JP2022540544A - Fine-scale eutectic textures, specifically alloys with nano-eutectic textures, and the manufacture of such alloys - Google Patents

Abstract

The present invention provides an alloy, particularly a light metal alloy, having an alloy composition having at least three components and a eutectic structure, wherein the eutectic structure is defined by the pseudo-eutectic point of the phase diagram of the alloy. (pE) field and obtained by cooling the alloy from the liquid state to the solid state under conditions resulting in the presence of at least 85 mol % of the eutectic structure in the alloy. Alloys also relate to methods for producing alloys of this type.

Description

The present invention relates to alloys, particularly light metal alloys, having an alloy composition having at least three components and a eutectic structure obtained by cooling the alloy from a liquid state to a solid state.

The present invention is a method for producing an alloy, particularly a light metal alloy, having a eutectic structure, the alloy having an alloy composition having at least three components, the alloy forming a eutectic structure The method further relates to a method in which the alloy is cooled to a solid state starting from the liquid state in order to

It is known that it can be advantageous if part of the alloy's structure is embodied with a eutectic structure to affect the casting properties or strength properties of the alloy. Binary cast alloys, ie, alloys having two components with a eutectic microstructure, are often used as alloys of industrial utility. These alloys are usually characterized by a eutectic point in their phase diagram, at which the liquid phase of the alloy and the two solid phases of the alloy are in thermodynamic equilibrium with each other, or the alloy is in the liquid phase. A direct transition from the liquid state to the solid state occurs upon cooling from , where a eutectic structure is formed. According to Gibbs' phase law f=NP+1 for solids at constant pressure, where f is the number of thermodynamic degrees of freedom, N is the number of components, and P is the number of equilibrium phases. corresponds to the number. A direct transition from the liquid phase to the solid phase thereby often results in the formation of a fine, lamellar texture.

Similarly, in the context of ternary alloy systems, attempts are also known to produce alloys with compositions close to the ternary eutectic point in order to improve strength properties using embodiments of the eutectic structure. . According to the Gibbs phase law f=N−P+1, it has 3 components and 4 phases, also corresponding to f=0 degrees of freedom. However, in order to produce a eutectic structure with significant fineness in alloy quantities that can be used in practice, high cooling rates are usually required to form this type of alloy, and for precipitation hardening, Additional combinations, matched with other elements, are often required to increase the strength of the alloy. Most often, cooling rates in the range of 50 K/s to 200 K/s are used for this purpose. However, the requirement for high cooling rates specifically limits the technical usefulness of this type of alloy to small-scale parts.

This is addressed by the present invention. The object of the present invention is to identify alloys with at least three components that have high strength and good deformability.

A further goal of the invention is to identify a method for producing alloys of this type.

With alloys of the type mentioned at the beginning, under conditions where the composition of the alloy is in the near-quasi-eutectic field of the phase diagram of the alloy, resulting in at least 85 mol % or at % of the eutectic structure in the alloy. and a eutectic structure is obtained, the object is achieved according to the invention.

The rationale for the present invention is that using an alloy composition having at least three constituents or elements at or near the pseudo-eutectic point in the phase diagram of the alloy results in a particularly finely sized or finely textured A refined eutectic structure may be embodied, which may in particular have a finer eutectic structure than an alloy with a selected composition at the "normal" eutectic point in the phase diagram. This is the discovery. Thus, in particular, characteristic texture spacings of the eutectic texture in the low micrometer range and especially in the nanometer range can be achieved, also referred to as nanoeutectic texture. In addition, the eutectic structure thereby formed usually constitutes the main or predominant microstructure, specifically the often small or negligible primary solidification phase and/or residual solidification phase. It has been shown that phases only occur near the eutectic point, or in fields near the eutectic point, specifically at the eutectic point, or not at all. This combination of the extremely fine microstructure and its predominant presence of the eutectic structure in the alloy embodies an alloy that has both high strength, specifically compressive strength, and outstanding deformability. . In notation, the pseudo-eutectic point is typically abbreviated by "e" or "pE" and the eutectic point is represented as abbreviated by "E".

Technically, in a ternary phase diagram, the liquidus and solidus, known from the binary phase diagram, typically correspond to the area of the curved surface and the binary phase area to the phase volume, respectively. . In a ternary phase diagram, the intersecting lines of the liquidus area typically form eutectic channels, also called liquidus boundary lines or first-variant lines, which terminate at the ternary eutectic point of the phase diagram. . The quasi-eutectic point thereby represents a point on the liquidus boundary forming a saddle point, i.e., the local extremum along the liquidus boundary and the shortest represents a perpendicular line.

The binary eutectic is also inconsistently referred to as the pseudo-eutectic point in the notation of the intersection of the content of the binary boundary system, or specifically the ternary, phase diagram. However, this type of nomenclature is not meant in the sense of the inventive concept, which is neither expressly intended nor implied by, or constituted by, the terminology "pseudo-eutectic point". Nor will it be done. Specifically, the pseudo-eutectic point is characterized in that its presence requires the addition or presence of at least a third component or element.

に対応する。この、三元共晶点Ｅに関する０の自由度と比較して増大した、擬共晶点ｐＥでの１の自由度は、共晶点で形成されるミクロ組織と比較してしばしば最大で数桁までより微細な、擬共晶点の領域における共晶ミクロ組織の実施形態の原因であると考えられる。 From the point of view of the Gibbs phase rule, the pseudo-eutectic point pE represents a local extremum along the liquidus boundary in a ternary alloy system, the extremum being one degree of freedom greater than the ternary eutectic E , and the number of degrees of freedom one less than monophasic solidification MC. According to the Gibbs phase law f=N−P+1, where f is the number of thermodynamic degrees of freedom, N is the number of components, and P is the number of equilibrium phases, this is

corresponds to This increased one degree of freedom at the pseudo-eutectic point pE compared to the zero degree of freedom with respect to the ternary eutectic point E is often at most a few degrees compared to the microstructure formed at the eutectic point. It is believed to account for the embodiment of the eutectic microstructure in the region of the pseudo-eutectic point, which is orders of magnitude finer.

Thus, for alloys with four components, the liquidus boundary corresponds to the two-dimensional area and the pseudo-eutectic point corresponds to the pseudo-eutectic line. For each higher component alloy with more than four components, the dimensionality of the relevant state regions increases as well. Therefore, within this document, the term "pseudo-eutectic point" refers to the pseudo-eutectic point in the phase diagram of a ternary alloy and the corresponding pseudo-eutectic line in the phase diagram of an alloy with four components, or It should be clearly understood as a general term meaning both the corresponding quasi-eutectic multidimensional regions in the phase diagram of alloys with more than four constituents. In this connection, therefore, in particular the terms “pseudo-eutectic point” and “pseudo-eutectic region” are used interchangeably. It is thereby to be understood that the pseudo-eutectic point of the ternary alloy system constitutes a specific embodiment.

は三元合金の状態図の擬共晶点に、または３つより多い成分の状態図の擬共晶点、具体的には擬共晶線もしくは擬共晶領域に当てはまる。 Thus, according to this description, specifically for the ternary alloy system, according to the Gibbs phase rule,

applies to the quasi-eutectic point of the phase diagram of a ternary alloy, or to the quasi-eutectic point of the phase diagram of more than three components, specifically to the quasi-eutectic line or region.

Thus, the pseudo-eutectic point, or the alloy composition at the pseudo-eutectic point of the phase diagram of an alloy having at least three components N, specifically follows the Gibbs phase law, with the number of degrees of freedom f being 0 and N− It is characterized by being between 1 and

The alloy composition is near the pseudo-eutectic point or in a field near the pseudo-eutectic point, specifically at the pseudo-eutectic point or at the saddle point representing said point, resulting in at least 85 mol % or at % It has been shown that the presence of a eutectic structure (expressed in mole percent and atomic percent, respectively) in the alloy is sufficient for both high strength and significant deformability of the alloy. It is preferred if at least 90 mol % or at %, particularly preferably at least 95 mol % or at % of the eutectic structure is present in the alloy. As a result, the advantageous properties of good deformability and at the same time high strength can be developed in a particularly pronounced way. Among other things, it is often possible to achieve up to 98 mol % or at %, as a result of which the mechanical properties of the alloy are determined virtually exclusively by the eutectic microstructure. A eutectic structure typically forms in a liquid-solid transformation or in solidification of an alloy.

High strength and significant deformability of the alloy are achievable both when the alloy is a ternary alloy and when the alloy comprises four components or at least five components. Specifically, the alloy may include multiple components, for example, in the form of other additional components for mixed crystal hardening and/or precipitation hardening, depending on the practical purpose. When the alloy is a ternary or quaternary alloy, it is particularly simple and feasible to embody an alloy with high strength and deformability.

The eutectic texture typically has an average characteristic texture spacing or phase content thereof, in particular an average lamellar spacing, of less than 3 μm. Particularly outstanding strength and deformability can thereby be achieved if the average spacing is less than 2 μm, in particular less than 1 μm. For example, this can be achieved if the alloy composition of the alloy is selected to be closer to the pseudo-eutectic stoichiometry. Particularly high intensities can thereby be achieved if the average spacing is less than 800 nm, in particular less than 600 nm. Additionally or alternatively, the average spacing of phase quantities can be affected by varying the cooling rate of the alloy while the alloy is solidifying.

It is advantageous if the alloy has residual solidification in an amount of at most 5 mol % or at %, preferably at most 3 mol % or at %, in particular preferably at most 2 mol % or at %. In this way, the aforementioned properties advantageously obtained by the eutectic structure are only slightly affected, or not at all, by the amount of residual solidification. The amount of residual solidification is set by selecting the alloy composition so that it is closer to the stoichiometric composition of the pseudo-eutectic point. Residual solidification typically occurs after the formation of the eutectic structure, where the residual amount of liquid phase solidifies in the form of a structure that is no longer eutectic, or the number or type of phases being formed is reduced by the end of eutectic solidification represents the amount of microstructure, which varies in The amount of residual solidification typically constitutes a limiting factor in the properties provided by the eutectic structure formed, and for that reason it is beneficial if residual solidification is kept as low as possible. Here, it is particularly beneficial if the residual solidification is not embodied in the form of a network or network, but rather preferably in the form of islands or units separated from each other, if present. be. Usually residual solidification is embodied in an amount of at least 1 mol % or at %, but preferably less.

It is advantageous for outstanding strength and deformability if the alloy has primary solidification in an amount of less than 10 mol % or at %, in particular less than 5 mol % or at %, preferably less than 3 mol % or at %. is. This allows a very dominant embodiment of the eutectic texture, or an embodiment of the eutectic texture with a high amount of texture, which accordingly can advantageously be achieved with the aforementioned properties. Primary solidification refers to the portion of the solidified microstructure that occurs immediately prior to the formation of the eutectic structure and does not solidify in the form of the eutectic structure, the primary solidification being the properties that can be achieved by embodiments of the eutectic structure. Although less relevant than the residual coagulation mentioned above in terms of limitations, this too should preferably be kept as low as possible. Generally, primary solidification is embodied in an amount of at least 1 mol % or at %, but preferably less.

For embodiments of high strength and particularly pronounced deformability, a primary solidification is formed which has or is made of a mixed crystal phase, in particular without or made of an intermetallic phase. useful if This appears to be an advantageous condition that applies to all alloy systems in order to achieve particularly practical strength and deformation properties.

The aforementioned amounts of residual solidification and/or primary solidification can be controlled or predetermined in conventional technical manner using thermodynamic calculations according to Schill-Gulliver. The Shile-Gulliver calculus or equation, also simply called the Shile calculus or equation, describes the distribution of alloy content in an alloy during solidification, where local equilibrium with respect to progressive solidification and negligence of diffusion in the solid phase are usually assumed. be done. Calculations of this type constitute routine technical instrumentation or textbook knowledge in the field of metallurgy and are presumed to be known to those skilled in the art. This is J. A. It is exemplified in the textbook "Solidification" by Dantzig et al. (ISBN: 978-2-940222-17-9).

It is advantageous if the alloy has a density of less than 8.0 g/cm ³ , in particular less than 7.5 g/cm ³ and preferably less than 6 g/cm ³ . Accordingly, the alloy may have a strength-to-weight ratio that is particularly advantageous for utility as a structural component. It is particularly beneficial if the alloy is embodied as a light metal alloy. A particularly high practical suitability of the alloy can thus be achieved. For this purpose it is advantageous if the alloy has a density of less than 5.0 g/cm ³ , in particular less than 3.0 g/cm ³ .

With respect to viable use as a practical material, it is advantageous if the alloy is a magnesium-based alloy, an aluminum-based alloy, a lithium-based alloy, or a titanium-based alloy.

It is advantageous if the alloy is a cast alloy. This allows a particularly viable production, in particular of structural parts, in particular having the aforementioned properties.

It has proven effective when the alloy is an Al--Mg alloy. The alloy may contain other alloying constituents, depending on the specific intended application. In this way, utility parts, in particular structural parts, which have particular relevance to practical situations, can be produced comprising or made from the alloy. It is particularly advantageous here if the alloy is an Al--Mg--Si alloy. Advantageously, the alloy may also contain zinc (Zn), specifically in an amount greater than 0.01 wt%, typically greater than 1 wt%. Therefore, the compressive strength of the alloy can be optimized. Thereby in most cases the alloy contains less than 15 wt%, in particular less than 10 wt%, preferably between 1.0 and 5.0 wt%, particularly preferably about 3.0 wt% zinc including.

The alloy is an Al-Cu-Li alloy, Al-Cu-Mg alloy, Mg-Li-Al alloy, Mg-Cu-Zn alloy, Al-Cu-Mg-Zn alloy, or Al-Mg-Si-Zn alloy In this case, a high practical suitability can be achieved, with both strength and pronounced deformability being particularly advantageous.

Alloys which exhibit particularly high strength and deformability and which have a high practical suitability, especially in the form of structural parts, are alloys containing 15.0% to 70.0% lithium, more than 0.0%, in particular More than 0.01%, preferably more than 0.05% aluminum, magnesium, and the balance comprising, and specifically made of, manufacturing-related impurities (in at %), the ratio of aluminum to magnesium ( at %) is between 1:6 and 4:6, which can be achieved in the case of magnesium-based alloys.

This type of Mg—Li—Al alloy has an alloy composition at or near the alloy composition in the pseudo-eutectic point in the Mg—Li—Al phase diagram. As a result, a finely textured or microscale eutectic microstructure can be achieved. The fine-scale microstructure is accompanied by high strength, in particular high compressive strength, and at the same time there is good deformability of the magnesium alloy at the corresponding aforementioned amounts of lithium in the magnesium alloy. The orientation composition or orientation line in the phase diagram is thereby specifically a ratio of aluminum to magnesium (in atomic percent, abbreviated as at%) of about 3:6, which is particularly uniform on the fine scale. Or because a uniform, fine lamellar, microstructure or morphology is found at this ratio.
In the range encompassing this ratio, especially at aluminum to magnesium ratios (in at%) of 1:6 to 4:6, a fine, particularly fine lamellar microstructure or morphology is also variably pronounced. are found to varying degrees, and the texture or morphology is usually associated with different significant magnitudes of strength, specifically compressive strength, and deformability or ductility of magnesium alloys. These special morphological features in the composition range described thus enable the formation of magnesium alloys having both high strength, specifically compressive strength, and good deformability. This magnesium alloy and the method for its manufacture, as well as raw materials, semi-finished products or parts, and their implementation as specific embodiments, are part of European Patent Application No. 19184999.1 and International Application PCT/ It has been filed with the European Patent Office and disclosed as part of EP 2020/058280, the disclosure of which is hereby incorporated in its entirety by the disclosure of this document. This is inter alia, as described in the above application, a Mg-Al-Li alloy of 30.0% to 60.0%, in particular 40% to 50%, preferably 45% to 50%, Particularly preferred is a case containing 45% to 48% lithium (at %). It is further advantageous if the Mg--Al--Li alloy contains more than 0.05%, in particular more than 0.1%, typically more than 1% aluminum (in at %). Thereby the Mg-Al-Li alloy has an aluminum to magnesium ratio (in at%) of 1.2:6 to 4:6, specifically 1.4:6 to 4:6, preferably 1.4:6 to 4:6. It has been shown that when it is 5:6 to 4:6 it can be embodied with a highly fine microstructure, in particular lamellar. a ratio of aluminum to magnesium (in at %) of 1.8:6 to 3.5:6, specifically 2:6 to 3.5:6, preferably 2.5:6 to 3.5:6 is advantageous for a pronounced fineness, or a fine, specifically lamellar, microstructure. Particularly high strengths, in particular compressive strengths, can thus be achieved. This is especially true at aluminum to magnesium ratios (at %) of 2.8:6 to 3.3:6, preferably about 3:6, where a very uniform fine morphology or microscopic You can get organization. To this end, the magnesium alloy (in at %) is 30.0% to 60.0% lithium and has a ratio of aluminum to magnesium (in at %) of 2.5:6 to 3.5. :6, in particular from 2.8:6 to 3.3:6, preferably about 3:6. In this regard, reference is made specifically to FIG. 1 of the aforementioned application document, in which the corresponding arrangement in the Mg—Li—Al phase diagram is schematically illustrated, the disclosure of which and related descriptions are It is also considered part of this document as appropriate. A particularly pronounced homogeneity is also achievable when the magnesium alloy thereby contains 40.0% to 60.0% lithium (in at %). As described in the aforementioned application, and where appropriate incorporated into this disclosure, the properties of Mg--Al--Li alloys are such that the amounts of calcium, rare earth metals, specifically yttrium, zinc, and/or silicon are It can be further optimized if additionally present according to the aforementioned application in the corresponding content range described in the application. For example, alloys of this type are Mg-20%Li-15%Al-1%Ca-0.5%Y (in wt%) or Mg-20%Li-24%Al-1%Ca-0.5 It can be embodied as %Y (in wt%).

Using methods of the type listed at the beginning, the composition lies in the field near the pseudo-eutectic point of the phase diagram of the alloy such that as the alloy cools or solidifies into the solid phase, the eutectic Other objects of the invention are achieved provided that the composition is provided such that the tissue is embodied in an amount of at least 85 mol % or at %. Thus, as discussed above, alloys can be embodied with high strength and significant deformability. Since the alloy composition is selected in a field near the pseudo-eutectic point, eutectic phase reactions or phase transformations occur as the alloy cools from the liquid state to the solid state or during the transition from liquid to solid. , the phase reaction or transformation embodies a eutectic microstructure, which has a particularly high degree of fineness, or fine structure, as the predominant microstructure content of the alloy.

The method according to the invention may be embodied within the alloy according to the invention, corresponding to or similar to the features, advantages, implementations and effects specifically described above. Good things should be understood. The same applies to the alloys according to the invention with respect to the process according to the invention.

The raw material, semi-finished product or part advantageously comprises an alloy, in particular is made out of it and realized. In accordance with the foregoing descriptions, features and advantages of the alloy according to the invention or of the alloy produced using the method according to the invention, raw materials, semi-finished products or parts formed with the alloy also have advantageously high strength. and have good deformability.

Additional features, advantages, and effects are obtained from the exemplary embodiments described below. In the drawings referenced thereby,

1 shows an illustration of a phase diagram for the Al—Mg—Si system with alloy compositions of exemplary alloys represented. 1 shows an illustration of a phase diagram for the Al—Mg—Si system with alloy compositions of exemplary alloys represented. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 3 shows optical microscope images of the exemplary alloys from FIGS. 1 and 2. FIG. 11 shows a yield stress diagram of the exemplary alloys from FIGS. 1, 2 and 10; FIG. FIG. 12 shows a yield stress diagram for the exemplary alloys from FIGS. 1, 2, and 11; FIG. 7 shows a yield stress diagram for the exemplary alloys from FIGS. 1, 2 and 6; FIG. FIG. 8 shows a yield stress diagram for the exemplary alloys from FIGS. 1, 2, and 7; FIG. 4 shows a yield stress diagram for the exemplary alloys from FIGS. 1 and 3; FIG. 5 shows a yield stress diagram for the exemplary alloys from FIGS. 1 and 4; FIG. 10 shows yield stress diagrams for the exemplary alloys from FIGS. 1, 2, 8, and 9; FIG. FIG. 13 shows a yield stress diagram for the exemplary alloys from FIGS. 1, 2, and 12; 1 shows an illustration of a phase diagram for the Al—Cu—Mg system with alloy compositions plotted for exemplary alloys. 22 shows an optical microscope image of the exemplary alloy from FIG. 21; FIG. 23 shows a yield stress diagram for the exemplary alloys from FIGS. 21 and 22; FIG. 1 shows an illustration of a phase diagram for the Mg—Al—Li system with alloy compositions of exemplary alloys represented. 25 shows an optical microscopy image of the exemplary alloy from FIG. 25 shows an optical microscopy image of the exemplary alloy from FIG. 25 shows an optical microscopy image of the exemplary alloy from FIG. FIG. 27 shows a yield stress diagram for the exemplary alloys from FIGS. 24 and 26; FIG. FIG. 28 shows a yield stress diagram for the exemplary alloys from FIGS. 24 and 27; FIG. 1 shows an illustration of a phase diagram for the Mg—Cu—Zn system with alloy compositions plotted for exemplary alloys. 31 shows an optical microscope image of the exemplary alloy from FIG. 30; 31 shows an optical microscope image of the exemplary alloy from FIG. 30; FIG. 33 shows a yield stress diagram for the exemplary alloys from FIGS. 30-32; FIG. 2 shows an electron microscope image of an exemplary alloy from the Al-Cu-Mg-Zn system. FIG. 35 shows a yield stress diagram for the exemplary alloy from FIG. 34; FIG. 2 shows a diagram of phase quantities of exemplary alloys from the Al—Mg—Si—Zn system. FIG. 37 shows a diagram of the solids content of the Shile-Gulliver calculation for the exemplary alloy from FIG.

として示され、それに応じて無次元の、変形度の関数としてＭＰａで降伏応力を描写する降伏曲線が、結果として算出された。 In the course of developing alloys according to the present invention, a series of tests were conducted using various alloy compositions of various alloy systems. Thereby, in each case, we select an alloy having an alloy composition at or near the pseudo-eutectic point field of the respective relevant phase diagram, and cool the alloy from the liquid state to the solid state by cooling the eutectic formed an organization. The microstructure was then examined using microscopy. In addition, various dilatometer test series and compression tests were performed at room temperature about 20° C. as a reference, where the change in length ΔL relative to the starting length L ₀ , i.e.

A yield curve plotting the yield stress in MPa as a function of the degree of deformation, denoted as , correspondingly dimensionless, was calculated as a result.

In the following, the alloy systems Al-Mg-Si, Al-Cu-Mg, Mg-Li-Al, Mg-Cu-Zn, Al-Cu-Mg-Zn, and Al- Test results for exemplary alloys from Mg--Si--Zn are shown in representative fashion.

Al-Mg-Si system:
Figures 1 and 2 show illustrations of the ternary phase diagram for the Al-Mg-Si system, and Figure 2 is an illustration of a segment from the phase diagram for the purpose of detailing the alloy composition range of interest. . Ten exemplary alloys from the Al-Mg-Si system were produced and studied. The alloy compositions of exemplary alloys from the Al-Mg-Si system are represented in Table 1 as Exemplary Alloy 1 through Exemplary Alloy 10 in weight percent and atomic percent, respectively, and correspond to reference numerals 1 through 10. , the numbers specifically represent the respective alloy compositions in the phase diagrams from FIGS.

As can be seen in the phase diagrams from FIGS. 1 and 2, exemplary alloys 8-10 each have a composition located in the field near the pseudo-eutectic point pE, where exemplary alloy 8 and 9 are positioned very close to the pseudo-eutectic point, and exemplary alloy 10 is positioned somewhat greater distance from the pseudo-eutectic point pE. Thereby, the alloy composition of exemplary alloy 9 is substantially at the pseudo-eutectic point pE. The pseudo-eutectic point pE is illustrated in FIG. 2 by the drawn reference line, where the pseudo-eutectic point pE is located at the intersection of the line of variation and the reference line in the direction of Al ₃ Mg ₂ is doing. It can also be seen in FIG. 2 that exemplary alloys 3-5 are located in the field near the eutectic point E of the phase diagram. Further, exemplary alloys 6 and 7 are provided as a comparison, and their compositions lie at a large distance from the pseudo-eutectic point pE, evident in FIG. , the alloy is positioned in the immediate vicinity of the liquidus boundary, but at a greater distance from both the pseudo-eutectic point pE and the eutectic point E, evident in FIG. .

Optical microscope images of exemplary alloys 1-10 are shown in FIGS. 3-12 to illustrate the respective microstructures. In FIGS. 13-20, yield stress diagrams are illustrated as a result of a dilatometric test series of an exemplary Al--Mg--Si alloy conducted at room temperature of about 20.degree. A yield stress curve is shown, where the yield stress in MPa is illustrated as a function of the degree of deformation. Each of the yield stress diagrams shows multiple yield stress curves from an alloy sample having an alloy composition corresponding to the alloy composition of one of Exemplary Alloys 1-10. Accordingly, each yield stress diagram represents the alloy composition of one of the exemplary alloys 1-10.

As can be seen in FIGS. 10-12, the microscopic images of exemplary alloys 8-10, which have an alloy composition in the field near or near the pseudo-eutectic point pE, are predominantly finely textured. It exhibits a eutectic texture on a fine scale. By comparison, microscopic images of exemplary alloys 4 and 5 can be seen in FIGS. 6 and 7, which alloys have alloy compositions near the eutectic point E. They exhibit a pronounced degree of eutectic texture, including a coarse texture, compared to the microstructures of exemplary alloys 8 and 9. Comparing these with the microscopic images of exemplary alloys 1 and 2 shown in FIGS. It can be seen to exhibit a coarse eutectic microstructure. Also shown in FIGS. 8 and 9 are microscopic images of exemplary alloys 6 and 7, which have an alloy composition in the far region of, or a large distance from, the pseudo-eutectic point pE. It can be seen that the eutectic texture is already present, but has a relatively coarse texture and is significantly less prevalent and less abundant. In addition, a high amount of residual coagulation is also evident, discernible in the form of thin channels in FIGS.

13-14 show yield stress diagrams for exemplary alloys 8 and 9 having alloy compositions in the field near or near the pseudo-eutectic point pE. It can be seen that both Exemplary Alloy 8 and Exemplary Alloy 9 have high strength, specifically compressive strength, and remarkable deformability with a yield stress between 300 and 400 MPa, here , illustrated in FIG. 13, notably exhibits a yield stress of up to 400 MPa. By comparison, the yield stress diagrams of exemplary alloys 4 and 5 can be seen in FIGS. 15 and 16, which alloys have alloy compositions near the eutectic point E. Exemplary alloys 4 and 5 also exhibit high strength and high deformability, at least subject to the individual specimens, where the yield stress is about 300 MPa, which is below that of exemplary alloys 8 and 9. , or consistently below with respect to exemplary alloy 5 illustrated in FIG. This result shows that an exemplary alloy with an alloy composition in the pseudo-eutectic pE field has its eutectic Correlating with the finding of exhibiting a particularly high degree of microtexturing of the structure, this also explains the higher strength and pronounced elasticity of the alloys in the field of the quasi-eutectic point.

In FIG. 20, the yield stress diagram of exemplary alloy 10 whose alloy composition is located at a somewhat greater distance from the pseudo-eutectic point pE is shown. Slightly lower yield stress values and especially greater variability between individual measurements are evident. 17 and 18, by comparison, exemplary alloy 1 and Exemplary alloy 2 is further shown to have significantly worse strength and deformation properties. In FIG. 19, yield stress diagrams corresponding to the alloy compositions of exemplary alloys 6 and 7 are additionally shown, whose alloy compositions are positioned at a relatively large distance from the alloy composition at the pseudo-eutectic point pE in the phase diagram. be The corresponding yield stress curves show a clearly reduced yield stress compared to that of alloy compositions closer to the pseudo-eutectic point pE, such as the yield stress curve shown in FIG. 13 for exemplary alloy 8.

It is clear that alloy compositions in the field near the pseudo-eutectic point pE correspond to a finely textured eutectic microstructure and correspondingly high strength and pronounced deformability.

In the detailed diagram, it can be seen that with respect to the line of variation or liquidus boundary that is in the _Al3Mg2 direction, the exemplary alloy 8 in the phase diagram from _FIG . ₂ is above the line in the Mg2Si region. It can be seen that for this reason, as the alloy cools from the liquid phase, solidification begins specifically with the undesirable formation of Mg ₂ Si, or primary solidification with an intermetallic Mg ₂ Si phase. It is formed. Primary solidification formed with intermetallic phases has been shown to have an adverse effect on both high strength and deformability embodiments. In order to achieve particularly advantageous strength and deformability, therefore, one generally seeks to keep or prevent primary solidification with or from intermetallic phases as little as possible. However, the primary solidification for Exemplary Alloy 8 is so slightly pronounced that it imposes no practical limit on mechanical properties. The microscopic image of exemplary alloy 8 in FIG. 10 shows extensive regions with fine eutectic texture, in this case formed with Al mixed crystal phase and Mg ₂ Si. Advantageously, residual solidification of the Al mixed crystal phase is also only slightly pronounced or almost non-existent. An effort is made to keep residual solidification as low as possible or to prevent it without impairing the advantageous strength and deformation properties achievable with a eutectic structure. Specifically, the residual solidification is embodied in the form of units that are not connected in a network or separated from each other, which likewise leads to advantageous embodiments of high strength and pronounced deformability. Increase. Exemplary Alloy 8 is thus confirmed to be well suited for controlling strength properties and deformability with both low residual solidification and low primary solidification based on the fine eutectic structure. This is true if the alloy composition is chosen such that it has or is made of a mixed crystal phase and a primary solidification with no intermetallics or phases is formed, i.e. for exemplary alloy 8, the primary It can be even more optimized if the solidification is located in the Al mixed crystal phase region.

This figure, and the accompanying description, of Exemplary Alloy 8 applies to Exemplary Alloy 9 as well. Exemplary alloy 9 has an alloy composition substantially at the pseudo-eutectic point pE. As can be seen in FIG. 11, Exemplary Alloy 9 also exhibits a fine eutectic texture with little residual solidification and primary solidification. The somewhat lower strength compared to Exemplary Alloy 8 is explained by the lower dissolution of Mg in the Al mixed crystal phase. Strength can be advantageously achieved by varying the amount of dissolved elements in the mixed crystal phase, but as noted above, primary solidification is preferably in the mixed crystal rather than in the intermetallic phase region. in the area.

By comparison, as can be seen in FIG. 12, Exemplary Alloy 10 also exhibits a fine eutectic structure, but with a greater amount of residual solidification in the form of the Al mixed crystal phase and Si, and also the residual solidification. It has a mesh shape. Due to the low Mg content, most of the Mg is bound in the form of Mg ₂ Si and as a result the mixed crystal hardening of the Al mixed crystal phase is very slightly pronounced. This corresponds to a lower yield stress in the yield stress diagram from FIG.

In a more detailed view of an exemplary alloy spaced from the pseudo-eutectic point pE in relation to alloy composition, an exemplary alloy in the field of the eutectic point E, illustrated in FIGS. It can be seen that 4 and 5 contain a low amount of primary solidification with a relatively coarse eutectic structure formed with two phases located near the primary solidification. The remaining predominant amount of eutectic structure is embodied as a ternary eutectic formed with the mixed crystal phases Al ₂ Si and Si. Mechanical properties, specifically strength and deformability, are adversely affected, specifically by a coarse binary eutectic structure or phase. A fine eutectic ternary texture is present to some extent locally, and the texture transitions to a significantly coarser texture at some locations. The difference between the microstructures of exemplary alloys having an alloy composition at or in a field of the pseudo-eutectic point pE compared to an alloy composition at or in a field of the eutectic point E is that Correlating with the finding of correspondingly improved strength and deformation properties of alloy compositions in fields at or near the pseudo-eutectic point pE.

In the associated microscopic images shown in FIGS. 8 and 9, it can be further seen that exemplary alloys 6 and 7 contain rough, polygonal primary solidifications. This is explained by the positioning of the relevant alloy compositions in the Mg ₂ Si region of the phase diagram, which results in the formation of significant Mg ₂ Si primary solidification. Between the coarse eutectic texture is also discernible a high amount of residual solidification, which is evident from the thin areas or channels in FIGS. Due to this organizational morphology, exemplary alloys 6 and 7 exhibit significantly reduced strength and yield stress, specifically associated with crack initiation and brittle fracture.

In FIG. 2, regions of particular interest for the Al—Mg—Si alloy embodiment are depicted as gray flat areas. This essentially refers to, or corresponds to, the aforementioned alloy compositions of exemplary alloys 8 and 9, but with mixed crystal phases embodied as primary solidification and specifically no intermetallic phases. It has a variation of alloy composition such as This allows for particularly high strength embodiments with pronounced deformability. A particularly advantageous range of implementation for this type of Al--Mg--Si alloy is therefore guaranteed when the Al--Mg--Si alloy is placed in a field near the pseudo-eutectic point in the Al--Mg--Si phase diagram. , where the alloy compositions in the phase diagram are placed on the side of the corresponding line of variation starting from the pseudo-eutectic point of the phase diagram of FIG. 2 above and pointing in the direction of increasing Al content.

Al-Cu-Mg system:
FIG. 21 shows an illustration of the ternary phase diagram for the Al—Cu—Mg system. One exemplary alloy from the Al-Cu-Mg system was produced and investigated. The relevant alloy composition, in weight percent and atomic percent, is represented as exemplary alloy 13 in Table 2 and corresponds to reference number 13, which specifically represents the alloy composition in the phase diagram from FIG. .

As can be seen in the phase diagram from FIG. 21, exemplary alloy 13 has an alloy composition located in the field near the pseudo-eutectic point pE. The associated microstructure is illustrated in FIG. 22 using optical microscope images. A very fine scale eutectic microstructure and a low amount of primary solidification formed with mixed crystal phases are evident. In FIG. 23, a yield stress diagram is shown as a result of a dilatometric test series for the Al—Cu—Mg exemplary alloy 13, where, as before, the yield stress in MPa is illustrated as a function of the degree of deformation. ing. It is clear that very high strength and yield stress are achieved. The elongation at break is also in the technically relevant range for this alloy system. Strength and deformability correspond to a fine eutectic microstructure and specifically to a low amount of primary solidification.

Mg-Al-Li system:
FIG. 24 shows an illustration of the ternary phase diagram for the Mg—Al—Li system. Three exemplary alloys from the Mg-Al-Li system were produced and investigated. The alloy compositions of exemplary alloys from the Mg—Al—Li system, in weight percent and atomic percent, respectively, are represented as exemplary alloys 14, 15, and 16 in Table 3, corresponding to reference numerals 14, 15, and 16. 24, and the numbers specifically represent the respective alloy compositions in the phase diagram from FIG.

As can be seen in the phase diagram from FIG. 24, exemplary alloys 14-16 each have an alloy composition located in the field near the pseudo-eutectic point pE. The pseudo-eutectic point pE is illustrated in FIG. 24 by the reference line drawn, where the pseudo-eutectic point pE is located at the intersection of the first line or liquidus boundary line and the reference line. there is The addition of CaY, specifically about 1 wt. can be feasibly stabilized without

In the phase diagram, exemplary alloys 14 and 15 are somewhat closer in the vicinity of the pseudo-eutectic point, and exemplary alloy 16 is somewhat further, where the alloy composition of exemplary alloy 14 is approximately It is positioned at the pseudo-eutectic point pE. According to currently available data, exemplary alloys 14-16 exist in the mixed crystal regime, specifically such that they form a body-centered cubic lattice bcc.

In Figures 25-27, the microstructures are each visualized using microscopic images. The tissue morphologies from FIGS. 25 and 26 show very fine-scale tissue embodiments that can no longer be resolved by optical microscopy. Grain boundaries recognizable thereby can be attributed to oxide impurities. The microstructure of exemplary alloy 16 was studied using scanning electron microscopy and is illustrated in FIG. Evident in FIG. 27 are, on the one hand, a thin grain boundary phase (whitish gray) identified as Al—Ca, and on the other hand, in the region surrounded by the grain boundary phase, specifically the center of said region. A pronounced microcrystalline texture or morphology in the section or within the mixed crystal phase, specifically visible in the image on the right from FIG. In the phase diagram from FIG. 24, the alloy composition of exemplary alloy 16 appears to be relatively far away from the first line and the pseudo-eutectic point pE. However, in this case, according to established technical knowledge, the gradient in the region of the body-centered cubic lattice bcc in the phase diagram, in which the exemplary alloy 16 is also located, is very flat and the three elements Mg , Al, and Li also exhibit high solubility in each other. Thus, it is possible to explain why such a wide field results that a high amount of advantageous fine-scale eutectic microstructure near the pseudo-eutectic point can be embodied.

28 and 29 show yield stress diagrams for exemplary alloys 15 and 16 as a result of a dilatometer test series, where, as before, yield stress in MPa is illustrated as a function of degree of deformation. 28 shows the yield stress curve for exemplary alloy 15, and FIG. 29 shows the yield stress curve for exemplary alloy 16. FIG. It is clear that both exemplary alloys have high strength and yield stress, as well as pronounced deformability, corresponding to the identified fine eutectic microstructures. The potential for further property optimization by heat treatment is also illustrated in FIG. 29 for exemplary alloy 16. FIG.

FIG. 29 depicts the yield curve of the alloy sample as-manufactured (as a casting) of Exemplary Alloy 16, depicted in solid line in FIG. 16-2 depicts the yield curve of an exemplary alloy sample immediately after heat treatment (of aging) of exemplary alloy 16, depicted by reference number 16-2. For this purpose, a sample of Exemplary Alloy 16 was subjected to a heat treatment at 330° C. for 3 hours and then a compression test was used to calculate the yield curve. A distinct effect of heat treatment on strength, specifically compressive strength, and deformability is evident, with the result that heat treatment is used to specifically improve compressive strength and deformability in an optimized manner. There is a setting potential for the intended use to obtain.

As explained earlier in this document, the alloy contains 15% to 70.0% lithium, more than 0.0%, in particular more than 0.01%, preferably more than 0.05% Containing and specifically consisting of (in at %) high amounts of aluminum, magnesium, and the balance manufacturing-related impurities, with a ratio of aluminum to magnesium (in at %) of 1:6 to 4:6 , in the case of magnesium-based alloys, has been found to be beneficial in achieving alloys with high practical suitability. Exemplary Alloy 16, as set forth within European Patent Application No. 19184999.1 and International Application PCT/EP2020/058280, both filed at the European Patent Office, is defined as It can be seen as a representative example. Here again, specific reference is made to FIG. 1 from each of these applications. In FIG. 24, the corresponding aluminum to magnesium ratio (in at %) of 1:6 is depicted with a dashed line. Thereby, the aforesaid aluminum to magnesium ratio range (in at % or mol %) of 1:6 to 4:6 is located to the left of this line in the phase diagram from FIG. Specifically, fields near the pseudo-eutectic point pE constitute a particular embodiment.

A particularly advantageous working range for Mg—Li—Al alloys that can be used as practical alloys, in particular for structural parts, is that the Mg—Li—Al alloys have a phase diagram of 1 :6 in the region between the line indicating the aluminum to magnesium ratio (in at%) and the first line or liquidus boundary line, specifically having the aforementioned Li content range, Guaranteed. This type of range is represented as a gray flat area in the state diagram from FIG.

the alloy composition is in the field of the pseudo-eutectic point pE, as was already the case previously within the exemplary alloys from the Al—Si—Mg system, and additionally preferably has mixed crystal phases, Or it becomes clear that the alloy composition is preferably chosen such that it contains a primary solidification made of it, ie the corresponding alloy composition is positioned in the mixed crystal region in the phase diagram.

Mg—Cu—Zn:
FIG. 30 shows an illustration of the ternary phase diagram for the Mg—Cu—Zn system. One exemplary alloy from the Mg--Cu--Zn system was produced and investigated. The relevant alloy composition, in weight percent and atomic percent, is represented as exemplary alloy 17 in Table 4 and corresponds to reference numeral 17, which specifically represents the alloy composition in the phase diagram from FIG.

As can be seen in the phase diagram from FIG. 30, exemplary alloy 17 has an alloy composition located in the field near the pseudo-eutectic point pE. The associated microstructure is illustrated in FIGS. 31 and 32 using optical microscope images. A very fine-scale eutectic microstructure at the limit of optical microscope resolution is evident. Here a relatively large amount of primary solidification can be seen. Therefore, it is advantageous with respect to high strength and deformability if an alloy composition even closer to the pseudo-eutectic point pE or closer to the first line or liquidus boundary line is chosen.

In FIG. 33, a yield stress diagram is shown as a result of a dilatometer test series for exemplary alloy 17, where, as before, yield stress in MPa is illustrated as a function of degree of deformation. Based on the significant amount of primary solidification clearly seen in the microscopic images, however, it is found that high strength and yield stress are achieved which can be further improved by choosing an alloy composition closer to the pseudo-eutectic point pE. it is obvious. In FIG. 33, there is shown the yield curve of Exemplary Alloy 17 immediately after manufacture (as a casting), denoted by reference numeral 17-1, and denoted by reference numeral 17-2. , the yield curve of exemplary alloy 17 after heat treatment is also shown. For this purpose, a sample of Exemplary Alloy 17 was subjected to heat treatment at 350° C. for 4 hours and then the yield curve was calculated by compression test. A clear impact of heat treatment on strength and deformability is evident and as a result there is potential to further optimize strength and deformability using heat treatment.

A study of quaternary alloy systems and quaternary eutectics was then also carried out. For this purpose, the alloy system Al--Cu--Mg--Zn and specifically Al--Mg--Si--Zn was considered.

Al—Cu—Mg—Zn:

For the alloy system Al--Cu--Mg--Zn, exemplary alloys in the field of pseudo-eutectic point pE were produced and studied. The alloy composition, in weight percent and atomic percent, is represented as Exemplary Alloy 18 in Table 5 and corresponds to reference numeral 18.

To study the eutectic microstructure, an electron microscope image of exemplary Alloy 18, shown in FIG. 34, was recorded. Specifically, there is a finely textured eutectic structure with texture dimensions in the nanometer range, clearly visible on the right side of the image from FIG. it is obvious.

This is the binary eutectic structure in a system with four components or elements and thus the increase in the thermodynamic degrees of freedom f mentioned at the outset, between 1 and 3 (quaternary eutectic).

In FIG. 34 substructures are discernible in the primary region (grey), where these are artifacts of the iso-stoichiometric structural transformation (bcc to fcc) in the solid state. They are negligible in terms of direct impact on strength and deformability. Relatively large amounts of primary solidification in the form of mixed crystal phases (light gray to whitish in color) and intermetallic secondary phases (black) in particular in the form of Laves phases can also be seen.

FIG. 35 shows a yield stress diagram as a result of a dilatometric test series for exemplary alloy 18. FIG. The yield curve before the heat treatment is completed, denoted by reference number 18-1, and the yield curve after the heat treatment is completed, denoted by reference number 18-2 are depicted, where the same as before. As such, the yield stress in MPa is illustrated as a function of the degree of deformation. Exemplary alloy 18 clearly exhibits very high strength with concurrent elongation at break, where deformability can be varied using heat treatment.

The significant primary solidification and secondary phases present should be considered brittle increasing factors, and therefore the amount of these should be reduced, e.g. or by bringing the alloy composition closer to the pseudo-eutectic point pE in the phase diagram.

Al—Mg—Si—Zn:
For the alloy system Al--Mg--Si--Zn, an exemplary alloy in the field of pseudo-eutectic point pE was studied by simulation. The alloy composition is represented as Exemplary Alloy 19 in Table 6 and corresponds to reference number 19.

As a result of the simulation, phase amounts are illustrated in FIG. 36 as a function of temperature for exemplary alloy 19. A direct solid-to-liquid transition is evident, corresponding to the embodiment of the eutectic structure. In FIG. 35 there is shown a corresponding illustration of the amount of solids as a function of temperature, determined using the Shile-Gulliver solidification calculation. The equilibrium and Schil-Gulliver solidification curves shown depict an alloy system exhibiting binary eutectic solidification with four components or elements. Correspondingly, therefore, as before, there is an increase in the thermodynamic degrees of freedom f from 1 to 3. In FIG. 37, the Shile-Gulliver calculations show a very small amount of primary solidification in the form of mixed crystal phases of less than 3 mol % or at % and in addition virtually no residual solidification.

Therefore, in addition to positioning the alloy composition in the field of the pseudo-eutectic pE, the amount of primary solidification and/or residual solidification is advantageously minimized to further increase or improve strength and deformation properties. It is equally clear that it is also possible to

Therefore, alloys according to the present invention having more than three components, which have a eutectic structure created by cooling from the liquid state to the solid state, have a finely textured eutectic structure, specifically can be advantageously embodied with a microstructure in the nanometer range, when the alloy composition of the alloy is located in the field at or near the pseudo-eutectic point in the phase diagram , constitutes the predominant or major phase or structure content in the alloy. Accordingly, the alloy can be embodied with advantageously high strength and significant deformability. This is especially true when embodied with very little primary and/or residual coagulation. In particular it is beneficial if a primary solidification is formed which has or is made of a mixed crystal phase, in particular without or made of an intermetallic phase, or where the alloy composition is selected in the corresponding region in the phase diagram. The alloys thus formed are therefore preferably adapted for specific purposes into strong and resilient members, in particular structural members, in particular in the automotive, aerospace and/or aerospace industries. It offers realization potential for its intended application in industry.
[Item 1]
An alloy having an alloy composition having at least three components and a eutectic structure, specifically a light metal alloy, wherein the eutectic structure is such that the composition of the alloy is above the pseudo-eutectic point (pE ), obtained by cooling said alloy from a liquid state to a solid state under conditions in which at least 85 mol % of the eutectic structure is present in said alloy.
[Item 2]
An alloy according to item 1, characterized in that the eutectic structure has an average spacing of phase quantities of the eutectic structure of less than 3 µm, preferably less than 1 µm.
[Item 3]
3. Alloy according to item 1 or 2, characterized in that it contains residual solidification in an amount of at most 5 mol %, preferably at most 3 mol %.
[Item 4]
An alloy according to any one of items 1 to 3, characterized in that it comprises primary solidification in an amount of less than 10 mol %, in particular less than 5 mol %.
[Item 5]
5. The alloy of item 4, wherein the primary solidification is formed with a mixed crystal phase.
[Item 6]
An alloy according to any one of items 1 to 5, characterized in that it has a density of less than 8 g/cm ³ .
[Item 7]
An alloy according to any one of items 1 to 6, characterized in that it is a magnesium-based alloy, an aluminum-based alloy, a lithium-based alloy or a titanium-based alloy.
[Item 8]
Alloy according to any one of items 1 to 7, characterized in that it is a cast alloy.
[Item 9]
Alloy according to any one of items 1 to 8, characterized in that it is an Al-Mg-Si alloy.
[Item 10]
10. Alloy according to item 9, characterized in that the Al-Mg-Si alloy contains between 0.01 wt% and 5.0 wt% zinc, in particular about 3.0 wt%.
[Item 11]
A method for producing an alloy having a eutectic structure, in particular an alloy according to any one of items 1 to 10, said alloy having an alloy composition having at least three components , the alloy is such that the alloy composition is in a field near the pseudo-eutectic point (pE) of the phase diagram of the alloy, resulting in the eutectic structure in an amount of at least 85 mol % during cooling to the solid phase. starting from the liquid state of the alloy and cooling to the solid state to form the eutectic structure under conditions provided with the alloy composition as embodied.
[Item 12]
A raw material, semi-finished product or structural material comprising the alloy according to any one of items 1 to 10 or obtainable using the method according to item 11.

Claims (11)

An alloy having an alloy composition having at least three components and a eutectic structure, in particular a light metal alloy, wherein said eutectic structure is such that said alloy composition is above the pseudo-eutectic point (pE ), obtained by cooling said alloy from a liquid state to a solid state under conditions in which at least 85 mol % of the eutectic structure is present in said alloy. 2. An alloy according to claim 1, characterized in that the eutectic structure has an average spacing of phase quantities of the eutectic structure of less than 3 [mu]m, preferably less than 1 [mu]m. 3. An alloy according to claim 1 or 2, characterized in that it contains residual solidification in an amount of at most 5 mol %, preferably at most 3 mol %. An alloy according to any one of the preceding claims, characterized in that it comprises primary solidification in an amount of less than 10 mol%, in particular less than 5 mol%. 5. The alloy of claim 4, wherein said primary solidification is formed with a mixed crystal phase. An alloy according to any one of the preceding claims, characterized in that it has a density of less than 8 g/cm ³ . An alloy according to any one of the preceding claims, characterized in that it is a magnesium-based alloy, an aluminum-based alloy, a lithium-based alloy or a titanium-based alloy. An alloy according to any one of the preceding claims, characterized in that it is an Al-Mg-Si alloy. 9. An alloy according to claim 8, characterized in that the Al-Mg-Si alloy contains between 0.01 wt% and 5.0 wt% zinc, in particular about 3.0 wt%. A method for producing an alloy, in particular an alloy according to any one of claims 1 to 9, having a eutectic structure, said alloy having an alloy composition having at least three components and said alloy is such that said alloy composition is in a field near the pseudo-eutectic point (pE) of said phase diagram of said alloy, resulting in said eutectic structure being at least 85 mol % during cooling to the solid phase. starting from a liquid state of said alloy and cooling to a solid state to form said eutectic structure under conditions provided said alloy composition as embodied in. Raw material, semi-finished product or structural material comprising the alloy according to any one of claims 1 to 9 or obtainable using the method according to claim 10.

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