CA2423459C - High strength magnesium alloy and its preparation method - Google Patents
High strength magnesium alloy and its preparation method Download PDFInfo
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- CA2423459C CA2423459C CA002423459A CA2423459A CA2423459C CA 2423459 C CA2423459 C CA 2423459C CA 002423459 A CA002423459 A CA 002423459A CA 2423459 A CA2423459 A CA 2423459A CA 2423459 C CA2423459 C CA 2423459C
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C23/00—Alloys based on magnesium
- C22C23/02—Alloys based on magnesium with aluminium as the next major constituent
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C23/00—Alloys based on magnesium
- C22C23/04—Alloys based on magnesium with zinc or cadmium as the next major constituent
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- 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
- C22F1/06—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of magnesium or alloys based thereon
Abstract
The present invention provides high strength magnesium alloys consisting essentially of 3 to 10 wt.% Zn, 0.25 to 3.0 wt.% Mn, and the balance of Mg and inevitable impurities, the high strength magnesium alloy further containing 1 to 6 wt.% Al, 0.1 to 4.0 wt.% Si, and 0.1 to 2.0 wt.% Ca, in order to provide a high strength magnesium alloy having an improved hardness and strength, and an excellent elongation at an ambient temperature. In addition, the present invention provides a method for preparing the high strength magnesium alloy characterized in that a Zn-Mn mother alloy is added to a magnesium melt by a fluxless melting method, and process conditions for working and heat-treating an obtained cast material.
Description
HIGH STRENGTH MAGNESIUM ALLOY AND ITS PREPARATION
METHOD
TECHNICAL FIELD
The present invention relates to high strength magnesium alloys and a method for preparing the same, and more specifically to magnesium alloys having, improved mechanical properties including strength, hardness, and elongation while having an improved formability, high strength and elongation and their economical processing method, by adding specific alloying elements or changing processing conditions including specific heat treatments.
BACKGROUND ART
Among magnesium alloys, Mg-Zn-based alloys exhibit a superior age hardening behavior. These alloys exhibit relatively high strength and ductility while having advantages in that it is easily processible and weldable.
On the other hand, Mg-Zn-based alloys have also disadvantages since it is difficult for them to be applied to a casting process such as die-casting because the addition of Zn to Mg tends to increase the formation of micro-pores during casting.
Furthermore, such Mg-Zn-based alloys have a limitation in improvement in strength because it is not easy to refine the microstructure by the addition of alloying elements or an over-heating heat treatment, compared to other magnesium alloys. This limitation has restricted their commercial applications.
In order to solve these problems, research efforts have been made to add alloying elements to an Mg-Zn binary alloy. Examples of these research efforts are as follows.
In 1947, J. P. Doan and G. Ansel have proposed a method in which Zr is added to an Mg-Zn-based alloy to refine the grain size of the alloy, thereby improving the strength of the alloy (J. P. Doan and G. Ansel, Trans.
AIME, vol. 171 (1947), pp. 286-295). In this method, however, there is a difficulty in adding Zr to magnesium melts because Zr has a high melting point.
Also, the addition of rare earth elements such as La, Ce or Nd, or Th to Mg-Zn alloy has been known. This method is known to have advantages in that it is possible to reduce micro-pore formation and improve strength at high temperatures and weldability. However, this method has significant cost disadvantages due to expensive alloying elements, compared to other magnesium alloys.
In 1987, W. Unsworth and J. F. King have reported that the ductility of Mg-Zn alloys might be improved by refining the Pl' phase, which is the major strengthening precipitation phase of Mg-Zn alloy, by the addition of Cu to the Mg-Zn alloy (W. Unsworth and J. F. King, Magnesium Technology, The Inst. of Metal, 1987, pp. 25-35). However, the addition of Zn and Cu involves a limited elongation of less than 10 % at ambient temperature even though the elongation may be slightly increased depending upon the added amounts of those elements.
The following Table 1 shows tensile properties of commercial cast alloys and wrought alloys.
METHOD
TECHNICAL FIELD
The present invention relates to high strength magnesium alloys and a method for preparing the same, and more specifically to magnesium alloys having, improved mechanical properties including strength, hardness, and elongation while having an improved formability, high strength and elongation and their economical processing method, by adding specific alloying elements or changing processing conditions including specific heat treatments.
BACKGROUND ART
Among magnesium alloys, Mg-Zn-based alloys exhibit a superior age hardening behavior. These alloys exhibit relatively high strength and ductility while having advantages in that it is easily processible and weldable.
On the other hand, Mg-Zn-based alloys have also disadvantages since it is difficult for them to be applied to a casting process such as die-casting because the addition of Zn to Mg tends to increase the formation of micro-pores during casting.
Furthermore, such Mg-Zn-based alloys have a limitation in improvement in strength because it is not easy to refine the microstructure by the addition of alloying elements or an over-heating heat treatment, compared to other magnesium alloys. This limitation has restricted their commercial applications.
In order to solve these problems, research efforts have been made to add alloying elements to an Mg-Zn binary alloy. Examples of these research efforts are as follows.
In 1947, J. P. Doan and G. Ansel have proposed a method in which Zr is added to an Mg-Zn-based alloy to refine the grain size of the alloy, thereby improving the strength of the alloy (J. P. Doan and G. Ansel, Trans.
AIME, vol. 171 (1947), pp. 286-295). In this method, however, there is a difficulty in adding Zr to magnesium melts because Zr has a high melting point.
Also, the addition of rare earth elements such as La, Ce or Nd, or Th to Mg-Zn alloy has been known. This method is known to have advantages in that it is possible to reduce micro-pore formation and improve strength at high temperatures and weldability. However, this method has significant cost disadvantages due to expensive alloying elements, compared to other magnesium alloys.
In 1987, W. Unsworth and J. F. King have reported that the ductility of Mg-Zn alloys might be improved by refining the Pl' phase, which is the major strengthening precipitation phase of Mg-Zn alloy, by the addition of Cu to the Mg-Zn alloy (W. Unsworth and J. F. King, Magnesium Technology, The Inst. of Metal, 1987, pp. 25-35). However, the addition of Zn and Cu involves a limited elongation of less than 10 % at ambient temperature even though the elongation may be slightly increased depending upon the added amounts of those elements.
The following Table 1 shows tensile properties of commercial cast alloys and wrought alloys.
Table 1 Properties of Commercial Magnesium Alloys Composition (%) Tensile Properties Tensi Yield le Stren Stren Elonga T Z gth gth tion Alloys Al Mn h Zn r Others (MPa) (Mpa) (%) Cast AZ91 C- 8.7 0.13 - 0.7 - - 145 275 6 EQ21A- - - - - 0. 1.5 Ag, 195 235 2 H K31 A- - - 3. - 0. - 105 220 8 WE54A- - - - - 0. 5.2 Y, 172 250 2 ZC63A- - 0.25- - 6.0 - 2.7 Cu 125 210 4 ZE63A- - - - 5.8 0. 2.6 RE 190 300 10 ZK61A- - - - 6.0 0. - 195 310 10 Die-cast AM60A- 6.0 0.13 - - - - 115 205 6 AS41A-F 4.3 0.35 - - - 1.0 Si 150 220 4 AZ91A,B 9.0 0.13 - 0.7 - - 150 230 3 Extrude AZ80A- 8.5 - - 0.5 - - 275 380 7 M 1 A-F - 1.2 - - - - 180 255 12 ZC71-F - 0.5-1.0 - 6.5 - 1.2 Cu 340 360 5 HM31A- - 1.2 3. - - - 230 290 10 ZK60A- - - - 5.5 0. - 305 365 11 Sheet AZ31 B- 3.0 - - 1.0 - - 220 290 15 HK31A- - - 3. - 0. - 200 255 9 HM21A- - 0.6 2. - - - 170 235 11 PE 3.3 - - 0.7 - - - - -Referring to Table 1, it can be found that commercial wrought alloys generally exhibit higher yield strength, tensile strength, and elongation compared to commercial cast alloys. Even in the case of commercial wrought alloys. However, it is difficult even for the existing commercial wrought alloys to obtain a combination of high strength and high elongation.
That is, high strength alloys exhibiting a high tensile strength exceeding 300 MPa have a drawback in that their elongation rarely exceed about 10%. It has also been reported that, in the case of Zn and Zr-added alloys exhibiting superior properties in terms of strength, there are many restrictions on preparing processes in adding Zr.
US Patent No. 4,997,622 discloses properties of magnesium alloys prepared by a rapid solidification processing method. According to This Patent, magnesium alloys prepared by a rapid solidification processing method exhibit improved yield strength, tensile strength and elongation.
However, the results of research made up to the present show that those alloys have high processing costs and limited applications, compared to the existing commercial alloys.
DISCLOSURE OF THE INVENTION
An object of the invention is to provide high strength magnesium alloys exhibiting improvements in refining of the microstructure and precipitation behavior, enhancements in mechanical properties such as hardness, strength, and elongation, and an improvement in formability in accordance with an addition of alloying elements, less expensive than those conventionally used, to an Mg-Zn-based alloy.
Another object of the invention is to provide a method for preparing high strength magnesium alloys exhibiting a definitely superior elongation, compared to strength, using an optimum heat treatment condition, and to provide an economical processing condition for the preparation method.
In accordance with one aspect, the present invention provides high strength magnesium alloys consisting essentially of 3- 10 wt.% Zn, 0.25 - 3.0 wt.% Mn, and the balance of Mg and inevitable impurities.
The magnesium alloy may further contain 1- 6 wt.%
Al. Also, the magnesium alloy may further contain 0.1 -4.0 wt.% Si, or a combination of 0.1 - 4.0 wt.% Si and 0.1 -2.0 wt.% Ca. Preferably, the content of Al is not more than the content of Zn.
Preferably, the content of Zn is 5.0 - 7.0 wt.%, the content of Mn is 0.75 - 2.0 wt.%, the content of Si is 1.5 - 3.0 wt.%, and the content of Ca is 0.3 - 1.0 wt.%.
The important feature of the present invention is to add Al, as an alloying element, to an Mg-Zn-based alloy so as to achieve a decrease in yield strength resulting in an improvement in formability and an enhancement in work hardening ability, thereby to provide a high strength magnesium alloy having a high strength and a high elongation.
According to one aspect of the present invention, there is provided a high strength magnesium alloy consisting essentially of: 3 to 10 wt.% Zn, forming a precipitation phase in the alloy; 0.25 to 3.0 wt.% Mn refining the precipitation phase; 1 to 6 wt.% Al; 0 to 4.0 wt.% Si; 0 to 2.0 wt.% Ca; and the balance of Mg and inevitable impurities, wherein the content of Al is not more than the content of the Zn.
That is, high strength alloys exhibiting a high tensile strength exceeding 300 MPa have a drawback in that their elongation rarely exceed about 10%. It has also been reported that, in the case of Zn and Zr-added alloys exhibiting superior properties in terms of strength, there are many restrictions on preparing processes in adding Zr.
US Patent No. 4,997,622 discloses properties of magnesium alloys prepared by a rapid solidification processing method. According to This Patent, magnesium alloys prepared by a rapid solidification processing method exhibit improved yield strength, tensile strength and elongation.
However, the results of research made up to the present show that those alloys have high processing costs and limited applications, compared to the existing commercial alloys.
DISCLOSURE OF THE INVENTION
An object of the invention is to provide high strength magnesium alloys exhibiting improvements in refining of the microstructure and precipitation behavior, enhancements in mechanical properties such as hardness, strength, and elongation, and an improvement in formability in accordance with an addition of alloying elements, less expensive than those conventionally used, to an Mg-Zn-based alloy.
Another object of the invention is to provide a method for preparing high strength magnesium alloys exhibiting a definitely superior elongation, compared to strength, using an optimum heat treatment condition, and to provide an economical processing condition for the preparation method.
In accordance with one aspect, the present invention provides high strength magnesium alloys consisting essentially of 3- 10 wt.% Zn, 0.25 - 3.0 wt.% Mn, and the balance of Mg and inevitable impurities.
The magnesium alloy may further contain 1- 6 wt.%
Al. Also, the magnesium alloy may further contain 0.1 -4.0 wt.% Si, or a combination of 0.1 - 4.0 wt.% Si and 0.1 -2.0 wt.% Ca. Preferably, the content of Al is not more than the content of Zn.
Preferably, the content of Zn is 5.0 - 7.0 wt.%, the content of Mn is 0.75 - 2.0 wt.%, the content of Si is 1.5 - 3.0 wt.%, and the content of Ca is 0.3 - 1.0 wt.%.
The important feature of the present invention is to add Al, as an alloying element, to an Mg-Zn-based alloy so as to achieve a decrease in yield strength resulting in an improvement in formability and an enhancement in work hardening ability, thereby to provide a high strength magnesium alloy having a high strength and a high elongation.
According to one aspect of the present invention, there is provided a high strength magnesium alloy consisting essentially of: 3 to 10 wt.% Zn, forming a precipitation phase in the alloy; 0.25 to 3.0 wt.% Mn refining the precipitation phase; 1 to 6 wt.% Al; 0 to 4.0 wt.% Si; 0 to 2.0 wt.% Ca; and the balance of Mg and inevitable impurities, wherein the content of Al is not more than the content of the Zn.
In accordance with another aspect, the present invention provides a method for preparing the high strength magnesium alloy, in which an addition of Mn to a magnesium melt is achieved by adding a Zn-Mn mother alloy to the magnesium melt.
Preferably, the high strength magnesium alloy is prepared in the form of a cast ingot by adding a Zn-Mn mother alloy having an Mn content of 10 - 20 wt.% (Zn-10 -20 wt.% Mn mother alloy) to the magnesium melt in a temperature range of 670 to 720 C, and adding Zn or Zn along with Al to the magnesium melt. Alternatively, the high strength magnesium alloy may be 5a preferably prepared in the form of a cast ingot by adding a Zn-10 - 20 wt.%
Mn mother alloy to the magnesium melt in a temperature range of 670 to 720 C, adding an Mg-Si mother alloy to the magnesium melt, and adding Zn or Zn along with Al and/or Ca to the magnesium melt.
Preferably, the cast ingot may be subsequently subjected to a homogenization process in a temperature range of 340 to 410 C for 6 to 12 hours, so that it is formed into a billet. The billet may be mechanically worked after being preheated in a temperature range of 150 to 400 C for 30 minutes to 2 hours.
More preferably, the worked or wrought body may be subjected to a primary aging process in a temperature range of 70 to 100 C for 24 to 96 hours, and then to a secondary aging process in a temperature range of 150 to 180 C for 48 hours or more.
Prior to the double aging process, a solution heat treatment may be carried out in a temperature range of 340 to 410 C for 6 to 12 hours.
Alternatively, a stretching of 3 to 7% may be carried out prior to the double aging process.
The reason why the composition range of each alloying element used in accordance with the present invention is limited to the above mentioned range is as follows:
Zinc (Zn): 3 - 10 wt.%
The maximum solid solubility limit of Zn in an Mg matrix is 6.2 wt.%
at 340 C. Where Zn is added in an amount of 3.0 wt.% or more to Mg matrix, it forms an acicular precipitation phase when it is subjected to a heat treatment, thereby exhibiting an age hardening behavior. Generally, the content of Zn is determined, based on the solid solubility limit thereof. When Zn is added in an amount of about 5.0 - 7.0 wt.% approximating to the maximum solid solubility limit thereof, it is possible to obtain a maximized age hardening behavior. At the Zn content of not more than 3.0 wt.%
corresponding to a minimum solid solubility limit at a general aging temperature, it is difficult to expect a desired precipitation strengthening phenomenon because of an insufficient formation of precipitation phase. On the other hand, when Zn is added. in an amount of 10.0 wt.% or more, precipitation of equilibrium phases at grain boundaries is promoted. As a result, degradation in mechanical properties may occur. Accordingly, the content of Zn is limited to a range of 3 - 10 wt.%, preferably a range of 5.0 -7.0 wt.%, in accordance with the present invention.
Manganese (Mn): 0.25 - 3.0 wt.%
The maximum solid solubility limit of Mn in an Mg matrix is about 2.2 wt.% at 650 C corresponding to the melting point of Mg. The solid solubility limit of Mn is rapidly decreased at lower temperatures, so that Mn may be present in the form of a-Mn in the Mg matrix. It is generally known that, in magnesium alloys, Mn contributes to an improvement in corrosion resistance when it is added in an amount of 0.1 wt.% or more. Where Mg is added for purposes other than the improvement in corrosion resistance, for example, a strengthening purpose, it may contribute to an improvement in the strength of the alloy product at its content of 0.25 - 2.0 wt.% even though the effect may be varied, depending on the matrix alloy of the alloy product. In particular, the inventors experimentally found that Mn existing in the wrought body serves to refine the precipitation phase of the Mg-Zn binary alloy when the wrought body is subjected to an aging process following a solution heat treatment, thereby providing effects of an improvement in strength and an improvement in elongation. Based on this fact, addition of Mn is made to strengthen the alloy in accordance with the present invention. In this regard, the minimum content of Mn is determined to be 0.25 wt.% in accordance with the present invention. Taking into consideration the maximum solid solubility limit of Mn and the processing method used, it is difficult to add a large amount of Mn using a general melting process. Where Mn is added in an amount of 3.0 wt.% or more, it is mainly present in the form of a-Mn in the matrix. Thus, the surplus amount of Mn does not contribute to the improvement in the properties of the alloy, but results in an undesirable result in terms of the preparation costs. Accordingly, the content of Mn is limited to a range of 0.25 - 3.0 wt.%, preferably a range of 0.75 - 2.0 wt.%, in accordance with the present invention.
Aluminum (Al): 1 - 6 wt.%
In an Mg matrix, AI exhibits a maximum solid solubility limit of about 12 wt.% at 437 C. It is known that an Mg17AI12 precipitation phase is formed in Mg-Al binary alloys in accordance with a heat treatment used. However, addition of Al according to the present invention is irrespective of the formation of such Mg AI12 precipitation phase. In accordance with the present invention, the addition of Al is adapted to improve the major strengthening phase in an Mg-Zn-Mn-based ternary alloy, that is, the Mg-Zn-based acicular precipitation phase. Accordingly, the content of Al is determined within a range involving limited formation of Mg-Al-based precipitation phases, taking into consideration the temperature range of heat treatment, such as the aging temperature range, and the content of Zn added as a major alloying element. The lower content limit of Al is determined to be 1.0 wt.% because the solid solubility limit of Al in the Mg matrix corresponds to about 1 wt.% in the aging temperature range. In order to limit the formation of an Mg-Al-based precipitation phase resulting from a surplus amount of Al exceeding the content of Zn, the upper content limit of Al is determined to be 6.0 wt.%. Meanwhile, where AI is added in an amount considerably more than the added amount of Zn, the possibility of the formation of the Mg-Al-based precipitation phase, that is, the Mg AI12 precipitation phase, is greatly increased. Such a precipitation phase may be coarsely precipitated at grain boundaries, and even interior of grains at a certain heat treatment temperature. Since this precipitation phase is very brittle in terms of strength, it provides a fracture path when the alloy is subjected to a fracture, thereby resulting in a degradation in strength. For this reason, it is desirable for the content of Al to be less than the content of Zn. The inventors experimentally found that Al present in the wrought body exhibits effects of refining acicular precipitation phases even without a solution heat treatment for the wrought body, and providing a remarkable improvement in tensile strength and, in particular, elongation, even though there is a slight decrease in yield strength. It was also found that an increase in the content of Al results in a decrease in yield strength while causing an increase in tensile strength.
Preferably, the high strength magnesium alloy is prepared in the form of a cast ingot by adding a Zn-Mn mother alloy having an Mn content of 10 - 20 wt.% (Zn-10 -20 wt.% Mn mother alloy) to the magnesium melt in a temperature range of 670 to 720 C, and adding Zn or Zn along with Al to the magnesium melt. Alternatively, the high strength magnesium alloy may be 5a preferably prepared in the form of a cast ingot by adding a Zn-10 - 20 wt.%
Mn mother alloy to the magnesium melt in a temperature range of 670 to 720 C, adding an Mg-Si mother alloy to the magnesium melt, and adding Zn or Zn along with Al and/or Ca to the magnesium melt.
Preferably, the cast ingot may be subsequently subjected to a homogenization process in a temperature range of 340 to 410 C for 6 to 12 hours, so that it is formed into a billet. The billet may be mechanically worked after being preheated in a temperature range of 150 to 400 C for 30 minutes to 2 hours.
More preferably, the worked or wrought body may be subjected to a primary aging process in a temperature range of 70 to 100 C for 24 to 96 hours, and then to a secondary aging process in a temperature range of 150 to 180 C for 48 hours or more.
Prior to the double aging process, a solution heat treatment may be carried out in a temperature range of 340 to 410 C for 6 to 12 hours.
Alternatively, a stretching of 3 to 7% may be carried out prior to the double aging process.
The reason why the composition range of each alloying element used in accordance with the present invention is limited to the above mentioned range is as follows:
Zinc (Zn): 3 - 10 wt.%
The maximum solid solubility limit of Zn in an Mg matrix is 6.2 wt.%
at 340 C. Where Zn is added in an amount of 3.0 wt.% or more to Mg matrix, it forms an acicular precipitation phase when it is subjected to a heat treatment, thereby exhibiting an age hardening behavior. Generally, the content of Zn is determined, based on the solid solubility limit thereof. When Zn is added in an amount of about 5.0 - 7.0 wt.% approximating to the maximum solid solubility limit thereof, it is possible to obtain a maximized age hardening behavior. At the Zn content of not more than 3.0 wt.%
corresponding to a minimum solid solubility limit at a general aging temperature, it is difficult to expect a desired precipitation strengthening phenomenon because of an insufficient formation of precipitation phase. On the other hand, when Zn is added. in an amount of 10.0 wt.% or more, precipitation of equilibrium phases at grain boundaries is promoted. As a result, degradation in mechanical properties may occur. Accordingly, the content of Zn is limited to a range of 3 - 10 wt.%, preferably a range of 5.0 -7.0 wt.%, in accordance with the present invention.
Manganese (Mn): 0.25 - 3.0 wt.%
The maximum solid solubility limit of Mn in an Mg matrix is about 2.2 wt.% at 650 C corresponding to the melting point of Mg. The solid solubility limit of Mn is rapidly decreased at lower temperatures, so that Mn may be present in the form of a-Mn in the Mg matrix. It is generally known that, in magnesium alloys, Mn contributes to an improvement in corrosion resistance when it is added in an amount of 0.1 wt.% or more. Where Mg is added for purposes other than the improvement in corrosion resistance, for example, a strengthening purpose, it may contribute to an improvement in the strength of the alloy product at its content of 0.25 - 2.0 wt.% even though the effect may be varied, depending on the matrix alloy of the alloy product. In particular, the inventors experimentally found that Mn existing in the wrought body serves to refine the precipitation phase of the Mg-Zn binary alloy when the wrought body is subjected to an aging process following a solution heat treatment, thereby providing effects of an improvement in strength and an improvement in elongation. Based on this fact, addition of Mn is made to strengthen the alloy in accordance with the present invention. In this regard, the minimum content of Mn is determined to be 0.25 wt.% in accordance with the present invention. Taking into consideration the maximum solid solubility limit of Mn and the processing method used, it is difficult to add a large amount of Mn using a general melting process. Where Mn is added in an amount of 3.0 wt.% or more, it is mainly present in the form of a-Mn in the matrix. Thus, the surplus amount of Mn does not contribute to the improvement in the properties of the alloy, but results in an undesirable result in terms of the preparation costs. Accordingly, the content of Mn is limited to a range of 0.25 - 3.0 wt.%, preferably a range of 0.75 - 2.0 wt.%, in accordance with the present invention.
Aluminum (Al): 1 - 6 wt.%
In an Mg matrix, AI exhibits a maximum solid solubility limit of about 12 wt.% at 437 C. It is known that an Mg17AI12 precipitation phase is formed in Mg-Al binary alloys in accordance with a heat treatment used. However, addition of Al according to the present invention is irrespective of the formation of such Mg AI12 precipitation phase. In accordance with the present invention, the addition of Al is adapted to improve the major strengthening phase in an Mg-Zn-Mn-based ternary alloy, that is, the Mg-Zn-based acicular precipitation phase. Accordingly, the content of Al is determined within a range involving limited formation of Mg-Al-based precipitation phases, taking into consideration the temperature range of heat treatment, such as the aging temperature range, and the content of Zn added as a major alloying element. The lower content limit of Al is determined to be 1.0 wt.% because the solid solubility limit of Al in the Mg matrix corresponds to about 1 wt.% in the aging temperature range. In order to limit the formation of an Mg-Al-based precipitation phase resulting from a surplus amount of Al exceeding the content of Zn, the upper content limit of Al is determined to be 6.0 wt.%. Meanwhile, where AI is added in an amount considerably more than the added amount of Zn, the possibility of the formation of the Mg-Al-based precipitation phase, that is, the Mg AI12 precipitation phase, is greatly increased. Such a precipitation phase may be coarsely precipitated at grain boundaries, and even interior of grains at a certain heat treatment temperature. Since this precipitation phase is very brittle in terms of strength, it provides a fracture path when the alloy is subjected to a fracture, thereby resulting in a degradation in strength. For this reason, it is desirable for the content of Al to be less than the content of Zn. The inventors experimentally found that Al present in the wrought body exhibits effects of refining acicular precipitation phases even without a solution heat treatment for the wrought body, and providing a remarkable improvement in tensile strength and, in particular, elongation, even though there is a slight decrease in yield strength. It was also found that an increase in the content of Al results in a decrease in yield strength while causing an increase in tensile strength.
Silicon (Si): 0.1 - 4.0 wt.%
Because silicon is hardly soluble in an Mg matrix, Si forms an Mg2Si phase, when it is added to the Mg matrix as an alloying element. Such a compound may provide a dispersion strengthening effect when its morphology and/or size is modified in the preparation and heat treatment procedures of the wrought body. The inventors experimentally found that a desired dispersion strengthening effect, as mentioned above, is obtained when Si is added to an Mg-Zn-Al-Mn-based quaternary alloy. When the content of Si is less than 0.1 wt.%, the intended effect of the Si addition can hardly be expected. At an Si content of more than 4.0 wt.%, coarse Mg2Si is formed, thereby resulting in a decrease in elongation. In this regard, the content of Si is limited to a range of 0.1 - 4.0 wt.%, preferably a range of 1.5 - 3.0 wt.%, in accordance with the present invention.
Calcium (Ca): 0.1 - 2.0 wt.%
In the case of an Mg alloy containing Si, it is possible to reduce the grain size of the alloy while improving the morphology of the Mg2Si phase, in accordance with an addition of Ca. To this end, Ca is added to an Si-containing Mg-Zn-Al-Mn alloy in accordance with the present invention. At a Ca content of less than 0.1 wt.%, it is hardly expected to observe the effect of improving the Mg2Si phase. On the other hand, when Ca is added in an amount of 2.0 wt.% exceeding considerably the maximum solid solubility limit of Ca in the Mg matrix at 516 C, that is, 1.34 wt.%, an Mg2Ca precipitation phase is formed. Due to such an Mg2Ca precipitation phase at grain boundaries, beside the effect of improving the Mg2Si phase, there is a degredation in strength. The inventors experimentally found that it is possible to more effectively control the morphology of the Mg2Si phase formed in the Mg-Zn-Al-Mn alloy at a Ca content of 0.3 - 1.0 wt.%, and thus to obtain improvements in strength and elongation. In this regard, the content of Ca is limited to a range of 0.1 - 2.0 wt.%, and preferably 0.3 - 1.0 wt.%, in accordance with the present invention.
In addition, the principal impurities of the Mg alloy should be appropriately limited because they mainly have fatally adverse affects on the corrosion resistance of the alloy, rather than on the mechanical properties of the alloy. Impurities generally known include Fe, Ni, and Cu. Although Cu has adverse effects on corrosion resistance in Mg-Al-based alloys widely used, it has no significant effect in the Mg-Zn-based alloy according to the present invention. In association with the Mg-Zn-based alloy according to the present invention, Fe and Ni are regarded as impurities to be limited in their contents. Typically, these impurities are conservatively limited to a maximum content of 0.005 wt.%. Adverse effects resulting from Fe may be eliminated by an addition of Mn. In the case of Mg alloys, adverse effects of Fe can be minimized by the reduction of the content ratio between Fe and Mn, Fe/Mn, to 0.032 or less. Since Mn is basically added in accordance with the present invention, it is possible to effectively eliminate adverse effects of Fe on corrosion resistance in so far as the content of Fe is less than the conservative limit. In the case of Mg alloys, the remaining impurities including Fe, Ni, and Cu are typically limited to a maximum content of 0.3 wt.% based on the total content thereof.
The preparation method of the present invention, which uses a fluxless melting method in addition to the above mentioned specific composition, has an important feature in that Mn is added in the form of a Zn-Mn mother alloy, taking into consideration of the fact that it is impossible to add Mn to molten magnesium using a method of directly melt Mn into the molten magnesium, because Mn has a very high melting point. At an early stage of the development of magnesium alloys, a method was used in which Mn is added in the form of a flux. Since molten magnesium involves a danger of burning when it is exposed to air, a flux is used which serves to shield the molten magnesium from air, thereby inhibiting the danger of burning. For such a flux, an Mn-added flux is conventionally used to achieve a desired addition of Mn. In this case, Mn is penetrated into the melt by diffusion. In this method, there is a limitation on the amount of Mn added.
Furthermore, it is difficult to control the content of impurities. Thus, this method involves many difficulties associated with the preparation of an intended alloy. After the generalization of fluxless melting methods, in that the surface of the melt is covered by protective gas, addition of Mn was achieved, mainly using a method in which Mn is added in the form of a Mg-Mn mother alloy. In accordance with this method, Mn is added to molten magnesium heated to a temperature, at which Mn melts directly, in a protective gas atmosphere capable of preventing the molten magnesium to burn. Thus, an Mg-Mn mother alloy is separately prepared. In the alloy preparation using the fluxiess melting method, a desired amount of Mn is added using the prepared Mg-Mn mother alloy. However, this method requires an expensive melting device configured to control the given atmosphere. Furthermore, a large amount of magnesium may be lost during the preparation of the mother alloy, because magnesium exhibits a high vapor pressure at a high temperature. Thus, the method using the Mg-Mn mother alloy involves an increase in the processing costs. After a series of experiments, the inventors found that, in the alloy preparation using the fluxless melting method, addition of Mn can be effectively achieved using a method in which Mn is added to molten magnesium in the form of a low-melting-point Zn-Mn mother alloy. In accordance with this method, it is possible to eliminate the burning possibility of the magnesium melt or a great loss of material. Accordingly, an economical preparation of magnesium alloys is possible. It is also possible to easily and conveniently control the content of impurities.
Preferably, the temperature of the magnesium melt, to which the Zn-Mn mother alloy is added, is limited to a range of 670 to 720 C, taking into consideration of the fact that although the melting point of magnesium is about 650 C, the magnesium melt secures a sufficient fluidity at a temperature of at least 670 C, and the fact that there is an increased possibility of burning at a temperature of the magnesium melt exceeding 720 C. The Zn-Mn mother alloy preferably has a Zn-10 - 20 wt.% Mn composition having an Mn content of 10 to 20 wt.% so that it is sufficiently melted in the above mentioned temperature range of the magnesium melt.
More preferably, a stirring process is carried out during the addition of the Zn-Mn mother alloy to the magnesium melt.
Si is added in the form of an Mg-Si mother alloy. Preferably, the temperature, at which the addition of Si is carried out, is limited to a range of 700 to 720 C, taking into consideration of a high melting point of the mother alloy and a desired inhibition of burning at the surface of the magnesium melt. In this case, it is more preferable to conduct a stirring process during the addition of the mother alloy.
For a shortage of zinc, Zn is added alone or along with Al. At this time, Ca may be selectively added. Preferably, the addition of Zn is carried out after a furnace cooling process conducted for the magnesium melt, in order to reduce the loss of Zn exhibiting a high vapor pressure at the alloy preparation temperature used. The furnace cooling may be carried out to a temperature of about 670 C, taking into consideration of the fluidity of the magnesium melt. It is more preferable to conduct a stirring process during the addition of those elements.
The resultant Mg melt is then cast to form an ingot. Preferably, the casting is conducted after the Mg melt is furnace-cooled to a temperature of 660 to 670 C, in order to inhibit the generation of heat from the Mg melt as much as possible.
Preferably, the cast alloy ingot prepared in accordance with the above mentioned method is subjected to a homogenization treatment in order to eliminate segregation of alloying elements possibly generated during the casting process, and non-uniformity of the wrought body, in terms of properties, resulting from the segregation. The homogenization treatment is carried out at a temperature of 340 to 410 C for 6 to 12 hours, taking into consideration of conditions capable of allowing precipitation phases resulting from the major alloying element, that is, Zn, to be sufficiently dissolved, and a desired thermal stability of the alloy.
Thereafter, the cast ingot is formed into a billet to be extruded.
Preferably, the billet is preheated at a temperature of 150 to 400 C for 30 minutes to 2 hours, and then subjected to mechanical working processes including extrusion, rolling, forging, swaging, and drawing, in the same temperature range. Generally, magnesium alloys do not have a desired workability at ambient temperature. In order to obtain a sound wrought body, accordingly, the magnesium alloy is subjected to hot mechanical working processes. In accordance with the present invention, the working temperature is experimentally determined to be within a range capable of securing the soundness of the wrought body.
Preferably, the wrought body is primarily aged at a temperature of 70 to 100 C for 24 to 96 hours, and secondarily aged, just after the primary aging process, at a temperature of 150 to 180 C for 48 hours or more. Such a double aging process is adapted to maximize the effect of the precipitation phase contributing to an improvement in strength by conducting the primary aging process at a temperature not higher than the G. P. zone solvus temperature of the predominant precipitation phase of Mg-Zn-based alloys, that is, the Pl' phase, and then conducting the secondary aging process at a temperature higher than the primary aging temperature. In accordance with the present invention, therefore, the primary aging temperature is limited to a range of 70 to 100 C slightly lower than the generally-known G. P. zone solvus temperature of the Pl' phase, and the primary aging time is determined to be a period of time enough to expect an improvement in hardness to a desired level. Also, the second aging temperature is limited to a range of 150 to 180 C. At a second aging temperature of less than 150 C, there is a problem associated with the execution of the aging process because a lot of time is required until a maximum hardness is obtained. At a second aging temperature exceeding 180 C, the maximum hardness obtained cannot reach a desired level even though it is rapidly obtained.
More preferably, the wrought body is subjected, prior to the double aging process, to a solution heat treatment for 6 to 12 hours at a temperature of 340 to 410 C corresponding to a temperature range, in which precipitation phases possibly generated during the working process can be present in the form of solid solutions, in order to maximize the effect of the precipitation phase contributing to an improvement in strength. The temperature range and period of the solution heat treatment are determined, based on the phase diagram of the Mg-Zn binary system, taking into consideration of conditions capable of allowing precipitation phases resulting from the major alloying element, that is, Zn, to be sufficiently dissolved, and a desired thermal stability of the alloy.
Meanwhile, it is more preferable to conduct a stretching process, prior to the double-aging process. Generally, the amount of stretching in a working process involving a heat treatment to strengthen a subject alloy is limited to a range from an elastic limit to a maximum strength limit, based on a strain measured by a tensile test for the alloy conducted before the heat treatment. In accordance with the present invention, therefore, the amount of stretching is limited to a range of 3 to 7 %.
In accordance with the present invention, it is possible to prepare an inexpensive high strength magnesium alloy having an improved elongation compared to the strength, compared to existing commercial wrought alloys.
That is, the magnesium alloy of the present invention can exhibit an elongation improved by two times or more over the maximum strength one of the existing commercial extruded alloys described in Table 1, that is, a ZC71 alloy, while maintaining a strength level similar to that of the ZC71 alloy.
In accordance with the present invention, it is also possible to prepare a high strength alloy without using expensive alloying elements having a difficulty in handling, for example, a radioactive element such as Th, or alloying elements involving a difficulty in addition in association with the preparation process, for example, Zr. In accordance with the present invention, it is possible to reduce the loss of materials, and thus to reduce the preparation cost, because Mn is added in the form of a Zn-Mn mother alloy, compared to conventional methods in which Mn is added in the form of an Mg-Mn mother alloy.
BRIEF DESCRIPTION OF THE DRAWINGS
The above objects, and other features and advantages of the present invention will become more apparent after a reading of the following detailed description when taken in conjunction with the drawings, in which:
Fig. la is a photomicrograph of an Mg-Zn-based binary alloy extruded product (Z6);
Figs. lb and lc are photomicrographs of respective extruded products of an Mg-Zn-based alloy added with Mn and an Mg-Zn-based alloy added with Al and Mn (ZM61 and ZAM621);
Figs. 1d and le are photomicrographs of respective extruded products of an Mg-Zn-based alloy added with Al, Mn and Si, and an Mg-Zn-based alloy added with Al, Mn, Si and Ca (ZAM631 + 2.5Si, and ZAM631 +2.5Si + O.4Ca);
Fig. 2 is a graph depicting the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during an aging process;
Fig. 3 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during a double aging process;
Fig. 4 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process following a solution heat treatment with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited when the same treatment is conducted;
Fig. 5 is a graph depicting respective age hardening behaviors of the Al + Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si + O.4Ca) exhibited during a double aging process;
Fig. 6 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature;
Fig. 7 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature;
Fig. 8 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature after the double aging process;
Fig. 9 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature after a double aging process;
Fig. 10 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a solution heat treatment with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted; and Fig. 11 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a 5% stretching process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted.
BEST MODE FOR CARRYING OUT THE INVENTION
Now, the high strength magnesium alloy according to the present invention will be described in more detail, with reference to the following examples.
Examples 1 to 11 In accordance with the present invention, alloy cast products were prepared which have rated compositions described in the following Table 2, respectively. In the preparation of each alloy cast product, melting of a raw alloy and alloying elements was achieved using a fluxless melting method in which a mixture gas of C02 + 0.5 % SF6 is sprayed over the surface of the melt at a flow rate of 2 f/minute. A steel crucible was also used in the melting process. Mn was added in the form of a Zn-15 wt.% Mn mother alloy to the melt at a temperature of 700 C. Thereafter, the melt was stirred for 5 minutes using a stirrer, and then cooled to 670 C in a furnace cooling fashion. The cooled melt was added with Zn alone or along with Al, and then stirred for 2 minutes. In the case involving an addition of Si, this Si was added in the form of an Mg-10 wt.% Si mother alloy to the melt. In this case, the resultant melt was then stirred at 720 C for 10 minutes. After the completion of the stirring process, the melt was cooled to 670 C in a furnace cooling fashion. The cooled melt was added with Zn alone or along with Al and/or Ca, and then stirred for 2 minutes. Thereafter, the melt was furnace-cooled to 660 C. Finally, the crucible was completely dipped in water maintained at ambient temperature. Thus, an alloy cast product was prepared.
Table 2 Rated Compositions of Alloys of Examples Rated Composition (wt%) Example Alloys Zn Al Mn Si Ca Mg 1 Z6 6 - - - - bal.
2 ZM60 6 - 0.5 - - bal.
3 ZM61 6 - 1 - - bal.
4 ZM62 6 - 1.5 - - bal.
5 ZAM611 6 1 1 - - bal.
6 ZAM621 6 2 1 - - bal.
7 ZAM631 6 3 1 - - bal.
8 ZAM641 6 4 1 - - bal.
9 ZAM661 6 6 1 - - bal.
Because silicon is hardly soluble in an Mg matrix, Si forms an Mg2Si phase, when it is added to the Mg matrix as an alloying element. Such a compound may provide a dispersion strengthening effect when its morphology and/or size is modified in the preparation and heat treatment procedures of the wrought body. The inventors experimentally found that a desired dispersion strengthening effect, as mentioned above, is obtained when Si is added to an Mg-Zn-Al-Mn-based quaternary alloy. When the content of Si is less than 0.1 wt.%, the intended effect of the Si addition can hardly be expected. At an Si content of more than 4.0 wt.%, coarse Mg2Si is formed, thereby resulting in a decrease in elongation. In this regard, the content of Si is limited to a range of 0.1 - 4.0 wt.%, preferably a range of 1.5 - 3.0 wt.%, in accordance with the present invention.
Calcium (Ca): 0.1 - 2.0 wt.%
In the case of an Mg alloy containing Si, it is possible to reduce the grain size of the alloy while improving the morphology of the Mg2Si phase, in accordance with an addition of Ca. To this end, Ca is added to an Si-containing Mg-Zn-Al-Mn alloy in accordance with the present invention. At a Ca content of less than 0.1 wt.%, it is hardly expected to observe the effect of improving the Mg2Si phase. On the other hand, when Ca is added in an amount of 2.0 wt.% exceeding considerably the maximum solid solubility limit of Ca in the Mg matrix at 516 C, that is, 1.34 wt.%, an Mg2Ca precipitation phase is formed. Due to such an Mg2Ca precipitation phase at grain boundaries, beside the effect of improving the Mg2Si phase, there is a degredation in strength. The inventors experimentally found that it is possible to more effectively control the morphology of the Mg2Si phase formed in the Mg-Zn-Al-Mn alloy at a Ca content of 0.3 - 1.0 wt.%, and thus to obtain improvements in strength and elongation. In this regard, the content of Ca is limited to a range of 0.1 - 2.0 wt.%, and preferably 0.3 - 1.0 wt.%, in accordance with the present invention.
In addition, the principal impurities of the Mg alloy should be appropriately limited because they mainly have fatally adverse affects on the corrosion resistance of the alloy, rather than on the mechanical properties of the alloy. Impurities generally known include Fe, Ni, and Cu. Although Cu has adverse effects on corrosion resistance in Mg-Al-based alloys widely used, it has no significant effect in the Mg-Zn-based alloy according to the present invention. In association with the Mg-Zn-based alloy according to the present invention, Fe and Ni are regarded as impurities to be limited in their contents. Typically, these impurities are conservatively limited to a maximum content of 0.005 wt.%. Adverse effects resulting from Fe may be eliminated by an addition of Mn. In the case of Mg alloys, adverse effects of Fe can be minimized by the reduction of the content ratio between Fe and Mn, Fe/Mn, to 0.032 or less. Since Mn is basically added in accordance with the present invention, it is possible to effectively eliminate adverse effects of Fe on corrosion resistance in so far as the content of Fe is less than the conservative limit. In the case of Mg alloys, the remaining impurities including Fe, Ni, and Cu are typically limited to a maximum content of 0.3 wt.% based on the total content thereof.
The preparation method of the present invention, which uses a fluxless melting method in addition to the above mentioned specific composition, has an important feature in that Mn is added in the form of a Zn-Mn mother alloy, taking into consideration of the fact that it is impossible to add Mn to molten magnesium using a method of directly melt Mn into the molten magnesium, because Mn has a very high melting point. At an early stage of the development of magnesium alloys, a method was used in which Mn is added in the form of a flux. Since molten magnesium involves a danger of burning when it is exposed to air, a flux is used which serves to shield the molten magnesium from air, thereby inhibiting the danger of burning. For such a flux, an Mn-added flux is conventionally used to achieve a desired addition of Mn. In this case, Mn is penetrated into the melt by diffusion. In this method, there is a limitation on the amount of Mn added.
Furthermore, it is difficult to control the content of impurities. Thus, this method involves many difficulties associated with the preparation of an intended alloy. After the generalization of fluxless melting methods, in that the surface of the melt is covered by protective gas, addition of Mn was achieved, mainly using a method in which Mn is added in the form of a Mg-Mn mother alloy. In accordance with this method, Mn is added to molten magnesium heated to a temperature, at which Mn melts directly, in a protective gas atmosphere capable of preventing the molten magnesium to burn. Thus, an Mg-Mn mother alloy is separately prepared. In the alloy preparation using the fluxiess melting method, a desired amount of Mn is added using the prepared Mg-Mn mother alloy. However, this method requires an expensive melting device configured to control the given atmosphere. Furthermore, a large amount of magnesium may be lost during the preparation of the mother alloy, because magnesium exhibits a high vapor pressure at a high temperature. Thus, the method using the Mg-Mn mother alloy involves an increase in the processing costs. After a series of experiments, the inventors found that, in the alloy preparation using the fluxless melting method, addition of Mn can be effectively achieved using a method in which Mn is added to molten magnesium in the form of a low-melting-point Zn-Mn mother alloy. In accordance with this method, it is possible to eliminate the burning possibility of the magnesium melt or a great loss of material. Accordingly, an economical preparation of magnesium alloys is possible. It is also possible to easily and conveniently control the content of impurities.
Preferably, the temperature of the magnesium melt, to which the Zn-Mn mother alloy is added, is limited to a range of 670 to 720 C, taking into consideration of the fact that although the melting point of magnesium is about 650 C, the magnesium melt secures a sufficient fluidity at a temperature of at least 670 C, and the fact that there is an increased possibility of burning at a temperature of the magnesium melt exceeding 720 C. The Zn-Mn mother alloy preferably has a Zn-10 - 20 wt.% Mn composition having an Mn content of 10 to 20 wt.% so that it is sufficiently melted in the above mentioned temperature range of the magnesium melt.
More preferably, a stirring process is carried out during the addition of the Zn-Mn mother alloy to the magnesium melt.
Si is added in the form of an Mg-Si mother alloy. Preferably, the temperature, at which the addition of Si is carried out, is limited to a range of 700 to 720 C, taking into consideration of a high melting point of the mother alloy and a desired inhibition of burning at the surface of the magnesium melt. In this case, it is more preferable to conduct a stirring process during the addition of the mother alloy.
For a shortage of zinc, Zn is added alone or along with Al. At this time, Ca may be selectively added. Preferably, the addition of Zn is carried out after a furnace cooling process conducted for the magnesium melt, in order to reduce the loss of Zn exhibiting a high vapor pressure at the alloy preparation temperature used. The furnace cooling may be carried out to a temperature of about 670 C, taking into consideration of the fluidity of the magnesium melt. It is more preferable to conduct a stirring process during the addition of those elements.
The resultant Mg melt is then cast to form an ingot. Preferably, the casting is conducted after the Mg melt is furnace-cooled to a temperature of 660 to 670 C, in order to inhibit the generation of heat from the Mg melt as much as possible.
Preferably, the cast alloy ingot prepared in accordance with the above mentioned method is subjected to a homogenization treatment in order to eliminate segregation of alloying elements possibly generated during the casting process, and non-uniformity of the wrought body, in terms of properties, resulting from the segregation. The homogenization treatment is carried out at a temperature of 340 to 410 C for 6 to 12 hours, taking into consideration of conditions capable of allowing precipitation phases resulting from the major alloying element, that is, Zn, to be sufficiently dissolved, and a desired thermal stability of the alloy.
Thereafter, the cast ingot is formed into a billet to be extruded.
Preferably, the billet is preheated at a temperature of 150 to 400 C for 30 minutes to 2 hours, and then subjected to mechanical working processes including extrusion, rolling, forging, swaging, and drawing, in the same temperature range. Generally, magnesium alloys do not have a desired workability at ambient temperature. In order to obtain a sound wrought body, accordingly, the magnesium alloy is subjected to hot mechanical working processes. In accordance with the present invention, the working temperature is experimentally determined to be within a range capable of securing the soundness of the wrought body.
Preferably, the wrought body is primarily aged at a temperature of 70 to 100 C for 24 to 96 hours, and secondarily aged, just after the primary aging process, at a temperature of 150 to 180 C for 48 hours or more. Such a double aging process is adapted to maximize the effect of the precipitation phase contributing to an improvement in strength by conducting the primary aging process at a temperature not higher than the G. P. zone solvus temperature of the predominant precipitation phase of Mg-Zn-based alloys, that is, the Pl' phase, and then conducting the secondary aging process at a temperature higher than the primary aging temperature. In accordance with the present invention, therefore, the primary aging temperature is limited to a range of 70 to 100 C slightly lower than the generally-known G. P. zone solvus temperature of the Pl' phase, and the primary aging time is determined to be a period of time enough to expect an improvement in hardness to a desired level. Also, the second aging temperature is limited to a range of 150 to 180 C. At a second aging temperature of less than 150 C, there is a problem associated with the execution of the aging process because a lot of time is required until a maximum hardness is obtained. At a second aging temperature exceeding 180 C, the maximum hardness obtained cannot reach a desired level even though it is rapidly obtained.
More preferably, the wrought body is subjected, prior to the double aging process, to a solution heat treatment for 6 to 12 hours at a temperature of 340 to 410 C corresponding to a temperature range, in which precipitation phases possibly generated during the working process can be present in the form of solid solutions, in order to maximize the effect of the precipitation phase contributing to an improvement in strength. The temperature range and period of the solution heat treatment are determined, based on the phase diagram of the Mg-Zn binary system, taking into consideration of conditions capable of allowing precipitation phases resulting from the major alloying element, that is, Zn, to be sufficiently dissolved, and a desired thermal stability of the alloy.
Meanwhile, it is more preferable to conduct a stretching process, prior to the double-aging process. Generally, the amount of stretching in a working process involving a heat treatment to strengthen a subject alloy is limited to a range from an elastic limit to a maximum strength limit, based on a strain measured by a tensile test for the alloy conducted before the heat treatment. In accordance with the present invention, therefore, the amount of stretching is limited to a range of 3 to 7 %.
In accordance with the present invention, it is possible to prepare an inexpensive high strength magnesium alloy having an improved elongation compared to the strength, compared to existing commercial wrought alloys.
That is, the magnesium alloy of the present invention can exhibit an elongation improved by two times or more over the maximum strength one of the existing commercial extruded alloys described in Table 1, that is, a ZC71 alloy, while maintaining a strength level similar to that of the ZC71 alloy.
In accordance with the present invention, it is also possible to prepare a high strength alloy without using expensive alloying elements having a difficulty in handling, for example, a radioactive element such as Th, or alloying elements involving a difficulty in addition in association with the preparation process, for example, Zr. In accordance with the present invention, it is possible to reduce the loss of materials, and thus to reduce the preparation cost, because Mn is added in the form of a Zn-Mn mother alloy, compared to conventional methods in which Mn is added in the form of an Mg-Mn mother alloy.
BRIEF DESCRIPTION OF THE DRAWINGS
The above objects, and other features and advantages of the present invention will become more apparent after a reading of the following detailed description when taken in conjunction with the drawings, in which:
Fig. la is a photomicrograph of an Mg-Zn-based binary alloy extruded product (Z6);
Figs. lb and lc are photomicrographs of respective extruded products of an Mg-Zn-based alloy added with Mn and an Mg-Zn-based alloy added with Al and Mn (ZM61 and ZAM621);
Figs. 1d and le are photomicrographs of respective extruded products of an Mg-Zn-based alloy added with Al, Mn and Si, and an Mg-Zn-based alloy added with Al, Mn, Si and Ca (ZAM631 + 2.5Si, and ZAM631 +2.5Si + O.4Ca);
Fig. 2 is a graph depicting the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during an aging process;
Fig. 3 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during a double aging process;
Fig. 4 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process following a solution heat treatment with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited when the same treatment is conducted;
Fig. 5 is a graph depicting respective age hardening behaviors of the Al + Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si + O.4Ca) exhibited during a double aging process;
Fig. 6 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature;
Fig. 7 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature;
Fig. 8 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature after the double aging process;
Fig. 9 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature after a double aging process;
Fig. 10 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a solution heat treatment with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted; and Fig. 11 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a 5% stretching process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted.
BEST MODE FOR CARRYING OUT THE INVENTION
Now, the high strength magnesium alloy according to the present invention will be described in more detail, with reference to the following examples.
Examples 1 to 11 In accordance with the present invention, alloy cast products were prepared which have rated compositions described in the following Table 2, respectively. In the preparation of each alloy cast product, melting of a raw alloy and alloying elements was achieved using a fluxless melting method in which a mixture gas of C02 + 0.5 % SF6 is sprayed over the surface of the melt at a flow rate of 2 f/minute. A steel crucible was also used in the melting process. Mn was added in the form of a Zn-15 wt.% Mn mother alloy to the melt at a temperature of 700 C. Thereafter, the melt was stirred for 5 minutes using a stirrer, and then cooled to 670 C in a furnace cooling fashion. The cooled melt was added with Zn alone or along with Al, and then stirred for 2 minutes. In the case involving an addition of Si, this Si was added in the form of an Mg-10 wt.% Si mother alloy to the melt. In this case, the resultant melt was then stirred at 720 C for 10 minutes. After the completion of the stirring process, the melt was cooled to 670 C in a furnace cooling fashion. The cooled melt was added with Zn alone or along with Al and/or Ca, and then stirred for 2 minutes. Thereafter, the melt was furnace-cooled to 660 C. Finally, the crucible was completely dipped in water maintained at ambient temperature. Thus, an alloy cast product was prepared.
Table 2 Rated Compositions of Alloys of Examples Rated Composition (wt%) Example Alloys Zn Al Mn Si Ca Mg 1 Z6 6 - - - - bal.
2 ZM60 6 - 0.5 - - bal.
3 ZM61 6 - 1 - - bal.
4 ZM62 6 - 1.5 - - bal.
5 ZAM611 6 1 1 - - bal.
6 ZAM621 6 2 1 - - bal.
7 ZAM631 6 3 1 - - bal.
8 ZAM641 6 4 1 - - bal.
9 ZAM661 6 6 1 - - bal.
10 ZAM631+2.5Si 6 3 1 2.5 - bal.
11 ZAM631+2.5Si+0.4Ca 6 3 1 2.5 0.4 bal.
In order to control the microstructure of each alloy cast product prepared as above, the alloy cast product was subjected to a homogenization process at a temperature of 340 to 410 C for 12 hours. The alloy cast product was then formed into a billet, which was, in turn, preheated at a temperature of 320 to 360 C for 30 minutes. The billet was then extruded by an extrusion machine, in which the temperature of the billet container and die was set to a temperature of 320 to 360 C. Thus, an extruded alloy product was prepared.
Figs. la, 1 b, and 1 c are photomicrographs of Z6, ZM61, and ZAM621 alloy extruded products prepared as mentioned above, respectively. Figs. 1 d and le are photomicrographs of ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloy extruded products, respectively. Referring to Figs. la to le, the Z6 alloy, which is a conventional magnesium alloy, has a grain size of about 22 pm, whereas the ZM61 and ZAM621 alloys according to the present invention have grain sizes of about 12 pm and about 8 pm, respectively. The ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloys have grain sizes of about 12 pm and about 6 pm, respectively.
Accordingly, it can be found that when Mn is added in an amount of 1 wt.% to the conventional Mg-Zn alloy, the microstructural grain size of the alloy is reduced by about 1/2. It can also be found that the conventional Mg-Zn alloy is added with 1 wt.% Mn and 2 wt.% Al, its grain size is reduced by about 1/3. In the case of the ZAM631 + 2.5Si + 0.4Ca alloy, that is, where 0.4 wt.% Ca is added to the ZAM631 + 2.5Si alloy, it has a grain size of about 6 pm corresponding to about 2/3 the grain size of the ZAM621 alloy.
Consequently, it can be found that in the alloys according to the present invention, a grain size reduction is obtained by about 2/3 for those added with Mn and Al and about 3/4 for those added with Si and Ca along with Mn and Al.
Fig. 2 is a graph depicting the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during an aging process.
For the Z6 alloy extruded product, a single aging process and a double aging process were carried out to obtain maximized improvements in hardness and strength, respectively. The Z6 alloy extruded product was primarily aged at 90 C for 48 hours, and then secondarily aged at 180 C for 384 hours under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. The measured age hardening behavior variation is depicted in Fig. 2. Referring to Fig. 2, it can be found that the alloy subjected to the double aging treatment exhibits an increase in maximum hardness and a reduction in the time taken to obtain a maximum hardness, compared to the alloy subjected to the single aging treatment.
Fig. 3 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during the double aging process. Each of the Z6, ZM61, and ZAM621 alloy extruded products was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 384 hours under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. The measured age hardening behavior variation is depicted in Fig. 3. Referring to Fig. 3, it can be found that the ZAM621 alloy prepared by adding 2 wt.% Al and 1 wt.% Mn to the Z6 alloy exhibits an improvement in hardness by about 35 % in an extruded state, and by about 20 % in a state subjected to a maximum double age treatment, compared to the Z6 alloy. However, the ZM61 alloy prepared by adding only Mn to the Z6 alloy exhibited no or little hardening effect during the aging period even though it exhibited a high hardness in an extruded state. Also, the maximum hardness of the ZM61 alloy was lower than that of the Z6 alloy.
Fig. 4 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process following a solution heat treatment with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited when the same treatment is conducted.
Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a solution heat treatment for 12 hours under the condition in which it was maintained at a temperature of 380 to 410 C. Following the solution heat treatment, each alloy was subjected to a double aging treatment. The age hardening behavior of each alloy resulting from the double aging treatment is depicted in Fig. 4. Referring to Fig. 4, it can be found that where the Z6 alloy is added with Mn alone or along with Al, it substantially exhibits an improvement in hardness during the age hardening process. The addition of the alloying element or elements contributed to an improvement in hardness by 10 % or more, based on the maximum hardness. In particular, the age hardening behavior of the ZM61 alloy prepared only by an addition of Mn was considerably different from the age hardening behavior exhibited under the heat treatment condition involving no solution heat treatment followed by the double aging treatment in that a remarkable improvement in hardness was obtained. The maximum hardness of the ZM61 alloy was similar to the ZAM621 alloy prepared by adding both AI and Mn.
Fig. 5 is a graph depicting respective age hardening behaviors of the Al + Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si + 0.4Ca) exhibited during a double aging process. After being extruded, each of the ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloys was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for a given period of time under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. Referring to Figs. 3 and 5, it can be found that the addition of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy results in an improvement in hardness by about 12 % and a considerable reduction in the time taken to obtain a maximum hardness under the condition in which the alloy is subjected to a maximum double aging treatment at 150 C.
Table 3 Alloys Yield Strength Tensile Elongation (Mpa) strength (%) (Mpa) ZAM631+2.5Si 143 271 10 ZAM631 +2.5Si+0.4Ca 150 371 16 Fig. 6 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature. Referring to Fig. 6, it can be found that the addition of Mn alone or along with Al to the Z6 alloy results in a considerable increase in the yield strength and maximum tensile strength exhibited in an extruded state. In a wrought state using an extrusion process, the alloy exhibited a superior elongation of 25 % or more. Detailed results are described in the above Table 3.
Fig. 7 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature. Referring to Fig. 7, it can be found that the addition of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy results in an increase in the maximum tensile strength exhibited in an extruded state. In a wrought state using an extrusion process, the alloy exhibited a superior elongation of 16 % or more. Detailed results are described in the above Table 3.
Table 4 Tensile Alloys Yield Strength Strength Elongation (Mpa) (Mpa) (%) ZAM631 +2.5Si 265 317 7 ZAM631 +2.5Si+0.4Ca 234 398 12 Fig. 8 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature after the double aging process. Each of the Z6, ZM61, and ZAM621 alloys was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 8. Referring to Fig. 8, it can be found that the case involving the double aging treatment exhibits an increase in the yield strength and maximum tensile strength of each alloy, compared to the case involving no double aging treatment. It can also be found that both cases exhibit similar elongations, respectively. The tensile properties of each alloy, exhibited after a tensile test conducted following the double aging treatment, are described in the above Table 4.
Fig. 9 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature after a double aging process. Each extruded alloy product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 24 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 9. Referring to Fig. 9, it can be found that the extruded products of the ZAM631 + 2.5Si and ZAM631 +2.5Si + 0.4Ca alloys respectively prepared by adding 2.5 wt.% Si and a mixture of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy obtain effects of considerably increased yield strength and maximum tensile strength, compared to those subjected to no double aging treatment. Detailed results are described in the above Table 4.
Referring to Table 4, it can be found that the ZM61 alloy prepared by adding Mn to the Z6 alloy exhibits a slight increase in tensile properties by virtue of the double aging treatment, compared to the Z6 alloy. The ZAM621 alloy prepared by adding both AI and Mn to the Z6 alloy exhibited a superior strength over the Z6 alloy, by virtue of the double aging treatment. In particular, the ZAM621 alloy exhibits a remarkable increase in maximum tensile strength. All alloys exhibit a superior elongation even after the double aging treatment.
Fig. 10 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature- after a double aging process following a solution heat treatment with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted. Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a solution heat treatment for 12 hours at a temperature of 380 to 410 C. Following the solution heat treatment, each alloy extruded product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 10. Referring to Fig. 10, it can be found that when the solution heat treatment is carried out prior to the double aging treatment, the ZM61 alloy exhibits a considerable increase in yield strength and maximum tensile strength while exhibiting an elongation similar to that of the Z6 alloy.
Although the ZAM621 alloy exhibited a reduced yield strength, compared to the ZM61 alloy, it was similar to the ZM61 alloy in terms of the maximum tensile strength. In particular, the ZMA621 alloy exhibited a considerable increase in elongation. The tensile properties of each alloy exhibited at ambient temperature after the double aging treatment following the solution heat treatment are described in the following Table 5.
Table 5 Tensile Alloys Yield Strength Strength Elongation (Mpa) (Mpa) (%) Meanwhile, a comparison was made between the case, in which each alloy extruded product is subjected to a double aging process just after the extrusion thereof, and the case in which the alloy extruded product is subjected to the double aging process following a solution heat treatment.
Based on the results of the comparison, it can be found that the ZM61 alloy exhibits a considerable increase in strength by virtue of the solution heat treatment followed by the double aging process whereas the Z6 and ZAM621 alloys exhibit a slight strength increase under the same treatment conditions. It can also be found that the Z6 and ZM61 alloys exhibit a considerable reduction in elongation by virtue of the solution heat treatment followed by the double aging process whereas the ZAM621 alloy exhibit a similar elongation even after the solution heat treatment.
Fig. 11 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a 5% stretching process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted. Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a 5 % stretching process.
Following the stretching process, each extruded alloy product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours.
The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 11. Referring to Fig. 11, it can be found that the ZAM621 alloy exhibits an improvement in strength, compared to the case involving no stretching process. The ZAM621 alloy also exhibited an elongation of 20 % or more. All alloys substantially exhibited an improvement in strength by virtue of the stretching process followed by the double aging process. In particular, when the ZAM621 alloy was subjected only to the stretching process while being subjected to the solution heat treatment, prior to the double aging process, in order to obtain an improvement in strength, it exhibited a strength similar to that of the ZM61 alloy while exhibiting a great increase in elongation. The tensile properties of each alloy exhibited at ambient temperature after the double aging process following the 5 %
stretching process are described in the following Table 6.
Table 6 Alloys Yield Strength Tensile Strength. Elongation (Mpa) (Mpa) (%) INDUSTRIAL APPLICABILITY
As apparent from the above description, the present invention provides a magnesium alloy having improvements in the hardness and strength at ambient temperature and an enhancement in elongation by adding Mn alone or along with Al to an Mg-Zn binary alloy while simultaneously adding Si alone or along with Ca to the Mg-Zn binary alloy to prepare a wrought body having a reduced grain size, and conducting a heat treatment or a working process involving a heat treatment for the wrought body.
Although the preferred embodiments of the invention have been disclosed for illustrative purposes, those skilled in the art will appreciate that various modifications, additions and substitutions are possible, without departing from the scope and spirit of the invention as disclosed in the accompanying claims.
In order to control the microstructure of each alloy cast product prepared as above, the alloy cast product was subjected to a homogenization process at a temperature of 340 to 410 C for 12 hours. The alloy cast product was then formed into a billet, which was, in turn, preheated at a temperature of 320 to 360 C for 30 minutes. The billet was then extruded by an extrusion machine, in which the temperature of the billet container and die was set to a temperature of 320 to 360 C. Thus, an extruded alloy product was prepared.
Figs. la, 1 b, and 1 c are photomicrographs of Z6, ZM61, and ZAM621 alloy extruded products prepared as mentioned above, respectively. Figs. 1 d and le are photomicrographs of ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloy extruded products, respectively. Referring to Figs. la to le, the Z6 alloy, which is a conventional magnesium alloy, has a grain size of about 22 pm, whereas the ZM61 and ZAM621 alloys according to the present invention have grain sizes of about 12 pm and about 8 pm, respectively. The ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloys have grain sizes of about 12 pm and about 6 pm, respectively.
Accordingly, it can be found that when Mn is added in an amount of 1 wt.% to the conventional Mg-Zn alloy, the microstructural grain size of the alloy is reduced by about 1/2. It can also be found that the conventional Mg-Zn alloy is added with 1 wt.% Mn and 2 wt.% Al, its grain size is reduced by about 1/3. In the case of the ZAM631 + 2.5Si + 0.4Ca alloy, that is, where 0.4 wt.% Ca is added to the ZAM631 + 2.5Si alloy, it has a grain size of about 6 pm corresponding to about 2/3 the grain size of the ZAM621 alloy.
Consequently, it can be found that in the alloys according to the present invention, a grain size reduction is obtained by about 2/3 for those added with Mn and Al and about 3/4 for those added with Si and Ca along with Mn and Al.
Fig. 2 is a graph depicting the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during an aging process.
For the Z6 alloy extruded product, a single aging process and a double aging process were carried out to obtain maximized improvements in hardness and strength, respectively. The Z6 alloy extruded product was primarily aged at 90 C for 48 hours, and then secondarily aged at 180 C for 384 hours under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. The measured age hardening behavior variation is depicted in Fig. 2. Referring to Fig. 2, it can be found that the alloy subjected to the double aging treatment exhibits an increase in maximum hardness and a reduction in the time taken to obtain a maximum hardness, compared to the alloy subjected to the single aging treatment.
Fig. 3 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited during the double aging process. Each of the Z6, ZM61, and ZAM621 alloy extruded products was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 384 hours under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. The measured age hardening behavior variation is depicted in Fig. 3. Referring to Fig. 3, it can be found that the ZAM621 alloy prepared by adding 2 wt.% Al and 1 wt.% Mn to the Z6 alloy exhibits an improvement in hardness by about 35 % in an extruded state, and by about 20 % in a state subjected to a maximum double age treatment, compared to the Z6 alloy. However, the ZM61 alloy prepared by adding only Mn to the Z6 alloy exhibited no or little hardening effect during the aging period even though it exhibited a high hardness in an extruded state. Also, the maximum hardness of the ZM61 alloy was lower than that of the Z6 alloy.
Fig. 4 is a graph depicting a comparison of respective age hardening behaviors of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited during a double aging process following a solution heat treatment with the age hardening behavior of the Mg-Zn-based binary alloy extruded product (Z6) exhibited when the same treatment is conducted.
Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a solution heat treatment for 12 hours under the condition in which it was maintained at a temperature of 380 to 410 C. Following the solution heat treatment, each alloy was subjected to a double aging treatment. The age hardening behavior of each alloy resulting from the double aging treatment is depicted in Fig. 4. Referring to Fig. 4, it can be found that where the Z6 alloy is added with Mn alone or along with Al, it substantially exhibits an improvement in hardness during the age hardening process. The addition of the alloying element or elements contributed to an improvement in hardness by 10 % or more, based on the maximum hardness. In particular, the age hardening behavior of the ZM61 alloy prepared only by an addition of Mn was considerably different from the age hardening behavior exhibited under the heat treatment condition involving no solution heat treatment followed by the double aging treatment in that a remarkable improvement in hardness was obtained. The maximum hardness of the ZM61 alloy was similar to the ZAM621 alloy prepared by adding both AI and Mn.
Fig. 5 is a graph depicting respective age hardening behaviors of the Al + Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si + 0.4Ca) exhibited during a double aging process. After being extruded, each of the ZAM631 + 2.5Si and ZAM631 + 2.5Si + 0.4Ca alloys was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for a given period of time under the condition in which a variation in age hardening behavior was periodically measured within the secondary aging period. Referring to Figs. 3 and 5, it can be found that the addition of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy results in an improvement in hardness by about 12 % and a considerable reduction in the time taken to obtain a maximum hardness under the condition in which the alloy is subjected to a maximum double aging treatment at 150 C.
Table 3 Alloys Yield Strength Tensile Elongation (Mpa) strength (%) (Mpa) ZAM631+2.5Si 143 271 10 ZAM631 +2.5Si+0.4Ca 150 371 16 Fig. 6 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature. Referring to Fig. 6, it can be found that the addition of Mn alone or along with Al to the Z6 alloy results in a considerable increase in the yield strength and maximum tensile strength exhibited in an extruded state. In a wrought state using an extrusion process, the alloy exhibited a superior elongation of 25 % or more. Detailed results are described in the above Table 3.
Fig. 7 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature. Referring to Fig. 7, it can be found that the addition of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy results in an increase in the maximum tensile strength exhibited in an extruded state. In a wrought state using an extrusion process, the alloy exhibited a superior elongation of 16 % or more. Detailed results are described in the above Table 3.
Table 4 Tensile Alloys Yield Strength Strength Elongation (Mpa) (Mpa) (%) ZAM631 +2.5Si 265 317 7 ZAM631 +2.5Si+0.4Ca 234 398 12 Fig. 8 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature after the double aging process. Each of the Z6, ZM61, and ZAM621 alloys was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 8. Referring to Fig. 8, it can be found that the case involving the double aging treatment exhibits an increase in the yield strength and maximum tensile strength of each alloy, compared to the case involving no double aging treatment. It can also be found that both cases exhibit similar elongations, respectively. The tensile properties of each alloy, exhibited after a tensile test conducted following the double aging treatment, are described in the above Table 4.
Fig. 9 is a graph depicting respective tensile properties of the Al +
Mn + Si-added Mg-Zn-based alloy extruded product (ZAM631 + 2.5Si) and Al + Mn + Si + Ca-added Mg-Zn-based alloy extruded product (ZAM631 +
2.5Si + 0.4Ca) exhibited at ambient temperature after a double aging process. Each extruded alloy product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 24 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 9. Referring to Fig. 9, it can be found that the extruded products of the ZAM631 + 2.5Si and ZAM631 +2.5Si + 0.4Ca alloys respectively prepared by adding 2.5 wt.% Si and a mixture of 2.5 wt.% Si and 0.4 wt.% Ca to the ZAM631 alloy obtain effects of considerably increased yield strength and maximum tensile strength, compared to those subjected to no double aging treatment. Detailed results are described in the above Table 4.
Referring to Table 4, it can be found that the ZM61 alloy prepared by adding Mn to the Z6 alloy exhibits a slight increase in tensile properties by virtue of the double aging treatment, compared to the Z6 alloy. The ZAM621 alloy prepared by adding both AI and Mn to the Z6 alloy exhibited a superior strength over the Z6 alloy, by virtue of the double aging treatment. In particular, the ZAM621 alloy exhibits a remarkable increase in maximum tensile strength. All alloys exhibit a superior elongation even after the double aging treatment.
Fig. 10 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature- after a double aging process following a solution heat treatment with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted. Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a solution heat treatment for 12 hours at a temperature of 380 to 410 C. Following the solution heat treatment, each alloy extruded product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours. The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 10. Referring to Fig. 10, it can be found that when the solution heat treatment is carried out prior to the double aging treatment, the ZM61 alloy exhibits a considerable increase in yield strength and maximum tensile strength while exhibiting an elongation similar to that of the Z6 alloy.
Although the ZAM621 alloy exhibited a reduced yield strength, compared to the ZM61 alloy, it was similar to the ZM61 alloy in terms of the maximum tensile strength. In particular, the ZMA621 alloy exhibited a considerable increase in elongation. The tensile properties of each alloy exhibited at ambient temperature after the double aging treatment following the solution heat treatment are described in the following Table 5.
Table 5 Tensile Alloys Yield Strength Strength Elongation (Mpa) (Mpa) (%) Meanwhile, a comparison was made between the case, in which each alloy extruded product is subjected to a double aging process just after the extrusion thereof, and the case in which the alloy extruded product is subjected to the double aging process following a solution heat treatment.
Based on the results of the comparison, it can be found that the ZM61 alloy exhibits a considerable increase in strength by virtue of the solution heat treatment followed by the double aging process whereas the Z6 and ZAM621 alloys exhibit a slight strength increase under the same treatment conditions. It can also be found that the Z6 and ZM61 alloys exhibit a considerable reduction in elongation by virtue of the solution heat treatment followed by the double aging process whereas the ZAM621 alloy exhibit a similar elongation even after the solution heat treatment.
Fig. 11 is a graph depicting a comparison of respective tensile properties of the cast products of the Mn-added Mg-Zn-based alloy extruded product (ZM61) and Al + Mn-added Mg-Zn-based alloy extruded product (ZAM621) exhibited at ambient temperature after a double aging process following a 5% stretching process with the tensile properties of the Mg-Zn-based binary alloy extruded product (Z6) exhibited at ambient temperature when the same treatment is conducted. Each of the Z6, ZM61, and ZAM621 alloy extruded products was first subjected to a 5 % stretching process.
Following the stretching process, each extruded alloy product was primarily aged at 70 C for 48 hours, and then secondarily aged at 150 C for 96 hours.
The tensile properties of each alloy, exhibited after the double aging process, are depicted in Fig. 11. Referring to Fig. 11, it can be found that the ZAM621 alloy exhibits an improvement in strength, compared to the case involving no stretching process. The ZAM621 alloy also exhibited an elongation of 20 % or more. All alloys substantially exhibited an improvement in strength by virtue of the stretching process followed by the double aging process. In particular, when the ZAM621 alloy was subjected only to the stretching process while being subjected to the solution heat treatment, prior to the double aging process, in order to obtain an improvement in strength, it exhibited a strength similar to that of the ZM61 alloy while exhibiting a great increase in elongation. The tensile properties of each alloy exhibited at ambient temperature after the double aging process following the 5 %
stretching process are described in the following Table 6.
Table 6 Alloys Yield Strength Tensile Strength. Elongation (Mpa) (Mpa) (%) INDUSTRIAL APPLICABILITY
As apparent from the above description, the present invention provides a magnesium alloy having improvements in the hardness and strength at ambient temperature and an enhancement in elongation by adding Mn alone or along with Al to an Mg-Zn binary alloy while simultaneously adding Si alone or along with Ca to the Mg-Zn binary alloy to prepare a wrought body having a reduced grain size, and conducting a heat treatment or a working process involving a heat treatment for the wrought body.
Although the preferred embodiments of the invention have been disclosed for illustrative purposes, those skilled in the art will appreciate that various modifications, additions and substitutions are possible, without departing from the scope and spirit of the invention as disclosed in the accompanying claims.
Claims (13)
1. A high strength magnesium alloy consisting essentially of:
3 to 10 wt.% Zn, forming a precipitation phase in the alloy;
0.25 to 3.0 wt.% Mn refining the precipitation phase;
1 to 6 wt.% Al;
0 to 4.0 wt.% Si;
0 to 2.0 wt.% Ca; and the balance of Mg and inevitable impurities, wherein the content of Al is not more than the content of the Zn.
3 to 10 wt.% Zn, forming a precipitation phase in the alloy;
0.25 to 3.0 wt.% Mn refining the precipitation phase;
1 to 6 wt.% Al;
0 to 4.0 wt.% Si;
0 to 2.0 wt.% Ca; and the balance of Mg and inevitable impurities, wherein the content of Al is not more than the content of the Zn.
2. The high strength magnesium alloy according to claim 1, wherein the Si content is 0.1 to 4.0 wt.%.
3. The high strength magnesium alloy according to claim 2, wherein the Ca content is 0.1 to 2.0 wt.%.
4. The high strength magnesium alloy according to claim 1, wherein the Zn content is 5.0 to 7.0 wt.%.
5. The high strength magnesium alloy according to claim 1, wherein the Mn content is 0.75 to 2.0 wt.%.
6. A method for preparing a high strength magnesium alloy according to any one of claims 1 to 5, wherein Mn is added in the form of a Zn-Mn mother alloy to a magnesium melt.
7. The method according to claim 6, comprising the steps of:
adding a Zn-10 - 20 wt.% Mn mother alloy to the magnesium melt at a temperature of 670 to 720°C;
adding Zn alone or along with Al to the magnesium melt; and forming the resultant magnesium melt into a cast body.
adding a Zn-10 - 20 wt.% Mn mother alloy to the magnesium melt at a temperature of 670 to 720°C;
adding Zn alone or along with Al to the magnesium melt; and forming the resultant magnesium melt into a cast body.
8. The method according to claim 6, comprising the steps of:
adding a Zn-10 - 20 wt.% Mn mother alloy to the magnesium melt at a temperature of 670 to 720°C;
adding an Mg-Si mother alloy to the magnesium melt;
adding Zn alone or along with Al and/or Ca to the magnesium melt; and forming the resultant magnesium melt into a cast body.
adding a Zn-10 - 20 wt.% Mn mother alloy to the magnesium melt at a temperature of 670 to 720°C;
adding an Mg-Si mother alloy to the magnesium melt;
adding Zn alone or along with Al and/or Ca to the magnesium melt; and forming the resultant magnesium melt into a cast body.
9. The method according to claim 6 or 7, wherein the step of adding Zn alone or along with Al is carried out under a condition in which the magnesium melt has been cooled in a furnace cooling fashion to reduce a loss of Zn exhibiting a high vapor pressure at a temperature at which the alloy is prepared.
10. The method according to claim 7 or 8, further comprising the steps of:
subjecting the cast body to a homogenization process at a temperature of 340 to 410°C for 6 to 12 hours;
forming the cast body into a billet;
preheating the billet at a temperature of 150 to 400°C for 30 minutes to 2 hours; and working the billet, thereby forming a wrought body.
subjecting the cast body to a homogenization process at a temperature of 340 to 410°C for 6 to 12 hours;
forming the cast body into a billet;
preheating the billet at a temperature of 150 to 400°C for 30 minutes to 2 hours; and working the billet, thereby forming a wrought body.
11. The method according to claim 10, further comprising the steps of:
primarily aging the wrought body at a temperature of 70 to 100°C for 24 to 96 hours; and secondarily aging the primarily-aged wrought body at a temperature of 150 to 180°C for 48 hours or more.
primarily aging the wrought body at a temperature of 70 to 100°C for 24 to 96 hours; and secondarily aging the primarily-aged wrought body at a temperature of 150 to 180°C for 48 hours or more.
12. The method according to claim 11, further comprising the step of:
conducting a solution heat treatment for the wrought body at a temperature of 340 to 410°C for 6 to 12 hours, prior to the double aging.
conducting a solution heat treatment for the wrought body at a temperature of 340 to 410°C for 6 to 12 hours, prior to the double aging.
13. The method according to claim 11, further comprising the step of:
conducting a stretching of 3 to 7% for the wrought body, prior to the double aging.
conducting a stretching of 3 to 7% for the wrought body, prior to the double aging.
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| PCT/KR2000/001072 WO2002027053A1 (en) | 2000-09-26 | 2000-09-26 | High strength magnesium alloy and its preparation method |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| CA2423459A1 CA2423459A1 (en) | 2002-04-04 |
| CA2423459C true CA2423459C (en) | 2009-09-15 |
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Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA002423459A Expired - Lifetime CA2423459C (en) | 2000-09-26 | 2000-09-26 | High strength magnesium alloy and its preparation method |
Country Status (10)
| Country | Link |
|---|---|
| EP (1) | EP1339888B1 (en) |
| JP (1) | JP3891933B2 (en) |
| CN (1) | CN100390313C (en) |
| AT (1) | ATE505567T1 (en) |
| AU (2) | AU7688400A (en) |
| CA (1) | CA2423459C (en) |
| DE (1) | DE60045848D1 (en) |
| IL (1) | IL154897A (en) |
| NO (1) | NO20031349L (en) |
| WO (1) | WO2002027053A1 (en) |
Families Citing this family (22)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| AU2002950563A0 (en) * | 2002-08-02 | 2002-09-12 | Commonwealth Scientific And Industrial Research Organisation | Age-Hardenable, Zinc-Containing Magnesium Alloys |
| CN1300357C (en) * | 2004-09-29 | 2007-02-14 | 上海交通大学 | Preparation of high-strength creep resistant deforming magnesium alloy |
| CN100410407C (en) * | 2005-10-19 | 2008-08-13 | 郑州大学 | Mg-Al-Si-Mn-Ca alloy and method for preparing same |
| JP2010047777A (en) | 2007-05-09 | 2010-03-04 | National Institute For Materials Science | Mg-BASED ALLOY |
| JP5289904B2 (en) * | 2008-11-18 | 2013-09-11 | 三協立山株式会社 | Method for producing extruded magnesium alloy profile |
| CN101985715B (en) * | 2010-12-08 | 2012-07-04 | 沈阳工业大学 | High performance cast magnesium alloy and preparation method thereof |
| KR101264219B1 (en) * | 2011-01-05 | 2013-05-14 | 주식회사 에너텍 | Mg alloy and the manufacturing method of the same |
| JP5590413B2 (en) * | 2011-03-22 | 2014-09-17 | 株式会社豊田自動織機 | High thermal conductivity magnesium alloy |
| CN103849798B (en) * | 2012-11-30 | 2017-11-07 | 沈阳工业大学 | A kind of cast Mg alloy with high strength and preparation method thereof |
| CN103436827B (en) * | 2013-09-04 | 2015-09-09 | 中南大学 | A kind of thermal treatment process of large-size high-strength wrought magnesium alloy forging |
| JP6560193B2 (en) * | 2014-03-28 | 2019-08-14 | 古河電気工業株式会社 | Magnesium alloy tube material and manufacturing method thereof, and stent using the same and manufacturing method thereof |
| CN104532094A (en) * | 2014-12-15 | 2015-04-22 | 镁联科技（芜湖）有限公司 | Casting magnesium alloy as well as preparation method and application thereof |
| CN105714168A (en) * | 2016-03-25 | 2016-06-29 | 中国兵器科学研究院宁波分院 | High-yield-strength magnesium alloy and preparation method thereof |
| CN105937005B (en) * | 2016-06-17 | 2019-12-06 | 东北大学秦皇岛分校 | Ageing strengthening magnesium alloy with uniformly distributed granular quasicrystal and rod-shaped phase and preparation method thereof |
| CN109022978A (en) * | 2017-05-19 | 2018-12-18 | 曹丹丹 | The preparation method of magnesium alloy |
| CN111218594A (en) * | 2018-11-26 | 2020-06-02 | 内蒙金属材料研究所 | Magnesium-beryllium rare earth die-casting alloy and preparation method thereof |
| CN110835702A (en) * | 2019-10-29 | 2020-02-25 | 天津东义镁制品股份有限公司 | Magnesium alloy template profile and preparation method thereof |
| CN113430403B (en) * | 2021-05-17 | 2022-05-31 | 中北大学 | Method for preparing high-strength and high-toughness rare earth magnesium alloy through pre-aging |
| CN114107712B (en) * | 2021-11-26 | 2022-12-09 | 西安交通大学 | Medical magnesium-based composite material bar and preparation method thereof |
| CN114318094A (en) * | 2021-12-20 | 2022-04-12 | 重庆大学 | Mn particle reinforced Mg-Zn composite material and preparation method thereof |
| CN114703388A (en) * | 2022-04-12 | 2022-07-05 | 重庆大学 | Method for refining Mn-containing Mg-Zn-Al series cast magnesium alloy grains |
| CN115044813A (en) * | 2022-04-29 | 2022-09-13 | 北京工业大学 | Low-cost high-strength magnesium alloy material and preparation method thereof |
Family Cites Families (11)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US1886251A (en) | 1930-11-20 | 1932-11-01 | Dow Chemical Co | Magnesium-manganese-zinc alloys |
| GB1149502A (en) | 1965-05-11 | 1969-04-23 | Birmetals Ltd | Improved heat-treatment for magnesium-base alloys |
| NO132492C (en) * | 1973-10-01 | 1975-11-19 | Nl Industries Inc | |
| FR2642439B2 (en) | 1988-02-26 | 1993-04-16 | Pechiney Electrometallurgie | |
| JPH0718364Y2 (en) * | 1989-02-03 | 1995-05-01 | タイガー魔法瓶株式会社 | Tableware dryer |
| US5087304A (en) * | 1990-09-21 | 1992-02-11 | Allied-Signal Inc. | Hot rolled sheet of rapidly solidified magnesium base alloy |
| JP3184516B2 (en) | 1990-10-18 | 2001-07-09 | 住友金属鉱山株式会社 | Magnesium alloy for galvanic anode |
| JP3204572B2 (en) * | 1993-06-30 | 2001-09-04 | 株式会社豊田中央研究所 | Heat resistant magnesium alloy |
| JPH07316713A (en) * | 1994-05-25 | 1995-12-05 | Kobe Steel Ltd | Production of magnesium-base alloy having high strength and high corrosion resistance and casting made of this alloy |
| JPH09310130A (en) * | 1996-05-21 | 1997-12-02 | Sumikou Boshoku Kk | Production of magnesium alloy for galvanic anode |
| JPH10204556A (en) * | 1997-01-13 | 1998-08-04 | Mitsui Mining & Smelting Co Ltd | Magnesium alloy with high flowability, and its production |
-
2000
- 2000-09-26 EP EP00966545A patent/EP1339888B1/en not_active Expired - Lifetime
- 2000-09-26 WO PCT/KR2000/001072 patent/WO2002027053A1/en active IP Right Grant
- 2000-09-26 DE DE60045848T patent/DE60045848D1/en not_active Expired - Lifetime
- 2000-09-26 IL IL154897A patent/IL154897A/en not_active IP Right Cessation
- 2000-09-26 AU AU7688400A patent/AU7688400A/en active Pending
- 2000-09-26 JP JP2002530813A patent/JP3891933B2/en not_active Expired - Lifetime
- 2000-09-26 CA CA002423459A patent/CA2423459C/en not_active Expired - Lifetime
- 2000-09-26 AU AU2000276884A patent/AU2000276884B2/en not_active Ceased
- 2000-09-26 CN CNB008199914A patent/CN100390313C/en not_active Expired - Lifetime
- 2000-09-26 AT AT00966545T patent/ATE505567T1/en not_active IP Right Cessation
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2003
- 2003-03-25 NO NO20031349A patent/NO20031349L/en not_active Application Discontinuation
Also Published As
| Publication number | Publication date |
|---|---|
| EP1339888B1 (en) | 2011-04-13 |
| WO2002027053A1 (en) | 2002-04-04 |
| CA2423459A1 (en) | 2002-04-04 |
| IL154897A (en) | 2009-12-24 |
| NO20031349D0 (en) | 2003-03-25 |
| EP1339888A1 (en) | 2003-09-03 |
| AU2000276884B2 (en) | 2005-09-29 |
| JP2004510057A (en) | 2004-04-02 |
| EP1339888A4 (en) | 2005-03-16 |
| NO20031349L (en) | 2003-05-23 |
| IL154897A0 (en) | 2003-10-31 |
| CN100390313C (en) | 2008-05-28 |
| ATE505567T1 (en) | 2011-04-15 |
| JP3891933B2 (en) | 2007-03-14 |
| CN1469937A (en) | 2004-01-21 |
| DE60045848D1 (en) | 2011-05-26 |
| AU7688400A (en) | 2002-04-08 |
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