WO2021258747A1 - 一种含Nb和Al的钛青铜合金带材及其制备方法 - Google Patents

一种含Nb和Al的钛青铜合金带材及其制备方法 Download PDF

Info

Publication number
WO2021258747A1
WO2021258747A1 PCT/CN2021/076741 CN2021076741W WO2021258747A1 WO 2021258747 A1 WO2021258747 A1 WO 2021258747A1 CN 2021076741 W CN2021076741 W CN 2021076741W WO 2021258747 A1 WO2021258747 A1 WO 2021258747A1
Authority
WO
WIPO (PCT)
Prior art keywords
titanium bronze
bronze alloy
alloy strip
alloy
strip
Prior art date
Application number
PCT/CN2021/076741
Other languages
English (en)
French (fr)
Inventor
唐宁
张镇凯
支月鹏
杨谏
武博
Original Assignee
宁波博威合金板带有限公司
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by 宁波博威合金板带有限公司 filed Critical 宁波博威合金板带有限公司
Priority to US17/423,698 priority Critical patent/US11913102B2/en
Priority to KR1020217025759A priority patent/KR102640850B1/ko
Publication of WO2021258747A1 publication Critical patent/WO2021258747A1/zh

Links

Images

Classifications

    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C9/00—Alloys based on copper
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/46—Roll speed or drive motor control
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/56—Elongation control
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/74—Temperature control, e.g. by cooling or heating the rolls or the product
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B45/00—Devices for surface or other treatment of work, specially combined with or arranged in, or specially adapted for use in connection with, metal-rolling mills
    • B21B45/004—Heating the product
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0268—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment between cold rolling steps
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • C22C1/03—Making non-ferrous alloys by melting using master alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C14/00—Alloys based on titanium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C27/00—Alloys based on rhenium or a refractory metal not mentioned in groups C22C14/00 or C22C16/00
    • C22C27/02—Alloys based on vanadium, niobium, or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C9/00—Alloys based on copper
    • C22C9/01—Alloys based on copper with aluminium as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/002—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working by rapid cooling or quenching; cooling agents used therefor
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/02—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working in inert or controlled atmosphere or vacuum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/08—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of copper or alloys based thereon
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B21B2003/005—Copper or its alloys
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B2265/00—Forming parameters
    • B21B2265/10—Compression, e.g. longitudinal compression

Definitions

  • the invention belongs to the technical field of copper alloy materials, and particularly relates to a titanium bronze alloy strip containing Nb and Al.
  • the titanium bronze alloy strip has excellent stability, especially the stability of mechanical properties at high temperatures.
  • the invention also relates to a preparation method of the titanium bronze alloy strip.
  • Titanium bronze alloy is a copper alloy with titanium as the main alloying element. It has high strength and excellent forming performance. It can be used to replace beryllium copper alloy in some applications.
  • Titanium bronze is a kind of amplitude modulation decomposition strengthening and aging precipitation strengthening alloy.
  • the main strengthening structure is the amplitude modulation decomposition structure and ⁇ -Cu 4 Ti phase.
  • the strengthening method of titanium bronze alloy is amplitude modulation decomposition strengthening.
  • the solid solution of Ti atoms in the copper matrix diffuses to form periodic Ti atom enrichment areas in the crystal grains, that is, the amplitude modulation decomposition structure. With the continuation of the time-effect process, the amplitude modulation decomposition structure gradually transforms into periodically arranged ⁇ -Cu 4 Ti phase.
  • the amplitude-modulated decomposition structure and ⁇ -Cu 4 Ti phase have poor stability at high temperatures and are prone to evolution, which will adversely affect the mechanical properties of the alloy.
  • the stability of material performance is very important. Good stability can ensure that the product will not fail quickly when sudden overload and high temperature occur during processing and application.
  • titanium bronze has high strength and excellent elastic properties, it has a wide range of application prospects in electric vehicles, 5G communication base stations and other fields. In these fields, especially in the field of electric vehicles, there are often instantaneous or continuous high-temperature operating conditions, and the temperature may reach 200°C or more.
  • the invention designs a Cu-Ti-Nb-Al system alloy by adding a certain amount of Nb and Al to the titanium bronze at the same time. Compared with conventional titanium bronze alloys, while ensuring excellent bending performance, the Cu-Ti-Nb-Al system alloy has significantly improved mechanical properties and stability at high temperatures, and the alloy strength has also been improved.
  • the technical problem to be solved by the present invention is: in view of the shortcomings of the prior art, how to ensure the excellent mechanical properties and bending performance of the titanium bronze alloy strip while making the alloy strip have optimized stability, especially at high temperatures Stability of mechanical properties.
  • a titanium bronze alloy strip containing Nb and Al the weight percentage composition of the titanium bronze alloy strip includes: 2.0-4.5wt% Ti, 0.005-0.4wt % Nb, 0.01-0.5wt% Al, the balance is Cu and unavoidable impurities.
  • the Ti content of the titanium bronze alloy strip of the present invention is 2.0-4.5 wt%.
  • the Ti content of the titanium bronze alloy strip is 2.5-4.0 wt%. More preferably, the Ti content of the titanium bronze alloy strip is 2.9-3.5 wt%.
  • Ti is the main strengthening element.
  • the amplitude-modulated decomposition structure is first formed by the diffusion of Ti atoms in the solid solution.
  • the strength of the copper alloy increases significantly; with the increase of the aging time, the needle-like ⁇ is gradually precipitated in the matrix.
  • ⁇ -Cu 4 Ti phase the aging strengthening effect gradually reaches its peak during this process; as the aging time is further extended, flaky ⁇ -Cu 4 Ti phase will be precipitated on the grain boundary, and its volume fraction will gradually increase with time It will eventually replace the ⁇ -Cu 4 Ti phase.
  • the strengthening effect of the copper alloy gradually decreases.
  • the amplitude modulation decomposition structure is a uniform nano-scale structure, and the ⁇ -Cu 4 Ti phase is also a nano-scale precipitation phase, which is dispersed in the matrix. Both of these structures can hinder the movement of grain boundaries and dislocations, making the copper alloy strong improve. By controlling the aging process, it means the formation of different microstructures, which can effectively control the overall properties of the alloy.
  • Nb or Al can be optionally added as a secondary alloying element in a titanium bronze alloy.
  • Nb when Nb is added alone, it can be dissolved in a small amount in the copper matrix, which slightly improves the strength of the alloy, but has little effect on other properties.
  • due to the high melting point of Nb its melting point is far beyond that of other alloying elements commonly used in copper and copper alloys.
  • Conventional production processes often fail to obtain beneficial effects.
  • Nb fails to dissolve into the copper matrix.
  • the solid solubility of Al in the copper matrix is about 8%.
  • the addition of Al will have a certain solid solution strengthening effect.
  • experiments have found that the addition of Al alone has no significant effect on the properties of titanium bronze.
  • 0.005-0.40wt% of Nb and 0.01-0.50wt% of Al are added to the titanium bronze alloy strip.
  • the inventor found that adding the amount of Nb and Al at the same time can significantly improve the strength of the titanium bronze alloy strip and the stability of its mechanical properties at high temperatures, and it can still ensure excellent bending performance.
  • a dispersed nano-scale intermetallic compound containing Nb and Al will be formed in the alloy matrix, which has a dispersion strengthening effect on the titanium bronze alloy. This strengthening effect is more obvious than adding Nb or Al alone to the improvement of the mechanical properties of the alloy.
  • These fine particles of Nb and Al-containing intermetallic compounds are dispersed in the alloy matrix, with a particle size of about 10nm to 10 ⁇ m.
  • the dispersed nanoparticles in the alloy will hinder the movement of dislocations and have the effect of dispersion strengthening, thereby improving the mechanical properties of the alloy.
  • intermetallic compounds containing Nb and Al are intermetallic compounds with high melting point and high stability. Its melting point can reach above 1900°C, and it will not interact with the copper matrix at high temperatures. Its strengthening effect still exists under temperature. Compared with the conventional titanium bronze alloy, the Cu-Ti-Nb-Al alloy of the present invention has significantly improved stability of the mechanical properties of the alloy at high temperatures.
  • the Nb content added to the titanium bronze alloy strip is less than 0.005wt% and the Al content is less than 0.01wt%, the number of intermetallic compound particles containing Nb and Al is less, and the mechanical properties of the alloy have no stability at high temperatures.
  • the improvement of the performance of the Cu-Ti-Nb-Al alloy of the present invention relative to the conventional titanium bronze alloy is mainly due to the dispersion strengthening effect of the high-stability nanoparticles.
  • the Nb content added to the titanium bronze alloy strip is greater than 0.40wt.% and the Al content is greater than 0.5wt%, the number of intermetallic compound particles containing Nb and Al in the alloy is too large, which is likely to occur during the preparation process.
  • the Nb content of the titanium bronze alloy strip of the present invention is 0.005-0.40 wt%, and the Al content is 0.01-0.5 wt%, and both elements need to be added at the same time. More preferably, the Nb content is 0.01-0.30 wt%, and the Al content is 0.05-0.3 wt%.
  • the number of Nb and Al-containing intermetallic compound particles with a particle size between 50-500nm in the titanium bronze alloy strip is not less than 1 ⁇ 10 5 /mm 2
  • the number of Al intermetallic compound particles is not higher than 1 ⁇ 10 3 particles/mm 2 .
  • the titanium bronze alloy of the present invention has a large number of dispersed and distributed fine granular intermetallic compounds containing Nb and Al inside the crystal grains.
  • the particle size of the intermetallic compound particles containing Nb and Al (the maximum size of the compound particle, the same below) is different in the number of particles between 50-500nm
  • the number of Nb and Al-containing intermetallic compound particles with a particle size of less than 1 ⁇ 10 5 particles/mm 2 and a particle size greater than 1 ⁇ m is not higher than 1 ⁇ 10 3 particles/mm 2 .
  • the dispersed nano-scale particles can pin the dislocations, effectively hinder the movement of the dislocations, limit the growth of crystal grains, and strengthen the alloy matrix.
  • due to the high stability of Nb and Al-containing intermetallic compounds at high temperatures their strengthening effect still exists at high temperatures.
  • the present invention has found that when the particle size of the intermetallic compound is too large, the agglomeration of the particles will increase, which will deteriorate the strength and bending performance of the material. Therefore, the number of intermetallic compound particles containing Nb and Al with a particle size greater than 1 ⁇ m is preferably not More than 1 ⁇ 10 3 pieces/mm 2 .
  • the inventor found that by controlling a certain number of nano-scale Nb and Al-containing intermetallic compound particles in the titanium bronze alloy matrix, the stability of the mechanical properties of the titanium bronze alloy at high temperatures can be further improved.
  • the test results of the finished product show: Cu-Ti-Nb-Al alloy has stable mechanical properties at high temperatures The performance has been significantly improved, and the conductivity of the alloy has also been improved. Therefore, the co-addition of Nb and Al can improve the stability of the mechanical properties of the titanium bronze alloy at high temperatures.
  • the average grain size of the titanium bronze alloy strip is less than or equal to 20 ⁇ m.
  • the metallographic phase of the conventional titanium bronze alloy without Nb and Al and only adding Nb or Al is shown in Figure 2-4: The average grain size is all 30 ⁇ m or more, except for a small amount of inclusions at the grain boundaries, there is nothing inside the grains Mass point.
  • the metallographic phase of the titanium bronze alloy containing Nb and Al of the present invention is shown in Figure 1: The average grain size is 18 ⁇ m, which is at least 40% lower than the prior art Cu-Ti alloy . During the alloy preparation process, the control of the grain size will directly affect the performance of the final product.
  • the crystal grain size is mainly controlled by adjusting the solution treatment temperature and time.
  • the processing time is shortened to a certain value, the allowable process error range will be drastically reduced, which will reduce the yield in production.
  • the growth of crystal grains is mainly accomplished by the migration of grain boundaries. Nano-scale Nb and Al-containing intermetallic compound particles stably exist in the matrix at high temperatures, which restrict the growth of the matrix grains by hindering the movement of the grain boundaries. Even if the solution time is long, the grain refinement effect is still very significant. This kind of grain refinement effect is very important to the improvement of product mechanical properties and yield.
  • the titanium bronze alloy strip has excellent high temperature stability. After the alloy strip is kept at 500°C for 1 hour in an atmospheric atmosphere, the hardness attenuation value H is less than 5%.
  • the high-temperature stability evaluation index of copper alloys is mainly the high-temperature softening temperature of copper alloys.
  • the national standard "GB/T33370-2016 Copper and Copper Alloy Softening Temperature Measurement Method" stipulates that after holding at a certain temperature for 1 hour, when the alloy hardness value drops to 80% of the original hardness, the corresponding holding temperature is the high temperature softening of copper alloy temperature.
  • the softening degree of the alloy is not linearly related to the holding temperature of the alloy.
  • the higher the temperature of the alloy the faster its performance changes.
  • the attenuation amplitude value of the alloy hardness at a certain holding temperature is used to characterize the stability of the mechanical properties of the titanium bronze alloy at high temperatures, which can more intuitively reflect the performance changes of the alloy at high temperatures, thereby facilitating the product processing process and Application design.
  • the hardness decay rate H of the conventional titanium bronze alloy is greater than 10% after being kept at 500°C for 1 hour in the atmosphere.
  • the hardness decay rate of the titanium bronze alloy of the present invention is much lower than that of the conventional titanium bronze alloy. This excellent high temperature stability enables the titanium bronze alloy strip to maintain stable performance in different processing and application scenarios, which is beneficial to expand the application range of the titanium bronze alloy strip.
  • one or more elements of Ni, Co, Fe, Sn, Mn, Si, Cr, Mg, B, Zr or Ag with a total weight percentage of not more than 0.50 wt% can be added to the titanium bronze alloy.
  • Ni, Co, Fe, Sn, Mn, Si, Cr, Mg, B will form intermetallic compounds with Nb and Al to further improve the stability of the strip, but adding too much of these elements will reduce the precipitation of CuTi Quantity, which will reduce the mechanical properties of the strip.
  • Zr and Ag can be dissolved in copper to increase the strength of the strip without reducing the conductivity. Therefore, the total amount of Ni, Co, Fe, Sn, Mn, Si, Cr, Mg, B, Zr or Ag and combinations thereof in the titanium bronze alloy strip of the present invention does not exceed 0.50 wt%.
  • the titanium bronze alloy strip of the present invention has a closed composition.
  • the balance of the titanium bronze alloy strip is Cu And inevitable impurities. If any element other than the above-mentioned elements is added, even if it is a small amount, it will have an adverse effect on the overall performance of the titanium bronze alloy strip, especially the bending performance, yield strength and high temperature stability.
  • the present invention also relates to a method for preparing a titanium bronze alloy strip containing Nb and Al as described above, including the following steps:
  • Thermal processing thermal processing of the ingot at a temperature of 700-980°C, to control the reduction of the cross-sectional area of the ingot thermal processing by not less than 75%;
  • Milling surface milling the surface of the material obtained by thermal processing
  • the first cold rolling control the cross-sectional area of the material to be reduced by not less than 70%;
  • the first aging select inactive atmosphere protection and keep it in the temperature range of 350-500°C for 0.5-24h;
  • the second aging select inactive atmosphere protection and keep it in the temperature range of 200-550°C for 1min-10h.
  • the casting method in step 1) is iron mold casting, horizontal continuous casting or vertical semi-continuous casting.
  • the hot working in step 2) is hot forging, hot rolling, or a combination of both.
  • the hot forging holding temperature is controlled at 700-980°C
  • the holding time is 1-12h
  • the initial forging temperature is controlled at 700-980°C. Free forging or die forging is used. When the temperature drops and deformation is difficult Reheating to increase the temperature of the billet.
  • the hot rolling holding temperature is controlled at 700-980°C, the holding time is 1-12h, the initial rolling temperature is controlled at 700-980°C, the hot rolling speed is 5-200m/min, and the final rolling
  • the temperature is above 500°C, the rolling reduction rate is controlled above 75%, and the on-line water cooling is performed after rolling.
  • the final rolling temperature is lower than 500°C, since the rolled piece is thin and long in the later stage of hot rolling, the temperature drop will cause the large temperature difference between the head and tail of the rolled piece, and the second phase will be precipitated, resulting in uneven structure, reducing material plasticity, and easy to generate cracks. So that it cracked.
  • multi-pass cold rolling is performed in step 6), wherein the deformation amount in a single pass is controlled at 5%-20%.
  • the crystal rotation promotes the proliferation of dislocations and the disordered arrangement of atoms.
  • the increased energy storage and lattice defects in the material are conducive to the progress of amplitude modulation decomposition or the precipitation of strengthening phases during the aging process, which can significantly increase the strength of the alloy.
  • the deformation amount in a single pass is controlled at 5%-20%, so that the force in the thickness direction during the rolling deformation is more uniform, which is beneficial to control the plate shape.
  • the solution treatment of step 5) and the intermediate cold rolling of step 6) are used as a step unit, and the step unit is repeated at least twice, wherein the intermediate cold-rolled material between the two adjacent solution treatments The cross-sectional area is reduced by ⁇ 30%.
  • the aging in step 7) is performed in an atmosphere containing hydrogen, nitrogen, argon, or any mixture of these gases.
  • the polishing and pickling steps for removing surface oxide scale are performed.
  • step 1) the vacuum smelting method is used.
  • the first step is to add electrolytic copper and Nb-containing master alloy in the smelting furnace at the same time to start smelting;
  • the second step after the electrolytic copper and Nb-containing master alloy are completely melted, add the Ti, Al-containing raw materials and optionally one or more raw materials containing one or more of Ni, Co, Fe, Sn, Mn, Si, Cr, Mg, B, Zr and Ag;
  • Step 3 All After the raw materials are melted, they are refined at 1300 ⁇ 50°C for 30-60min.
  • the melting point of Nb is as high as 2469°C, and its solid solubility in Cu is very low.
  • Adding the Nb-containing master alloy and electrolytic copper to the melting furnace at the same time can maximize the smelting time of Nb, thereby promoting the melting of Nb. If the smelting time of Nb is too short, elemental Nb particles with larger sizes are likely to appear in the ingot, which affects the quality of the ingot.
  • the refining process in step 1) will directly affect the stability of the mechanical properties of the titanium bronze strip of the present invention at high temperatures, and a suitable refining time is helpful for the generation of nano-level Nb and Al-containing intermetallic compounds, and It is beneficial to the dispersion and distribution of nano-level Nb and Al-containing intermetallic compound particles in the ingot.
  • the refining time is too short, a sufficient amount of Nb and Al-containing intermetallic compounds cannot be generated; if the refining time is too long, the nano-level Nb and Al-containing intermetallic compound particles are prone to agglomeration and growth, which will affect The properties of the final alloy.
  • the Nb-containing master alloy may be a Cu-Nb master alloy or a Nb-Ti master alloy
  • the Ti-containing and Al-containing raw material may be pure Ti, pure Al or a Ti and/or Al-containing master alloy, containing Ni
  • One or more of the raw materials of one or more of, Co, Fe, Sn, Mn, Si, Cr, Mg, B, Zr, and Ag may be simple substances of these elements or intermediate alloys containing these elements.
  • step 7) and step 9) the alloy is aged twice.
  • the main purpose of the first aging is to form an amplitude-modulated decomposition structure, increase the precipitation of ⁇ -Cu 4 Ti phase, and achieve a strengthening effect.
  • cold deformation will produce a large number of movable dislocations in the alloy. These dislocations are more likely to move at high temperatures and will greatly Affect the stability of the mechanical properties of the alloy at high temperatures.
  • the second time efficiency effectively reduces the density of movable dislocations in the alloy caused by the last cold rolling, thereby improving the stability of the structure and performance of the titanium bronze strip at room temperature and high temperature.
  • the above steps 1)-9) must be carried out in the order shown. If you change the order of the steps shown or reduce one or more of the above steps, or replace one or more of the above steps with other steps, it will affect the overall performance of the titanium bronze alloy strip, especially at high temperatures. The stability of mechanical properties has a significant impact.
  • the high-temperature stability of the titanium bronze alloy strip containing Nb and Al of the present invention is excellent: after being kept at 500° C. in the atmosphere for 1 hour, the alloy hardness decay rate H ⁇ 5%.
  • the titanium bronze alloy strip containing Nb and Al of the present invention can realize the ratio of the bending radius parallel to the rolling direction (that is, the good direction) to the strip thickness R 1 / T ⁇ 0.5, which is perpendicular to the rolling direction ( That is, the ratio of the bending radius in the bad direction to the thickness of the strip R 2 /T ⁇ 1.0.
  • This excellent bending performance enables the titanium bronze alloy strip to undergo severe bending in different directions at the same time, making it suitable for the preparation of small and complex-shaped terminals for consumer electronics and other connector-related industries. .
  • the "strip" is a common material form in the field, and its thickness is usually not greater than 1 mm.
  • Fig. 1 is the metallographic structure of the Cu-Ti-Nb-Al alloy strip according to the present invention.
  • Figure 2 shows the metallographic structure of a Cu-Ti alloy strip in the prior art.
  • Figure 3 shows the metallographic structure of a Cu-Ti-Nb alloy strip in the prior art.
  • Figure 4 shows the metallographic structure of a Cu-Ti-Al alloy strip in the prior art.
  • Fig. 5 is a scanning electron micrograph of an intermetallic compound containing Nb and Al in a Cu-Ti-Nb-Al alloy strip according to the present invention.
  • the first step adding electrolytic copper and Cu-Nb master alloy in the smelting furnace at the same time, and start smelting;
  • Step 2 After the electrolytic copper and Cu-Nb master alloy are completely melted, add pure Ti, pure Al and selected from Ni, Co, Fe, Sn, Mn, Si, Cr, Mg, B, Zr and The simple substance of the optional element of Ag;
  • the third step after all the raw materials are melted, refining at 1300 ⁇ 50°C for 30-60min. After smelting, a rectangular ingot is cast by a vertical semi-continuous casting method.
  • the ingot is kept at 800-950°C for 1-12h and then hot rolled, the hot rolling speed is 50-120m/min, the single pass processing rate of rolling is controlled between 10-30%, and the final rolling temperature is 650°C Above, after hot rolling, it is water-cooled on-line, and face milling is performed after hot rolling.
  • the first cold rolling is carried out, and the total cold rolling processing rate is controlled above 80%.
  • the solution temperature is 700-950°C
  • the holding time is 1-100s
  • the cooling rate is 10-250°C/s.
  • the solid solution is performed twice, the solution temperature is 700-950°C, the holding time is 1-100s, and the cooling rate is 10°C/s-250°C/s.
  • the intermediate cold rolling is performed again, the rolling rate is controlled at 10-60%, and the deformation of a single pass is controlled at 5-20%.
  • the rolling rate can be in the range of 5-99% according to the actual product specifications. Internal change, and the number of solution treatment and intermediate cold rolling can be once or more than twice.
  • the first aging is then carried out in an atmosphere containing hydrogen, nitrogen, argon, or any mixture of these gases.
  • the aging temperature is 400°C and the holding time is 4h.
  • the final cold rolling is carried out, and the rolling rate is controlled at 10-30%. It should be noted that although a specific rolling rate is involved in the final cold rolling step here, the rolling rate can vary in the range of 5-80% according to the actual product specifications.
  • the second aging is carried out in an atmosphere containing hydrogen, nitrogen, argon, or any mixture of these gases.
  • the aging temperature is 350°C and the holding time is 4h.
  • the number of Nb and Al-containing intermetallic compound particles with a particle size between 50-500 nm and a particle size> 1 ⁇ m in the alloy was measured, and the mechanical properties, electrical conductivity, bending properties and the resistance of the alloy strips obtained were tested. Stability of mechanical properties at high temperature.
  • the thickness specification of the finished product of Example 12 is 0.15mm, and the specific process is as follows:
  • Step 1 Add electrolytic copper and Cu-Nb master alloy into the smelting furnace at the same time to start smelting;
  • Step 2 After the electrolytic copper and Cu-Nb master alloy are completely melted, add pure Ti, pure Al and pure Co in sequence;
  • the ingot is kept at 930°C for 8 hours and then hot rolled.
  • the hot rolling speed is 110m/min, the single pass processing rate is 30%, and the final rolling temperature is above 650°C.
  • online water cooling is performed.
  • hot rolling Perform face milling.
  • the first cold rolling was carried out, and the total cold rolling processing rate was 90%.
  • the solution temperature is 700°C
  • the holding time is 80s
  • the cooling rate is 100°C/s.
  • the solid solution is performed twice, the solution temperature is 950°C, the holding time is 5s, and the cooling rate is 200°C/s.
  • the intermediate cold rolling is performed again, the rolling rate is controlled at 20%, and the deformation of a single pass is controlled at 5%.
  • the first aging is then carried out in an atmosphere containing a mixture of hydrogen and argon.
  • the aging temperature is 400°C and the holding time is 4h.
  • the final cold rolling is performed, the rolling rate is 20%, and the final thickness is 0.15mm.
  • the second aging is performed in an atmosphere containing a mixture of hydrogen and argon, the aging temperature is 350° C., and the holding time is 4 hours to obtain the finished material.
  • the room temperature tensile test is carried out on the electronic universal testing machine in accordance with "GB/T228.1-2010 Metallic Material Tensile Test Part 1: Room Temperature Test Method".
  • the sample adopts a rectangular cross-section proportional sample with a proportionality factor of 5.65.
  • the yield strength of the strips of the inventive examples and comparative examples given in Table 1 below is the yield strength of the parallel rolling direction.
  • the conductivity test is in accordance with "GB/T3048-2007 Test Method for Electrical Properties of Wires and Cables Part 2: Resistivity Test of Metallic Materials", expressed in %IACS.
  • the bending performance is measured by the following method: take a long strip of copper alloy strip along the rolling direction (that is, the good direction), and take a long strip perpendicular to the rolling direction (that is, the bad direction), the sample width is 10mm, and then use A 90° V-shaped punch with different radii at the tip is used to bend the long strip, and then use a stereo microscope to observe the outer surface of the bend, and use the minimum bending radius/thickness (R/T) that does not produce cracks on the surface. Indicates bending performance. When the R/T value is 0, it means that the minimum bending radius R is 0, and the bending performance is the best.
  • the average grain size is measured in accordance with the test method of "YS/T 347-2004 Copper and Copper Alloy Average Grain Size Measurement Method".
  • the stability test of mechanical properties at high temperature is carried out with reference to "GB/T33370-2016 Copper and Copper Alloy Softening Temperature Determination Method".
  • the sample is kept in an air atmosphere of 500°C for 1 hour and then air-cooled to test the hardness of the sample.
  • the hardness attenuation rate H (%) of the sample after being kept at a certain temperature and high temperature indicates the stability of the mechanical properties of the sample at high temperature.
  • the specific operation method is as follows: take a cross-section parallel to the rolling direction of the copper alloy strip, and observe its structure with a 25 ⁇ m ⁇ 40 ⁇ m (1000 ⁇ m 2 ) rectangle as the basic unit; select 10 rectangles at different positions in the field of view, and measure each In the rectangle, the number of particles with a particle size between 50-500nm and the number of particles with a particle size greater than 1 ⁇ m are counted, and finally the average value is taken as the judgment basis.
  • the particle size is defined as the maximum size of the particles.
  • the copper alloys of all the examples have achieved yield strength ⁇ 900 MPa, electrical conductivity ⁇ 10% IACS, and alloy bending.
  • Excellent performance that is, the ratio of bending radius to strip thickness (R 1 / T) ⁇ 0.5, perpendicular to the rolling direction (ie bad direction), the ratio of bending radius to strip thickness (R 2 /T) ⁇ 1.0.
  • the hardness decay rate H of the alloy samples of Examples 1-20 is less than 5%.
  • Examples 1-20 and Comparative Examples 1-10 reflect the effects of different Nb and Al contents and the number of Nb and Al-containing intermetallic compound particles on the overall performance of the titanium bronze alloy strip. At the same time, Examples 1-20 also show that adding one or more optional elements selected from Si, Zn, Co, Fe, Sn, Mn, Mg, Cr, B, Ag and Zr in a reasonable small amount is important for alloy strength and High temperature stability can be improved to a certain extent.
  • composition, number of Nb and Al-containing intermetallic compound particles and performance test results of the titanium bronze alloy strips of Examples 1-20 and Comparative Examples 1-10 are shown in Table 1.
  • Comparative examples 8-10 show that although the hardness decay rate H ⁇ 5%, due to the excessive content of Al and/or Nb, the yield strength and bending performance of the titanium bronze alloy are adversely affected. Especially when Al and Nb are excessive at the same time, they will agglomerate into large precipitate particles, which is not conducive to improving the strength of the alloy. At the same time, it will increase the risk of cracking during bending (R 1 /T and R 2 /T of Comparative Example 10 are both large ).

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Conductive Materials (AREA)

Abstract

本发明公开了一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材的重量百分比组成包括2.00-4.50wt％的Ti，0.005-0.4wt％的Nb，和0.01-0.5wt％的Al，余量为Cu和不可避免的杂质。优选地，该钛青铜合金带材的微观组织中粒径在50-500nm之间的含Nb和Al的金属间化合物粒子的数量不低于1×105个/mm2，粒径大于1μm的含Nb和Al的金属间化合物粒子的数量不高于1×103个/mm 2。在确保优异的折弯性能的情况下，该钛青铜合金带材具有优异的稳定性，尤其是在高温下力学性能的稳定性。本发明还涉及该钛青铜合金带材的制备方法。

Description

一种含Nb和Al的钛青铜合金带材及其制备方法 技术领域
本发明属于铜合金材料技术领域，特别涉及一种含有Nb和Al的钛青铜合金带材。该钛青铜合金带材具有优异的稳定性，尤其是在高温下力学性能的稳定性。本发明还涉及该钛青铜合金带材的制备方法。
背景技术
随着消费电子以及其它连接器相关行业产品的小型化与多功能化的快速发展，设计者们需要选用强度更高、成形性能更加优异的铜合金材料来制造其中的接触件，以满足其终端产品的轻薄化和小型化的设计需求。在现有铜合金体系中，作为高强度、高传导代表的铍铜合金能够满足上述性能要求，但因成本及含铍材料在加工过程中产生剧毒物质的问题限制了材料的使用。钛青铜合金是以钛为主要合金元素的铜合金，其强度较高并且成形性能较优异，在一些应用场合可以用于替代铍铜合金。
钛青铜是一种调幅分解强化加时效析出强化型合金，主要强化组织为调幅分解组织与β`-Cu 4Ti相。在时效处理早期，钛青铜合金的强化方式为调幅分解强化，铜基体中固溶的Ti原子经扩散，在晶粒内形成周期性的Ti原子富集区，即调幅分解组织。随时效过程的继续，调幅分解组织逐渐转变为周期性排列的β`-Cu 4Ti相。然而，调幅分解组织与β`-Cu 4Ti相在高温下稳定性较差，容易发生演变，从而对合金的力学性能造成不利影响，温度越高其性能的恶化越快。在材料加工以及应用过程中，材料性能的稳定性至关重要，良好的稳定性能确保产品在加工与应用过程中出现突然过载以及高温时不会快速失效。由于钛青铜具有高强度以及优异的弹性性能，其在电动汽车、5G通讯基站等领域有广泛的应用前景。在这些领域，尤其是电动汽车领域中往往会出现瞬时或持续的高温工况，温度可能达到200℃以上。如果开发材料时，未考虑材料在高温下的力学性能稳定性以及材料在高温工况下使用后性能的变化情况，这会使得由该材料制备的元器件在高温工况下的使用寿命存在不确定性，并且元器件甚至存在突然失效的风险，从而造成较大的安全隐患。因此在设计钛青铜合金材料体系时仅仅针对常规的强度、电导率、加工性能等进行调控并不能完全满足材料后续的各种加工与应用场景的要求。在考虑常规性能指标的同时，还应考虑钛青铜合金材料性能的稳定性，尤其是在高温下力学性能的稳定性。
截止目前，经本发明人检索，在现有技术中没有发现关于钛青铜合金带材在高温下的力学性能稳定性的研究。
发明内容
本发明通过在钛青铜中同时添加一定量的Nb和Al，设计了一种Cu-Ti-Nb-Al体系合金。相对于常规钛青铜合金，在确保优异的折弯性能的同时，该Cu-Ti-Nb-Al体系合金在高温下的力学性能稳定性得到明显提升，并且合金强度也得到提升。
本发明所要解决的技术问题是：针对现有技术的不足，如何在确保钛青铜合金带材的优异力学性能及折弯性能的同时，使合金带材具有优化的稳定性、尤其是在高温下力学性能的稳定性。
本发明解决上述技术问题所采用的技术方案为：一种含有Nb和Al的钛青铜合金带材，该钛青铜合金带材的重量百分比组成包括：2.0-4.5wt％的Ti，0.005-0.4wt％的Nb，0.01-0.5wt％的Al，余量为Cu和不可避免的杂质。
本发明在钛青铜合金带材中添加2.0-4.5wt％的Ti。Ti有助于改善钛青铜合金的力学性能。当添加的Ti含量小于2.0wt％时，钛青铜合金带材虽然具有较高的电导率但得不到理想的力学性能，从而应用受限。当添加的Ti含量超过4.5wt％时，过高的Ti含量会降低合金带材的电导率并显著劣化其加工性能，尤其劣化折弯性能。因此，本发明的钛青铜合金带材的Ti含量为2.0-4.5wt％。作为优选，钛青铜合金带材的Ti含量为2.5-4.0wt％。进一步优选，钛青铜合金带材的Ti含量为2.9-3.5wt％。
本发明中Ti是主要强化元素，在时效过程中首先通过固溶体中Ti原子的扩散形成调幅分解组织，此时铜合金的强度显著升高；随时效时间的增加，基体中逐渐析出针状的β`-Cu 4Ti相，在此过程中时效强化效果逐渐达到峰值；随着时效时间的进一步延长，在晶界上会析出片状β-Cu 4Ti相，其体积分数随时间延长而逐渐增大，最终代替β`-Cu 4Ti相，在此过程中铜合金的强化效果逐渐下降。调幅分解组织为均匀的纳米尺度的组织，β`-Cu 4Ti相也为纳米级沉淀相，弥散分布在基体中，这两种组织均能阻碍晶界和位错的移动，使铜合金强度提高。通过控制时效过程，是指形成不同的微观组织，可有效调控合金的综合性能。
现有技术表明，在钛青铜合金中可以任选地添加少量的Nb和Al中的任一者作为次要的合金化元素。一方面，单独添加Nb时，其可以少量固溶在铜基体中，对合金的强度略有提升，但对其他性能影响不明显。但是由于Nb熔点较高，其熔点远远超 出铜及铜合金中常用的其他合金化元素，采用常规工艺生产往往并不能获得有益的效果，反而因为Nb未能固溶到铜基体中而影响合金的应用性能。另一方面，Al在铜基体中的固溶度在8％左右，理论上Al的添加会有一定的固溶强化效果。然而，试验发现单独添加Al对钛青铜性能实际上并没有明显的影响。
本发明在钛青铜合金带材中添加0.005-0.40wt％的Nb和0.01-0.50wt％的Al。本发明人发现，同时添加所述量的Nb、Al对钛青铜合金带材的强度及其在高温下力学性能的稳定性有显著提升效果，并且仍然能够保证优异的折弯性能。实验发现，同时添加Nb与Al后，在合金基体中会形成弥散分布的纳米尺度的含Nb和Al的金属间化合物，对钛青铜合金有弥散强化的效果。这种强化效果相比单独添加Nb或Al对合金力学性能的提升都要明显。这些细小颗粒状的含Nb和Al的金属间化合物，弥散分布在合金基体中，其粒径约为10nm至10μm。合金中弥散分布的纳米颗粒会对位错的运动有阻碍作用，具有弥散强化的效果，从而提升合金的力学性能。
更为重要的是，含Nb和Al的金属间化合物是高熔点高稳定性的金属间化合物，其熔点可达到1900℃以上，并且在高温下不会与铜基体发生界面反应，因而在较高温度下其强化作用仍然存在。本发明的Cu-Ti-Nb-Al合金与常规钛青铜合金相比，合金在高温下力学性能的稳定性得到显著提升。
当钛青铜合金带材中添加的Nb含量小于0.005wt％、Al含量小于0.01wt％时，生成的含Nb和Al的金属间化合物粒子的数量较少，合金在高温下力学性能的稳定性无明显改善。本发明的Cu-Ti-Nb-Al合金相对常规钛青铜合金性能的提升主要归因于高稳定性纳米粒子的弥散强化作用。然而，当钛青铜合金带材中添加的Nb含量大于0.40wt.％、Al含量大于0.5wt％时，合金中含Nb和Al的金属间化合物粒子的数量过多，在制备的过程中易发生粒子的团聚，最终对合金的性能(尤其是屈服强度和折弯性能)产生不利影响。因此，本发明的钛青铜合金带材的Nb含量为0.005-0.40wt％、Al含量为0.01-0.5wt％，两种元素需同时添加。进一步优选，Nb含量为0.01-0.30wt％、Al含量为0.05-0.3wt％。
作为优选，该钛青铜合金带材中粒径在50-500nm之间的含Nb和Al的金属间化合物粒子的数量不低于1×10 5个/mm 2，粒径大于1μm的含Nb和Al的金属间化合物粒子的数量不高于1×10 3个/mm 2。如图5的扫描电镜照片所示，本发明的钛青铜合金的晶粒内部存在大量弥散分布的细小颗粒状的含Nb和Al的金属间化合物。经研究，在本发明的钛青铜合金带材中有利的是：含Nb和Al的金属间化合物粒子的粒径(化 合物粒子的最大尺寸，下同)在50-500nm之间的粒子的数量不低于1×10 5个/mm 2，且粒径大于1μm的含Nb和Al的金属间化合物粒子的数量不高于1×10 3个/mm 2。弥散分布的纳米尺度粒子能钉扎位错，有效阻碍位错的运动，限制晶粒的长大，强化合金基体。重要的是，由于含Nb和Al的金属间化合物在高温下的高稳定性，其强化作用在高温下仍然存在。本发明发现当金属间化合物粒子粒径过大，由于粒子团聚作用增大，反而会恶化材料的强度及折弯性能，因此粒径大于1μm的含Nb和Al的金属间化合物粒子的数量优选不高于1×10 3个/mm 2。本发明人发现通过在钛青铜合金基体中控制一定数量的一定纳米尺度的含Nb和Al的金属间化合物粒子，可进一步提升钛青铜合金在高温下的力学性能稳定性。
申请人希望强调：Nb和Al的协同作用为本发明中Cu-Ti合金体系高温力学性能的稳定性提升的最主要因素。实验发现：在Cu-Ti合金体系中，单独添加Nb时，合金强度有所提升，但对合金在高温下的力学性能稳定性无改善；单独添加Al时，合金各方面性能无明显改善；同时添加Nb和Al时，Cu-Ti-Nb-Al合金基体中生成弥散分布的含Nb和Al的金属间化合物粒子，成品测试结果显示：Cu-Ti-Nb-Al合金在高温下的力学性能稳定性得到明显提升，同时合金电导率也得到提升。因此Nb和Al的共同添加，能够提升钛青铜合金在高温下的力学性能稳定性。
该钛青铜合金带材的平均晶粒度≤20μm。不含Nb和Al、仅添加Nb或Al的常规钛青铜合金的金相如图2-4所示：平均晶粒尺寸均为30μm以上，除晶界处有少量夹杂物之外，晶粒内部无任何质点。相比之下，在经相同工艺后，本发明的含有Nb和Al的钛青铜合金的金相如图1所示：平均晶粒尺寸为18μm，较现有技术的Cu-Ti合金降低了至少40％。在合金制备过程中，晶粒度大小的控制将会直接影响最终产品的性能。在常见的铜合金制备过程里，主要通过固溶处理温度与时间的调节来调控晶粒度大小。然而，当处理时间缩短至一定值后，允许的工艺误差范围会急剧降低，从而会降低生产中的成品率。晶粒的长大主要通过晶界的迁移完成，纳米尺度的含Nb和Al的金属间化合物粒子在高温下稳定存在于基体中，其通过阻碍晶界的运动从而限制基体晶粒的长大。即使固溶时间较长，其晶粒细化效果仍然很显著，这种细晶效果对产品力学性能与成品率的提升都至关重要。
如上所述，该钛青铜合金带材具有优异的高温稳定性。该合金带材在大气气氛中在500℃下保温1h后，硬度衰减值H小于5％。现有技术中铜合金高温稳定性能的评价指标主要为铜合金高温软化温度。国标《GB/T33370-2016铜及铜合金软化温度的 测定方法》中规定，在一定温度下保温1h后，合金硬度值下降到原始硬度的80％时所对应的保温温度为铜合金的高温软化温度。然而合金的软化程度与合金保温温度并非呈线性关系，通常合金所处温度越高其性能变化越快。随着产品加工工艺与应用的日益复杂化，仅仅考虑合金的高温软化温度可能并不能满足产品设计与应用的需求。本发明中，采用在一定保温温度下合金硬度的衰减幅度值表征钛青铜合金在高温下的力学性能稳定性，能更直观地反应合金在高温下的性能变化情况，从而有利于产品加工过程以及应用的设计。常规钛青铜合金在大气气氛中在500℃下保温1h后，合金的硬度衰减率H大于10％。本发明的钛青铜合金的硬度衰减率远低于常规钛青铜合金。这种优异的高温稳定性使得该钛青铜合金带材能够在不同的加工和应用场景中保持性能的稳定，有利于扩大钛青铜合金带材的应用范围。
作为优选，该钛青铜合金中还可以添加总重量百分比不超过0.50wt％的Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr或Ag中的一种或多种元素。其中Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B会与Nb和Al形成金属间化合物从而可以进一步提高带材的稳定性，但加入过多的这些元素会降低CuTi沉淀相的数量，从而会降低带材的力学性能。Zr、Ag能够固溶在铜中，从而提高带材强度，但又不会降低导电性。因此，本发明的钛青铜合金带材中的Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr或Ag及其组合的总量不超过0.50wt％。
需要指出的是，本发明的钛青铜合金带材具有封闭式的组成。除了以上提及的必要元素Ti、Nb、Al以及任选元素Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr或Ag以外，该钛青铜合金带材的余量为Cu和不可避免的杂质。如果添加上述元素以外的任何元素，即使是微量，也将对钛青铜合金带材的综合性能、尤其是折弯性能、屈服强度和高温稳定性带来不利的影响。
本发明还涉及如上所述的一种含有Nb和Al的钛青铜合金带材的制备方法，包括以下步骤：
1)熔铸：采用真空或气氛保护的熔炼方法，在1200-1400℃将铜合金原料熔化；
2)热加工：在700-980℃温度下对铸锭进行热加工，控制铸锭热加工的横断面面积缩减不低于75％；
3)铣面：将热加工获得的材料铣面；
4)第一次冷轧：控制材料的横断面面积缩减不低于70％；
5)固溶处理：将冷轧后的材料加热至700-950℃的温度并且保温1-100s的时间， 随后进行水冷或气冷处理，冷却速度控制在10-250℃/s；
6)中间冷轧：控制材料的横断面面积缩减5-99％；
7)第一次时效：选择非活性气氛保护在350-500℃的温度范围内保温0.5-24h；
8)最终冷轧：控制横断面面积缩减5-80％。
9)第二次时效：选择非活性气氛保护在200-550℃的温度范围内保温1min-10h。
作为优选，步骤1)中的铸造方式为铁模铸造、水平连续铸造或垂直半连续铸造。
作为优选，步骤2)中的热加工为热锻、热轧、或两者的结合。
进一步优选，在上述热锻中，热锻保温温度控制在700-980℃，保温时间1-12h，初始锻造温度控制在700-980℃，采用自由锻或模锻，当温度降低、变形困难时回炉加热以提高坯料温度。
又进一步优选，在上述热轧中，热轧保温温度控制在700-980℃，保温时间1-12h，初始轧制温度控制在700-980℃，热轧速度5-200m/min，并且终轧温度在500℃以上，轧制压下率控制在75％以上，轧制后在线水冷却。如果终轧温度低于500℃，由于热轧后期轧件薄而长，温降大使轧件头尾与中间温差大，会有第二相析出，产生组织不均匀，降低材料塑性，容易生成裂纹以致开裂。作为优选，在步骤6)中进行多道次冷轧，其中单道次变形量控制在5％-20％。
轧制过程中晶体转动促进位错的增殖和原子的错乱排布，材料中增加的储能和晶格缺陷等有利于时效过程中调幅分解的进行或强化相的析出，可显著提高合金的强度。单道次变形量控制在5％-20％，使得轧制变形中沿厚度方向上的受力更均匀，有利于控制板型。
作为优选，以步骤5)的固溶处理和步骤6)的中间冷轧为一个步骤单元，重复实施该步骤单元至少两次，其中相邻两次固溶处理之间的中间冷轧的材料的横断面面积缩减≥30％。
作为优选，在步骤7)中时效在包含氢气、氮气、氩气、或这几种气体的任何混合物的气氛中进行。
作为优选，在固溶处理之后和/或在时效之后进行用以去除表面氧化皮的研磨、酸洗工序。
需要对以上方法中的关键步骤进行如下说明：
在步骤1)中，采用真空熔炼方法，第一步：在熔炼炉中同时加入电解铜和含Nb中间合金，开始熔炼；第二步：电解铜和含Nb中间合金完全熔化后，依次加入含Ti、 含Al原料和任选的含Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag中一者或多者的一种或多种原料；第三步：所有原材料熔化后，在1300±50℃下精炼30-60min。Nb熔点高达2469℃，且在Cu中固溶度很低，将含Nb中间合金与电解铜同时加入熔炼炉中可使得Nb的熔炼时间最大化，从而促进Nb的熔化。若Nb的熔炼时间过短，在铸锭中易出现尺寸较大的单质Nb颗粒，影响铸锭的品质。需要特别强调，步骤1)中的精炼工序会直接影响本发明钛青铜带材高温下的力学性能稳定性，合适的精炼时间有助于纳米级别的含Nb和Al的金属间化合物的生成，并有利于纳米级别的含Nb和Al的金属间化合物粒子在铸锭中的弥散分布。如果精炼时间过短，则不能生成足够量的含Nb和Al的金属间化合物；如果精炼时间过长，则纳米级别的含Nb和Al的金属间化合物粒子易发生聚集长大，这都将影响最终合金的性能。
在步骤1)中，含Nb中间合金可以是Cu-Nb中间合金或Nb-Ti中间合金，含Ti、含Al原料可以是纯Ti、纯Al或者含Ti和/或Al的中间合金，含Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag中一者或多者的一种或多种原料可以是这些元素的单质或含这些元素的中间合金。
在步骤7)、步骤9)中，对合金进行两次时效处理。第一次时效的主要目的是形成调幅分解组织，增加β`-Cu 4Ti相的析出，达到强化效果。第一次时效后为了进一步强化合金，需要对合金进行一次冷轧工艺处理，但是冷变形会在合金内部产生大量可动位错，这些位错在高温下更易发生移动，将会极大程度地影响合金在高温下力学性能的稳定性。第二次时效能有效降低最后一次冷轧在合金中造成的可动位错的密度，从而提高钛青铜带材在室温以及高温下组织与性能的稳定性。
以上步骤1)-9)必须按所示顺序进行。如果调换所示步骤的顺序或者减少上述步骤中的一个或多个步骤或者用其他步骤替换上述步骤中的一个或多个步骤，都将对钛青铜合金带材的综合性能、尤其是在高温下力学性能的稳定性产生显著影响。
本发明的有益效果
与现有技术相比，本发明的优点在于：
(1)本发明的含有Nb和Al的钛青铜合金带材的高温稳定性能优异：在大气气氛中在500℃下保温1h后，合金硬度衰减率H＜5％。
(2)本发明的含有Nb和Al的钛青铜合金带材可以实现平行于轧制方向(即好方向)的弯曲半径与带材的厚度比R 1/T≤0.5，垂直于轧制方向(即坏方向)的弯曲半径与带材的厚度比R 2/T≤1.0。这种优异的折弯性能使得该钛青铜合金带材能够在不同 方向同时经受严苛的折弯成形，从而适合于制备用于消费电子以及其它连接器相关行业的小型化且具有复杂形状的端子。
在本文中，所述“带材”是本领域中常见的材料形式，其厚度通常不大于1mm。
除非另外指出，否则说明书和权利要求书中使用的表示成分的量、化学和力学性质、工艺条件等的所有数字应理解为在所有情况下都由术语“约”修饰。因此，除非相反地指出，说明书和所附权利要求书中阐述的数值参数是可取决于通过本文示例性实施方案寻求获得的期望性质而变化的近似值。至少每个数值参数应当根据有效数字的数值和普通的舍入方法来解释。
尽管阐述示例性实施方案的宽泛范围的数值范围和参数是近似值，但尽可能精确地报告在具体实施例中阐述的数值。然而，任何数值固有地含有由在它们各自的测试测量中发现的标准偏差所必然产生的某些误差。在整个说明书和权利要求书中给出的每个数值范围将包括落入这样的较宽数值范围内的每个较窄的数值范围，如同在本文中也明确地写出这样的较窄的数值范围。此外，实施例中报告的任何数值可用于限定本文公开的较宽组成范围的上端点或下端点。
附图说明
图1为根据本发明的Cu-Ti-Nb-Al合金带材的金相组织。
图2为现有技术的Cu-Ti合金带材的金相组织。
图3为现有技术的Cu-Ti-Nb合金带材的金相组织。
图4为现有技术的Cu-Ti-Al合金带材的金相组织。
图5为根据本发明的Cu-Ti-Nb-Al合金带材中含Nb和Al的金属间化合物的扫描电镜图片。
具体实施方式
以下结合附图和实施例对本发明作进一步详细描述。
设计了20个实施例和10个对比例合金。每一种合金根据合金原料添加量的要求(参见下表1)采用前述两步添加合金原料的熔炼方法，第一步：在熔炼炉中同时加入电解铜和Cu-Nb中间合金，开始熔炼；第二步：电解铜和Cu-Nb中间合金完全熔化后，根据表1成分依次加入纯Ti、纯Al以及选自Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag的可选元素的单质；第三步：所有原材料熔化后，在1300±50℃下 精炼30-60min。熔炼后，通过垂直半连续铸造法铸造出矩形铸锭。
将铸锭在800-950℃下保温1-12h后进行热轧，热轧速度为50-120m/min，轧制单道次加工率控制在10-30％之间，终轧温度在650℃以上，热轧后在线水冷却，热轧后进行铣面。
随后进行第一次冷轧，冷轧总加工率控制在80％以上。
第一次冷轧后进行固溶，固溶温度为700-950℃，保温时间为1-100s，冷却速度10-250℃/s。
固溶后进行中间冷轧，轧制率控制在30-60％，单道次变形量控制在5-20％。
中间冷轧后进行二次固溶，固溶温度为700-950℃，保温时间为1-100s，冷却速度10℃/s-250℃/s。
二次固溶后再次进行中间冷轧，轧制率控制在10-60％，单道次变形量控制在5-20％。
需要说明的是：尽管上面的中间冷轧步骤中涉及特定的轧制率和两次固溶处理与两次中间冷轧，但是根据实际成品规格要求，轧制率可在5-99％的范围内变化，并且固溶处理与中间冷轧次数可为一次或两次以上。
随后进行第一次时效，在包含氢气、氮气、氩气、或这几种气体的任何混合物的气氛中进行，时效温度为400℃，保温时间4h。
第一次时效后进行最终冷轧，轧制率控制在10-30％。需要说明的是：尽管此处最终冷轧步骤中涉及特定的轧制率，但是根据实际成品规格要求，轧制率可在5-80％的范围内变化。
最后进行第二次时效，在包含氢气、氮气、氩气、或这几种气体的任何混合物的气氛中进行，时效温度为350℃，保温时间4h。
需要说明的是，在第一次和第二次时效过程中尽管使用了特定的气体气氛，但是应理解还可使用其他的非活性气体作为保护气氛。
随后测量了合金中粒径在50-500nm之间和粒径>1μm的含Nb和Al的金属间化合物粒子的数量，并且测试所得的合金带材的力学性能、导电性能、折弯性能和在高温下的力学性能稳定性。
需要说明的是，为了避免使得本申请的说明书过于冗长，以下记载实施例12的详细工艺参数作为示例。尽管没有记载其他实施例的详细工艺参数，但是应当理解，本领域技术人员根据本说明书的公开内容已经足以实施本申请请求保护的发明并且 这样的公开内容也能够充分支持权利要求所要求的保护范围。
实施例12的成品厚度规格为0.15mm，具体工艺如下：
根据实施例12的合金原料添加量配料并进行熔炼。第一步：在熔炼炉中同时加入电解铜和Cu-Nb中间合金，开始熔炼；第二步：电解铜和Cu-Nb中间合金完全熔化后，依次加入纯Ti、纯Al和纯Co；第三步：所有原材料熔化后，在1300℃下精炼45min。熔炼后，通过垂直半连续铸造法铸造出矩形铸锭。
将铸锭在930℃下保温8h后进行热轧，热轧速度为110m/min，轧制单道次加工率30％，终轧温度在650℃以上，热轧后在线水冷却，热轧后进行铣面。
随后进行第一次冷轧，冷轧总加工率为90％。
第一次冷轧后进行固溶，固溶温度为700℃，保温时间为80s，冷却速度100℃/s。
固溶后进行中间冷轧，轧制率控制在55％，单道次变形量控制在20％。
中间冷轧后进行二次固溶，固溶温度为950℃，保温时间为5s，冷却速度200℃/s。
二次固溶后再次进行中间冷轧，轧制率控制在20％，单道次变形量控制在5％。
随后进行第一次时效，在包含氢气和氩气的混合物的气氛中进行，时效温度为400℃，保温时间4h。
第一次时效后进行最终冷轧，轧制率20％，最终厚度为0.15mm。
最后在包含氢气和氩气的混合物的气氛中进行第二次时效，时效温度为350℃，保温时间4h，得到成品材料。
测试标准：
室温拉伸试验按照《GB/T228.1-2010金属材料拉伸试验第1部分：室温试验方法》在电子万能力学性能试验机上进行，试样采用比例系数为5.65的矩形横截面比例试样。下表1中给出的本发明实施例和对比例的带材屈服强度为平行轧制方向的屈服强度。
导电率测试按照《GB/T3048-2007电线电缆电性能试验方法第2部分：金属材料电阻率试验》，用％IACS表示。
折弯性能采用如下方法测定：将铜合金带材沿轧制方向(即好方向)取长条样、垂直于轧制方向(即坏方向)取长条样，试样宽度为10mm，随后采用90°、尖端具有不同半径的V形冲头对长条样进行折弯，然后采用体视显微镜观测折弯处外表面，以表面不产生裂纹的最小弯曲半径/板厚(R/T)来表示折弯性能。当R/T值为0，表示最小弯曲半径R为0，折弯性能最优异。
平均晶粒度按照《YS/T 347-2004 铜及铜合金 平均晶粒度测定方法》的测试方法进行测定。
在高温下力学性能的稳定性测试参考《GB/T33370-2016铜及铜合金软化温度的测定方法》进行，将样品在500℃空气气氛下，保温1h后空冷，测试样品硬度。样品在某一温度高温保温后相比原始样品的硬度衰减率H(％)表征样品在高温下力学性能的稳定性，相同温度下硬度衰减率H越低，其在高温下力学性能的稳定性越好。
通过金相显微镜观察合金晶粒尺寸与金属间化合物粒子分布情况。使用扫描电镜观察合金中的金属间化合物粒子并对其大小数量进行统计。具体操作方式如下：取平行于铜合金带材轧制方向的截面，以25μm×40μm(1000μm 2)的矩形作为基本单元对其组织进行观察；选取视野中不同位置的10个矩形，对每个矩形中粒径在50-500nm之间的粒子数量以及粒径大于1μm的粒子数量进行统计，最后取其平均值作为判断依据，其中粒径的定义为颗粒物的最大尺寸。
根据实施例1-20可以发现，本发明通过对Ti、Nb、Al含量的合理控制，所有实施例的铜合金均实现了屈服强度≥900MPa，导电率≥10％IACS的性能，同时合金折弯性能优异即平行于轧制方向(即好方向)弯曲半径与带材的厚度比(R 1/T)≤0.5，垂直于轧制方向(即坏方向)弯曲半径与带材的厚度比(R 2/T)≤1.0。经500℃高温保温试验后，发现实施例1-20的合金样品的硬度衰减率H＜5％。
实施例1-20及对比例1-10反映了不同Nb和Al含量和含Nb和Al的金属间化合物粒子数量对钛青铜合金带材综合性能的影响。同时，实施例1-20也表明，合理少量添加选自Si、Zn、Co、Fe、Sn、Mn、Mg、Cr、B、Ag和Zr中的一种或多种任选元素对于合金强度以及高温稳定性可起到一定的改善。
实施例1-20和对比例1-10的钛青铜合金带材的成分、含Nb和Al的金属间化合物粒子数量及性能测试结果见表1。
虽然对比例1-5的钛青铜合金带材的屈服强度和折弯性能满足要求，但是由于没有添加Nb和Al(对比例1)或没有同时添加Nb和Al(对比例2-5)，在基体中不存在含Nb和Al的金属间化合物粒子，因而硬度衰减率H均较高(H>10％)。对比例6和7虽然同时添加了Nb和Al，但是对比例6添加的Nb含量不足，对比例7添加的Al含量不足，这都不能产生足量的含Nb和Al的金属间化合物粒子，其强化效果不明显，因而硬度衰减率H仍然较高(H>10％)
对比例8-10表明：尽管硬度衰减率H<5％，但是由于Al和/或Nb含量过多，钛 青铜合金的屈服强度和折弯性能受到不利影响。特别是Al和Nb同时过量时，会团聚成大的析出物粒子，不利于合金强度提升，同时会在弯曲时增加开裂的风险(对比例10的R 1/T和R 2/T都较大)。
Figure PCTCN2021076741-appb-000001
Figure PCTCN2021076741-appb-000002

Claims (10)

  1. 一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材的重量百分比组成包括2.0-4.5wt％的Ti，0.005-0.40wt％的Nb和0.01-0.50wt％的Al，余量为Cu和不可避免的杂质。
  2. 根据权利要求1所述的一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材的重量百分比组成包括：2.5-4.0wt％的Ti，优选为2.9-3.5wt％的Ti；和/或0.01-0.3wt％的Nb；和/或0.05-0.3wt％的Al。
  3. 根据权利要求1或2所述的一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材中粒径在50-500nm之间的含Nb和Al的金属间化合物粒子的数量不低于1×10 5个/mm 2，粒径大于1μm的含Nb和Al的金属间化合物粒子的数量不高于1×10 3个/mm 2。
  4. 根据权利要求1或2所述的一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材在大气气氛中在500℃下保温1h后，硬度衰减率H小于5％。
  5. 根据权利要求1或2所述的一种含Nb和Al的钛青铜合金带材，其特征在于：(1)该钛青铜合金带材平行于轧制方向的弯曲半径与带材的厚度比R 1/T≤0.5，垂直于轧制方向的弯曲半径与带材的厚度比R 2/T≤1.0；和/或(2)该钛青铜合金带材的屈服强度大于900MPa，导电率为10-20％IACS。
  6. 根据权利要求1或2所述的一种含Nb和Al的钛青铜合金带材，其特征在于：该钛青铜合金带材的重量百分比组成还包括总量0-0.50wt％的选自Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag中的一种或多种。
  7. 制备根据权利要求1-6中任一项所述的一种含Nb和Al的钛青铜合金带材的方法，其特征在于包括以下步骤：
    1)熔铸：采用真空或气氛保护的熔炼方法，在1200-1400℃将铜合金原料熔化；
    2)热加工：在700-980℃温度下对铸锭进行热加工，控制铸锭热加工的横断面面积缩减不低于75％；
    3)铣面：将热加工获得的材料铣面；
    4)第一次冷轧：控制材料的横断面面积缩减不低于70％；
    5)固溶处理：将冷轧后的材料加热至700-950℃的温度并且保温1-100s的时间，随后进行水冷或气冷处理，冷却速度控制在10-250℃/s；
    6)中间冷轧：控制材料的横断面面积缩减5-99％；
    7)第一次时效：选择非活性气氛保护在350-500℃的温度范围内保温0.5-24h；
    8)最终冷轧：控制横断面面积缩减5-80％；
    9)第二次时效：选择非活性气氛保护在200-550℃的温度范围内保温1min-10h。
  8. 根据权利要求7所述的方法，其中满足以下一者或多者：
    步骤1)中的铸造方式为铁模铸造、水平连续铸造或垂直半连续铸造；
    步骤2)中的热加工为热锻、热轧、或两者的结合；
    在步骤3)中将材料上下铣面0.5-2.0mm，以去除表面缺陷；
    在步骤6)中进行多道次冷轧，其中单道次变形量控制在5-20％；
    以步骤5)的固溶处理和步骤6)的中间冷轧为一个步骤单元，重复实施该步骤单元至少两次，其中相邻两次固溶处理之间的中间冷轧的材料的横断面面积缩减≥30％；和
    在步骤7)和/或步骤9)中时效在包含氢气、氮气、氩气、或这几种气体的任何混合物的气氛中进行。
  9. 根据权利要求7或8所述的方法，其中在步骤1)中，熔炼过程共分三步，第一步：在熔炼炉中同时加入电解铜和含Nb中间合金，开始熔炼；第二步：电解铜和含Nb中间合金完全熔化后，依次加入含Ti、含Al原料和任选的含Ni、Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag中一者或多者的一种或多种原料；第三步：所有原材料熔化后，在1300±50℃下精炼30-60min，完成后铸造出铸锭。
  10. 根据权利要求9所述的方法，其中含Nb中间合金是Cu-Nb中间合金或Nb-Ti中间合金，含Ti、含Al原料是纯Ti、纯Al或者含Ti和/或Al的中间合金，含Ni、 Co、Fe、Sn、Mn、Si、Cr、Mg、B、Zr和Ag中一者或多者的一种或多种原料是这些元素的单质或含这些元素的中间合金。
PCT/CN2021/076741 2020-06-24 2021-02-18 一种含Nb和Al的钛青铜合金带材及其制备方法 WO2021258747A1 (zh)

Priority Applications (2)

Application Number Priority Date Filing Date Title
US17/423,698 US11913102B2 (en) 2020-06-24 2021-02-18 Titanium-copper alloy strip containing Nb and Al and method for producing same
KR1020217025759A KR102640850B1 (ko) 2020-06-24 2021-02-18 엔비(Nb) 및 에이엘(Al)을 함유하는 티타늄 구리 합금 스트립 및 그의 제조 방법

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
CN202010586105.5A CN111534714B (zh) 2020-06-24 2020-06-24 一种含Nb和Al的钛青铜合金带材及其制备方法
CN202010586105.5 2020-06-24

Publications (1)

Publication Number Publication Date
WO2021258747A1 true WO2021258747A1 (zh) 2021-12-30

Family

ID=71971254

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/CN2021/076741 WO2021258747A1 (zh) 2020-06-24 2021-02-18 一种含Nb和Al的钛青铜合金带材及其制备方法

Country Status (4)

Country Link
US (1) US11913102B2 (zh)
KR (1) KR102640850B1 (zh)
CN (1) CN111534714B (zh)
WO (1) WO2021258747A1 (zh)

Families Citing this family (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN111534714B (zh) * 2020-06-24 2021-08-31 宁波博威合金板带有限公司 一种含Nb和Al的钛青铜合金带材及其制备方法
CN113802027B (zh) * 2021-09-18 2022-07-15 宁波博威合金板带有限公司 一种钛青铜及其制备方法

Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2011118400A1 (ja) * 2010-03-25 2011-09-29 Jx日鉱日石金属株式会社 高強度チタン銅板及びその製造方法
JP2015091603A (ja) * 2014-11-21 2015-05-14 新日鐵住金株式会社 銅合金の製造方法
CN110923499A (zh) * 2019-12-27 2020-03-27 宁波博威合金材料股份有限公司 一种含Ce和B的钛青铜合金带材及其制备方法
CN111534714A (zh) * 2020-06-24 2020-08-14 宁波博威合金板带有限公司 一种含Nb和Al的钛青铜合金带材及其制备方法

Family Cites Families (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2004244646A (ja) * 2002-12-17 2004-09-02 Nikko Metal Manufacturing Co Ltd 抵抗体用圧延銅合金箔及びその製造方法
JP2005113259A (ja) * 2003-02-05 2005-04-28 Sumitomo Metal Ind Ltd Ｃｕ合金およびその製造方法
JP3731600B2 (ja) * 2003-09-19 2006-01-05 住友金属工業株式会社 銅合金およびその製造方法
EP2180071A1 (en) * 2007-07-27 2010-04-28 Materials Solution Inc. Copper alloy material
CN104278171B (zh) * 2014-09-16 2016-09-07 中南大学 一种CuTi系弹性铜合金及其制备方法

Patent Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2011118400A1 (ja) * 2010-03-25 2011-09-29 Jx日鉱日石金属株式会社 高強度チタン銅板及びその製造方法
JP2015091603A (ja) * 2014-11-21 2015-05-14 新日鐵住金株式会社 銅合金の製造方法
CN110923499A (zh) * 2019-12-27 2020-03-27 宁波博威合金材料股份有限公司 一种含Ce和B的钛青铜合金带材及其制备方法
CN111534714A (zh) * 2020-06-24 2020-08-14 宁波博威合金板带有限公司 一种含Nb和Al的钛青铜合金带材及其制备方法

Also Published As

Publication number Publication date
CN111534714A (zh) 2020-08-14
US20220341001A1 (en) 2022-10-27
CN111534714B (zh) 2021-08-31
KR20220000900A (ko) 2022-01-04
US11913102B2 (en) 2024-02-27
KR102640850B1 (ko) 2024-02-27

Similar Documents

Publication Publication Date Title
CN108642363B (zh) 一种高强高塑共晶高熵合金及其制备方法
CN111485132B (zh) 一种综合性能优异的铜合金带材及其制备方法
CN110923499B (zh) 一种含Ce和B的钛青铜合金带材及其制备方法
EP3363922B1 (en) Cobalt silicide-containing copper alloy
TWI387657B (zh) Cu-Ni-Si-Co based copper alloy for electronic materials and method of manufacturing the same
CN113355554A (zh) 一种铜铬合金带材及其制备方法
EP2692879B1 (en) Cu-co-si-based copper alloy strip for electron material, and method for manufacturing same
US20110027122A1 (en) Cu-Ni-Si-Co-Cr System Alloy for Electronic Materials
WO2021258747A1 (zh) 一种含Nb和Al的钛青铜合金带材及其制备方法
CN108384986B (zh) 一种铜合金材料及其应用
CN113106290B (zh) 一种高性能锡磷青铜带材及其制备方法
CN111996411A (zh) 一种高强高导铜合金材料及其制备方法和应用
CN111621668B (zh) 一种镍硅系铜合金带材及其制备方法
CN114657410B (zh) 一种高强高导铜铁系合金及其制备方法
WO2014109211A1 (ja) 電子・電気機器用銅合金、電子・電気機器用部品及び端子
WO2012133651A1 (ja) 銅合金および銅合金の製造方法
JP6246456B2 (ja) チタン銅
CN108796296A (zh) 一种铜合金及其应用
CN115198138B (zh) 一种铜合金带材及其制备方法
CN114277280B (zh) 一种析出强化型锡黄铜合金及其制备方法
CN112281023B (zh) 一种具有优异折弯性的铜合金材料及其制备方法和应用
CN112501472B (zh) 一种高性能铜合金带材及其制备方法
JP2016211077A (ja) チタン銅
CN115874080B (zh) 一种铜基合金材料及其制备方法和应用
CN115094266B (zh) 一种高强导电弹性铜合金及其制备方法

Legal Events

Date Code Title Description
121 Ep: the epo has been informed by wipo that ep was designated in this application

Ref document number: 21829662

Country of ref document: EP

Kind code of ref document: A1

NENP Non-entry into the national phase

Ref country code: DE

122 Ep: pct application non-entry in european phase

Ref document number: 21829662

Country of ref document: EP

Kind code of ref document: A1