JP2007077486A - Aluminum alloy sheet for forming - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing a high Mg Al-Mg based alloy sheet having improved formability by controlling the existing state of Al-Mg based intermetallic compounds inevitably generated in the production process. <P>SOLUTION: The Al-Mg based aluminum alloy sheet contains 6.0 to 15.0% Mg, and the balance Al with inevitable impurities. Regarding Al-Mg based precipitates in the crystal grains observed by a transmission electron microscope of 50,000 magnifications and also identified by an electron beam probe microanalyzer in the structure of the central part in the sheet thickness of the sheet, the average grain size is ≤100 nm, and the average density is 0.1 to 10<SP>3</SP>pieces/μm<SP>2</SP>, thus the balance of strength/ductility is increased, and its press formability is improved. <P>COPYRIGHT: (C)2007,JPO&INPIT

Description

  The present invention relates to a high Mg content Al—Mg-based aluminum alloy plate having high formability.

  As is well known, various aluminum alloy plates (hereinafter referred to as “Al”) have been conventionally used for transportation equipment such as automobiles, ships, aircraft or vehicles, machines, electrical products, architecture, structures, optical equipment, and components and parts of equipment. Is also widely used depending on the characteristics of each alloy.

  In many cases, these aluminum alloy plates are formed by press molding or the like, and are used as members and parts for the above-described applications. In this respect, from the viewpoint of high formability, among the Al alloys, an Al-Mg Al alloy having an excellent balance between strength and ductility is advantageous.

  For this reason, with regard to Al-Mg-based Al alloy sheets, examination of component systems and optimization of manufacturing conditions have been conventionally performed. As this Al-Mg based Al alloy, for example, JIS A 5052, 5182 and the like are typical alloy component systems. However, even this Al—Mg-based Al alloy is inferior in ductility and inferior in formability compared to cold-rolled steel sheets.

  For this reason, with regard to Al-Mg-based Al alloy sheets, examination of component systems and optimization of manufacturing conditions have been conventionally performed. As this Al-Mg based Al alloy, for example, JIS A 5052, 5182 and the like are typical alloy component systems. However, even this Al—Mg-based Al alloy is inferior in ductility and inferior in formability compared to cold-rolled steel sheets.

  On the other hand, when the Al-Mg-based Al alloy is increased in Mg content to a high Mg content exceeding 6%, preferably 8%, the strength ductility balance is improved. However, it is difficult to industrially manufacture such a high Mg Al—Mg alloy by an ordinary manufacturing method in which an ingot cast by DC casting or the like is subjected to hot rolling after soaking. The reason for this is that Mg is segregated in the ingot during casting, and the normal hot rolling significantly reduces the ductility of the Al—Mg alloy, so that cracking is likely to occur.

  On the other hand, it is also difficult to hot-roll high-Mg Al—Mg alloys at low temperatures while avoiding the above-described temperature range where cracks occur. This is because, in such low-temperature rolling, the deformation resistance of the high-Mg Al—Mg-based alloy material is remarkably increased, and the product size that can be produced is extremely limited by the current rolling mill capability.

  In addition, a method of adding a third element such as Fe or Si has been proposed in order to increase the allowable Mg content of a high Mg Al—Mg alloy. However, when the content of these third elements is increased, a coarse intermetallic compound is easily formed, and the ductility of the aluminum alloy plate is lowered. For this reason, there is a limit to the increase in the Mg content allowable amount, and it was difficult to contain an amount of Mg exceeding 8%.

  For this reason, various proposals have heretofore been made for producing high-Mg Al—Mg-based alloy plates by a continuous casting method such as a twin roll type. In the twin roll type continuous casting method, molten aluminum alloy is poured from a refractory hot water supply nozzle between a pair of rotating water-cooled copper molds (twin rolls) and solidified. Immediately after that, it is reduced and rapidly cooled to form an aluminum alloy thin plate. As this twin roll type continuous casting method, the Hunter method, the 3C method and the like are known.

  The cooling rate of the twin roll type continuous casting method is 1 to 3 orders of magnitude higher than that of the conventional DC casting method or belt type continuous casting method. For this reason, the obtained aluminum alloy sheet has a very fine structure and is excellent in workability such as press formability. Moreover, the aluminum alloy plate having a relatively thin plate thickness of 1 to 13 mm is obtained by casting. For this reason, steps such as hot rough rolling and hot finish rolling can be omitted as in the case of a conventional DC ingot (thickness 200 to 600 mm). Furthermore, ingot homogenization may be omitted.

An example in which the structure of a high Mg Al—Mg alloy plate manufactured using such a twin-roll type continuous casting method is defined for the purpose of improving formability has been proposed. For example, an aluminum alloy sheet for automobiles with excellent mechanical properties is proposed in which the average size of Al-Mg based intermetallic compounds of Al-Mg based alloy sheets with a high Mg content of 6-10% is 10 μm or less. (See Patent Document 1). In addition, an aluminum alloy sheet for automobile body sheets, in which the number of Al-Mg intermetallic compounds of 10 μm or more is 300 pieces / mm ² or less and the average crystal grain size is 10 to 70 μm has been proposed (Patent Document 2). See).
JP 7-252571 A (full text) JP-A-8-165538 (full text)

  As described in Patent Documents 1 and 2, the Al—Mg intermetallic compound that crystallizes during casting is likely to be a starting point of fracture during press molding. Therefore, in order to improve the press formability of high-Mg Al-Mg-based alloy sheets, these Al-Mg-based intermetallic compounds (also referred to as β-phase, Al-Mg-based compounds) are disclosed in Patent Documents 1 and 2. As described above, it is effective to reduce the size or the size of coarse particles.

  However, Patent Documents 1 and 2 commonly suppress the Al-Mg intermetallic compound (β phase) that is crystallized during casting by increasing the cooling rate (casting rate) in the casting process. . However, the higher the Mg content, the more difficult it is to reduce the β phase of the Al—Mg-based alloy sheet to such an extent that it does not adversely affect the press formability only by controlling the cooling rate in the casting process.

  That is, even if the cooling rate (casting rate) in the twin-roll continuous casting method is increased to suppress the Al-Mg intermetallic compound that crystallizes during casting, in the subsequent process, continuous casting is performed. In addition to subsequent cooling to room temperature, the temperature of the plate ingot or sheet is 400 ° C or higher, such as homogenization heat treatment before cold rolling, intermediate annealing during cold rolling, solution treatment after cold rolling, etc. The process of heating the plate or cooling the heated plate-shaped ingot or sheet is selectively included in the process design. In these thermal history processes, there is a possibility that an Al—Mg-based intermetallic compound called β phase is generated.

  Therefore, it is difficult to simply suppress the generation of Al-Mg intermetallic compounds. For example, even if an Al-Mg intermetallic compound is present, the existence form of the Al-Mg intermetallic compound, etc. It can be said that there is a need for a technology that improves the press formability of Al-Mg alloy plates with high Mg by controlling the above.

  The present invention has been made to solve such problems, and its purpose is to improve the formability by controlling the presence of Al-Mg intermetallic compounds that are inevitably generated in the production process. Another object of the present invention is to provide a method for producing a high Mg Al-Mg alloy plate.

In order to achieve this object, the gist of the aluminum alloy sheet for molding of the present invention is an Al-Mg aluminum alloy sheet containing Mg: 6.0 to 15.0% in mass%, the balance being Al and inevitable impurities. The Al-Mg-based precipitates in the crystal grains, which are observed with a transmission electron microscope of 50000 times and identified by the electron probe microanalyzer, show the average grain size in the center of the thickness of the plate. And the average density is 0.1 piece / μm ² or more and 10 ³ piece / μm ² or less.

  The present inventors present the presence of Al-Mg-based precipitates (Al-Mg-based intermetallic compounds) called β-phase in the high-Mg Al-Mg-based alloy sheet produced by twin-roll continuous casting. About form, it discovered that a moldability could be improved by making it precipitate finely to a nano level.

As described above, the β-phase nano-level fine precipitation means that the average grain size of the Al-Mg-based precipitates in the crystal grains is 100 nm or less, and the average density is 50,000 times the TEM observation of the plate structure. It is 0.1 piece / μm ² or more and 10 ³ piece / μm ² or less.

It has been difficult to observe such a β-phase nanoprecipitation state. This is because the Al—Mg-based precipitate (Al ₃ Mg ₂ ) constituting the β phase has a nano-level size and was difficult to identify even with the transmission electron microscope (FE-TEM).

  On the other hand, the present inventors have identified the nanometer-scale fineness of the β phase by means of the later-described identification using EPMA (electron probe microanalyzer) and observation within the crystal grains using FE-TEM (transmission electron microscope). The observation of the precipitation state was made possible.

  In addition, the inventors have also found that the state of the microstructure of the β-phase nano-level fine precipitation state defines the formability, which is a macro characteristic of the plate.

  In the present invention, on the basis of these facts, the nano-phase fine precipitation state of the β phase in the high Mg Al-Mg alloy plate is defined as described above, and the high Mg Al-Mg alloy plate is formed. It improves the performance.

(Nano-level fine precipitation state of β phase)
Fig. 1 shows the microstructure in the crystal grains (50,000-times FE-TEM photograph). Al-Mg-based precipitates in the crystal grains of the high-Mg Al-Mg-based alloy plate of Invention Example 1 of the examples described later. (β phase). In FIG. 1, the black dots dispersed and present are Al—Mg-based precipitates (β phase) in the crystal grains. In FIG. 1, black streaks or lines are dislocations introduced into the crystal grains. In FIG. 1, the β-phase with black dots has an average particle diameter of 24 nm which is 100 nm or less, and the average density is within the range of the present invention of 0.65 particles / μm ² .

As supported by the examples described later, in the present invention, in order to improve the moldability, the nano-level β phase in the crystal grains has an average grain size of 100 nm or less, an average density of 0.1 particles / μm ² or more, 10 ³ The number of pieces / μm ² or less. By finely precipitating the β phase at the nano level as in the present invention, the β phase coarsened with a grain size exceeding 100 nm is not present in the crystal grains.

On the other hand, if the β phase is coarsened with an average particle diameter exceeding 100 nm or an average density exceeding 10 ³ particles / μm ² , the moldability is lowered. On the other hand, if the average density of the β phase is less than 0.1 pieces / μm ² , the effect of these fine precipitates cannot be obtained and the moldability is lowered.

The β phase with an average particle size of 100 nm or less is the surface of the Al-Mg alloy plate that has been subjected to electro-etching after 0.05 to 0.1 mm of the thickness center of the plate (both in the plate thickness direction and in the cold rolling direction of the plate) ) Is observed by 50,000 times FE-TEM (transmission electron microscope) of the tissue and identified by an X-ray spectrometer (EDX). The structure observation by FE-TEM at the center of the plate thickness is performed so that the total area of the observation field is 4 μm ² or more at one center of the plate thickness, and this is placed at an appropriate distance in the longitudinal direction of the plate Average the observed results.

  The diameter of the β phase is the equivalent circle diameter of the β phase per piece, and the equivalent circle diameter is measured for all the β phases in the field of view and averaged to obtain the average particle diameter at one observation point. Further, the average particle diameter per one place is averaged at five places to obtain the average particle diameter referred to in the present invention.

The average density of β phases is the number of β phases per 1 μm ² of the total area of the observation field. The average density referred to in the present invention is obtained by averaging the number per one observation at five locations.

(Average crystal grain size)
It is preferable as a precondition for improving the formability that the average crystal grain size on the surface of the Al alloy plate is refined to 100 μm or less. By making the crystal grain size fine or small within this range, press formability is ensured or improved. When the crystal grain size becomes larger than 100 μm, the press formability is remarkably deteriorated, and defects such as cracks and rough skin during forming tend to occur. On the other hand, even if the average crystal grain size is too small, SS (stretcher strain) marks, which are peculiar to 5000 series Al alloy plates, are generated during press molding. From this point of view, the average crystal grain size is 20 μm or more. It is preferable to do.

  The crystal grain size referred to in the present invention is the maximum diameter of crystal grains in the longitudinal (L) direction of the plate. The crystal grain size is measured by a line intercept method in the L direction by observing the surface of the Al alloy plate that has been mechanically polished by 0.05 to 0.1 mm and then electrolytically etched using a 100 × optical microscope. 1 The measurement line length is 0.95mm, and the total measurement line length is 0.95 x 15mm by observing a total of 5 fields with 3 lines per field.

(Chemical composition)
The significance of each alloy element and the reason for its limitation in the chemical composition of the Al alloy sheet of the present invention will be described below. The Al alloy sheet of the present invention has a chemical component composition containing, by mass%, Mg: 6.0 to 15.0%, the balance being Al and inevitable impurities.

(Mg: 6.0-15.0%)
Mg is an important alloy element that increases the strength and ductility of the Al alloy sheet. If the Mg content is too small, the strength and ductility are insufficient, the characteristics of a high Mg Al—Mg-based Al alloy are not obtained, and the formability is insufficient. On the other hand, if the Mg content is too large, the precipitation of Al-Mg compounds increases even if the production method and conditions are controlled. As a result, moldability is significantly reduced. In addition, the work hardening amount is increased and the cold rollability is also lowered. Therefore, Mg is in the range of 6.0 to 15.0%, preferably more than 8% and 14% or less.

(Fe: 1.0% or less, Si: 0.5% or less)
Fe and Si are impurities that should be regulated to the smallest possible amount. Fe and Si are produced in a large amount in the amount of Al-Mg compounds composed of Al-Mg- (Fe, Si) and the like, and the amount of compounds other than Al-Mg compounds such as Al-Fe and Al-Si. When the Fe content exceeds 1.0% and the Si content exceeds 0.5%, the amount of these compounds becomes excessive, which significantly impairs fracture toughness and formability. As a result, the moldability is significantly reduced. Therefore, Fe is regulated to 1.0% or less, preferably 0.5% or less, and Si is regulated to 0.5% or less, preferably 0.3% or less.

  In addition, Mn, Cu, Cr, Zr, Zn, V 2, Ti, B, etc. are also impurity elements, and it is better that the content is small. However, for example, Mn, Cr, Zr, and V have the effect of refining the rolled plate structure, and Ti and B have the effect of refining the cast plate (ingot) structure. Cu and Zn also have the effect of improving strength. For this reason, it may be included with the aim of these effects, and it is allowed to contain one or more of these elements within a range that does not impair the formability that is a characteristic of the plate of the present invention. These allowable amounts are, respectively,% by mass, Mn: 0.3% or less, Cr: 0.3% or less, Zr: 0.3% or less, V: 0.3% or less, Ti: 0.1% or less, B: 0.05% or less, Cu: 1.0% or less, Zn: 1.0% or less.

(Production method)
Below, the manufacturing method of the Al-Mg type | system | group Al alloy plate in this invention is demonstrated.
The high-Mg Al-Mg-based Al alloy plate of the present invention may be an ordinary production method in which an ingot cast by DC casting or the like is subjected to hot rolling after soaking, but in this ordinary production method, As described above, it is difficult to industrially manufacture high-Mg Al-Mg Al alloy plates efficiently.

  Therefore, when industrially producing the high Mg Al-Mg-based Al alloy sheet of the present invention, currently, continuous casting such as twin roll type, cold rolling and annealing without hot rolling are performed. It is preferable to use a plate having a thickness of 0.5 to 3 mm manufactured in combination.

  In this regard, at the time of the twin roll type continuous casting, the molten aluminum alloy containing high Mg is poured into a pair of rotating twin rolls, and the cooling rate of the twin rolls is set to 100 ° C./s or more. It is preferably produced by continuous casting in the range of 1 to 13 mm. Furthermore, in order to reliably achieve higher press formability, it is preferable that the twin roll surface is not lubricated during the continuous casting.

(Double roll type continuous casting)
As a continuous casting method, there are a belt caster type, a propel type, a block caster type, etc. in addition to the twin roll type. However, in order to increase the cooling rate at the time of casting a high Mg Al—Mg-based Al alloy plate as described later, twin-roll continuous casting is preferable.

  This twin roll type continuous casting, as described above, between the twin rolls such as a pair of rotating water-cooled copper molds, from the hot water supply nozzle made of refractory material, Al alloy molten metal having the above composition is poured and solidified, and Then, between the twin rolls, the Al alloy thin plate is obtained by reducing and quenching immediately after the solidification.

(Double roll lubrication)
At this time, as the twin roll, it is desirable to use a roll whose surface is not lubricated by a lubricant. Conventionally, oxide powder (alumina powder, zinc oxide powder, etc.), SiC powder, graphite powder, In general, a lubricant (release agent) such as oil or molten glass is applied to or flowed down on the twin roll surface. However, when these lubricants are used, the cooling rate becomes slow and the required cooling rate cannot be obtained. For this reason, the crystal grains become coarse, and the formability of the high Mg Al—Mg alloy plate decreases.

  In addition, when these lubricants are used, cooling unevenness is likely to occur due to the uneven concentration and thickness of the lubricant on the twin roll surface, and the solidification rate tends to be insufficient depending on the part of the plate. For this reason, as the Mg content increases, macro segregation and micro segregation increase, and it becomes more likely that it becomes difficult to make the formability of the Al-Mg alloy plate uniform.

(Double roll cooling rate)
For example, in order to reduce the average grain size of high-Mg Al-Mg alloy plates even if the thickness of the cast plate is in the range of relatively thin plates of 1 to 13 mm, the cooling of the casting by this twin roll is used. The cooling rate must be as high as possible (solidification rate) of at least 100 ° C / s. When the above lubricant is used, even if the cooling rate is theoretically high, the actual or actual cooling rate tends to be substantially less than 100 ° C./s. For this reason, the average crystal grain size of the high Mg Al—Mg alloy plate cannot be made fine, and the press formability is significantly reduced.

Since this cooling rate is difficult to measure directly, a method known from the dendrite arm spacing (Dendrite secondary branch interval, DAS) of the cast plate (ingot) (for example, Light Metal Society, 8.20 1988) Published in “Methods of measuring aluminum dendrite arm spacing and cooling rate”). That is, the average distance d between adjacent dendrite secondary arms (secondary branches) in the cast structure of the cast plate was measured using the intersection method (number of fields of view of 3 or more, number of intersections of 10 or more). Using d, the following formula is obtained: d = 62 × C ^−0.337 (where d: dendrite secondary arm interval mm, C: cooling rate ° ^C./s ).

(Twin roll casting thickness)
The thickness of the thin plate continuously cast by twin rolls is in the range of 1 to 13 mm. And, preferably, a thin plate thickness of 1 mm or more and less than 5 mm. Continuous casting with a thickness of less than 1 mm is difficult due to casting limitations such as pouring between twin rolls and controlling the roll gap between twin rolls. On the other hand, when the plate thickness is 13 mm, or more strictly, the plate thickness exceeds 5 mm, the cooling rate of the casting becomes extremely slow, and the overall intermetallic compounds such as Al-Mg system become coarse or a large amount of crystallization occurs. Tend to. As a result, there is a high possibility that the press formability is significantly lowered.

(Twin roll pouring temperature)
The pouring temperature when pouring Al alloy molten metal into the twin rolls is preferably set to the liquidus temperature + 30 ° C. or lower. When the pouring temperature exceeds the liquidus temperature + 30 ° C, the casting cooling rate described later becomes small, and all intermetallic compounds such as the Al-Mg system may become coarse or crystallize in large quantities. As a result, the strength-elongation balance is lowered, and the press formability may be significantly lowered. In addition, the rolling effect of the twin rolls is reduced, the number of center defects increases, and the basic mechanical properties of the Al alloy plate itself may be deteriorated.

(Twin roll speed)
The peripheral speed of the pair of rotating twin rolls is preferably 1 m / min or more. If the peripheral speed of the twin rolls is less than 1 m / min, the contact time between the molten metal and the mold (twist roll) becomes long, and the surface quality of the cast thin plate may deteriorate. In this respect, the higher the peripheral speed of the twin rolls, the better, and the preferable peripheral speed is 30 m / min or more.

(Reduction by twin rolls)
In the present invention, selectively or as needed, the length of the plate-shaped ingot is 1 m by the twin rolls with respect to the plate-shaped ingot solidified between the twin rolls after pouring into the twin rolls. Casting may be performed while applying a rolling load of 300 tons or more, that is, 300 tons / m or more.

  Due to the load of the rolling load, the gas generated during pouring or during solidification is easily released from the inside of the plate-shaped slab. For this reason, even in the case of a high Mg Al—Mg alloy having a wide solidification temperature range of about 100 ° C., no gas stays in the slab structure, and the resulting voids are suppressed. Then, due to a synergistic effect with the subsequent cold rolling, it is possible to suppress casting defects such as voids to a range that does not affect molding characteristics such as elongation of the manufactured plate.

  This effect due to the load of the rolling load is of course also affected by the thickness of the casting and the casting conditions, but in the range of relatively thin plates with a casting thickness of 1 to 13 mm, the reduction is 300 ton / m or more. Demonstrated by load. In addition, 300 ton / m or more is a rolling load (ton) per 1 m in the longitudinal direction of the plate-shaped ingot.

(Homogenization heat treatment)
Homogenization heat treatment (also referred to as soaking) is essential before hot rolling in order to suppress the segregation of Mg in ingots cast by DC casting or the like. Further, even a plate-like ingot produced by a twin roll type continuous casting method with relatively little Mg segregation is preferably applied before cold rolling in order to suppress Mg segregation.

  The homogenization heat treatment is performed at a temperature of 400 ° C or higher and below the liquidus temperature for the required time. This time is approximately 1 second (1 s) or less when homogenizing heat treatment using a continuous heat treatment furnace for thin ingots produced by the twin roll continuous casting method. In addition, when the ingot cast by DC casting or the like is subjected to homogenization heat treatment using a batch heat treatment furnace, 1 to 10 hours (1 to 10 hours) is a standard. By this homogenization heat treatment, the degree of Mg segregation is reduced, and the degree of Mg segregation can be suppressed within the scope of the present invention.

  In the homogenization heat treatment, if the heating rate and the cooling rate are low in the course of both the heating and cooling of the ingot, there is a possibility that an Al-Mg intermetallic compound is generated. In particular, the temperature range where Al-Mg-based intermetallic compounds are likely to occur is the range where the temperature of the ingot center is 200 ° C to 400 ° C when the temperature is raised, and the homogenization heat treatment temperature is 100 ° C when it is cooled. Range.

  For this reason, when selectively performing such a homogenization heat treatment, the temperature of the ingot center is reduced during heating to the homogenization heat treatment temperature in order to suppress the generation of Al-Mg intermetallic compounds. The average rate of temperature rise in the range from 200 ° C to 400 ° C is preferably 5 ° C / s or more. In cooling from the homogenization heat treatment temperature, the average cooling rate in the range from the homogenization heat treatment temperature to 100 ° C. is preferably 5 ° C./s or more.

(Hot rolling)
An ingot cast by DC casting or the like is cooled to a hot rolling temperature after the homogenization heat treatment or hot rolled as it is. This hot rolling condition may be a conventional method. On the other hand, the plate-shaped ingot by the twin roll type continuous casting method is cold-rolled without being hot-rolled either online or offline.

(Cold rolling)
In cold rolling, a plate-shaped ingot produced by a twin-roll continuous casting method, and in an ingot cast by DC casting or the like, the hot-rolled hot-rolled plate has a thickness of 0.5 to 3 mm. By cold rolling, the cast structure is processed.

  In this regard, when the thickness of the cold-rolled plate is thick, it is preferable that intermediate annealing is performed in the middle of cold rolling so that the cold rolling rate in the final cold rolling is 60% or less. Note that the degree of work organization in cold rolling depends on the cold rolling rate of cold rolling, and the cast structure may remain due to the above-mentioned texture control, but this hinders formability and mechanical properties. Tolerable range.

(Final annealing)
The Al alloy cold-rolled sheet is preferably finally annealed at 400 ° C. to the liquidus temperature. If the annealing temperature is less than 400 ° C., there is a high possibility that the solution effect will not be obtained, and there is a high possibility that the amount of β phase will increase. For this reason, nano-level fine precipitation is not possible with the β phase in the crystal grains having an average particle size of 100 nm or less and an average density of 0.1 / μm ² or more and 10 ³ / μm ² or less. For this reason, there is a high possibility that the elongation of the high-Mg Al—Mg-based alloy plate is lowered, the strength-ductility balance is lowered, and the press formability is lowered. The final annealing temperature is preferably 450 ° C. or higher.

  Further, after this final annealing, it is necessary to cool at a temperature range of 500 to 300 ° C. at an average cooling rate as fast as possible of 10 ° C./s or more. If the average cooling rate after the final annealing is slow and less than 10 ° C./s, a large amount of β phase precipitates during the cooling process. As a result, nano-level fine deposition cannot be performed. For this reason, there is a high possibility that the elongation of the high-Mg Al—Mg-based alloy plate is lowered, the strength-ductility balance is lowered, and the press formability is lowered. For this reason, the average cooling rate is preferably 15 ° C./s or more.

(Additional annealing)
Furthermore, if the additional annealing shown below is performed again during cooling after the final annealing, it is effective to improve the strength-ductility balance by making the β phase in the crystal grains into the above-mentioned specific nano-level fine precipitates.

  In the temperature range of 150 ° C to 250 ° C during cooling after final annealing, the β-phase in the crystal grains is cooled to room temperature after additional annealing that is held for 0 s to 10 min. The fine precipitation of the level can be achieved, and the strength-ductility balance is improved. When the additional annealing temperature exceeds 250 ° C. or when the holding time exceeds 10 minutes, the amount of β phase in the crystal grains increases, so that the strength-ductility balance decreases and press formability decreases. Preferably, the addition annealing temperature is 160 ° C. or more and 220 ° C. or less, or the addition annealing time is 2 min or more and 8 min or less. The cooling rate after the additional annealing is 10 ° C./s or more, preferably 15 ° C./s or more.

  Further, in the temperature range of 50 ° C. or more and less than 150 ° C. during cooling after the final annealing, the same effect can be obtained by performing additional annealing that is held for 30 minutes or more and 10 hours or less. When the additional annealing holding time exceeds 10 hours, the amount of β phase in the crystal grains increases, so that the strength-ductility balance decreases and press formability decreases. Further, when the additional annealing time is less than 30 min or the additional annealing temperature is 50 ° C. or lower, these effects cannot be obtained. The temperature is preferably 60 ° C or higher and 120 ° C or lower.

  Examples of the present invention will be described below. Conditions shown in Table 2 for Al-Mg-based Al alloy melts (Invention Examples A to D, Comparative Examples E and F) having various chemical composition compositions shown in Table 1 by the twin roll continuous casting method and the DC casting method described above. In each ingot plate thickness was cast.

  In the case of the twin roll continuous casting method, each Al alloy sheet ingot was soaked under the conditions shown in Table 2 and then cold rolled to a thickness of 1.0 mm without hot rolling. In the case of DC casting, each aluminum alloy ingot is soaked under the conditions shown in Table 2 and then rolled to a plate thickness of 4.0 mm at a start temperature of 480 ° C and an end temperature of 350 ° C. And then cold rolled to a thickness of 1.0 mm. In addition, the intermediate annealing during these cold rolling was not performed.

  Each cold-rolled sheet was subjected to final annealing and additional annealing in a continuous annealing furnace at the temperatures, cooling conditions, and additional annealing conditions shown in Table 2. In addition, addition annealing used the batch annealing furnace or the continuous annealing furnace depending on annealing conditions. In the case of a batch annealing furnace, after winding at the additional annealing temperature, it was immediately inserted into the batch annealing furnace to perform additional annealing.

  At the time of twin roll continuous casting, the peripheral speed of the twin roll is 70 m / min, and the pouring temperature when pouring the Al alloy molten metal into the twin roll is constant at the liquidus temperature + 20 ° C in each example. The roll surface was not lubricated.

The thus obtained high-Mg Al-Mg-based Al alloy sheet after the final annealing was obtained at any measurement location, 10 locations at a distance of 100 mm or more in the longitudinal direction (rolling direction). A sample was taken from the center of the plate thickness, and the average particle diameter (nm) and average density (pieces / μm ² ) of the Al—Mg-based precipitates in the crystal grains were measured by the measurement method described above. Table 3 shows the measurement results. The FE-TEM used was a field emission transmission electron microscope (HF-2000) manufactured by Hitachi, Ltd.

  In both the inventive examples and the comparative examples, except for Comparative Example 13, the average crystal grain size on the surface of the obtained Al alloy plate was in the range of 30 to 60 μm. In Comparative Example 13, the average crystal grain size on the surface of the obtained Al alloy plate exceeded 100 μm.

  Further, specimens were collected from the center part of the plate thickness, and the mechanical properties of each specimen and the average value of the balance between strength and ductility [tensile strength (TS: MPa) × total elongation (EL:%)] (MPa%) In addition, properties such as moldability were measured and evaluated. These results are also shown in Table 3.

  The tensile test was performed according to JIS Z 2201, and the shape of the test piece was a JIS No. 5 test piece, and the test piece was manufactured so that the longitudinal direction of the test piece coincided with the rolling direction. The crosshead speed was 5 mm / min, and the test piece was run at a constant speed until the test piece broke.

  As a material test evaluation of formability, a bulging test in a plane strain state was performed as an evaluation of the bulging property, and a burring test was performed as an evaluation of stretch flangeability.

  The overhang test uses a ball head overhang punch with a diameter of 101.6 mm, applies R-303P as a lubricant to a test piece with a length of 180 mm and a width of 110 mm, and overhangs at a forming speed of 4 mm / s and a wrinkle holding load of 200 kN. A molding test was performed, and the height (mm) when the test piece was cracked was measured.

In the burring test, a 10 mm diameter hole is punched into a square plate with a side of 100 mm. Then, using a 60 ° conical punch with a diameter of 25 mm, using a burrs on the upper surface (die surface) side, using rust preventive oil as a lubricating oil, conducting a burring test with a wrinkle holding force of 4.0 tons and a punch speed of 10 m / min The punch was stopped when the edge of the punched hole was broken, and the burring rate (λ) was determined from the hole inner diameter (d _s ) after the fracture and the initial hole diameter (d ₀ ) before the molding test by the following formula.
λ: (d _s −d ₀ ) / d ₀ × 100

  About the hole internal diameter after a fracture | rupture, it measured in the rolling direction and the direction perpendicular | vertical to a rolling direction, respectively, calculated | required each burring rate, and took the average, and was taken as the burring rate of each sample. Further, the burring test was performed three times for each sample, and the average value was finally set as the burring rate (λ%). These results are also shown in Tables 2 and 3.

As shown in Tables 1 and 2, Invention Examples 1 to 14 are examples of high Mg Al-Mg based Al alloy plates having compositions within the scope of the present invention of A to D in Table 1, and within the preferable production condition range. It is manufactured. Therefore, as shown in Table 3, Invention Examples 1 to 14 show that the Al-Mg-based precipitates in the crystal grains of the structure in the center of the plate thickness of the plate have an average particle size of 100 nm or less and an average density of 0.1 / It is not less than μm ^{2 and not} more than 10 ³ / μm ² . In Invention Examples 1 to 14, the β phase coarsened with a particle size exceeding 100 nm was not the result of the structure observation. As a result, Invention Examples 1 to 14 have high strength ductility balance, limit overhang height, and high λ, and are excellent in press formability.

  On the other hand, as shown in Tables 1 and 2, Comparative Examples 17 to 23 are examples of high-Mg Al-Mg-based Al alloys having compositions within the scope of the present invention of A and B in Table 1, but preferred production. Manufactured outside the range of conditions. For this reason, as shown in Table 3, the average grain size or average density of the Al—Mg-based precipitates in the crystal grains of the structure in the central portion of the plate thickness is outside the range of the present invention. As a result, the strength ductility balance is low and the press formability is poor. Note that the limit overhang height in the overhang test in Table 3 shows that even if the difference between the inventive example and the comparative example is about 2 mm, the difference in formability in the actual overhang forming becomes larger. The difference of about 2mm is important.

  Comparative Examples 17 and 18 are examples of alloys having a composition within the scope of the present invention of A in Table 1, but no additional annealing was performed after the final annealing.

  Comparative Example 19 is an alloy example having a composition within the scope of the present invention of A in Table 1, but the additional annealing temperature is too high after the final annealing.

  Although Comparative Example 20 is an alloy example having a composition within the range of the present invention of A in Table 1, the additional annealing temperature after the final annealing is too low.

  Comparative Example 21 is an alloy example having a composition within the range of the present invention of A in Table 1, but the additional annealing time after the final annealing is too long.

  Comparative Example 22 is an example of an alloy having a composition within the scope of the present invention of A in Table 1, but no additional annealing is performed after the final annealing.

  Comparative Example 23 is an alloy example having a composition within the scope of the present invention of A in Table 1, but the cooling rate after the final annealing is too low.

On the other hand, Comparative Examples 15 and 16 are examples of high-Mg Al—Mg-based Al alloys having compositions outside the scope of the present invention of E 1 and F 2 in Table 1, although manufactured within the range of preferable manufacturing conditions. For this reason, the strength ductility balance is low and the press formability is poor.
Comparative Example 15 uses an alloy of E 2 whose Mg content is too low below the lower limit.
Comparative Example 16 uses an F 2 alloy whose Mg content exceeds the upper limit and is too high.

  Therefore, from these, the critical significance for strength, ductility, strength-ductility balance, and formability of the Al—Mg-based precipitate definition in the crystal grains of the present invention and preferable production conditions within this range can be understood.

  As described above, according to the present invention, it is possible to improve the elongation and strength ductility balance of aluminum alloys including high-Mg Al—Mg alloys, and to improve the formability. As a result, it can be applied to aluminum alloy plate applications that require formability, such as automobiles, ships, airplanes, vehicles, and other transport equipment, machines, electrical products, architecture, structures, optical equipment, and members and parts of equipment. Can be expanded.

It is explanatory drawing which shows the structure | tissue photograph (invention example) in the crystal grain of an Al-Mg type alloy plate.

Claims (3)

Mass-%, Mg: 6.0 to 15.0% Al-Mg-based aluminum alloy plate consisting of Al and inevitable impurities, a 50000 times transmission electron microscope of the structure at the center of the plate thickness Al-Mg-based precipitates in the crystal grains, which are observed by the electron beam probe microanalyzer, are average particle diameter of 100 nm or less, average density is 0.1 / μm ² or more, 10 ³ / μm An aluminum alloy sheet for molding, characterized in that it is ² or less.   The aluminum alloy sheet for forming according to claim 1, wherein the Mg content is more than 8% and 14% or less.   The aluminum alloy plate is an element other than Mg, Fe: 1.0% or less, Si: 0.5% or less, Mn: 0.3% or less, Cr: 0.3% or less, Zr: 0.3% or less, V: 0.3% or less, Ti The aluminum alloy sheet for forming according to claim 1 or 2, which is suppressed to: 0.1% or less, Cu: 1.0% or less, and Zn: 1.0% or less.

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