KR20080024554A - Cu-based alloy having high strength and high plasticity and method for preparing the same - Google Patents

Abstract

A Cu-based alloy having high strength and high elongation and a method for manufacturing the same are provided, wherein the Cu-based alloy maintains high strength and high elongation, and can be manufactured by an environmentally friendly, simple and economical process without using hazardous elements such as beryllium. A Cu-based alloy having high strength and high elongation comprises: 5 to 15 atomic wt.% of Zr; 1 to 5 atomic wt.% of Al; and the balance of Cu, and the Cu-based alloy has a composite microstructure consisting of: a superlattice phase in which a Cu5Zr layer and a Zr layer are repeated; and a Cu crystalline phase. The Cu-based alloy has a volume ratio of the superlattice phase to the Cu crystalline phase of 3:7 to 5:5. The superlattice phase comprises 15 to 25 wt.% of the Zr layer and 75 to 85 wt.% of the Cu5Zr layer. The Cu crystalline phase comprises 94 to 98 wt.% of Cu and 2 to 6 wt.% of Al. The Cu-based alloy has yield strength of 1.0 to 1.6 GPa, and an elongation of 7.0 to 15.0%. The Cu-based alloy has a formula of Cu86Zr11Al13.

Description

High strength and high elongation copper alloy and its manufacturing method {Cu-based alloy having high strength and high plasticity and method for preparing the same}

1 is a graph showing the X-ray diffraction pattern (XRD) results for a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ according to an embodiment of the present invention.

2A is a graph showing a scanning electron microscope (SEM) photograph of a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ according to an embodiment of the present invention.

FIG. 2B is a high resolution transmission electron microscope (HRTEM) photograph of the S and S ′ regions of FIG. 2A and a corresponding restriction field in the copper alloy having the formula of Cu ₈₆ Zr ₁₁ Al ₃ according to one embodiment of the present invention. It is a diffraction diagram.

FIG. 2C is a high-resolution transmission electron microscope (HRTEM) photograph of the region B of FIG. 2A and a corresponding limited viewing diffraction diagram in the copper alloy having the chemical formula Cu ₈₆ Zr ₁₁ Al ₃ according to one embodiment of the present invention. .

FIG. 2D is a transmission electron microscope (TEM) image of a portion (bottom photo) and the remaining region of the S region of FIG. 2A in the copper alloy having the chemical formula Cu ₈₆ Zr ₁₁ Al ₃ according to one embodiment of the present invention. .

FIG. 3 is a high resolution transmission electron microscope (HRTEM) photograph of both ends of region S of FIG. 2A and a corresponding limitation in the copper alloy having the formula of Cu ₈₆ Zr ₁₁ Al ₃ according to one embodiment of the present invention. It is a visual diffraction diagram.

4A is a graph illustrating stress-strain measurement results of a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ according to an embodiment of the present invention.

FIG. 4B shows a scanning electron microscope (SEM) photograph of a fracture shape of a specimen subjected to a compression test in a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ according to one embodiment of the present invention.

5 is a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ in accordance with an embodiment of the present invention, a superstructure (S region) of the superlattice phase (S region), a superstructure phase and a structure consisting of a Cu phase (S '+ B) It is a graph showing the result of measuring the hardness of the region).

6A to 6D are scanning electron microscope (SEM) images observed while increasing a compressive strain in a copper alloy having a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ according to an embodiment of the present invention.

The present invention relates to a high-strength and high elongation copper alloy and a method for manufacturing the same, and more particularly, it has high strength and high elongation characteristics, and does not use harmful elements such as beryllium. It relates to a high strength and high elongation copper alloy manufacturable by a process and a method for producing the same.

Copper alloys are excellent in ductility and inexpensive in terms of material cost, but their specific strength is low so that they are subject to various limitations in application as a structural material. Accordingly, various reinforcement methods have been devised to increase the strength of such copper alloys, and examples thereof include solid solution strengthening, precipitation hardening, and dispersion strengthening.

Solid solution strengthening refers to a phenomenon in which a solute atom is dissolved in a lattice of solvent atoms to form a solid solution and cause a distortion of the lattice, thereby strengthening the material by forming a stress field in the vicinity of the solute atom (e.g. 48845). In addition, there is a strengthening phenomenon due to the insoluble second phase finely dispersed in the matrix, which is classified into precipitation hardening and dispersion hardening according to which method the dispersed second phase is introduced. In other words, precipitation hardening means strengthening when the second phase is formed by precipitation from a supersaturated solid solution (for example, Korean Patent Application Publication No. 1984-7753), and dispersion hardening means that the second phase is not precipitated from solid solution. It refers to the strengthening phenomenon when formed by a process (for example, powder metallurgy method or internal oxidation method, etc.) (for example, Korean Patent Publication No. 1983-7874).

The double precipitation hardening method is an important method for reinforcing copper alloys. Among the alloys used commercially, copper-beryllium-based alloys as disclosed in Korean Patent Application Publication No. 1996-701230 are typical precipitation hardening alloys. However, beryllium metal is one of the representative environmentally harmful metals, and the method of reinforcing copper alloy using such beryllium metal has many limitations in its application.

In addition, in order to further improve the strength of the copper alloy in recent years, a study on a copper-based microcomposite having a grain size of μm has been conducted (MK Lee, SM Hong, GH Kim, KH Kim, and WW Kim, Met. Mater. Int. 10 (2004), 313-319; PD Funkenbusch, and TH Courtney, Scripta Metall. 15 (1981), 1349-1354; C. Biselli, and DG Morris, Acta Mater. 44 (1996), 493- 504), the smaller the grains are, the more limited the generation and movement of dislocations is, so that the hardening of the work is difficult and there is a problem that sudden fracture is caused by local shear stress.

Therefore, studies have been conducted on composites reinforced with the second phase as a method for suppressing brittle fracture (G. Frommeyer, and G. Wassermann, Acta Metall. 23 (1975), 1353-1360; H. Gleiter, Prog. Sci. 33 (1989), 223; Y. Sakai, K. Inoue, and H. Maeda, Acta Metall. Mater. 43 (1995), 1517-1522; SI Hong, MA Hill, Y. Sakai, JT Wood, and JD Embury, Acta Metall. Mater. 43 (1995), 3313-3323; HS Kim, J. Kor. Inst. Met. & Mater. 38 (2000), 357-365) As a common method of properly distributing two phases in a matrix, a phase separation of the matrix phase and the second phase is induced by adding a main element constituting the matrix and a metal element having a positive mixing heat. The materials developed in this way are composed of grains of 占 퐉 or ㎚ size in the base material to maintain high strength and to generate a large number of shear bands in the second phase present in the alloy, thereby effectively suppressing the propagation of cracks. As a result, high elongation can be exhibited.

However, the above-described prior arts have limitations in effectively improving the strength and elongation of copper alloys, causing other environmental pollution problems, and having many problems in terms of process simplicity and economy.

Accordingly, the present invention is to solve the above-mentioned problems of the prior art, and has a high strength and high elongation characteristics, and can be manufactured by an environmentally friendly, simple and economical process without using harmful elements such as beryllium. An object of the present invention is to provide a high strength and high elongation copper alloy and a method of manufacturing the same.

The present invention to achieve the above technical problem,

5 to 15 atomic percent Zr; 1 to 5 atomic percent Al; And balance Cu,

A superlattice phase in which a Cu ₅ Zr layer and a Zr layer are repeated; And it provides a high strength and high elongation copper alloy having a composite microstructure composed of a Cu crystal phase.

According to a preferred embodiment of the present invention, the volume ratio of the superlattice phase and the Cu crystal phase is 3: 7 to 5: 5.

According to a preferred embodiment of the present invention, the superlattice phase comprises 15 to 25% by weight of the Zr layer and 75 to 85% by weight of the Cu ₅ Zr layer.

According to a preferred embodiment of the present invention, the Cu crystal phase comprises 94 to 98% by weight of Cu and 2 to 6% by weight of Al.

According to a preferred embodiment of the present invention, the yield strength of the copper alloy is 1.0 to 1.6 GPa.

According to a preferred embodiment of the present invention, the elongation of the copper alloy is 7.0 to 15.0%.

According to a preferred embodiment of the present invention, the copper alloy has a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ .

In addition, the present invention to achieve the above other technical problem,

Weighing 5 to 15 atomic percent Zr, 1 to 5 atomic percent Al and the balance Cu to produce an ingot by an arc melting process;

Repeating melting and cooling the ingot a plurality of times for uniform mixing of the Zr, Al and Cu; And

Casting the copper mold by suction casting method

It provides a method for producing a high strength and high elongation copper alloy comprising a.

According to a preferred embodiment of the present invention, the arc melting process and the suction casting method is carried out under high purity argon atmosphere.

According to a preferred embodiment of the present invention, the melting of the melting and cooling step is carried out at a temperature of 2500 to 3000 ℃, cooling is carried out at a temperature of 10 to 30 ℃.

According to one preferred embodiment of the present invention, the melting and cooling steps are performed three to four times.

Hereinafter, the present invention will be described in more detail.

According to the present invention, it is possible to produce a copper alloy having a higher strength and elongation than conventional copper alloy reinforcing methods, the manufacturing process is also environmentally friendly, simple and economical.

Copper alloys according to the present invention comprise 5 to 15 atomic percent Zr; 1 to 5 atomic percent Al; And a balance of Cu, the superlattice phase in which a Cu ₅ Zr layer and a Zr layer are repeated; And a composite microstructure composed of Cu crystal phases.

As described above, many studies have been made on methods of improving strength by inducing phase separation using a second phase as a method for suppressing brittle fracture of a copper alloy. Unlike the conventional reinforcement methods, the inventors found that when the alloy solidifies, a high-strength crystal phase is formed first, and the remaining liquid phase forms a fine process structure through a process reaction. The material can be prepared, and the composite material thus produced has both high primary and fine process structures at the same time, so that the strength is high and the absorbing energy can be effectively absorbed, so that the elongation is not greatly reduced.

In the copper alloy according to the present invention, the structure corresponding to the high-strength primary crystal is a superlattice phase in which the Cu ₅ Zr layer and the Zr layer are repeated, the structure corresponding to the fine process structure is a Cu crystal phase, and other fine-sized Al-containing nanocrystal phases. This includes.

As described above, in order to secure high strength and excellent elongation at the same time, the volume ratio of the superlattice phase and the Cu crystal phase is preferably 3: 7 to 5: 5. When the superlattice phase is excessively out of the above-described volume ratio, the elongation is relatively not excellent, and when the Cu crystal phase is excessively present, there is a problem in that high strength properties cannot be secured.

On the other hand, the superlattice phase comprises 15 to 25% by weight of the Zr layer, 75 to 85% by weight of the Cu ₅ Zr layer, the Cu crystal phase preferably comprises 94 to 98% by weight of Cu and 2 to 6% by weight of Al. .

As described above, the copper alloy according to the present invention has excellent strength due to the presence of a superlattice phase, and has a good elongation due to a complicated microstructure in which cracks are hard to propagate. The yield strength is 1.0 to 1.6 GPa, and the elongation is excellent. Is 7.0 to 15.0%.

According to a preferred embodiment of the present invention, the copper alloy may have a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ .

In addition, the present invention, in order to achieve the above another technical problem, by measuring the 5 to 15 atomic% Zr, 1 to 5 atomic% Al and the balance of Cu to produce an ingot by the arc melting process; Repeating melting and cooling the ingot a plurality of times for uniform mixing of the Zr, Al and Cu; And it provides a method for producing a high strength and high elongation copper alloy comprising the step of casting a copper mold by suction casting method.

The arc melting process and the suction casting process may be carried out under conventional conditions as known in the art, but it is particularly preferable to carry out under high purity argon atmosphere. This is because when oxygen or nitrogen gas other than argon is present when the specimen is dissolved, it may react with the molten metal of the specimen to generate an oxide or nitride, and these impurities may lower the mechanical properties of the copper alloy according to the present invention.

In addition, the melting of the melting and cooling step is carried out at a temperature of 2500 to 3000 ℃, cooling is performed at a temperature of 10 to 30 ℃, it is preferable that the melting and cooling steps are carried out three to four times.

If the specimen is heated to a temperature lower than the above-mentioned temperature, the specimen may not be uniformly mixed, and impurities may be generated in the material, and it is impossible to heat the specimen to a temperature higher than the above-mentioned temperature because of limitations of the melting machine. In addition, when the cooling temperature using the cooling water is higher than 30 ℃ it is not preferable because the cooling rate of the specimen can be lowered to lower the physical properties of the alloy according to the present invention. On the other hand, if the number of melting of the specimen is less than three times, since uniform mixing of the elements is not performed, impurities may occur in a part of the specimen.If the number of melting is more than four, residual stress field is generated inside the specimen due to thermal stress. Occurs, which is undesirable because it can increase the brittleness of the specimen.

Hereinafter, the present invention will be described in more detail with reference to Examples, but the following Examples are only intended to help the understanding of the present invention, and are not intended to limit the scope of the present invention.

Example. Preparation of Copper Alloys According to the Invention

After measuring the metallic elements of Cu (99.99%), Zr (99.7%) and Al (99.999%) in the amounts of 86%, 11% and 3% by weight, respectively, the specimen was subjected to high purity (99.9999%) argon atmosphere. Ingots were prepared by performing arc melting at 2500 ° C. for several seconds.

Thereafter, the components were remelted four times so that the components could be uniformly mixed. The remelting was repeated for several seconds at 2500 ° C., followed by rapid cooling to room temperature. Subsequently, Cu ₈₆ Zr ₁₁ Al ₃ specimens having diameters of 1, 2 and 3 mm were cast using suction casting. A copper mold was used for suction casting, and cooling water was flowed around the copper mold so that the specimen could be rapidly cooled.

Evaluation example. Measurement of Properties of Copper Alloys According to the Present Invention

Evaluation Example 1 Phase Analysis Through X-Ray Diffraction Pattern

The X-ray diffraction pattern (XRD, Cu-Kα) was analyzed by selecting a specimen having a diameter of 2 mm among the specimens prepared in the above example as a representative specimen.

FIG. 1 is a graph showing the X-ray diffraction pattern results. From the pattern results in FIG. 1, it is understood that the main crystal phases constituting the alloy are Cu and Cu ₅ Zr, and a small amount of Cu ₉ Al ₄ is present. Could. It is expected that Cu and Cu ₅ Zr, which are marked high intensity peaks, exist as large crystal phases, and that Cu ₉ Al _{4, which} is represented as low intensity halo peaks, is present as small nanocrystalline phases. In addition, the presence of halo peaks, which are not easy to analyze, suggests that traces of nanocrystalline phases of different compositions are also present.

Evaluation Example 2 Analysis of Microstructure and Composition by SEM, HRTEM and EDS Pattern

On the other hand, for the images obtained by XRD analysis of Evaluation Example 1, Scanning Electron Microscopy (SEM), High Resolution Transmission Electron Microscopy (HRTEM) and Energy Dispersion Spectroscopy (EDS) The microstructure and composition, such as the actual shape, size and distribution of these phases were analyzed. The specimens used were fabricated by argon ion milling (6kV, 11.5 °).

FIG. 2A is a microstructure of Cu ₈₆ Zr ₁₁ Al ₃ alloy observed by SEM, in which the prepared alloy has acicular (15 × 2 μm) primary (S region) and lamellar process structures (S ′ and The area B).

Figure 2b is a high resolution photograph of the S and S 'region obtained from the HRTEM and the corresponding limited viewing diffraction diagram, it can be seen that the region consists of a superlattice (two different phases) appearing repeatedly. This superlattice phase is observed in the general cast structure is a very unusual result, as a result of the limited field diffraction diagram and EDS composition analysis, layer 1 and layer 2 of Figure 2b is composed of Zr and Cu ₅ Zr phase 2: 8 I can see that it is.

FIG. 2C shows a high resolution photograph of the B region (light portion) of FIG. 2A and the corresponding limited viewing diffraction diagram, which phase is Cu. However, it was found that Cu solid solution was found from the fact that the lattice constant measured from the high resolution photograph was slightly larger than the pure Cu lattice constant and that the trace amount of Al (~ 4 at.%) Was detected from the TEM-linked EDS analysis. .

On the other hand, Figure 2d is a TEM image of a portion (bottom of the photo) and the remaining area of the initial (S region) of Figure 2a, except for the region is generated as a result of the process reaction when considering the structure (lamella) Judging. Each composition for regions 3 to 6 of FIG. 2D was analyzed using EDS and High Angle Anular Dark Field (HAADF) analysis, and the results are shown in Table 1 below.

domain Cu (at.%) Zr (at.%) Al (at.%) Prediction 3 80.8 18.0 1.2 _{(Cu 5 Zr) 0.8 Zr 0} .2 4 79.3 20.7 - _{(Cu 5 Zr) 0.8 Zr 0} .2 5 95.1 - 4.9 α-Cu 6 96.0 - 4.0 α-Cu

As can be seen from the results of Table 1, it can be confirmed that the remaining region except for the primary crystal is composed of a superlattice phase and a Cu crystal phase made of a process reaction.

FIG. 3 is a high resolution photograph of both ends of the superlattice phase and its corresponding limited viewing diffraction diagram, which was observed to consist mainly of nanocrystalline phases. As a result of the analysis of the limited field diffraction diagram and the fast Fourier transform (FFT) of high resolution photographs, this region was composed of 5-10 nm Cu ₁₀ Zr ₇ and Cu ₉ Al ₄ nanocrystals, and the volume fraction was very small. All. These results are in good agreement with those of the XRD pattern described above.

Evaluation Example 3 Measurement of Strength and Elongation

Compression test was performed to measure the strength and elongation of the copper alloy according to the present invention, the results are shown in Figure 4a. The maximum strengths from rod-shaped specimens with diameters of 1, 2 and 3 mm were 1.60, 1.57 and 1.33 GPa, respectively, and the plastic strains were 7.2, 10.9 and 13.0%, respectively. In the case of relatively slow cooling ingot specimen (diameter 15 mm), the maximum strength and plastic strain were 1.00 GPa and 12.0%, respectively.

As can be seen from the experimental results, it can be seen that as the size of the cast specimen increases, the strength decreases and the elongation increases, indicating that the mechanical properties of the alloy are significantly influenced by the cooling rate, thereby changing the mechanical properties. .

In addition, as can be seen from Figure 4b, the fracture surface of the specimen forms an angle of about 47 ^o with the compression direction, it can be seen that the fracture of the specimen occurs in the direction of the maximum shear stress.

Cu-Be alloys, which are representative examples of precipitation hardening copper alloys, show slight differences depending on the composition and manufacturing method, but exhibit a yield strength of about 0.3 to 1.2 GPa and an elongation between 1 and 25% (WF Gale, and TC Totemeier, Smithells Metals Reference Book 8th ed. (Elsevier, England, 2004); ASM metals handbook, Vol. 2 Properties and Selection: Nonferrous Alloys and Pure Metals 9th ed. (ASM International, Ohio, 1986)). Table 2 below shows the results of comparing the mechanical properties obtained by performing compression experiments on the copper alloys of the examples with those obtained by performing tensile tests on conventional Cu-Be alloys.

Furtherance shape Common name Yield strength, σ _ys (MPa) Elongation, ε _p (%) Reference ^* Example (1 mm specimen) Cast (cast) - 1596 7.2-13.0 - Cu _{97 .44} Be _{2 .56} Workpiece (wrought) - 865 2.6 One _{Cu 98 .0 Be 1 .7 Co 0} .3 Processing material C17000 1070 4.0-11.0 2 _{Cu 97 .4 Be 1 .9 Co 0} .4 Pb 0 .3 Processing material C17300 1210 3.0-10.0 2 _{Cu 97 .0 Co 2 .5 Be 0} .5 Processing material C17500 690 10.0 ~ 25.0 2 _{Cu 97 .0 Co 1 .65 Ag 1} .00 Be 0 .35 Processing material C17600 690 10.0 ~ 25.0 2 _{Cu 99 .0 Cr 0 .8 Be 0} .06 Casting material C81400 250 11.0 2 _{Cu 97 .0 Co 1 .5 Ag 1} .0 Be 0 .4 Casting material C81800 515 8.0 2 _{Cu 97 .0 Co 2 .5 Be 0} .5 Casting material C82000 515 6.0 2 _{Cu 98 .0 Ni 1 .5 Be 0} .5 Casting material C82200 515 7.0 2 _{Cu 98 .0 Be 1 .7 Co 0} .3 Casting material C82400 1000 1.0 2 _{Cu 97 .2 Be 2 .0 Co 0} .5 Si 0 .25 Casting material C82500 1035 1.0 2 _{Cu 97 .0 Be 2 .4 Co 0} .5 Casting material C82600 1070 1.0 2 _{Cu 86 .6 Be 2 .6 Co 0} .5 Si 0 .3 Casting material C82800 1070 1.0 2 _{Cu 97 .0 Be 2 .0 Co 1} .0 Casting material Be-Cu 21C 725 5.0 2 Cu ₆₈ Ni _{30 .8 .2 .0} Be ₁ Casting material Be-Cu-Ni 72C 310 15.0 2

Reference 1: W. F. Gale, and T. C. Totemeier, Smithells Metals Reference Book 8th ed. (Elsevier, England, 2004)

Reference 2: ASM metals handbook, Vol. 2 Properties and Selection: Nonferrous Alloys and Pure Metals 9th ed. (ASM International, Ohio, 1986)

As can be seen from the results of Table 2, although considering the difference between the manufacturing method and the property evaluation method, the alloy according to the present invention shows a relatively excellent mechanical properties compared to the conventional Cu-Be alloy, which is the present invention It is judged that the alloy has a special microstructure.

In addition, the hardness of the phases constituting the alloy was measured through nanoindentation experiments to determine the degree to which the microstructure of the copper alloy according to the present invention contributes to the overall strength of the alloy.

5 is a result of measuring the hardness of the supercrystalline phase (S region) of the superlattice phase and the process structure (S '+ B region) composed of the superlattice phase and the Cu phase, and exhibits high hardness values of about 7.0 GPa and 4.1 GPa, respectively. It was. On the other hand, the relationship between the hardness value and the yield strength measured by the nano-indentation holds the following equation (PM Rice, and RE Stoller, Symposium Q; MRS (2000 Fall)), based on this As a result of calculating the yield strength, it was found that the maximum yield strengths of the superlattice and the process tissue were about 1.9 GPa and 1.1 GPa, respectively.

Δσ _y [MPa] = 274 ΔH _N [GPa]

(Equation 1, Δσ _y represents the amount of change in yield strength (unit: MPa), ΔH _N represents the amount of change in hardness value (unit: GPa) measured by nanoindentation.)

Therefore, considering that the maximum yield strength of the copper alloy according to the present invention is about 1.6 GPa, it can be seen that the phase contributing to the high strength of the alloy according to the present invention is a superlattice phase. Generally, in the nanostructured material or composite material including the nanocrystalline phase, the portion contributing to the increase in strength is the nanocrystalline phase, but the volume fraction occupied by the nanocrystalline phase in the copper alloy according to the present invention is very small, so that the overall strength of the alloy We do not believe it will significantly affect the growth.

Evaluation Example 4 Plastic Fracture Behavior of the Copper Alloy According to the Present Invention

Plastic fracture behavior of crystalline materials can be explained by the generation and propagation of cracks. Therefore, in order to analyze the plastic fracture behavior of the copper alloy according to the present invention, after applying different strains to the specimen, the formation and propagation paths of the cracks were observed at the strain rate.

6A is a SEM photograph observed after deformation of the rod-shaped specimen of the copper alloy according to the present invention having a diameter of 2 mm to yield strength (about 2% compressive strain), and it can be seen that a large number of microcracks were generated in the superlattice phase. . In order to observe the propagation process of the generated cracks, the specimen was deformed by 4%, and the result is shown in FIG. 6B.

As can be seen from FIGS. 6A and 6B, the cracks generated in the superlattice phase (S region) reach the interface of the process phase (S '+ B region), and then do not pass through the process phase but along the interface. After returning, it stops at the end of the superlattice. Therefore, since the crack propagates along a distant path, it can absorb a lot of plastic strain energy. The propagated cracks are connected to interfacial cracks on adjacent superlattices as the strain increases, and these results can be seen in FIG. 6C. Therefore, the interface failure caused by the shear stress occurs in several places, the shear failure occurs macroscopically as shown in Figure 4b.

Figure 6d is a photograph taken from the Cu phase present in the process structure after the specimen is broken, it can be seen that there is a very fine (~ 0.3 ㎛) dimples over a wide range of the specimen. Therefore, it can be seen that the plastic strain energy applied to the specimen is also absorbed by the deformation of the highly ductile Cu phase. In conclusion, the reason why the copper alloy according to the present invention does not show sudden breakage is related to its unique microstructure, and it is judged that the plastic deformation energy is absorbed in the following path.

1) Preferential multiple cracking occurs in high-strength superlattice phase (S).

2) The crack formed propagates along the interface after reaching the interface between the superlattice phase (S) and the process phase (S '+ B).

3) Crack connection between adjacent superlattices.

4) Shear failure after interface failure by shear stress.

According to the present invention, there is provided a high-strength and high elongation copper alloy having a high strength and high elongation characteristics and can be manufactured by an environmentally friendly, simple and economical process without using harmful elements such as beryllium, and a manufacturing method thereof. can do.

Claims (11)

5 to 15 atomic percent Zr; 1 to 5 atomic percent Al; And balance Cu, A superlattice phase in which a Cu ₅ Zr layer and a Zr layer are repeated; And a composite microstructure composed of a Cu crystal phase. The high-strength and high elongation copper alloy according to claim 1, wherein the volume ratio of the superlattice phase and the Cu crystal phase is 3: 7 to 5: 5. The high-strength and high elongation copper alloy according to claim 1, wherein the superlattice phase comprises 15 to 25 wt% of a Zr layer and 75 to 85 wt% of a Cu ₅ Zr layer. The high strength and high elongation copper alloy according to claim 1, wherein the Cu crystal phase comprises 94 to 98% by weight of Cu and 2 to 6% by weight of Al. The high strength and high elongation copper alloy according to claim 1, wherein the yield strength of the copper alloy is 1.0 to 1.6 GPa. The high strength and high elongation copper alloy according to claim 1, wherein the elongation of the copper alloy is 7.0 to 15.0%. The high strength and high elongation copper alloy of claim 1, wherein the copper alloy has a chemical formula of Cu ₈₆ Zr ₁₁ Al ₃ . Weighing 5 to 15 atomic percent Zr, 1 to 5 atomic percent Al and the balance Cu to produce an ingot by an arc melting process; Repeating melting and cooling the ingot a plurality of times for uniform mixing of the Zr, Al and Cu; And Casting the copper mold by suction casting method Method for producing a high strength and high elongation copper alloy comprising a. The method of claim 8, wherein the arc melting process and the suction casting method are performed under a high purity argon atmosphere. The method of claim 8, wherein the melting and cooling steps are performed at a temperature of 2500 to 3000 ° C., and cooling is performed at a temperature of 10 to 30 ° C. 10. The method of claim 8, wherein the melting and cooling steps are performed three to four times.

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