High-conductivity and high-strength copper alloy thermal mechanical treatment strengthening process

Abstract

The invention discloses a high-conductivity high-strength copper alloy thermal mechanical treatment strengthening process which comprises the following steps: firstly, heating a Cu-Ni-Be-Zr alloy cast ingot to 890-980 DEG C until the core part of the cast ingot reaches a set heating temperature, then compressing the cast ingot, carrying out water cooling at room temperature, and finally carrying out aging treatment and air cooling on the cast ingot at 480 DEG C * 3-480 DEG C * 8 hours; the Cu-Ni-Be-Zr alloy comprises the following components in percentage by weight: 1.6 to 2.5 percent of Ni, 0.25 to 0.5 percent of Be, 0.05 to 0.2 percent of Zr, not more than 0.1 percent of impurities and the balance of Cu. According to the method, the cast ingot of the Cu-Ni-Be-Zr alloy is directly subjected to thermal mechanical treatment and then subjected to aging, the process steps of forging, solid solution treatment and cold deformation in a traditional process are omitted, the performance of the treated Cu-Ni-Be-Zr alloy is basically consistent with that of the alloy obtained after the traditional production process is adopted, and therefore the Cu-Ni-Be-Zr alloy is obtained. The thermomechanical treatment combines the advantages of hot working and mechanical deformation, and the microstructure and performance of the alloy can be more efficiently improved.

Description

TECHNICAL FIELD

[0001] The application relates to a strengthening process of a copper alloy, in particular to a high-conductivity high-strength copper alloy thermomechanical treatment strengthening process. BACKGROUND

[0002] The Cu-Ni-Be-Zr alloy is widely used in the fields of aerospace, electronic equipment, ocean engineering and mechanical manufacturing as a high-performance copper-based material due to excellent strength, electrical conductivity, corrosion resistance and thermal stability.

[0003] Heat treatment is a core link for regulating the performance of the Cu-Ni-Be-Zr alloy, and research mainly focuses on the optimization of the solid solution and aging process of the alloy. The solid solution can make other alloy elements fully dissolve into the copper matrix, and the aging treatment can make the alloy elements precipitate in the form of a second phase, thereby improving the hardness and strength of the alloy. The traditional production process of the Cu-Ni-Be-Zr alloy is: smelting, casting, forging, solid solution, cold deformation, aging and service. Due to the market demand for production cost control, the short-process, low-cost and rapid production technology of the Cu-Ni-Be-Zr alloy has become particularly urgent. SUMMARY

[0004] The application aims to provide a high-conductivity high-strength copper alloy thermomechanical treatment strengthening process

[0005] Technical scheme: A high-conductivity high-strength copper alloy thermomechanical treatment strengthening process comprises the following steps:

[0006] Step S10: heating the Cu-Ni-Be-Zr alloy ingot to 890-980 DEG C, until the core of the ingot reaches the set heating temperature;

[0007] Step S20: compressing the ingot treated in step S10, and then water cooling at room temperature;

[0008] Step S30: aging the ingot treated in step S20 at 480 DEG C for 3-8 hours, and then air cooling;

[0009] The composition of the Cu-Ni-Be-Zr alloy is as follows: 1.6-2.5wt% of Ni, 0.25-0.5wt% of Be, 0.05-0.2wt% of Zr, not more than 0.1wt% of impurities, and the balance of Cu.

[0010] Further, the composition of the Cu-Ni-Be-Zr alloy is as follows: 1.8wt% of Ni, 0.38wt% of Be, 0.08wt% of Zr, not more than 0.1wt% of impurities, and the balance of Cu.

[0011] Further, the strain rate is 0.01s -1 ~ 10s -1 The ingot is compressed, and the true strain is 0.2~0.7.

[0012] Further, the temperature is 890~920℃, and the strain rate is 0.01s -1 ~ 0.1s -1 The ingot is compressed, and the temperature is 920~950℃, and the strain rate is 0.01s -1 ~ 1s -1 The ingot is compressed, and the temperature is 950~980℃, and the strain rate is 0.01s -1 ~ 10s -1 The ingot is compressed.

[0013] Further, the ingot is aged at 480℃ for 6h.

[0014] Further, the Cu-Ni-Be-Zr alloy ingot is in a cylindrical or cubic shape.

[0015] Further, the Cu-Ni-Be-Zr alloy is prepared by smelting electrolytic copper with a purity of ≥99.9%, nickel plate, beryllium copper alloy and sponge zirconium under the protection of argon, and is poured into an ingot.

[0016] Beneficial effects: the ingot of the Cu-Ni-Be-Zr alloy is directly subjected to the heat mechanical treatment, and then is aged, so that the forging + solid solution + cold deformation process steps in the traditional process are omitted; after the heat mechanical treatment strengthening process, the performance of the Cu-Ni-Be-Zr alloy is basically consistent with that of the alloy after the traditional production process, which indicates that the heat mechanical treatment combines the advantages of heat processing and mechanical deformation, can more efficiently improve the microstructure and performance of the alloy, and provides a new and reliable process way for the practical application of the Cu-Ni-Be-Zr alloy. BRIEF DESCRIPTION OF DRAWINGS

[0017] Figure 1 The flow stress curve of the Cu-Ni-Be-Zr alloy under different deformation temperatures (890, 920, 950, 980℃) and different strain rates (0.01, 0.1, 1, 10s -1 );

[0018] Figure 2 The metallographic structure of the Cu-Ni-Be-Zr alloy after hot compression;

[0019] Figure 3 The IPF graph of the Cu-Ni-Be-Zr alloy after hot compression;

[0020] Figure 4corresponding pole figure; Figure 3 corresponding pole figure;

[0021] Figure 5 KAM map of Cu-Ni-Be-Zr alloy after hot compression;

[0022] Figure 6 GOS map of Cu-Ni-Be-Zr alloy after hot compression;

[0023] Figure 7 hardness of Cu-Ni-Be-Zr alloy under different hot compression;

[0024] Figure 8 NRB compound in TEM bright field image;

[0025] Figure 9 line scanning EDS spectrum of NRB compound;

[0026] Figure 10 STEM image and surface scanning EDS spectrum image of NRB compound. DETAILED DESCRIPTION

[0027] The application will be further illustrated below in conjunction with the accompanying drawings and specific examples.

[0028] A high-conductivity high-strength copper alloy thermomechanical treatment strengthening process, specifically comprising the following steps:

[0029] Step S10: heat a Cu-Ni-Be-Zr alloy ingot, for example in a cylindrical shape, to 890-980℃, until the center of the ingot reaches the set heating temperature.

[0030] Step S20: compress the ingot after step S10 treatment at a strain rate of 0.01s -1 -10s -1 , with a true strain of 0.2-0.7, and then water cooling at room temperature.

[0031] heat to 890-920℃, compress the ingot at a strain rate of 0.01s -1 -0.1s -1 , heat to 920-950℃, compress the ingot at a strain rate of 0.01s -1 -1s -1 , heat to 950-980℃, compress the ingot at a strain rate of 0.01s -1 -10s -1 , compress the ingot.

[0032] Step S30: age the ingot after step S20 treatment at 480℃×3h-480℃×8h, and then air cool.

[0033] The Cu-Ni-Be-Zr alloy has a composition of Ni 1.6-2.5 wt%, Be 0.25-0.5 wt%, Zr 0.05-0.2 wt%, impurities less than 0.1 wt%, and the balance of Cu. The Cu-Ni-Be-Zr alloy is prepared by smelting and casting ingots under argon protection, using electrolytic copper with a purity of ≥99.9%, nickel plate, beryllium copper alloy, and sponge zirconium as raw materials. The ingots can be regular or irregular in shape, and the regular shape can be cylindrical or cubic.

[0034] Example 1

[0035] A Cu-Ni-Be-Zr alloy thermomechanical treatment strengthening process, specifically comprising the following steps:

[0036] Step S10: heat the Cu-Ni-Be-Zr alloy ingot to 890°C, so that the core of the ingot reaches 890°C;

[0037] Step S20: compress the ingot treated in step S10 at a strain rate of 0.01 s -1 , and then water cool at room temperature.

[0038] Step S30: age the ingot treated in step S20 at 480°C for 3h, and then air cool.

[0039] Example 2

[0040] A Cu-Ni-Be-Zr alloy thermomechanical treatment strengthening process, specifically comprising the following steps:

[0041] Step S10: heat the Cu-Ni-Be-Zr alloy ingot to 920°C, so that the core of the ingot reaches 920°C;

[0042] Step S20: compress the ingot treated in step S10 at a strain rate of 0.1 s -1 , and then water cool at room temperature.

[0043] Step S30: age the ingot treated in step S20 at 480°C for 4h, and then air cool.

[0044] Example 3

[0045] A Cu-Ni-Be-Zr alloy thermomechanical treatment strengthening process, specifically comprising the following steps:

[0046] Step S10: heat the Cu-Ni-Be-Zr alloy ingot to 950°C, so that the core of the ingot reaches 950°C;

[0047] Step S20: compress the ingot treated in step S10 at a strain rate of 1 s -1The ingot after step S10 is compressed with a true strain of 0.5, and then water cooled at room temperature;

[0048] Step S30: the ingot after step S20 is aged at 480℃ for 6h, and then air cooled.

[0049] Example 4

[0050] A Cu-Ni-Be-Zr alloy thermomechanical treatment strengthening process, specifically comprising the following steps:

[0051] Step S10: the Cu-Ni-Be-Zr alloy ingot is heated to 980℃, until the ingot core reaches 980℃;

[0052] Step S20: the ingot after step S10 is compressed at a strain rate of 10s -1 The ingot after step S10 is compressed with a true strain of 0.7, and then water cooled at room temperature;

[0053] Step S30: the ingot after step S20 is aged at 480℃ for 8h, and then air cooled.

[0054] In Examples 1-4, the composition of the Cu-Ni-Be-Zr alloy ingot is Ni 1.8wt%, Be 0.38wt%, Zr 0.08wt%, impurities less than 0.1wt%, and the balance of Cu. The ingot is cast into a cylindrical ingot.

[0055] Figure 1 (a)-(d) are the flow stress curves of the Cu-Ni-Be-Zr alloy under different deformation temperatures (890, 920, 950, 980℃) and different strain rates (0.01, 0.1, 1, 10s -1 ) conditions, (a) strain rate 0.01s -1 , (b) strain rate 0.1s -1 , (c) strain rate 1s -1 , (d) strain rate 10s -1 At the initial stage of thermal deformation, a large number of dislocations are generated in the alloy due to plastic deformation. These dislocations accumulate or entangle with each other in the crystal, and then form dislocation walls or subgrain boundaries. With the continuous progress of thermal deformation, the effect of dynamic softening mechanism (including dynamic recovery and dynamic recrystallization) gradually increases. With the further deepening of material deformation, the dynamic softening phenomenon becomes more and more significant. At this time, the growth trend of flow stress gradually slows down and shows a downward trend.

[0056] Table 1 is the flow stress curves of the Cu-Ni-Be-Zr alloy under different deformation temperatures (890, 920, 950, 980℃) and different strain rates (0.01, 0.1, 1, 10s -1) under the condition of peak stress (MPa). The peak stress of the alloy shows a regularity of increasing with the decrease of temperature or the increase of strain rate. When the temperature decreases, the atomic diffusion ability weakens, and the dynamic softening mechanism is difficult to fully play a role, and work hardening dominates, resulting in the increase of peak stress. When the strain rate increases, the deformation amount per unit time increases, and the dislocations cannot effectively coordinate the deformation through dynamic recovery and dynamic recrystallization and other ways, and the dislocation density increases rapidly, thereby increasing the peak stress.

[0057] Table 1

[0058]

[0059] Figure 2 (a)-(g) are the metallographic structures of Cu-Ni-Be-Zr alloy after hot compression, (a) 890℃ and strain rate 0.01s -1 , (b) 920℃ and strain rate 0.01s -1 , (c) 950℃ and strain rate 0.01s -1 , (d) 980℃ and strain rate 0.01s -1 , (e) 980℃ and strain rate 0.1s -1 , (f) 980℃ and strain rate 1s -1 , (g) 980℃ and strain rate 10s -1 .

[0060] In the hot deformation metallographic Figure 2 (d) at 980℃ and strain rate 0.01s -1 , different characteristics from the previous three are presented. At this time, the grain size of the alloy increases significantly, and the atomic diffusion ability in high temperature environment is significantly improved, and the grains gradually grow by swallowing the surrounding small grains. At the same time, there are some small recrystallized structures in the structure. The occurrence of dynamic recrystallization is closely related to factors such as deformation temperature, strain rate and deformation degree. Under this condition, although the strain rate is low, the higher temperature provides enough driving force for recrystallization nucleation and growth, so that dynamic recrystallization occurs in part of the region. However, compared with Figure 2 (a, b, c), the texture in this figure is not obvious. Higher temperature promotes the progress of the recrystallization process, and the newly generated recrystallized grains disrupt the orientation distribution of the original crystal, thereby weakening the texture characteristics. Figure 2 (e, f, g) shows the metallographic structure deformed at 980℃ with strain rates of 0.1, 1, and 10s -1 , respectively. In Figure 2In (e, f), obvious texture characteristics are presented, the grains are distributed in orientation, and there are some fine recrystallized structures. With the increase of the strain rate, the deformation amount of the alloy in unit time increases, the dislocations cannot move and coordinate sufficiently, the stress concentration in the crystal is intensified, and more dislocations are caused to provide more positions for recrystallization nucleation. Meanwhile, the heat generated in the deformation process cannot be dissipated in time under the higher strain rate, which further increases the temperature in the alloy, and is beneficial to the recrystallization. Under the influence of the strain rate, the deformation of the crystal is mainly concentrated in certain specific directions, and then obvious texture is formed. In (g), the size of the recrystallized grains is obviously larger than that in (e, f). Figure 2 In (g), the size of the recrystallized grains is obviously larger than that in (e, f). Figure 2 In (e, f), the size of the recrystallized grains is very uneven. Under the high strain rate of 10 s -1 , a large amount of heat generated in the deformation process makes the temperature in the alloy rise rapidly, and the recrystallization process can be carried out quickly, and the grains have more time and energy to grow. In addition, Figure 2 In (g), the texture is not obvious, and the grain size is relatively uniform. Under the high strain rate, a large amount of energy input makes the deformation in the alloy more uniform, and the recrystallization process occurs synchronously in the whole structure, so that the texture characteristics are weakened, and the grain size tends to be consistent. Figure 2 In (d, e, f), the grain size is very uneven. In (d), Figure 2 In (d), under the high temperature and low strain rate, the recrystallization process competes with the grain growth process, and the grains in some regions have grown, while the recrystallization in some regions has just begun, resulting in a large difference in grain size. In (e, f), Figure 2 In (e, f), the difference in strain rate causes different degrees of deformation and recrystallization processes in different regions, and then causes uneven distribution of grain size.

[0061] Figure 3 are IPF maps of the Cu-Ni-Be-Zr alloy after hot compression, (a) 890℃ and strain rate 0.01s -1 , (b) 920℃ and strain rate 0.01s -1 , (c) 950℃ and strain rate 0.01s -1 , (d) 980℃ and strain rate 0.01s -1 , (e) 980℃ and strain rate 0.1s -1 , (f) 980℃ and strain rate 1s -1 , (g) 980℃ and strain rate 10s -1 .

[0062] Figure 4 are pole figures corresponding to Figure 3 , (a) 890℃ and strain rate 0.01s -1(b) 920 °C and strain rate 0.01 s -1 (c) 950 °C and strain rate 0.01 s -1 (d) 980 °C and strain rate 0.01 s -1 (e) 980 °C and strain rate 0.1 s -1 (f) 980 °C and strain rate 1 s -1 (g) 980 °C and strain rate 10 s -1 .

[0063] In Figure 3 (a) IPF map at low temperature and low strain rate (890 °C, 0.01 s -1 ), elongated grains along the compression direction were observed, with a length to width ratio > 6:1, showing typical compression deformation characteristics. Grain boundaries were blurred, and there were a large number of dislocation tangles. Figure 4 (a) showed that the high intensity region was concentrated in the

[100] crystal direction, with a texture intensity of 23.2 mrd, indicating a high degree of preferred orientation of the crystal direction. Figure 3 (b, c) at medium temperature and low strain rate (920-950 °C, 0.01 s -1 ), elongated grains were still dominant, and the proportion of green areas (close to 101 direction) gradually decreased, Figure 4 (b, c) showed that the high intensity region was still concentrated in the

[100] crystal direction, and the texture intensity decreased from 22.37 mrd to 12.12 mrd, and the crystal direction distribution tended to be multi-directional. Figure 3 (d) at high temperature and low strain rate (980 °C, 0.01 s -1 ), elongated grains became thick. Figure 4 (d) showed that the high intensity region shifted to both sides. Figure 3 (e) (980 °C, 0.1 s -1 ), the green area disappeared, and the proportion of red areas (close to 001) increased. Figure 4 (e) had a lower significant peak, with a texture intensity < 7.29 mrd, and the orientation was relatively uniform. Figure 3 (f) (980 °C, 1 s -1 ), the grains were more uniform, and deformation grains of different directions appeared. Figure 4 (f) had a lower significant peak, with a texture intensity of 10.72 mrd. Figure 5 (g) at high temperature and high strain rate (980 °C, 10 s -1 ), a small amount of elongated grains coexisted with equiaxed grains. Figure 5 (g) showed that the high intensity region appeared on both sides of the

[100] crystal direction, with a texture intensity of 11.26 mrd.

[0064] The microstructure and texture evolution of materials during hot compression processing are synergistically regulated by temperature and strain rate. Low temperature and low strain rate (890℃, 0.01s) -1 Under certain conditions, limited dynamic recovery leads to extensive dislocation entanglement, causing grains to elongate dramatically along the compression direction and form a strong texture

[001] . At medium temperatures (920–950 °C), dynamic recovery is promoted, dislocations partially annihilate through climb and cross-slip, grains fragment and adjust their orientation, and the texture intensity decreases. At high temperatures and low strain rates (980 °C, 0.01 s⁻¹), dynamic recovery is further enhanced. -1 Under these conditions, dynamic recrystallization proceeds fully, grains nucleate and grow through grain boundary bowing mechanisms, and the crystal orientation preferentially aligns and re-enhances in the

[001] direction; at high temperatures (980℃, 0.1~1s), the strain rate is high. -1 High storage energy is generated, accelerating recrystallization, resulting in rapid grain refinement and randomized orientation. However, if the strain rate is increased to 10 s⁻¹, the effect is not significant. -1 Overall, a greater number of equiaxed grains appeared. Generally, increased temperature promotes dynamic recovery by enhancing atomic diffusion, while strain rate regulates the degree of crystal orientation preference by affecting the accumulation of stored energy and recrystallization kinetics.

[0065] Figure 5 KAM diagram of Cu-Ni-Be-Zr alloy after hot compression, (a) at 890℃ and strain rate of 0.01s. -1 (b) 920℃ and strain rate 0.01s -1 (c) 950℃ and strain rate 0.01s -1 (d) 980℃ and strain rate 0.01s -1 (e) 980℃ and strain rate 0.1s -1 (f) 980℃ and strain rate 1s -1 (g) 980℃ and strain rate 10s -1 .

[0066] exist Figure 5 In (a, b, c, d), for the same strain rate (0.01 s⁻¹), -1 The KAM graph shows the KAM values ​​as the processing temperature increases from 890℃ to 980℃. The trend of the KAM values ​​is as follows: Figure 6 (h), the KAM value decreases significantly with increasing temperature. While... Figure 6 In (d, e, f, g), at a hot working temperature of 980℃, the increase in KAM value is not significant with the increase in strain rate.

[0067] Figure 6 The GOS diagram of Cu-Ni-Be-Zr alloy after hot compression is shown in (a) at 890℃ and a strain rate of 0.01s. -1 (b) 920℃ and strain rate 0.01s -1(c) 950℃ and strain rate 0.01s -1 (d) 980℃ and strain rate 0.01s -1 (e) 980℃ and strain rate 0.1s -1 (f) 980℃ and strain rate 1s -1 (g) 980℃ and strain rate 10s -1 .

[0068] exist Figure 6 In (a, b, c, d), for the same strain rate (0.01 s⁻¹), -1 The GOS diagram shows the process temperature increased from 890℃ to 980℃. Its recrystallization accounts for... Figure 7 (h) decreased from 5.2% to 1.1%. The overall recrystallized grain percentage decreased. And... Figure 7 In (d, e, f, g), the recrystallization percentage increased from 1.1% to 50.2%, as shown in Figure 6(h). From the perspective of the competition mechanism between dynamic recovery and dynamic recrystallization, dynamic recovery and dynamic recrystallization are two important softening mechanisms in the hot working process. Dynamic recovery mainly reduces dislocation density through dislocation climb and cross-slip, while dynamic recrystallization softens the material by forming new distortion-free grains. The competition between the two depends on the stacking fault energy, deformation temperature, and strain rate of the material.

[0069] Variations in temperature and strain rate affect the degree of recrystallization by influencing the relative rates of dynamic recovery and dynamic recrystallization, as well as the accumulation of dislocation density. The presence of Be further promotes dynamic recovery and inhibits dynamic recrystallization by increasing stacking fault energy, especially under high temperature and low strain rate conditions. However, under high temperature and high strain rate conditions, the increase in strain rate may exceed the capacity for dynamic recovery, leading to dynamic recrystallization becoming dominant.

[0070] Figure 8 The hardness of Cu-Ni-Be-Zr alloy under different hot compression conditions is given by: (a) strain rate 0.01 s⁻¹ -1 (b) Strain rate 0.1 s -1 (c) Strain rate 1s -1 (d) Strain rate 10s -1 .

[0071] Figure 8 This indicates that under high temperature (980℃) conditions, the softening mechanisms inside the material, such as grain coarsening and complete dynamic recrystallization, greatly weaken the material's hardness, making the effect of strain rate on hardness difficult to manifest. However, under relatively low temperature conditions, strain rate can interact with the material's internal processes such as dynamic recovery and recrystallization, thereby affecting the material's hardness to varying degrees.

[0072] The samples were heated to 980℃ and subjected to hot compression at different strain rates, followed by aging treatment at 480℃ for 3 hours. The resulting alloy exhibited high hardness, as shown in Table 2 for the hardness after aging under different processes. On one hand, dynamic recrystallization during hot compression produces fine grains, which can improve the alloy's hardness. On the other hand, solution treatment and aging promote the formation of precipitates, while the dislocations introduced during thermomechanical treatment can make the precipitates finer and more dispersed, enhancing the precipitation strengthening effect. Furthermore, thermomechanical treatment integrates multiple steps to reduce energy consumption and time, improving efficiency, especially at a strain rate of 10 s⁻¹. -1 At a deformation temperature of 980℃, the microstructure exhibits good uniformity and enhanced adaptability. The strain rate of the Cu-Ni-Be-Zr alloy at 980℃ is 1 s⁻¹. -1 and 10s -1 After hot compression and aging treatment, the alloy achieved higher hardness than traditional heat treatment methods, proving the feasibility of the thermomechanical treatment method of this invention.

[0073] Table 2

[0074]

[0075] Figure 8 (a) and (b) are bright-field images of NRB compounds, (c) is a selected area electron diffraction of NRB compound (b), and (d) is a high-resolution TEM image of selected area b.

[0076] Figure 8 In (a), it is clearly visible that a large number of γ″ phases are encapsulated within NRB compounds in the Cu-Ni-Be-Zr alloy. From the perspective of secondary precipitation, this phenomenon has its inherent mechanism. During the aging process of the material, primary precipitation occurs first to form NRB compounds, while the subsequent precipitation of the γ″ phase within the NRB compounds is a secondary precipitation phenomenon. In dark-field images... Figure 9 (b) shows a large number of secondary γ″ phases in the selected area high-resolution image. Selected area electron diffraction... Figure 10 (c) Diffraction spots of the γ″ phase are also present, further confirming the existence of the secondary γ″ phase. From the perspective of secondary precipitation theory, the uneven distribution of solute atoms within the NRB compound is a key factor in the precipitation process.

[0077] Figure 8 This is the line scan EDS spectrum of the NRB compound. Figure 9 The images show STEM images of the NRB compound and EDS images of the surface scan.

[0078] NRB compounds themselves contain large amounts of elements such as Cu, Ni, Be, and Zr, such as Figure 8 and Figure 8As shown. In the early stages of aging, the alloy first forms NRB compounds through primary precipitation. As a secondary phase rich in Ni, Zr, etc., NRB preferentially segregates Cu, Ni, Be, etc., due to its chemical properties and crystal structure. With prolonged aging, the NRB compounds gradually grow, and the distribution of solute atoms within them becomes inhomogeneous. This compositional inhomogeneity provides the material basis for the subsequent secondary precipitation of the γ″ phase—the solute atoms required for the γ″ phase (such as Be) can meet the nucleation concentration conditions in locally enriched regions within the NRB, explaining the solute distribution mechanism of the secondary γ″ phase precipitation. From the perspective of the synergistic mechanism of interfacial energy and dislocations, the NRB compounds are coherent with the matrix and are circular, resulting in a low interfacial energy. In crystallography, a low interfacial energy provides favorable conditions for dislocation movement; therefore, dislocations easily cut through the NRB compounds, such as... Figure 8 As shown in (a) and (c), the NRB compound exhibits obvious dislocation cut marks in the high-resolution images. These correspond to additional diffraction spots caused by dislocation motion. ​ (b) and scattered diffraction spots ​ (d) The movement and presence of dislocations cause lattice distortion, which lowers the energy barrier required for the nucleation of the γ″ phase and provides favorable nucleation sites for the secondary precipitation of the γ″ phase.

[0079] Based on the above, we can conclude that:

[0080] 1. The suitable hot pressing deformation processing temperature for Cu-Ni-Be-Zr alloy is below 920℃, when the alloy is suitable for low strain rate processing; above 920℃, the alloy can be hot pressed deformation processing at both high and low strain rates.

[0081] 2. The texture of the Cu-Ni-Be-Zr alloy after hot compression is as follows: <100> Copper texture, at a strain rate of 0.01 s⁻¹ -1 At that time, the texture strength decreased as the processing temperature increased from 890℃ to 980℃. At a processing temperature of 980℃, the texture strength decreased as the strain rate increased from 0.01 s⁻¹. -1 Improved to 10 seconds -1 The texture strength showed no significant change. At a strain rate of 0.01 s⁻¹... -1 Below this point, as the processing temperature increases from 890℃ to 980℃, the dynamic recovery mechanism becomes dominant. At a processing temperature of 980℃, as the strain rate increases from 0.01 s⁻¹... -1 Improved to 10 seconds -1 The dynamic recovery mechanism still plays a role, but as the strain rate increases, the dynamic recrystallization mechanism gradually becomes dominant.

[0082] 3. Analysis of grain orientation and texture evolution using IPF plots and pole figures revealed that the microstructure is synergistically regulated by temperature and strain rate. The overall texture is... <100> Copper texture, at a strain rate of 0.01 s⁻¹-1 At that time, the texture strength decreased as the processing temperature increased from 890℃ to 980℃. At a processing temperature of 980℃, the texture strength decreased as the strain rate increased from 0.01 s⁻¹. -1 Improved to 10 seconds -1 The texture strength showed no significant change. The competition mechanism between dynamic recrystallization and dynamic recovery in the Cu-Ni-Be-Zr alloy during hot compression was investigated using KAM and GOS plots. This was achieved at a strain rate of 0.01 s⁻¹. -1 Below this point, as the processing temperature increases from 890℃ to 980℃, the dynamic recovery mechanism becomes dominant. At a processing temperature of 980℃, as the strain rate increases from 0.01 s⁻¹... -1 Improved to 10 seconds -1 The dynamic recovery mechanism still plays a role, but as the strain rate increases, the dynamic recrystallization mechanism gradually becomes dominant.

[0083] 4. Analysis of hardness curves at different processing rates and temperatures, using theories of dynamic recovery, high-temperature softening, and work hardening, reveals the following: At 890℃, hardness increases significantly with increasing strain rate, indicating that dislocation strengthening dominates and high strain rate inhibits softening. Between 920℃ and 950℃, the alloy hardness fluctuates little, with dynamic recovery and strain rate effects offsetting each other. At 980℃, hardness shows no significant fluctuation, indicating that high-temperature softening dominates the alloy hardness, rendering strain rate ineffective.

[0084] 5. The strain rate of the Cu-Ni-Be-Zr alloy at 980℃ is 1 s. -1 and 10s -1 After hot compression and aging treatment, the alloy achieved higher hardness than traditional heat treatment methods, proving the feasibility of the thermomechanical treatment method of this invention. Thermomechanical treatment combines the advantages of hot working and mechanical deformation, enabling more efficient improvement of the alloy's microstructure and properties, and providing a new and reliable process for the practical application of Cu-Ni-Be-Zr alloys.

[0085] 6. Precipitation mechanism of secondary γ″ phase in NRB compounds and primary γ″ phase: The precipitation of secondary γ″ phase is mainly due to solute atom segregation and dislocation-assisted nucleation. The misfit degree between the γ″ phase inertial plane and the matrix interface was calculated to be 0.02%, confirming the γ″ phase / matrix interface as a semi-coherent interface. The dislocation spacing D caused by the interfacial misfit between the γ″ phase and the matrix is ​​6.4 nm, and the critical shear stress τ is increased due to the dislocations generated by the interfacial misfit. c Reduced by 40%. (110) γ″1 ⊥(101) γ″2

[110] γ″1 ⊥

[101] γ″2 (110) γ″1 / / (100) γ″2

[001] γ″1 / /

[011] γ″2 In the phase relationship between these two new phases and the parent phase: (110)p⊥(101)p orientation, the interfacial atomic matching degree reaches 87%, and the additional interfacial energy Δγ caused by mismatch only increases by 0.03 J / m 2 The orientation matching degree of (110)p / / (100)p is only 64%, and Δγ increases to 0.12 J / m. 2 This difference results in the former having higher thermodynamic stability, making it the dominant orientation.

Claims

1. A thermomechanical strengthening process for high-conductivity, high-strength copper alloys, characterized in that... Includes the following steps: step S10: Heat the Cu-Ni-Be-Zr alloy ingot to 890-980℃ until the core of the ingot reaches the set heating temperature; Step S20: Compress the ingot processed in step S10, and then cool it with water at room temperature; Step S30: The ingot processed in step S20 is subjected to aging treatment at 480℃ for 3 hours to 480℃ for 8 hours, and then air-cooled. The composition of the Cu-Ni-Be-Zr alloy is Ni 1.6-2.5wt%, Be 0.25-0.5wt%, Zr 0.05-0.2wt%, impurities not exceeding 0.1wt%, and balance Cu.

2. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 1, characterized in that: The composition of the Cu-Ni-Be-Zr alloy is 1.8wt% Ni, 0.38wt% Be, 0.08wt% Zr, impurities not exceeding 0.1wt%, and balance Cu.

3. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 1, characterized in that: A strain rate of 0.01 s⁻¹ was used. -1 ~10s -1 For ingot compression, the true strain is 0.2–0.

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4. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 3, characterized in that: Heating to 890–920℃, with a strain rate of 0.01 s⁻¹ -1 ~0.1s -1 The ingot is compressed; heated to 920–950°C, with a strain rate of 0.01 s⁻¹. -1 ~1s -1 The ingot is compressed; heated to 950–980°C, with a strain rate of 0.01 s⁻¹. -1 ~10s -1 Compression of the ingot.

5. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 1, characterized in that: The ingots are subjected to aging treatment at 480℃ for 6 hours.

6. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 1, characterized in that: The Cu-Ni-Be-Zr alloy ingots are cylindrical or cubic.

7. The thermomechanical strengthening process for high-conductivity, high-strength copper alloys according to claim 1, characterized in that: The Cu-Ni-Be-Zr alloy is made from electrolytic copper, nickel plates, beryllium copper alloy and sponge zirconium with a purity of ≥99.9% as raw materials, and is melted and cast into ingots under argon protection.

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