preparing a Cu-15Ni-8Sn alloy is provided, comprising, by weight percentage: Ni 14.5~15.5%, Sn 7.5~8.5%, Co 0.1~1.0%, V 0.05~0.5%, Zr 0.02~0.2%, and Ce. The alloy composition is 0.01~0.15%, with the balance being Cu and unavoidable impurities. The preparation method includes the following steps: Step S1, batching, melting and casting according to the alloy composition ratio to obtain an ingot; homogenizing the ingot to obtain a homogenized ingot; Step S2, hot rolling the homogenized ingot to obtain a hot-rolled alloy; solution-cooling and water-quenching the hot-rolled alloy sequentially to obtain a solution-treated alloy; Step S3, cold-rolling the solution-treated alloy sequentially through first-stage cold rolling, first-stage aging, second-stage cold rolling, and second-stage aging to obtain an aging alloy; the deformation amount of the first-stage cold rolling is 30~55%, the temperature of the first-stage aging is 400~450℃, the deformation amount of the second-stage cold rolling is 10~35%, the temperature of the second-stage aging is 350~420℃, and the deformation amount of the first-stage cold rolling is greater than that of the second-stage cold rolling, and the temperature of the first-stage aging is greater than that of the second-stage aging; Step S4, tension annealing the aging alloy to obtain a Cu-15Ni-8Sn alloy.

[0031] Existing technologies mostly focus on improving strength through precipitation strengthening and grain refinement, and on controlling the grain boundary characteristic distribution (GBCD), especially for Σ3. n The increase in the proportion of special grain boundaries and its effect on high-temperature stress relaxation resistance have rarely been studied. Research indicates that Σ3 n Special grain boundaries (such as twin boundaries) possess low interfacial energy and excellent resistance to intergranular failure, effectively hindering grain boundary sliding and dislocation movement at high temperatures. They are crucial microstructural factors for improving the stress relaxation resistance of alloys. However, there are no reports in the current technology of actively incorporating grain boundary engineering into the stress relaxation control of Cu-15Ni-8Sn alloys, and even less research on precisely controlling Σ3 through multi-stage deformation heat treatment. n A special grain boundary ratio process is needed. Therefore, developing a new process that can synergistically control microalloying and grain boundary characteristic distribution is of great significance for obtaining Cu-15Ni-8Sn alloys with ultra-high stress relaxation resistance.

[0032] This invention first involves batching, smelting, and casting according to the alloy composition ratio to obtain an ingot. The ingot is then homogenized to eliminate dendritic segregation and obtain a homogeneous equiaxed crystal structure, resulting in a homogenized ingot. Subsequently, the homogenized ingot is hot-rolled to break up the as-cast structure and provide good plasticity for further processing, yielding a hot-rolled alloy. The hot-rolled alloy is then sequentially solution-treated and water-quenched to allow the alloying elements to fully dissolve in the copper matrix, preparing for subsequent aging precipitation strengthening, resulting in a solution-treated alloy.

[0033] The solid solution alloy is then subjected to a series of processes: primary cold rolling, primary aging, secondary cold rolling, and secondary aging. The primary cold rolling involves a large deformation to introduce a suitable amount of uniform dislocations, providing nucleation sites for subsequent precipitates and simultaneously supporting Σ3... n The formation of twin boundaries accumulates strain energy. After high-temperature aging treatment of the cold-rolled sheet, Co atoms begin to diffuse, altering vacancy concentration and diffusion kinetics. Initial precipitation of V and Zr nucleates on dislocation lines, while partial recrystallization and grain boundary migration occur, inducing Σ32... n Formation of twin boundaries. Cold rolling of the sheet material after primary aging, with a small deformation, breaks the existing dislocation network, introduces new dislocations, and makes the dislocation distribution more uniform. Simultaneously, it activates more Zr and V atoms to accumulate on the dislocations, further promoting the evolution of special grain boundaries. The sheet material after secondary cold rolling undergoes low-temperature aging treatment to promote the full precipitation of the Co-V-Zr composite nanophase. Utilizing the vacancy consumption effect of Co and the grain boundary pinning effect of Zr, the already formed Σ32 is stabilized. n Special grain boundaries are used to ultimately construct a stable microstructure with high dislocation density, multiphase dispersion, and a high proportion of special grain boundaries, resulting in an aging alloy. Finally, the aging alloy is subjected to tension annealing. Through thermo-mechanical coupling, the residual stress introduced by the previous cold rolling process is further eliminated, preventing stress concentration from becoming a potential relaxation source. This further stabilizes the dislocation configuration inside the alloy and inhibits high-temperature migration of grain boundaries, resulting in a Cu-15Ni-8Sn alloy.

[0034] This invention first introduces grain boundary engineering control through multi-stage deformation heat treatment. Appropriate cold deformation combined with intermediate aging treatment can induce low-energy special grain boundaries (especially Σ3). n The formation of twin boundaries. The multi-stage process of this invention breaks dislocation tangles through intermediate aging, promoting dislocation recombination and grain boundary migration, thus contributing to the formation of Σ3. n The nucleation and growth of twin boundaries provide favorable conditions. This invention not only focuses on precipitation strengthening, but also uses grain boundary engineering as the main control method in Cu-15Ni-8Sn alloys, ultimately obtaining a high proportion of Σ3 in the alloy. nSpecial grain boundaries. These grain boundaries have low interfacial energy and good resistance to intergranular failure. They can effectively hinder grain boundary sliding and dislocation movement at high temperatures, significantly suppress the nucleation and growth of discontinuous precipitates, and thus greatly improve the stress relaxation resistance of the alloy.

[0035] Based on the above-mentioned grain boundary engineering process, this invention further introduces a Co-V-Zr-Ce quaternary microalloying system. By changing the vacancy concentration and diffusion kinetics (Co), adjusting the precipitation kinetics to suppress coarsening (V), suppressing grain boundary pinning (Zr), and purifying the melt and refining the grains (Ce), a four-in-one strengthening process is achieved, realizing comprehensive and multi-scale coordinated control from intragranular dislocations to grain boundaries and precipitates. Based on the above-mentioned grain boundary engineering control, the yield strength and high-temperature microstructure stability of the alloy are further improved.

[0036] Specifically, in the composition of the Cu-15Ni-8Sn alloy, Co (cobalt) is the core element for resisting stress relaxation. Co atoms can effectively consume vacancies in the matrix, reduce the atomic diffusion rate at high temperatures, and generate a strong pinning effect through their interaction with dislocations, significantly inhibiting dislocation climb and slip, thus fundamentally delaying the occurrence of stress relaxation. Vanadium (vanadium) can regulate precipitation kinetics and inhibit precipitation coarsening. Vanadium can change the growth mode of amplitude-modulated decomposition products, promote the formation of finer, more dispersed precipitates, and improve the alloy's resistance to softening and its structural stability. Zirconium (zirconium) can form phases such as Ni4SnZr, acting as a modifier to refine grains and occupy discontinuous precipitation sites at grain boundaries. Zr atoms have a strong inhibitory effect on dislocation climb (DP), effectively suppressing the formation and growth of DP. Ce (cerium), as a rare earth element, has a strong purifying and microalloying effect. It can react with impurity elements such as O and S in the melt, purify the grain boundaries, and improve the grain boundary bonding force. At the same time, the addition of trace amounts of Ce can further refine the as-cast grains and improve the corrosion resistance and thermal stability of the alloy.

[0037] The preparation method of this invention is a process of Co-V-Zr-Ce quaternary microalloying combined with multi-stage cold rolling and graded aging grain boundary engineering. For the first time, grain boundary engineering has significantly increased the proportion of special grain boundaries in Cu-15Ni-8Sn alloys, and established a correlation between grain boundary characteristic distribution and stress relaxation resistance. Furthermore, while existing technologies consider grain boundary engineering to mainly rely on stacking fault energy and thermodynamic conditions to induce twin boundaries, this invention actively introduces the alteration of vacancy concentration and diffusion kinetics of Co, as well as the grain boundary pinning effect of Zr, into the grain boundary engineering regulation. By altering vacancy concentration and diffusion kinetics, the recrystallization grain boundary migration is delayed, and Zr segregation is used to stabilize special grain boundaries, resulting in a higher proportion of Σ3. n Special grain boundaries and ultra-high resistance to stress relaxation.

[0038] In summary, this invention introduces grain boundary engineering, which significantly improves stress relaxation resistance, and induces a high proportion of Σ3 through multi-stage deformation heat treatment. n Using special grain boundaries as the core technology, the distribution of grain boundary characteristics is actively controlled through a series of processes including primary cold rolling, primary aging, secondary cold rolling, and secondary aging, resulting in Σ3 n The proportion of special grain boundaries is significantly increased. Simultaneously, the combination of Co vacancy consumption and dislocation pinning effects, V precipitation kinetic regulation, Zr grain boundary strengthening, and Ce purification effect greatly suppresses dislocation movement and grain boundary sliding in the alloy at high temperatures. Furthermore, the alloy exhibits excellent grain boundary stability, with a high proportion of Σ3 formed through multi-stage deformation heat treatment. n Special grain boundaries possess low interfacial energy and resistance to intergranular failure. Combined with the segregation and pinning of grain boundaries by Zr and the purifying effect of Ce, this significantly inhibits grain boundary migration and the nucleation and growth of discontinuous precipitation at high temperatures, ensuring the long-term stability of the grain boundary region. Furthermore, the use of a Co-V-Zr-Ce quaternary microalloying system achieves ultra-high resistance to stress relaxation while maintaining a good balance between strength and ductility.

[0039] The smelting and casting processes can employ conventional methods in the art. For example, according to the component ratio, electrolytic copper, electrolytic nickel, pure tin, Cu-Co master alloy, Cu-V master alloy, Cu-Zr master alloy, and Cu-Ce master alloy are used as raw materials for vacuum smelting and casting into ingots. A power gradient heating method is used, first baking the material at a lower power for 3-5 minutes, then increasing the power to 1300℃±10℃, holding at that temperature for 8-10 minutes, and then casting at 1200℃±10℃ to obtain the ingot.

[0040] In a preferred embodiment, in step S3, the holding time for the first-stage aging is 0.5~3 hours; and / or the holding time for the second-stage aging is 1~4 hours. These conditions facilitate the introduction of high-density and uniformly distributed dislocation entanglements within the grains during the large-deformation first-stage cold rolling process, providing sufficient nucleation sites for subsequent precipitates, while also providing Σ3 n The formation of twin boundaries requires the accumulation of necessary strain energy. Subsequently, during first-order aging, Co atoms reduce the atomic diffusion rate by consuming vacancies, and V and Zr undergo initial precipitation at dislocation tangles, forming nanoscale clusters or GP regions. Simultaneously, the accumulated strain energy drives partial grain boundary migration, inducing Σ3+ in the previously dislocation tangled regions. nThe nucleation and growth of twin boundaries form a preliminary special grain boundary network. During the second-stage cold rolling of the alloy after first-stage aging, a smaller deformation is used. This breaks the recovered dislocation network and introduces new mobile dislocations, making the dislocation distribution more uniform. Furthermore, the newly introduced dislocations provide additional enrichment sites for V and Zr atoms, activating more solute atoms to segregate at dislocations and grain boundaries, further promoting the evolution and expansion of special grain boundaries during the second-stage aging process. Finally, during the second-stage aging process, the remaining Co, V, and Zr atoms are fully desoluble from the supersaturated matrix, precipitating fine, dispersed Co-V-Zr composite nanophases near the previously formed dislocation network and grain boundaries. These nanophases effectively pin dislocations through the Orowan mechanism, while the continuous consumption of vacancies by Co and the segregation of Zr at grain boundaries jointly stabilize the formed Σ3 n Special grain boundaries. Through the synergistic effect of the above multi-stage processes, a composite nanophase consisting of a high-density dislocation substructure, a diffusely distributed structure, and a high proportion of Σ32 grain boundaries is finally constructed. n A three-dimensional stable microstructure composed of special grain boundaries.

[0041] For similar reasons, in a preferred embodiment, in step S3, the deformation of the first-stage cold rolling is 35-45%; and / or the temperature of the first-stage aging is 410-430°C, and the holding time is 0.5-1.5h; and / or the deformation of the second-stage cold rolling is 15-25%; and / or the temperature of the second-stage aging is 370-390°C, and the holding time is 1-2h.

[0042] In a preferred embodiment, in step S3, the deformation amount of the first-stage cold rolling is greater than that of the second-stage cold rolling, with a difference of 10-30%; and / or the temperature of the first-stage aging is greater than that of the second-stage aging, with a difference of 20-60℃. These conditions facilitate the phased control of dislocation configuration through a gradient design of the two-stage cold rolling deformation. Specifically, the first-stage cold rolling employs a larger deformation amount, introducing high-density and uniformly distributed dislocation tangles within the grain, providing sufficient nucleation sites for subsequent precipitates, and simultaneously creating Σ3 n The formation of twin boundaries requires the accumulation of necessary strain energy. Secondary cold rolling uses a smaller deformation amount, maintaining a deformation difference of over 10% between primary and secondary cold rolling. This effectively reduces the risk of damage to the Σ3 crystals formed during primary aging caused by excessive deformation in the secondary stage. n The problem of special grain boundary networks can be addressed by introducing an appropriate amount of new dislocations to disrupt the recovered dislocation network, making the dislocation distribution more uniform. This provides additional nucleation sites for the supplementary precipitation of Co-V-Zr composite nanophases and activates more solute atoms to segregate at dislocations and grain boundaries. If the deformation difference is less than 10%, it may lead to excessive secondary cold rolling deformation or insufficient primary cold rolling deformation. Excessive secondary cold rolling deformation can easily destroy the already formed Σ3 nSpecial grain boundaries lead to a significant decrease in their proportion, while insufficient primary cold rolling deformation fails to accumulate enough strain energy. If the deformation difference exceeds 30%, the secondary cold rolling deformation will be too low (<10%), potentially hindering the effective introduction of new dislocations to activate atomic enrichment and promote further grain boundary evolution. Similarly, the gradient design of the two-stage aging temperatures is crucial. Primary aging uses a higher temperature to promote Co atom diffusion, altering vacancy concentration and diffusion kinetics, and inducing initial precipitation of V and Zr at dislocation tangles, forming nanoscale clusters or GP regions. Simultaneously, the accumulated strain energy drives grain boundary migration, inducing Σ3+. n The nucleation and growth of twin boundaries form a preliminary special grain boundary network. Second-stage aging uses a lower temperature, maintaining a temperature difference of over 20°C between first-stage and second-stage aging. This promotes the full precipitation of the Co-V-Zr composite nanophase. Simultaneously, the continuous consumption of vacancies by Co and the segregation of Zr at the grain boundaries contribute to the stabilization of the already formed Σ32. n Special grain boundaries. If the temperature difference between the two aging stages is less than 20°C, the second-stage aging temperature is relatively too high, which can easily lead to the formation of Σ3 grain boundaries. n Special grain boundaries degrade and migrate, while precipitated phases tend to coarsen, weakening the strengthening effect. If the temperature difference exceeds 60℃, the secondary aging temperature is too low, potentially leading to insufficient precipitation kinetics of the Co-V-Zr composite nanophase, making it difficult to achieve full precipitation strengthening and grain boundary stabilization. By synergistically combining the deformation difference and temperature difference, it is beneficial to reduce the stress relaxation rate of the alloy to below 3.0%, while obtaining ≥60% Σ3 n Special grain boundary ratio and ≤15% discontinuous precipitation volume fraction.

[0043] In a preferred embodiment, in step S1, the homogenization treatment temperature is 850~890℃, the holding time is 8~10h, and the atmosphere is argon; and / or after homogenization treatment, the obtained material is cooled in the furnace to 450~550℃ and then air-cooled to obtain a homogenized ingot. Prolonged high-temperature holding of the ingot in an inert environment promotes the full diffusion of solute atoms, effectively eliminating dendritic segregation in the ingot and obtaining a homogeneous matrix. A slow cooling rate helps to gradually release the thermal stress inside the material, reducing crack formation caused by rapid cooling, and ultimately obtaining a homogenized ingot with a uniform and defect-free microstructure. This provides a stable microstructure basis for subsequent hot rolling and multi-stage deformation heat treatment, and is beneficial for improving the alloy's final high-temperature stress relaxation resistance.

[0044] High-temperature plastic deformation directly breaks down coarse dendrites and casting defects in the ingot, significantly improving segregation and obtaining fine recrystallized grains, thereby further improving the alloy's plasticity and reducing internal stress. In a preferred embodiment, in step S2, the initial rolling temperature of hot rolling is 840~860℃, the holding time is 1~2h, and the final rolling temperature is 750~800℃; and / or the cumulative deformation of hot rolling is ≥85%, preferably ≥90%; and / or the solution temperature is 840~890℃, the holding time is 2~6h; and / or the water quenching cooling rate is 30~60℃ / s. These conditions are conducive to promoting the full dissolution of alloying elements such as Co, V, Zr, and Ce in the copper matrix to form a supersaturated solid solution. At the same time, the faster cooling rate further reduces the unexpected coarse precipitation of second-phase particles during cooling, thus providing good solution conditions for subsequent multi-stage deformation processing.

[0045] For similar reasons, in a preferred embodiment, the solution treatment temperature is 850~870℃, and the holding time is 3~5h; and / or the tension annealing temperature is 300~400℃, the holding time is 1~2h, and the applied longitudinal tension is 5~20MPa. These conditions facilitate the effective elimination of residual stress introduced by the preceding multi-stage deformation process through thermo-mechanical coupling, thereby further stabilizing the microstructure and significantly improving the alloy's high-temperature stress relaxation resistance.

[0046] By controlling the proportions of microalloying elements in the alloy, it is beneficial to fully leverage the synergistic effect of quaternary microalloying. In a preferred embodiment, the Cu-15Ni-8Sn alloy, by weight percentage, comprises Ni 14.5~15.5%, Sn 7.5~8.5%, Co 0.3~0.8%, V 0.1~0.4%, Zr 0.05~0.2%, Ce 0.02~0.1%, with the balance being Cu and unavoidable impurities; and / or 0.3%≤Co+V+Zr+Ce≤1.8%. This composition, combined with a multi-stage deformation heat treatment process, can more stably induce a high proportion of Σ3 n The formation of a special grain boundary structure and the reduction of discontinuous precipitation volume fraction enable the alloy to achieve better stress relaxation resistance under high-temperature service conditions.

[0047] The Cu-15Ni-8Sn alloy prepared by this invention has a special grain boundary network structure formed by deformation heat treatment. In a preferred embodiment, in the Cu-15Ni-8Sn alloy, Σ3 nThe proportion of grain boundaries in the total grain boundary length is ≥60%; and / or the volume fraction of discontinuous precipitates in the Cu-15Ni-8Sn alloy is ≤15%. The Cu-15Ni-8Sn alloy utilizes the low interfacial energy and excellent resistance to intergranular failure inherent in its high proportion of special grain boundaries, combined with the grain boundary pinning effect of microalloying elements such as Zr, to effectively hinder grain boundary sliding and dislocation movement under high-temperature conditions, thereby significantly suppressing the nucleation and growth of discontinuous precipitates. Simultaneously, controlling the volume fraction of discontinuous precipitates within a low range further ensures microstructural stability. This optimization of the microstructure directly improves the alloy's resistance to stress relaxation under high-temperature service conditions, maintaining stable mechanical properties even at high temperatures.

[0048] In some embodiments, the stress relaxation rate of the Cu-15Ni-8Sn alloy of the present invention is ≤3.0% after a 50-hour stress relaxation test at 150°C and with an initial stress of 50% of the yield strength, in accordance with GB / T26007~2017 standards.

[0049] In another typical embodiment of the present invention, a Cu-15Ni-8Sn alloy is also provided, which is obtained according to the above-described preparation method of Cu-15Ni-8Sn alloy. Its internal microstructure is controlled by multi-stage deformation heat treatment and microalloying, forming a high proportion of Σ3 n The unique grain boundary network structure effectively hinders dislocation movement during high-temperature service, significantly improving the alloy's microstructural stability at high temperatures. Combined with quaternary microalloying, this alloy maintains high strength while exhibiting excellent resistance to stress relaxation, meeting the requirements for the mechanical properties of elastic components under extreme operating conditions.

[0050] The present application will be further described in detail below with reference to specific embodiments, which should not be construed as limiting the scope of protection claimed in the present application.

[0051] Example 1

[0052] Step S1: Prepare the alloy according to the alloy composition ratio in Table 1, use vacuum melting, power gradient heating (low power 8kW for 4min, high power 25kW for 1300℃), hold for 8min and then cast at 1200℃ to obtain an ingot; perform homogenization treatment at 860℃ for 9h, with furnace temperature fluctuation ≤5℃, argon protection, cool with the furnace to 500℃ and then air cool to obtain a homogenized ingot;

[0053] Step S2: The homogenized ingot is hot rolled to 850°C and held for 1.5 hours. Multiple hot rolling passes are performed, with a cumulative deformation of 92%. The final rolling temperature is 760°C, and the intermediate billet thickness is 3 mm to obtain the hot-rolled alloy. The hot-rolled alloy is then solution-treated at 860°C for 4 hours and then water-quenched at a cooling rate of 50°C / s to obtain the solution-treated alloy.

[0054] Step S3: The solid solution alloy is subjected to first-stage cold rolling, first-stage aging, second-stage cold rolling, and second-stage aging in sequence. First-stage cold rolling: deformation amount 40% to 1.8 mm; first-stage aging: 420℃×1h; second-stage cold rolling: deformation amount 20% to 1.2 mm; second-stage aging: 380℃×1.5h to obtain the aging alloy.

[0055] Step S4: The aged alloy is subjected to tension annealing at 350℃ for 1.5h with a longitudinal tension of 10MPa and then cooled in the furnace to further stabilize the microstructure, thus obtaining the Cu-15Ni-8Sn alloy.

[0056] Examples 2 to 5

[0057] The difference from Example 1 is that the composition of Cu-15Ni-8Sn is different, as detailed in Table 1.

[0058] Examples 6 to 9

[0059] The difference from Example 1 is that the conditions for homogenization, hot rolling, solution treatment, and tension annealing are different, as detailed in Table 2.

[0060] Examples 10 to 13

[0061] The difference from Example 1 is that the conditions in step S3 are different, as detailed in Table 3.

[0062] Comparative Example 1

[0063] The difference from Example 1 is that the alloy composition is the conventional Cu-15Ni-8Sn, which does not contain Co, V, Zr, or Ce.

[0064] Comparative Example 2

[0065] The difference from Example 1 is that in step S3, the solid solution alloy is subjected to single-stage cold rolling (60% deformation) and single aging (400℃×2h) to obtain an aged alloy.

[0066] Comparative Example 3

[0067] Existing microalloyed alloys include: Cu-15Ni-8Sn-0.2Co.

[0068] Comparative Example 4

[0069] Traditional Cu-15Ni-8Sn alloys include: Cu-15Ni-8Sn.

[0070] Performance testing:

[0071] The above embodiments and comparative examples were subjected to the following tests, and the results are shown in Table 4.

[0072] Stress relaxation rate: Under the conditions of 150℃ and initial stress of 50% of yield strength, a 50-hour stress relaxation test was conducted according to the methods specified in GB / T 26007-2017 "Copper Alloy Strips and Foils for Elastic Elements and Connectors" and GB / T 39152-2020 "Test Method for Bending Stress Relaxation of Copper and Copper Alloys". The stress change curve over time was recorded and the stress relaxation rate was calculated.

[0073] Yield strength: The room temperature tensile test was conducted according to GB / T 228.1-2021 "Metallic materials - Tensile testing - Part 1: Test method at room temperature". The tensile rate was 0.5 mm / min. The specimen was processed along the rolling direction. The test was conducted using an AG-X Plus 100 kN universal tensile testing machine. The average value of three parallel tests was taken.

[0074] Σ3 n Specific grain boundary ratios: determined using electron backscatter diffraction (EBSD) technique. After mechanical polishing, the surface stress layer was removed by ion thinning or vibration polishing. Data acquisition was performed on a field emission scanning electron microscope (such as JSM-7800F) equipped with an EBSD probe, with an accelerating voltage of 20 kV, a working distance of 15 mm, a step size of 0.1–0.5 μm, and a scanning area ≥200 μm × 200 μm. The acquired regional samples were analyzed, and Σ3 was identified based on the coincidence position lattice (CSL) model. n Special grain boundaries (n=1,2,3…), statistical Σ3 n The proportion of special grain boundaries in the total number of grain boundaries.

[0075] Discontinuous precipitate volume fraction: Metallographic samples were observed using scanning electron microscopy (SEM) in backscattered electron (BSE) mode, with an accelerating voltage of 15–20 kV and a working distance of 8–10 mm. Images were acquired from at least 10 randomly selected fields of view (magnification of 1000–5000x). Grayscale identification and area statistics were performed on the discontinuous precipitate regions. The proportion of the discontinuous precipitate area to the total area of ​​the field of view was used as an approximate value of the volume fraction, and the average value of each field of view was taken.

[0076] Table 1

[0077]

[0078] Table 2

[0079]

[0080] Table 3

[0081]

[0082] Table 4

[0083]

[0084] The stress relaxation curve of Example 1 is shown below. Figure 1 The stress relaxation curves for Comparative Example 3 are shown below. Figure 2 The stress relaxation curves for Comparative Example 4 are shown below. Figure 3 It can be seen that the alloy of the present invention has the best stress relaxation resistance.

[0085] Comparative Example 1 is a conventional Cu-15Ni-8Sn, without Co, V, Zr, or Ce, but employs the same multi-stage deformation heat treatment process as Example 1. The lack of microalloying elements leads to vacancy depletion and insufficient grain boundary pinning, resulting in an increased volume fraction of discontinuous precipitation and a significantly worse stress relaxation rate than Example 1. However, its Σ3 n The specific grain boundary ratio is comparable to that of Example 1, confirming that the multi-stage deformation heat treatment process itself induces a high proportion of Σ3. n Special grain boundaries make a significant contribution.

[0086] Although Comparative Example 2 contains a Co-V-Zr-Ce quaternary microalloying composition, step S3 uses single-stage cold rolling and single-time aging, without secondary cold rolling and secondary aging. Due to the lack of crucial steps such as intermediate aging to break dislocation entanglement and secondary cold rolling to optimize dislocation distribution, Σ3... n The proportion of special grain boundaries was only 45%, the volume fraction of discontinuous precipitation was as high as 25%, and the stress relaxation rate was 4.8%, significantly inferior to all other embodiments. This confirms that the multi-stage deformation heat treatment process induces a high proportion of Σ3 n Key role in special grain boundaries.

[0087] Comparative Example 3 is an existing microalloyed alloy, Cu-15Ni-8Sn-0.2Co, containing only Co and without employing a multi-stage deformation heat treatment process. It lacks the synergistic effect of V, Zr, and Ce, and Σ3+. n Active regulation of special grain boundaries, its Σ3 n With a special grain boundary ratio of only 52%, a discontinuous precipitation volume fraction of 18%, and a stress relaxation rate of 3.23%, it is still significantly inferior to the embodiments of the present invention.

[0088] Comparative Example 4 is a traditional Cu-15Ni-8Sn alloy, which contains neither microalloying elements nor employs a multi-stage deformation heat treatment process. It lacks grain boundary pinning and special grain boundary control, Σ3 nWith only 38% special grain boundaries, a discontinuous precipitation volume fraction as high as 35%, and a lack of effective pinning of dislocation movement, it has the highest stress relaxation rate and the worst overall performance.

[0089] As can be seen from the above, compared with the comparative example, this invention employs a multi-stage deformation heat treatment process—first-stage cold rolling, first-stage aging, second-stage cold rolling, and second-stage aging—combining Co-V-Zr-Ce microalloying, which significantly reduces the stress relaxation rate and discontinuous precipitation volume fraction of the Cu-15Ni-8Sn alloy. EBSD analysis shows that this process improves the Σ3... n The significantly increased twin boundary ratio confirms the effectiveness of this process in controlling grain boundary engineering in Cu-15Ni-8Sn alloys. Furthermore, it can be seen that the overall effect is better when all process parameters are within the preferred range of this invention.

[0090] 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 protection of the present invention.

Claims

1. A method for preparing a Cu-15Ni-8Sn alloy, characterized in that, The composition, by weight percentage, includes Ni 14.5-15.5%, Sn 7.5-8.5%, Co 0.1-1.0%, V 0.05-0.5%, Zr 0.02-0.2%, Ce 0.01-0.15%, with the balance being Cu and unavoidable impurities; the preparation method includes the following steps: Step S1: Prepare, melt, and cast the alloy according to the alloy composition ratio to obtain an ingot; then homogenize the ingot to obtain a homogenized ingot. Step S2: The homogenized ingot is hot-rolled to obtain a hot-rolled alloy; the hot-rolled alloy is then subjected to solution treatment and water quenching to obtain a solution-treated alloy. Step S3: The solid solution alloy is subjected to first-stage cold rolling, first-stage aging, second-stage cold rolling, and second-stage aging in sequence to obtain an aging alloy; the deformation amount of the first-stage cold rolling is 30~55%, the temperature of the first-stage aging is 400~450℃, the deformation amount of the second-stage cold rolling is 10~35%, the temperature of the second-stage aging is 350~420℃, and the deformation amount of the first-stage cold rolling is greater than the deformation amount of the second-stage cold rolling, and the temperature of the first-stage aging is greater than the temperature of the second-stage aging; Step S4: The aging alloy is subjected to tension annealing to obtain the Cu-15Ni-8Sn alloy.

2. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1, characterized in that, In step S3 The heat preservation time for the first-stage aging process is 0.5~3 hours; and / or The heat preservation time for the second-stage aging is 1 to 4 hours.

3. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, In step S3 The deformation amount of the first-stage cold rolling is 35-45%; and / or The first-stage aging temperature is 410~430℃, and the holding time is 0.5~1.5h; and / or The deformation amount of the secondary cold rolling is 15-25%; and / or The temperature for the secondary aging process is 370~390℃, and the holding time is 1~2h.

4. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, In step S3 The deformation amount of the first-stage cold rolling is greater than that of the second-stage cold rolling, with a difference of 10-30%; and / or The temperature of the first-stage aging is greater than the temperature of the second-stage aging, with a difference of 20~60℃.

5. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, In step S1 The homogenization treatment is performed at a temperature of 850~890℃, for a holding time of 8~10 hours, in an argon atmosphere; and / or After the homogenization treatment, the resulting material is cooled in the furnace to 450~550℃ and then air-cooled to obtain the homogenized ingot.

6. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, In step S2 The hot rolling temperature is 840~860℃, the holding time is 1~2h, and the final rolling temperature is 750~800℃; and / or The cumulative deformation of the hot rolling is ≥85%; and / or The solution treatment temperature is 840~890℃, and the holding time is 2~6h; and / or The cooling rate of the water quenching is 30~60℃ / s.

7. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, The solution treatment temperature is 850~870℃, and the holding time is 3~5h; and / or The tension annealing temperature is 300~400℃, the holding time is 1~2h, and the applied tension is 5~20MPa.

8. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, By weight percentage, the Cu-15Ni-8Sn alloy comprises 14.5-15.5% Ni, 7.5-8.5% Sn, 0.3-0.8% Co, 0.1-0.4% V, 0.05-0.2% Zr, 0.02-0.1% Ce, with the balance being Cu and unavoidable impurities; and / or 0.3%≤Co+V+Zr+Ce≤1.8%.

9. The method for preparing the Cu-15Ni-8Sn alloy according to claim 1 or 2, characterized in that, In the Cu-15Ni-8Sn alloy, Σ3 n Grain boundaries account for ≥60% of the total grain boundary length; and / or In the Cu-15Ni-8Sn alloy, the volume fraction of discontinuous precipitated phases is ≤15%.

10. A Cu-15Ni-8Sn alloy, characterized in that, The Cu-15Ni-8Sn alloy is obtained by any one of claims 1 to 9.

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