AU2005324597B2 - Wrought magnesium alloy having excellent formability and method of producing same - Google Patents

Abstract

Disclosed is a wrought magnesium alloy having excellent strength and extrusion or rolling formability, and a method of producing the same. The wrought magnesium alloy comprises 0.1-1.5 at % group IIIa, 1.0-4.0 at % group IIIb, 0.35 at % or less of one selected from the group consisting of groups IIa, IVa, VIIa, IVb, and a mixture thereof, 1.0 at % or less of group IIb, and a balance of Mg and unavoidable impurities and thus has a second phase composite microstructure. The wrought magnesium alloy of the present invention has high strength, toughness, and formability in addition to the electromagnetic wave shield ability of magnesium. Accordingly, the wrought magnesium alloy is a material useful to portable electronic goods, such as notebook personal computers, mobile phones, digital cameras, camcorders, CD players, PDA, or MP3 players, automotive parts, such as engine room hoods, oil pans, or inner panel of door, or structural parts for airplane.

Description

WO 2006/075814 PCTiKR20051000697 Description WROUGHT MAGNESIUM ALLOY HAVING EXCELLENT FORMABILITY AND METHOD OF PRODUCING SAME Technical Field The present invention relates to a wrought magnesium alloy, which contains a second phase consisting of an intermetallic compound, thereby having excellent strength, formability, and corrosion resistance. More particularly, the present invention pertains to a wrought magnesium alloy, which comprises 0.1 1.5 at% group IIIa, 1.0 at% group IIIb, 0.35 at% or less of one selected from the group consisting of groups IIa, IVa, Vila, IVb, and a mixture thereof, 1.0 at% or less group IIb, and a balance of Mg and unavoidable impurities and which thus has a second phase composite microstructure consisting of an intermetallic compound, and a method of producing the same.

Background Art Having a density of 1.74 1.95 g/cm 3 or so, a relatively low specific gravity that is 2/3 of that of aluminum, excellent specific strength and machinability, magnesium alloys have been developed as light structural materials for airplanes and automobiles.

However, since magnesium has a hexagonal close packed (HCP) lattice crystal structure, formability is very low, and thus, the application is limited to a field in which the forming is achieved using a casting process. Particularly, its practical use is limited due to problems, such as severe oxidation of melts, reduction of strength at high temperatures, and low corrosion resistance. Effort has been made to avoid the above disadvantages, thus enabling stable dissolution in atmospheric air, while employing a sulfur hexafluoride (SF 6 gas, a carbon dioxide gas, an argon gas and the like, and production of a plate by Direct Chilled casting process.

Of the magnesium alloys, a Mg-Zn alloy shows an excellent age hardening behavior, and is advantageous in that since a microstructure is refined through heat treatments, strength and ductility significantly increase and it is easy to work and weld.

On the other hand, it is disadvantageous in that since micropores are formed as a casting process due to the addition of Zn, it is difficult to apply the Mg-Zn alloy to a casting process, such as die-casting. Additionally, it is difficult to desirably improve strength because it is grown as the coarse grain. To overcome the above disadvantages, studies have been made to improve formability using a grain boundary slip, in which some alloy elements are added to Mg-Zn binary alloys to refine grains. With respect to this, J. P. Doan and G. Ansel suggest a method of improving the strength of an alloy, in which Zr is added to refine grains constituting a Mg-Zn alloy P. Doan and G. Ansel, WO 2006/075814 PCT/KR2005/000697 Trans, AIME, vol. 171 (1947), pp. 286-295). However, since Zr has a high melting point and a low solubility to Mg at room temperature, it mostly exists at a grain boundary, thus acting as a fracture initiation site when external stress is applied. In this regard, it becomes possible to conduct a plastic working process using the ductility of a single-phase solid solution after alloy materials, including aluminum and zinc or manganese, such as AZ31B or AM20, are developed. However, even though their microstructures have a single-phase solid solution and thus have excellent ductility, they are disadvantageous in that since a strain hardening ability is poor and it is difficult to prevent the grain growth, formability is poor due to anisotropy. A technology, in which different portions are heated at different temperatures to achieve a warm working process, is suggested so as to avoid the above disadvantage. However, the technology is problematic in that the different heating temperatures of the different portions significantly increase the production cost of a press mold. As an alternative method, a thixo-molding method, in which preliminarily flake shaped powder is compacted at high temperatures in a region where liquid and solid phases coexist, is suggested.

However, this method is disadvantageous in that the powder is expensive, and in that it is difficult to apply an electroplating process because the powder pressed material has a porous structure. Magnesium has low corrosion resistance, and thus, it is necessary to treat a surface of magnesium, but undesirably, a gas phase plating process or an electroless plating process requires chemical and treatment costs that are much higher than the electroplating process. However, products having high porosity and low density, such as die-casting and thixo-molding products, are difficult to apply to a wetplating process because corrosion occurs due to chemicals soaked into the pores.

Furthermore, Korean Pat. Laid-Open Publication No.2003-0048412 discloses an alloy, which contains 3.0 10.0 wt% Zn, 0.25 3.0wt% Mn, Al, Si, and Ca. However, even though the alloy containing Zn in an amount of 2 or more has high strength, it has a disadvantage in that free zinc (Zn) readily forms a low melting point euteetic phase. For example, if Mg Zn3, having a low melting point that is less than 350°C, exists, corrosion resistance is low. And the plate is easily cracked at both sides during a rough rolling process for breaking the coarse dendrite structure, so draw-ability is poor because of high anisotropy. Korean Pat. Laid-Open Publication No. 2002-0078936 Pat. No. 6471797) discloses a method of improving strength and formability using a Mg-Zn-Y eutectic ternary alloy quasi-crystalline phase which contains 1 at% Zn and 0.1 3 at% Y. However, this method is disadvantageous in that the amount of Zn must be enough so as to desirably assure a quasi-crystalline phase effect. The composition of a cast product is not uniform because a specific gravity between zinc and magnesium is significantly different. The micro-pores at the grain boundary reduce corrosion resistance, and tears form at sides of a plate during the hot rolling process.

WO 2006/075814 PCT/KR2005/000697 On the other hand, in Korean Pat. Application No.10-2003-0044997 which has been made by the inventor of the present invention, tears caused by an ununiform composition during rolling process is reduced by reducing the amount of Zn. However, since the second phase improving plasticity is a low melting point eutectic phase that is formed at grain boundary after the matrix is formed, the second phases are dispersed through the break down rolling after solidification. Accordingly, it is difficult to uniformly disperse them. Hence, the initial rolling process must be repeated a few times within a reduction ratio of 5 10 and the process must be conducted within a reduction ratio range of 15 20 after the cast structure is broken so as to get a good quality without side cracks.

Many other patents disclose a method of producing a light magnesium alloy strip or powder having high strength, in which an amorphous structure is produced through the Rapid Solidification process. Korean Pat. Laid-Open Publication No. 1990-0004953, entitled 'high strength magnesium alloy', Korean Pat. Laid-Open Publication No.

1993-846, entitled 'magnesium alloy having high strength', Japanese Pat. Laid-Open Publication No. H05-70880, entitled 'magnesium alloy material having high strength and method of producing the same', Japanese Pat. Laid-Open Publication No.

H06-41701, entitled 'amorphous magnesium alloy having high strength and method of producing the same', Japanese Pat. Laid-Open Publication No. H07-54026, entitled 'magnesium alloy having high strength and method of producing the same', U.S. Pat.

4675157, 4765954, 4853035, 4857109, 4938809, 5071474, 5078806, 5078807, 5087304, 5129960, and 5316598, EP No. 0,361,136A1, and French Pat. No. 2,688,233 disclose the formation of an amorphous structure through a Rapid Solidification process. Since the cooling rate must be conducted at 10 5 10 7 0C IS to form the amorphous structure, the patents are useful to produce powder or a thin strip, but not to produce a common plate shape. Accordingly, an ingot, which is produced by compacting amorphous powder under the recrystallization temperature, is employed in order to conduct a rolling or a press forming.

Furthermore, U.S. Pat. Nos. 637040, 3391034, 4116731, 4194908, and 5059390, and English Pat. No. 2095288 disclose the fact that some rare-earth elements are used to prevent the grain growth or the grain boundary slip at high temperatures while their eutectic phases exist at grain boundary so as to improve creep resistance. However, the eutectic phase mostly has a coarse microstructure that is incoherent to a matrix microstructure, and thus, formability is insufficiently improved. As well, Japanese Pat.

Laid-Open Publication Nos. H7-109538A, U.S. Pat. Nos. 5693158, 5800640 and 6395224 disclose a method of producing goods having low crack sensitivity, in which Sr, Li or B is employed and heat treatments are conducted to refine a particle size of crystals of a cast product. However, these patents are useful to cast products, but 486061 00 cannot be directly applied to wrought products. Japanese Pat. Laid-Open Publication No.

0 HIO-147830A discloses the use of 6 12 wt% Y and 1 6 wt% Gd, and hot forging and O subsequent aging processes to improve creep resistance to be applied to engine parts.

SHowever, the patent cannot be applied to wrought products because the product cost MC 5 significantly increases due to the use of a lot expensive elements, and the coarse intermetallic compounds are incoherent to the matrix. Furthermore, a method of improving formability, in which an excessive amount of Li is employed to change the lattice structure t of a matrix microstructure into a body-centered cubic lattice, is suggested. However, this C method is not useful to casing materials when considering a Galvanic reaction of Li and an Cc, V) lo increased cost due to the use of an excessive amount of Li.

SDisclosure of Invention Technical Problem Accordingly, the present invention has been made keeping in mind the above problems occurring in the prior art, and the present invention seeks to provide a wrought magnesium alloy which contains the intermetallic compound coherent to a matrix microstructure and which has a second phase composite microstructure, thereby improving elongation and anisotropy to assure excellent formability and corrosion resistance. In order to accomplish the above, an alloy consisting of three or more elements is used to activate a slip plane. Additionally, in order to activate the slip plane according to increasing in temperature, lia and IIIb groups are added together to reduce stacking fault energy and to improve corrosion resistance of the matrix microstructure. Furthermore, fine intermetallic compound particles dispersed during extrusion and rolling processes are employed to improve strain hardening ability and formability.

Technical Solution In order to accomplish the above, the present invention provides a wrought magnesium alloy having excellent formability and plating properties, which comprises 0.1 at% group Ila, 1.0 4.0 at% group IIIb, 0.35 at% or less of one selected from the group consisting of groups Ha, IVa, Vila, IVb, and a mixture thereof, 1.0 at% or less of group Ilb, and a balance of Mg and impurities and which thus has a second phase intermetallic compound.

486061 00 In one aspect, the present invention seeks to provide a wrought magnesium alloy 0 Shaving excellent formability and plating properties, which comprises 0.1-1.5 at% of a first o essential element selected from the group liIa elements Sc, Y and lanthanides and mixtures Sthereof, 1.0-4.0 at% of a second essential element selected from the group IIIb Al, B and mixture thereof, and 0.65 at% or less of an element selected from the group Ilb Zn, Cd and mixture thereof, 0.35 at% or less of one selected from the group consisting of group IIa Ca and Sr, group IVa Ti, Zr and Hf, group VIIa Mn, group IVb Si and Ge, and mixtures thereof, and a balance of Mg and unavoidable impurities, and thus contains a second phase CN of finely precipitated intermetallic compounds.

In another aspect, the present invention seeks to provide a method of producing a C, wrought magnesium alloy, comprising: preparing a magnesium alloy cast billet, 0.1-1.5 at% of a first essential element selected from the group IIIa elements Sc, Y and lanthanides and mixtures thereof, 1.0-4.0 at% of a second essential element selected from the group IIIb Al, B and mixture thereof, 1i and 0.65 at% or less of an element selected from the group IIb Zn, Cd and mixture thereof, 0.35 at% or less of one selected from the group consisting of group IIa Ca and Sr, group IVa Ti, Zr and Hf, group VIIa Mn, group IVb Si and Ge, and mixtures thereof, and a balance of Mg and unavoidable impurities, and thus contains a second phase of finely precipitated intermetallic compounds; subjecting the magnesium alloy cast billet to diffusion annealing at a temperature of 250-450 0

C;

reheating the diffusion-annealed magnesium alloy cast billet in a heat treatment furnace of 250-400 0 C, extruding the reheated billet, and then rolling the extruded billet.

Brief Description of the Drawings FIG. 1 illustrates a box sample formed using a wrought magnesium alloy sheet, according to the present invention; FIG. 2 illustrates a cup-shaped sample formed using the wrought magnesium alloy sheet, according to the present invention; WO 2006/075814 PCTIKR20051/000697 [11] FIG. 3 illustrates the box sample formed using AZ31 sheet; [12] FIG. 4 illustrates a microstructure of a material of No. 1 in Table 1, which is cast and then diffusion annealed at 400 0 C for 5 hours; [13] FIG. 5 illustrates a microstructure of an extruded material according to the present invention, which is annealed; and [14] FIG. 6 illustrates a mnicrostructure of a rolled sheet according to the present invention.

Best Mode for Carrying Out the Invention The present invention is characterized in that the fine 2nd phase precipitates, which is coherent to matrix microstructure, is formed in the solid solution microstructure having excellent ductility, thereby making grains fine and improving formability.

When the grains are made fine, the strength of most materials increases. The reason is that a dislocation moves along the specific slip plane in the course of plastic deformation of the metal in such a way that the dislocation does not directly move from one grain to another grain. But direction of dislocation changes its route because of the grain boundary barrier effect. Accordingly, since the grain boundaries act as barriers in the movement of the dislocation, dislocations are pile up at a grain boundary, thereby preventing deformation. The high temperature stable phase must be capable of being formed in order to make the grain fine, and desired solid solubility must be assured at high temperatures in order to be coherent to a matrix microstructure. Furthermore, a size difference between elements of a matrix metal and atoms must be about 15 so as to assure a desirable matrix reinforcement effect. Many studies have been made of the effect of an intermetallic compound on a solid solution. Particularly, a matrix reinforcement effect caused by the dispersion of fine intermetallic compound particles is well known in the metallurgy engineering (Mechanical Metallurgy, 2nd ed., George E.

Dieter, McGraw-Hill, 1981, pp. 221-227). The intermetallic compound has a high melting point and strong bonding strength, thus having high hardness and thermally stable. Because of the finely dispersed second phase particles, these alloys are much more resistant to recrystallization and grain growth than single-phase alloys. However, if the intermetallic compound has a microstructure that is incoherent to the matrix microstructure, it acts as a fracture initiation site and thus has increased strength, but elongation or total ductility is reduced even though the matrix microstructure has ductility.

[16] If the second phase of conventional magnesium alloy is not a high melting point phase in the matrix microstructure. The 2 d phase is low melting point eutectic phase during the solidification instead of the precipitates. Therefore, the eutectic phase is mostly incoherent to the matrix microstructure. It is scarcely an atomic match for the WO 2006/075814 PCT/KR2005/000697 matrix microstructure, thus effectively preventing grain growth or over-aging.

However, disadvantageously, it reduces the formability of a material or acts as a fracture initiation site. And thus, these type alloys are unsuitable for wrought magnesium alloy. Even though a duplex microstructure is formed, the movement of the dislocation is ineffectively prevented if the second phase is not strong, resulting in undesirably improved anisotropy or strength.

[17] It is known that group IIIa elements employed in the present invention readily form intermetallic compounds having a cubic lattice and thus have high a matrix reinforcement effect and ductility. Alan Russel and Karl Gschneidner Jr. of the Ames Laboratory of Iowa State University, which is affiliated with the U.S. Department of Energy, reported that an intermetallic compound formed by the group IIIa has a B2 cubic lattice, such as CsCI, unlike a B27, B33, or DO orthorhombic lattice of a conventional intermetallic compound. And thus, it has excellent ductility (Nature Materials, 2, Sep. 2003, PP 587-590). Currently, it has been reported that many intermetallic compounds containing group IIIa elements are coherent to the magnesium matrix, and it is presumed that the ductility of the intermetallic compound is caused by stacking faults.

[18] Furthermore, many researchers, including A.P. Tsai, ascertained the fact that since a quasi-crystal intermetallic compound formed by group IIIa elements has high adhesion energy and Young's modulus, it is a substance having high strength and ductility.

Based on the above fact, many studies have been made to apply the quasi-crystal intermetallic compound to a structural material. Particularly, in the magnesium alloy field, Japan and Korea have taken the lead in the studies of a Mg-Zn-Y alloy containing quasi-crystal particles (Materials Science and Engineering A300, 2001, pp.312-315; Acta Materialia 50 (2002) pp.

2 3 4 3- 2 356; Materials transactions vol. 42, No. 10 (2001) pp.2144-2147; TMS 2002 conference, Magnesium Technology 2002, pp.

14 1-150; Journal of Alloys and Compounds 342 (2002) pp.445-450).

[19] The above studies have proven the following fact. Since zinc separation occurs during a dissolution process due to the high Zn content (4 at% or more), the composition is ununiform. And free Zn forms a low melting point eutectic phase, thus undesirably causing a side crack during rolling process. However, Y, the group IIIa elements, forms a icosahedral quasi-crystal phase in conjunction with Mg and Zn, and thus, the phase reinforces a matrix while being coherent to the matrix, thereby effectively preventing grain growth at a high temperature until 400 oC Particularly, A.

Inoue of Japan confirmed using a high resolution electron microscope (HREM) that in a magnesium alloy, which is produced through an RSP process and contains 2 at% Y and 1 at% Zn, the ABACAB type of stacking faults are formed every 6 periods (Scripta Materialia 49 (2003) pp.

4 1 7-422; Philosophical Magazine Letters vol.82 WO 2006/075814 PCT/KR2005/000697 (2002) pp.

54 3 55 1; Acta Materialia vol. 50 (2002) pp.38 4 5-3857).

The stacking faults are formed because a stacking order of a closely packed side is changed unlike a normal stacking order, and it is known that they are mostly formed due to plastic deformation. It is difficult to form stacking faults if stacking fault energy is high, and thus, strain hardening required as a press material is not high. Accordingly, since pure aluminum or copper has high stacking fault energy, energy supplied during a room temperature process is mostly converted into heat. Thus, it is difficult to accumulate internal deformations, and a driving force for nucleation is reduced during recrystallization. However, in the magnesium alloy of the present invention, the group IIIb and IIIa elements are alloyed with magnesium acting as a matrix element, thereby reducing the stacking fault energy of the intermetallic compound to provide ductility.

Additionally, fine second phases promote nucleation during a reheating process to make fine grains. Intermetallic compound particles prevent grain growth at a recrystallization temperature or higher.

[21] Based on the above description, the present inventor came to a conclusion that when a group IIIa is alloyed with magnesium to form a solid solution having low stacking fault energy, when a group IIIb is added to the solid solution to increase a solid-solution strengthening effect, and when a group IIb and other miniaturized elements are added to form a structure which contains an intermetallic compound coherent thereto, it is possible to create a material having excellent strain hardening ability, fineness through recrystallization by heat treatment, and improved anisotropy.

[22] Hereinafter, a detailed description will be given of elements and compositions of the wrought magnesium alloy according to the present invention.

[23] The group IIIa, that is, an essential element in the present invention, includes Sc, Y, lanthanides, and actinides. In this regard, it is preferable that Sc, Y, or lanthanides be employed alone or in combination instead of actinides radiating radioactive rays. They are solid-solved in Mg, thus reducing a c/a ratio to increase ductility and reducing the stacking fault energy to increase the driving force for nucleation by recrystallization.

Furthermore, particles, which exist in a form of Mg RE at high temperatures during a solidification process, form prism-shaped plate particles having a HCP structure, that is, a DO9 lattice structure, such as Mg RE or Mg 7RE at about 550 0 C through a 19 3 17 5 peritectic transformation. (RE is an abbreviated form of rare-earth elements belonging to the group IIIa) Thereby, the particles have a high reinforcement effect and are coherent to the matrix, and consequently, they do not act as the fracture initiation site.

After a rolling process, the particles may be compacted into a rod, a sphere, or a cube.

[24] In the present invention, a eutectic phase, which is not solid-solved after a diffusion heat-treatment, is finely dispersed during extrusion and rolling processes, thereby preventing grain growth during heat-treatment and acting as a site for nucleation by re- WO 2006/075814 PCTIKR20051/000697 crystallization. When the amount of the group Ia is less than 0.1 the second phase is formed in an insufficient amount. When the amount is more than 1.5 a fineness effect is saturated, and consequently, elongation is reduced and a production cost increases. This is the reason why that amount is limited.

The group IIIb includes B, Al, Ga, In, and Ti. Since Ga, In, and T1 having a low melting point form a low melting point eutectic phase, it is preferable to employ only Al or a mixture of B and Al. The group 1111b forms a fine deposit and thus contributes to reinforcement of the matrix. Al is used as a main alloy element. Since B has a low solid solubility to magnesium and forms a high melting point compound, such as B Y, 2 B3Y 2 or B 5

Y

3 it is employed in conjunction with Al in an amount of 0.010 or less so as to make the fine grains.

[26] In the present invention, Al of the group IIIb is solid-solved in Mg to increase corrosion resistance and to prevent the growth of a dendrite microstructure, thereby making a cast microstructure fine. Furthermore, since Al forms fine cubes, such as Al2 RE or Al RE during the solidification process and increases ductility of the matrix mi- 3 crostructure, it is possible to produce goods having high strength and excellent ductility. When the amount of Al is less than 1.0 it is difficult to assure the desirable reinforcement effect. When the amount is more than 4.0 since an unstable rod- or plate-shaped A]2Mg 3 or Al12Mg17 phase is enlarged in a grain boundary, even though room temperature strength is high, high temperature strength and corrosion resistance are reduced. This is the reason why that amount is limited.

[27] In order to make the fine grains and to help form the intermetallic compound, 0.35 or less of group IIa, group IVa, group VIIa, or group IVb is selectively employed alone or in combination, and 1.0 or less of group IIb is employed alone or in combination.

[28] The group IIa, group IVa, and group VIIa are used as a supplemental agent of the group IIIa and group III1b. In the group IIa, it is preferable to use Ca and Sr. Since Be, Ba, and Ra make toxic gases, they can be used only if a special ventilation device is adopted. Ca and Sr are particularly useful to make a fine cast structure in the casting a billet having a diameter of 200 mm or more in the present invention, and form diskshaped particles, such as (Mg, A1)2Ca, thereby improving a reinforcement effect.

[29] In the group Va, Ti, Zr, and Hf are most frequently employed, and Rf is added, using a protection device in unavoidable cases, because of the emission of radioactive rays. The group IVb makes a cast microstructure fine, and Si and Ge are most frequently employed because a melting point is high and it is easy to handle. A grain fining effect depends on the amount of each element added. That is, Zr, Si, and Ca make the grains have fine sizes of microns corresponding to the reciprocal of 52, 19, and 15 microns.

WO 2006/075814 PCTIKR20051/000697 Mn of the group VIIa is a cheap alloy element, prevents the formation of Al12Mg17 and A12Mg 3 phases, and promotes the formation of high temperature cubic Al2 Y to contribute to the fining of the grains and the improvement of corrosion resistance. Tc and Re of the group VIIa are costly and thus are used in unavoidable cases.

[31] The group IIa, group IVa, group VIIa, and group IVb elements have low solid solubility to magnesium, and thus, if they are excessively added, segregation occurs or coarse particles having high brittleness are formed when a cooling rate is low after a casting process. Accordingly, the amount is limited to 0.35 or less.

[32] The group IIb includes Zn, Cd, and Hg. Since Hg is toxic to humans when breathing, use of Hg is limited, and it is used in conjunction with an additional protection device. When Zn and Cd are added alone or in combination, a stacking fault structure is formed in a magnesium matrix microstructure to bring about strain hardening, and Zn and Cd are smoothly solid-solved with the group IIIa and group IIIb elements to promote the formation of cubic particles, such as (Mg, Zn) RE, Zn6Mg 2

RE,

or (Mg, Zn)17RE 3 However, the excessive amount of Zn and Cd increases gas solid solubility, thereby reducing corrosion resistance or plating workability and brining about the occurrence of hot tear and gravity separation phenomena. Hence, the amount is limited to 1.0 or less, and preferably, 0.65 or less.

[33] Hereinafter, a method of producing a plate using a magnesium alloy slab according to the present invention is described in detail through the following example which is set forth to illustrate, but is not to be construed as the limit of the present invention.

[34] A magnesium raw material is melted, and an alloy or a master alloy is added to the molten magnesium in a mixed gas atmosphere of SF and Ar or CO or an Ar gas 62 atmosphere while being blocked from contact with atmospheric air. Generally, a slab for a magnesium alloy plate is produced through mold casting, Direct Chilled casting, continuous casting, or strip casting processes.

In the present example, a mold, in which a cavity having a thickness of 30 mm, a width of 250 mm, and a height of 400 mm is formed, is preheated in a heating furnace heated to about 200 0 C. A molten magnesium alloy is poured into the mold at 710 760C, and then machined so as to remove surface defects from a cast product.

[36] Diffusion annealing is conducted at 250 450 0 C so that the duration time is 1 min/mm or more with respect to the thickness of the slab. When the heating temperature is less than 250 0 C or the duration time is less than 1 min/mm, the inside of the slab is insufficiently heated, thus forming cracks on a surface or on an edge during a rolling process. It is preferable to heat the slab at 350 400 0 C so as to reduce the diffusion time. When the heating temperature is more than 450 0 C, a free low melting point eutectic phase may be formed during the diffusion annealing. At this stage, the eutectic phase may be remelted and thus separated from the slab. Accordingly, the WO 2006/075814 PCT/KR2005/000697 molten eutectic phase may cling to a rolling roll. When the amount of the alloy element is large, the duration time and the heating temperature increase to improve workability.

[37] Initial coarse rolling is conducted once or more in a reduction ratio of 20 or less each time so as to fracture a coarse cast microstructure of a material, which is subjected to diffusion annealing, and to remove fine segregation. After the completion of a rolling process, a process annealing is conducted once or more at 200 450 0 C so that the duration time is 1 min/mm or more with respect to the thickness of a slab.

When the heating temperature is less than 200oC or the duration time is less than 1 min/mm, the inside of the slab is insufficiently heated, thus forming cracks on a surface or on an edge during the rolling process. In the initial coarse rolling, when a reduction ratio is more than 20 cracks may be formed at a grain boundary of a cast microstructure. At this stage, a surface temperature of a rolling roll must be maintained at 50 150 0 C so as to prevent the formation of fine surface cracks caused by the qu enched slab while the slab is in contact with the roll. When the temperature of the rolling roll is more than 150 0 C, delamination, in which a portion of a rolling material clings to the rolling roll and then delaminates, occurs during the rolling process, thus roughening the surface of the slab. If the plate is not excessively cooled after the initial coarse rolling, it is possible to conduct the rolling process again without reheating.

[38] When the cast microstructure of the slab is fractured, a second rolling process is repeatedly conducted in a reduction ratio of 50 or less each time until the desired thickness is gained. At this stage, the reduction ratio depends on the capacity of a motor of a rolling mill, a heat emitting state of a plate during a reduction process, elastic deformation of the rolling roll, and flatteness of the plate. It is preferable that a second process annealing be repeatedly conducted at 200 450 0 C each time while a duration time is maintained at I min/mm or more during the second rolling process.

However, in the second rolling process, a rolled microstructure becomes fine, causing crack resistance. Furthermore, in some cases, it is possible to conduct cold rolling.

Thus, annealing is not necessarily conducted every rolling process.

[39] After the final rolling process is completed, the final annealing is conducted at 180 350 0 C while a duration time is maintained at 1 min/mm or more, which depends on the thickness, strength, and elongation of the plate. When the annealing temperature is high and the time is long, elongation increases but strength is reduced. Particularly, when the annealing temperature is more than 350°C, undesirably, yield strength is significantly reduced.

Hereinafter, a detailed description will be given of extrusion using a magnesium alloy billet according to the present invention.

[41] A magnesium raw material is melted, and an alloy raw material or a master alloy is added to the molten magnesium raw material in a mixed gas atmosphere of SF and 6 WO 2006/075814 PCTIKR2005/0100697 Ar or CO, or an Ar gas atmosphere while being blocked from contact with atmospheric air. Subsequently, a molten magnesium alloy is poured into a mold, having a diameter of 185 mmn and a length of 650 mm, at 710 7600C to form a billet, and is then processed so as to remove surface defects. Needless to say, it is possible to conduct a continuous casting in addition to a mold casting.

[42] Diffusion annealing is conducted at 250 450°C while a duration time is maintained at 1 min/mm or more with respect to the diameter of the billet so as to fracture a coarse cast microstructure of a cast material and to remove fine segregation.

When a heating temperature is less than 2500C or the duration time is less than 1 min/ mm, stress is concentrated on a grain boundary, and consequently, alligatoring may occur, cracking the material in a direction of the extrusion. It is preferable to heat the material at 350 400( so as to reduce a diffusion time. When the heating temperature is more than 450°C, a free low melting point eutectic phase may be re-melted during the diffusion annealing and thus be separated from the material. When the amount of the alloy element is large, the duration time and the heating temperature increase to improve workability.

[43] The diffusion-annealed material is reheated in a heating furnace at 250 400C to be extruded. An extruder has an extrusion speed of a maximum of 20 m/min at an extrusion pressure of 850 MPa or more. If the extrusion is conducted at 500 MPa, the extrusion speed is significantly reduced to 3 4 m/min. A temperature of a container is 300 4500C. When the temperature is less than 300C, many surface cracks are formed. When the temperature is more than 4500C, high temperature cracks or deformations are significantly formed during the extrusion process. The container is heated at about 3500C, and an extrusion ratio is typically 10 100. Additionally, in the present invention, the material may be wound in a coil form during the extrusion process, and thus, it is possible to conduct reciprocating rolling.

[44] If the billet is very large or the cast microstructure is coarse, a first extrusion is conducted to fracture the cast microstructure and to disperse a second phase, and a second extrusion is then conducted. After the first extrusion, it is preferable to conduct a process annealing at 200 450C while a duration time is maintained at 1 min/mm or more. However, during the first extrusion, the microstructure is made fine, causing crack resistance, and the reheating is implemented in the container. Hence, annealing is not necessarily conducted.

After the final rolling process is completed, if the material is rolled into a plate, the final annealing is conducted at 180 350C while a duration time is maintained at 1 min/mm or more, which depends on a thickness, strength, and elongation of the plate.

When the annealing temperature is high and the time is long, elongation increases but strength is reduced. Particularly, when the annealing temperature is more than 3500C, WO 2006/075814 PCT/KR2005/000697 undesirably, yield strength is significantly reduced. Needless to say, when the plate and coil are annealed, the heat treatment may be implemented using a rapidly heating device, such as a heater, employing a gas nozzle, or an induction heater, instead of the furnace. At this stage, since a heating rate is high, it is necessary to set the annealing temperature higher. In this regard, the annealing temperature may deviate from the above range, without departing from the scope and concept of the invention.

[46] As shown in following Table 1 and 2, wrought magnesium alloys of the present invention were rolled to obtain test results. They were tested after being rolled into plates having a width of 150 mm and a thickness of 1 mm.

[47] Rectangular molds, which had a width of 80 mm, a length of 100 mm, and a depth of 45 mm, were formed, and edge cracks of the molds were observed, thereby achieving a forming test. Samples having an area of 80 mm X 50 mm were hung on a nylon thread as a hanger, and immersed in 200 cc of 2 HC1 aqueous solution in a beaker. Thereby, gases, generated from the samples, were dissolved in the solution. At this stage, weight reduction was measured, thereby achieving evaluation of corrosion resistance. The evaluation of formability is as follows. O means that forming is achieved without cracks and local reduction of a thickness, A means that cracks are not formed but a thickness deviation locally occurs, and x means that formability is very poor because of the formation of cracks. In the evaluation of characteristics of a wet plating process, O means a state that plating thickness and adhesion of a plated surface are excellent. A means a state that adhesion is fair, pinhole is not observed, and plating thickness is ununiform. x means a state in which the pinholes are observed or plating layer comes off the surface somewhere in the specimen.

[48] Table 1 No Chemical component(at%)The Process Mechanica ResultsO: Note balance includes impurities and Mg (dimension 1 excel-lent Group Group Group Group of mold, properties A:fairx:p liIa IIIb Ha, IVa, IIb and unit is oor VIIa, mm) IVb 1 Y 0.2Sc Al 2.00 Mn Zn 0.40 30 x 250 x 1 T. 4 F. 0 5 C.R. 0.11 0.13 400Mold 270MPaY 3.56P. O casting .224MPa 3 El 22% 2 Y Al 1.50 Mn Zn Dia. T. 4 F. 0 5 C.R. 0.50Nd 0.15 0.50Cd 185Billet 257MPaY. 3.2 6 P. O 0.01 0.10 casting 213MPaEl WO 2006/075814 WO 206/05814PCTIK-R20051000697 1 1 20% 1 3 Y 0.50 Al 3.00 Ca Zn 0.30 Dia. T. 4 F. 0 5 C.R. 10 1.S.

0.OlMn 185Billet 376MPaY. 3.0 6 P. 0 0.10 casting 312MPaEl 21% 4 Y Al 2.50 Sr Zn 0.50 Dia. T. 4 F. 0 5 C.R. 1 0 1I.S.

0.45Sc 0.O2Mn 3O5Billet 355MPaY. 3.3 6 P. 0 0.10 0.2OZr casting 287MPaE1 0.08 19% Y Al 3.00 Si Zn Dia. T. 4 F. 0 5 C.R. 1 0 1I.S.

1.OOLa 0.lOMn 0.4OCd 185Billet 346MPaY. 2.9 6 P. 0 0.05 0.15 0.10 casting 287MPaE1 18% 6 Y Al 3.50 Zr 0.05 Zn 0.50 Dia. T. 4 F. 0 5 C.R. 10 1S.

0.4ONd 185Billet 4O5MPaY. 3.4 6 P. 0 0.03 casting 336MPaEl 18% 7 Y 0.30 Al Zr 0.10 Zn 0.50 Dia. T. 4 F. 0 5 C.R. 101I.5.

2.50B 185Billet 331MPaY. 3.2 6 P. 0 0.005 casting 275MPaEl 23% Table 2 8 Y0. 15 Al 2.00 Ca Zn 0.30 Dia. T. 4 F. 0 5 C.R. 101.S.

0.lOMn 185Billet 245MPaY. 3.8 6 P. 0 0.10 casting 2O3MPaEl 18% 9 Y 0.25 Zr 0.80 Zn 1.55 D360 x T. 4 F. 0 5 C.R. 11

C.S.

tl2OD.C. 285MPaY. 3.8 7 failed casting 253MPaEl Ni .16% Y0. 15 AlO0.90 Zn 0.75 30 x250 x T. 4 F. 0 5 C.R. 11

C.S.

400Mold 261MPaY. 5.2layer casting 2OSMPaEl off 18% 11 Al 2.54 M~n Zn 0.30 30 x250 x T. 4 F. x 5 C.R. 11

C.S.(

486061 00

U

0.09 400Mold 265 MpaY. 4.19pinhole AZ31) casting 185MpaEl .18% Tensile strength 2 Y Yield point 3 El. Elongatin formability, 5 CR.: corrosion weight loss rate, 6p.: plating, 7 Failed Ni: impossible to form a Ni plating, 8 Layer off: the plating layer comes off the surface, 9 Pinhole: formation of micro pinholes in the plating layer,

I

0 LS.: present inventive sample CS.: comparative sample (No.1,9,10,11 rolling speed: 1.6 m/min, reduction ratio: 15 during a initial coarse rolling and then 20 45 No. 2-8 rolling speed: 16 20 m/min, reduction ratio: 30 67% after extrusion) When evaluating corrosion resistance, weights of the beakers, in which the samples were contained, were measured every five minutes for 60 min using a precision scale having an allowable margin of error of 1/1000 g to calculate a slope of weight reduction, thereby completing the evaluation of corrosion resistance. The higher slope brings about increased weight reduction, resulting in poor corrosion resistance.

In Table 1, since a magnesium alloy of No. 11, which is produced according to a conventional method, has poor formability, cracks are formed during the forming process as shown in FIG. 3. In the wet plating process, activation treatment is conducted in liquid, an electroplating process, such as a copper cyanide plating, a copper sulfate plating, or a nickel plating, is then implemented, and subsequently, a final plating, such as a chromium plating or a precious metal plating, is implemented. At this stage, if the pinholes are formed or a plating layer comes off surface, the reliability of corrosion resistance is significantly reduced.

Industrial Applicability As described above, in the present invention, a fine second phase intermetallic compound is dispersed so as to significantly improve the poor formability and corrosion resistance of a conventional magnesium plate. Thereby, the magnesium plate has excellent properties as a structural material, and consequently, it is possible to apply the magnesium plate to structural materials used in portable electronic products, automobiles, or airplanes.

As used herein, except where the context requires otherwise, the term "comprise" 486061 00 and variations of the term, such as "comprising", "comprises" and "comprised", are not intended to exclude other additives, components, integers or steps.

Reference to any prior art in the specification is not, and should not be taken as, an acknowledgment, or any form of suggestion, that this prior art forms part of the common 0 5 general knowledge in Australia or any other jurisdiction or that this prior art could reasonably be expected to be ascertained, understood and regarded as relevant by a person Sskilled in the art.

Claims (2)

1. A wrought magnesium alloy having excellent formability and plating Sproperties, which comprises 0.1-1.5 at% of a first essential element selected from the group cr Ilia elements Sc, Y and lanthanides and mixtures thereof, 1.0-4.0 at% of a second essential element selected from the group IIIb Al, B and mixture thereof, and 0.65 at% or less of an element selected from the group IIb Zn, Cd and mixture thereof, 0.35 at% or less of one selected from the group consisting of group IIa Ca and Sr, group IVa Ti, Zr and Hf, group 1 VIIa Mn, group IVb Si and Ge, and mixtures thereof, and a balance of Mg and unavoidable I impurities, and thus contains a second phase of finely precipitated intermetallic compounds. N o 2. A method of producing a wrought magnesium alloy, comprising: preparing a magnesium alloy cast billet, 0.1-1.5 at% of a first essential element selected from the group Ilia elements Sc, Y and lanthanides and mixtures thereof, 1.0-4.0 at% of a second essential element selected from the group IIIb Al, B and mixture thereof, and 0.65 at% or less of an element selected from the group lib Zn, Cd and mixture thereof, is 0.35 at% or less of one selected from the group consisting of group IIa Ca and Sr, group IVa Ti, Zr and Hf, group Vila Mn, group IVb Si and Ge, and mixtures thereof, and a balance of Mg and unavoidable impurities, and thus contains a second phase of finely precipitated intermetallic compounds; subjecting the magnesium alloy cast billet to diffusion annealing at a temperature of

250-450 0 C; reheating the diffusion-annealed magnesium alloy cast billet in a heat treatment furnace of 250-400 0 C, extruding the reheated billet, and then rolling the extruded billet. 3. A wrought magnesium alloy according to claim 1, substantially as hereinbefore described with reference to any one of the examples. 4. A method of producing a wrought magnesium alloy according to claim 2, substantially as hereinbefore described with reference to any one of the examples.