ALLOY HAVING FINE-SCALE EUTECTIC, IN PARTICULAR NANOEUTECTIC, STRUCTURE AND PRODUCTION OF SUCH AN ALLOY

The invention relates to an alloy, in particular a light metal alloy, having an alloy composition with at least three components and a eutectic structure that is obtained by cooling the alloy from a liquid state to a solid state, under the condition that a composition of the alloy lies in a field around a pseudoeutectic point (pE) of a phase diagram of the alloy, so that at least 85 mol % eutectic structure is present in the alloy. The alloy also relates to a method for producing an alloy of this type.

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Description

The invention relates to an alloy, in particular a light metal alloy, having an alloy composition with at least three components and a eutectic structure that is obtained by cooling the alloy from a liquid state to a solid state.

The invention furthermore relates to a method for producing an alloy, in particular a light metal alloy, with a eutectic structure, wherein the alloy has an alloy composition with at least three components and wherein the alloy is cooled, starting from a liquid state, to a solid state of the alloy in order to form the eutectic structure.

It is known that it can be advantageous if a portion of a structure of an alloy is embodied with a eutectic structure in order to influence casting properties or strength properties of an alloy. Binary casting alloys, that is, alloys with two components which have eutectic microstructures, are often used as technical application alloys. These alloys are normally characterized by a eutectic point in their phase diagram, at which point a liquid phase of the alloy and two solid phases of the alloy are in thermodynamic equilibrium with one another, or at which a direct transition from a liquid state to a solid state takes place when the alloy is cooled from the liquid phase, wherein a eutectic structure is formed. According to Gibbs' phase rule for solids at a constant pressure f=N−P+1 with a number of thermodynamic degrees of freedom f, a number of components N, and a number of equilibrium phases P, this corresponds to a number of degrees of freedom of f=0. The direct transition from liquid phase to solid phase thereby often results in a formation of a fine and lamellar structure.

Analogously, in relation to ternary alloy systems, attempts to create alloys with compositions close to a ternary eutectic point have also become known, in order to improve strength properties with an embodiment of a eutectic structure. According to Gibbs' phase rule f=N−P+1, this analogously likewise corresponds, with three components and four phases, to f=0 degrees of freedom. However, high cooling rates are normally required to form alloys of this type, in order to create a eutectic structure with pronounced fineness at alloy amounts which can be used in applications, and a coordinated additional combination with other elements for precipitation hardening is often necessary to increase a strength of the alloy. In most cases, cooling rates in a range of 50 K/s to 200 K/s are used for this purpose. However, the requirement of high cooling rates in particular limits a technical usability of alloys of this type to small-scale parts.

This is addressed by the invention. The object of the invention is to specify an alloy having at least three components that has high strength and good deformability.

A further goal of the invention is to specify a method for producing an alloy of this type.

The object is attained according to the invention if, with an alloy of the type named at the outset, the eutectic structure is obtained under the condition that a composition of the alloy lies in a field around a pseudoeutectic point of a phase diagram of the alloy, so that at least 85 mol % or at % eutectic structure is present in the alloy.

The basis of the invention is the finding that, with a composition of an alloy having at least three components or elements that are at or in proximity to a pseudoeutectic point in the phase diagram of the alloy, a particularly fine-scale or finely structured eutectic structure can be embodied which, in particular, can have a finer eutectic structure than an alloy having a composition chosen that is at the “usual” eutectic point in the phase diagram. In particular, characteristic structural spacing of the eutectic structure in the low micrometer range and specifically in the nanometer range can thus be realized, also referred to as a nanoeutectic structure. In addition, it has been shown that the eutectic structure thereby formed normally constitutes a principal or dominant microstructure, and in particular that often only a small or negligibly small primary solidification phase and/or residual solidification phase, or none at all, occurs in proximity to or in a field around, in particular at, a eutectic point. This combination of an extraordinarily fine microstructure of the eutectic structure and the dominant presence thereof in the alloy enables the alloy to be embodied with both high strength, in particular compressive strength, and also pronounced deformability. In representations, the pseudoeutectic point is typically denoted such that it is abbreviated by “e” or “pE” and the eutectic point abbreviated by “E”.

Technically, in a ternary phase diagram, the liquidus line and solidus line known from the binary phase diagram typically respectively correspond to curved surface areas and binary phase areas correspond to phase volumes. In the ternary phase diagram, the intersecting lines of liquidus areas typically form eutectic channels, also referred to as liquidus boundary lines or monovariant lines, which end in a ternary eutectic point of the phase diagram. The pseudoeutectic point thereby represents a point on the liquidus boundary line that forms a saddle point, that is, represents a local extreme along the liquidus boundary line and a minimum perpendicular thereto—in relation to the bounding single-phase fields.

Sometimes, in representations of two-component boundary systems or content intersections of, in particular ternary, phase diagrams, binary eutectics are also inconsistently referred to as pseudoeutectic points. However, a terminological designation of this type is not in the sense of the present concept, and it is explicitly not intended or signified by, nor is it comprised by, the nomenclature “pseudoeutectic point” in this document. In particular, the pseudoeutectic point is characterized in that the existence thereof requires an addition or a presence of at least a third component or a third element.

In terms of the Gibbs' phase rule, the pseudoeutectic point pE represents in the ternary alloy system a local extreme along the liquidus boundary line, which extreme has a number of degrees of freedom that is 1 greater than the ternary eutectic E and a number of degrees of freedom that is 1 less than a single-phase solidification MC. With the Gibbs' phase rule f=N−P+1 with a number of thermodynamic degrees of freedom f, a number of components N and a number of equilibrium phases P, this corresponds to:


f(E)=3−4+1=0


f(pE)=3−3+1=1


f(MC)=3−2+1=2

This increased degree of freedom of 1 at the pseudoeutectic point pE compared to the degree of freedom of 0 for the ternary eutectic point E is considered to be the cause of the embodiment of the finer, often by up to several orders of magnitude, eutectic microstructure in the region of the pseudoeutectic point compared to the microstructure formed at the eutectic point.

Accordingly, for an alloy having four components, the liquidus boundary line corresponds to a two-dimensional area and the pseudoeutectic point corresponds to a pseudoeutectic line. For respective higher-component alloys having more than four components, a dimensionality of associated state regions increases analogously. Within the scope of this document, the designation “pseudoeutectic point” is therefore specifically to be understood as a general term that signifies both a pseudoeutectic point in a phase diagram of a ternary alloy and also a corresponding pseudoeutectic line in a phase diagram of an alloy having four components, or a corresponding pseudoeutectic multi-dimensional region in a phase diagram of an alloy having more than four components. Thus, in this regard, the designations “pseudoeutectic point” and “pseudoeutectic region” in particular are used synonymously. It should be understood that a pseudoeutectic point of a ternary alloy system thereby constitutes a specific embodiment.

In accordance with this explanation, in particular for the ternary alloy system, the following thus holds true for a pseudoeutectic point of a phase diagram of a ternary alloy or for a pseudoeutectic point, in particular a pseudoeutectic line or a pseudoeutectic region, of a phase diagram with more than three components, according to Gibbs' phase rule:


f(E)<f(pE)<f(MC),


therefore:


0<f(pE)<N−1.

The alloy composition at the pseudoeutectic point or the pseudoeutectic point of the phase diagram of the alloy having at least three components N is thus in particular characterized in that, according to Gibbs' phase rule, the number of degrees of freedom f lies between 0 and N−1.

It has been shown that it is sufficient for both high strength and also for pronounced deformability of the alloy if the alloy composition lies in proximity to or in a field around, in particular at, the pseudoeutectic point or the saddle point representing said point, so that at least 85 mol % or at % (stated in mole percent and atomic percent, respectively) eutectic structure is present in the alloy. It is preferable if at least 90 mol % or at %, particularly preferably at least 95 mol % or at %, eutectic structure is present in the alloy. As a result, the advantageous properties of high strength with simultaneously good deformability can be developed in a particularly pronounced manner. Specifically, up to 98 mol % or at % can often be attained thereby, so that the mechanical properties of the alloy are virtually solely determined by the eutectic microstructure. The eutectic structure typically forms in a liquid-solid phase transformation or in a solidification of the alloy.

High strength and pronounced deformability of the alloy are achievable both if the alloy is a ternary alloy and also if the alloy comprises four components or at least five components. In particular, the alloy can comprise a plurality of components, for example in the form of other added components for mixed crystal hardening and/or precipitation hardening, depending on the application objective. It is particularly simple and feasible to embody the alloy with high strength and deformability if the alloy is a ternary alloy or quaternary alloy.

The eutectic structure normally has average characteristic structural spacing or an average spacing of the phase amounts thereof, in particular lamellae, of less than 3 μm. A particularly pronounced strength and deformability can be achieved if the average spacing is thereby less than 2 μm, in particular less than 1 μm. This can be achieved, for example, if the alloy composition of the alloy is selected such that it is in greater proximity to the stoichiometric composition of the pseudoeutectic point. Particularly high strengths 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 the phase amounts can be influenced by varying a cooling speed of the alloy during the solidification of the alloy.

It is advantageous if the alloy has a residual solidification at an amount of maximally 5 mol % or at %, preferably maximally 3 mol % or at %, in particular preferably maximally 2 mol % or at %. In this manner, the aforementioned properties advantageously obtained with the eutectic structure are influenced only insignificantly or not at all by the amount of residual solidification. The residual solidification amount can be set by selecting the alloy composition such that it is in greater proximity to the stoichiometric composition of the pseudoeutectic point. Residual solidification typically denotes that microstructure amount in which, after the formation of the eutectic structure, a residual amount of liquid phase solidifies in the form of a structure that is no longer eutectic, or the number or type of the phases being formed changes at the end of the eutectic solidification. The amount of residual solidification typically constitutes a factor which limits the properties effected by the eutectic structure formed, for which reason it is beneficial if the residual solidification is kept as small as possible. Here, it is particularly beneficial if the residual solidification is not embodied in a network-like manner, or with a form of a network structure, but rather is preferably embodied, where present, in the form of islands or units separated from one another. Normally, the residual solidification is embodied at an amount of at least 1 mol % or at %, but can preferably also be smaller.

It is expedient for a pronounced strength and deformability if the alloy has a primary solidification at an amount less than 10 mol % or at %, in particular less than 5 mol % or at %, preferably less than 3 mol % or at %. This enables a very dominant embodiment of the eutectic structure, or an embodiment of the eutectic structure at a high structural amount with aforementioned properties that can accordingly be advantageously achieved. The primary solidification, which denotes that part of the solidified microstructure which, immediately preceding the formation of the eutectic structure, does not solidify in the form of a eutectic structure, is, with regard to a limitation of the properties that can be attained with the embodiment of the eutectic structure, less relevant than the aforementioned residual solidification, but should also preferably be kept as small as possible. Normally, the primary solidification is embodied at an amount of at least 1 mol % or at %, but can preferably also be smaller.

For an embodiment of high strength and particularly pronounced deformability, it is beneficial if the primary solidification is formed having or being made of a mixed crystal phase and, in particular, not having or not being made of intermetallic phase. This appears to be an advantageous criterion that applies to all alloy systems, in order to achieve particularly application-friendly strength and deformability properties.

The amount of aforementioned residual solidification and/or primary solidification can, in a customary technical manner, be controlled or predetermined using a thermodynamic calculation according to Scheil-Gulliver. The Scheil-Gulliver calculation or equation, sometimes also simply referred to as the Scheil calculation or equation, describes a distribution of an alloy amount in an alloy during a solidification, wherein a local equilibrium on a progressing solidification front and a disregarded diffusion in solid phase are normally assumed. A calculation of this type constitutes a customary technical tool or textbook knowledge in the field of metallurgy and is presumed to be known to a person skilled in the art. This is exemplified in the textbook “Solidification” by J. A. Dantzig et al., (ISBN: 978-2-940222-17-9).

It is beneficial if the alloy has a density of less than 8.0 g/cm3, in particular less than 7.5 g/cm3, preferably less than 6 g/cm3. The alloy can thus have a particularly advantageous strength-to-weight ratio with regard to an application, especially as a structural part. It is particularly beneficial if the alloy is embodied as a light metal alloy. A particularly high application suitability of the alloy can thus be achieved. It is advantageous if, for this purpose, the alloy has less than 5.0 g/cm3, in particular less than 3.0 g/cm3.

For a feasible use as application material, it is beneficial if the alloy is a magnesium-based alloy, aluminum-based alloy, lithium-based alloy, or titanium-based alloy.

It is advantageous if the alloy is a casting alloy. This enables a particularly feasible production, specifically of structural parts having aforementioned properties in particular.

It has proven effective if the alloy is an Al—Mg alloy. Depending on the precise intended application, the alloy may comprise other alloy components. In this manner, application parts having particular relevance to practical situations, in particular structural parts, can be produced having or being made of the alloy. Here, it is particularly beneficial if the alloy is an Al—Mg—Si alloy. Advantageously, the alloy can also comprise zinc (Zn), in particular at an amount greater than 0.01 wt %, typically greater than 1 wt %. A compressive strength of the alloy can thus be optimized. In most cases, the alloy thereby comprises less than 15 wt %, in particular less than 10 wt %, preferably between 1.0 wt % and 5.0 wt %, particularly preferably approximately 3.0 wt %, zinc.

A high application suitability, to which both strength and also pronounced deformability are particularly advantageous, can be achieved if 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.

An alloy with high application suitability, specifically in the form of a structural part, that exhibits particularly high strength and deformability can be achieved if the alloy is a magnesium-based alloy comprising, in particular being made of (in at %):

15.0% to 70.0% lithium,
greater than 0.0%, in particular greater than 0.01%, preferably greater than 0.05%, aluminum,
magnesium and production-related impurities as a remainder,
wherein a ratio of aluminum to magnesium (in at %) is 1:6 to 4:6.

An Mg—Li—Al alloy of this type has an alloy composition in a field around or in proximity to or at an alloy composition of a pseudoeutectic point in the Mg—Li—Al phase diagram, so that a finely structured or micro-scale eutectic microstructure can be attained. The fine-scale microstructure is accompanied by a high strength, in particular a high compressive strength, there being at the same time a good deformability of the magnesium alloy at corresponding aforementioned amounts of lithium in the magnesium alloy. An orientation composition or orientation line in the phase diagram is thereby in particular an aluminum-to-magnesium ratio (in atomic percent, abbreviated by at %) of approx. 3:6, since a particularly homogeneous fine-scale, or homogeneous fine lamellar, microstructure or morphology is found at this ratio. In a range encompassing this ratio, above all at an aluminum-to-magnesium ratio (in at %) of 1:6 to 4:6, the fine, in particular fine lamellar, microstructure or morphology is also found in a varyingly pronounced degree, which is accordingly normally accompanied by varying pronounced magnitudes of strength, in particular a magnitude of compressive strength, as well as deformability or ductility of the magnesium alloy.

Because of these special morphological characteristics in the stated composition range, a formation of a magnesium alloy that has both a high strength, in particular compressive strength, and also a good deformability is thus enabled. This magnesium alloy and a method for the production thereof as well as a realization as a feedstock, semi-finished product or part, and also specific embodiments thereof were filed and disclosed in the European Patent Office as part of European Patent Application No. 19184999.1 and also as part of International Application No. PCT/EP2020/058280, the disclosures of which are hereby included in their entirety in the disclosure of this document. This applies specifically where, as is stated in said applications, the Mg—Al—Li alloy comprises (in at %) 30.0% to 60.0%, in particular 40% bis 50%, preferably 45% bis 50%, particularly preferably 45% to 48%, lithium. It is further advantageous if the Mg—Al—Li alloy comprises (in at %) greater than 0.05%, in particular greater than 0.1%, normally greater than 1% aluminum. It has thereby been shown that the Mg—Al—Li alloy can be embodied with a, in particular lamellar, microstructure with a high degree of fineness if the ratio of aluminum to magnesium (in at/o) is 1.2:6 to 4:6, in particular 1.4:6 to 4:6, preferably 1.5:6 to 4:6. It is beneficial to a pronounced fineness or a fine, in particular lamellar, microstructure if the ratio of aluminum to magnesium (in at %) is 1.8:6 to 3.5:6, in particular 2:6 to 3.5:6, preferably 2.5:6 to 3.5:6. A particularly high strength, in particular compressive strength, can thus be achieved. This holds especially true at an aluminum-to-magnesium ratio (in at %) of 2.8:6 to 3.3:6, preferably approximately 3:6, at which a very homogeneous fine morphology or microstructure is obtainable. To this end, it is particularly advantageous if the magnesium alloy (in at %) is 30.0% to 60.0/lithium and an aluminum-to-magnesium ratio (in at %) is 2.5:6 to 3.5:6, in particular 2.8:6 to 3.3:6, preferably approximately 3:6. In this regard, reference is made in particular to FIG. 1 of the aforementioned application documents, in which a corresponding arrangement in an Mg—Li—Al phase diagram is schematically illustrated, and the disclosure of which as well as the related description are accordingly also to be considered part of this document. A particularly pronounced homogeneity is also achievable if the magnesium alloy thereby comprises (in at %) 40.0% to 60.0% lithium. As stated in the aforementioned applications and also incorporated accordingly into the present disclosure, the properties of the Mg—Al—Li alloy can be further optimized if amounts of calcium, rare earth metals, in particular yttrium, zinc, and/or silicon are additionally present according to the aforementioned applications at corresponding content ranges stated in the aforementioned applications. For example, an alloy of this type can be embodied as Mg—20% Li—15% Al—1% Ca—0.5% Y (in wt %) or Mg—20% Li—24% Al—1% Ca—0.5% Y (in wt %).

The other object of the invention is attained with a method of the type named at the outset, under the condition that the composition is provided such that it lies in a field around a pseudoeutectic point of a phase diagram of the alloy, so that the eutectic structure is embodied at an amount of at least 85 mol % or at % when the alloy is cooled to the solid phase or solidifies. As stated above, the alloy can thus be embodied with high strength and pronounced deformability. Because the alloy composition is selected in a field around the pseudoeutectic point, a eutectic phase reaction or phase transformation takes place when the alloy is cooled from the liquid state to the solid state, or during the liquid-to-solid transition, which phase reaction or transformation embodies the eutectic microstructure with a particularly high degree of fineness or fine structuring as a principal microstructure amount of the alloy.

It should be understood that the method according to the invention may be embodied correspondingly or analogously to the features, advantages, implementations, and effects that are described, in particular as described above, within the scope of an alloy according to the invention. The same also applies to the alloy according to the invention with regard to a method according to the invention.

A feedstock, semi-finished product or part is advantageously realized having, in particular being made of, an alloy according to the invention or such that it is obtainable using a method according to the invention for producing an alloy according to the invention. In accordance with the foregoing explanations, features, and effects of the alloy according to the invention or of an alloy produced with a method according to the invention, a feedstock, semi-finished product or part formed with an alloy also has an advantageously high strength and good deformability.

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

FIG. 1 and FIG. 2 show phase diagram illustrations of an Al—Mg—Si system in which alloy compositions of exemplary alloys are indicated;

FIG. 3 through FIG. 12 show optical microscope images of exemplary alloys from FIG. 1 and FIG. 2;

FIG. 13 through FIG. 20 show yield stress diagrams of exemplary alloys from FIG. 1 through FIG. 12;

FIG. 21 shows a phase diagram illustration of an Al—Cu—Mg system with an alloy composition of an exemplary alloy being drawn;

FIG. 22 shows optical microscope images of the exemplary alloy from FIG. 21;

FIG. 23 shows a yield stress diagram of the exemplary alloy from FIG. 21 and FIG. 22;

FIG. 24 shows a phase diagram illustration of an Mg—Al—Li system in which alloy compositions of exemplary alloys are indicated;

FIG. 25 and FIG. 27 show optical microscope images of exemplary alloys from FIG. 24:

FIG. 28 and FIG. 29 show yield stress diagrams of exemplary alloys from FIG. 24 through FIG. 27;

FIG. 30 shows a phase diagram illustration of an Mg—Cu—Zn system with an alloy composition of an exemplary alloy being drawn;

FIG. 31 and FIG. 32 show optical microscope images of the exemplary alloy from FIG. 30;

FIG. 33 shows a yield stress diagram of the exemplary alloy from FIG. 30 through FIG. 32;

FIG. 34 shows electron microscope images of an exemplary alloy from an Al—Cu—Mg—Zn system;

FIG. 35 shows a yield stress diagram of the exemplary alloy from FIG. 34;

FIG. 36 shows a phase amount diagram of an exemplary alloy from an Al—Mg—Si—Zn system;

FIG. 37 shows a solid amount diagram of a Scheil-Gulliver calculation for the exemplary alloy from FIG. 36.

In the course of a development of the alloy according to the invention, series of tests were conducted with different alloy compositions of various alloy systems. In each case, alloys were thereby chosen with an alloy composition in the field of or around a pseudoeutectic point of a respectively related phase diagram, and a eutectic structure was formed by cooling the alloy from a liquid state to a solid state. The microstructure was then examined by means of microscopy. In addition, various dilatometric test series and compression tests were conducted at room temperature, approximately 20° C., as a standard, wherein yield curves which depict a yield stress, in MPa, as a function of a degree of deformation, illustrated as the amount of length change ΔL relative to a starting length L0, that is

Δ ⁢ L L 0 ,

and correspondingly dimensionless, were calculated as a result.

Below, test results for exemplary alloys from the alloy systems Al—Mg—Si, Al—Cu—Mg, Mg—Li—Al, Mg—Cu—Zn, Al—Cu—Mg—Zn, and Al—Mg—Si—Zn are shown in a representative manner in order to illustrate the aforementioned concept on a broad basis.

Al—Mg—Si system:

FIG. 1 and FIG. 2 show illustrations of a ternary phase diagram of an Al—Mg—Si system, wherein FIG. 2 is a segment illustration from the phase diagram for the purpose of showing the relevant alloy composition range in detail. Ten exemplary alloys from the Al—Mg—Si system were produced and examined. The alloy compositions of the exemplary alloys from the Al—Mg—Si system are respectively indicated in percentage by weight and atomic percent as exemplary alloy 1 through exemplary alloy 10 in Table 1 and correspond to the reference numerals 1 through 10, which in particular denote the respective alloy composition in the phase diagram from FIG. 1 and FIG. 2.

TABLE 1 Ten exemplary alloys from the Al—Mg—Si alloy system. Al Mg Si Exemplary alloy 1 wt % 68.10 9.00 22.90 at % 68.04 9.98 21.98 Exemplary alloy 2 wt % 71.10 9.10 19.80 at % 70.94 10.08 18.98 Exemplary alloy 3 wt % 78.10 6.30 15.60 at % 78.04 6.99 14.97 Exemplary alloy 4 wt % 81.50 5.00 13.50 at % 81.48 5.55 12.97 Exemplary alloy 5 wt % 82.00 6.00 12.00 at % 81.85 6.65 11.51 Exemplary alloy 6 wt % 82.00 10.50 7.50 at % 81.30 11.56 7.14 Exemplary alloy 7 wt % 82.70 10.00 7.30 at % 82.03 11.01 6.96 Exemplary alloy 8 wt % 84.90 10.90 4.20 at % 84.03 11.98 3.99 Exemplary alloy 9 wt % 85.80 9.60 4.60 at % 85.05 10.56 4.38 Exemplary alloy 10 wt % 86.50 7.20 6.30 at % 86.03 7.95 6.02

As can be seen in the phase diagram from FIG. 1 and FIG. 2, the exemplary alloys 8 through 10 each have compositions which are arranged in a field around a pseudoeutectic point pE, wherein the exemplary alloys 8 and 9 are positioned very close to the pseudoeutectic point and the exemplary alloy 10 is positioned at a somewhat greater distance from the pseudoeutectic point pE. The alloy composition of the exemplary alloy 9 thereby virtually lies at the pseudoeutectic point pE. The pseudoeutectic point pE is illustrated in FIG. 2 by a drawn reference line, wherein the pseudoeutectic point pE is located at the intersection of the monovariant line in the direction of Al3Mg and the reference line. In FIG. 2, it can also be seen that the exemplary alloys 3 through 5 are arranged in a field around a eutectic point E of the phase diagram. Furthermore, the exemplary alloys 6 and 7 are provided as comparisons, the compositions of which are located at a large distance from the pseudoeutectic point pE, evident in FIG. 2, as well as the exemplary alloys 1 and 2 which, though positioned in direct proximity to a liquidus boundary line, are positioned at a greater distance from both the pseudoeutectic point pE and also the eutectic point E, evident in FIG. 1.

In FIG. 3 through FIG. 12, optical microscope images of the exemplary alloys 1 through 10 are shown in order to illustrate a respective microstructure. In FIG. 13 through FIG. 20, yield stress diagrams are illustrated as the results of dilatometric test series of the Al—Mg—Si exemplary alloys which were conducted at room temperature, approximately 20° C. Yield stress curves are shown, wherein a 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 alloy specimens with an alloy composition corresponding to the alloy composition of one of the exemplary alloys 1 through 10. Each yield stress diagram thus represents an alloy composition of one of the exemplary alloys 1 through 10.

As can be seen in FIG. 10 through FIG. 12, the microscope images of the exemplary alloys 8 through 10, which have alloy compositions in proximity to or in a field around the pseudoeutectic point pE, show a dominant finely structured or fine-scale eutectic structure. By comparison, microscope images of the exemplary alloys 4 and 5 can be viewed in FIG. 6 and FIG. 7, which alloys have an alloy composition in proximity to the eutectic point E. These show a pronounced degree of a eutectic structure which comprises a course structure compared to the microstructures of the exemplary alloys 8 and 9. If one compares these with the microscope images of the exemplary alloys 1 and 2 shown in FIG. 3 and FIG. 4, the alloy compositions of which are located at a large distance from, but in the region of, a liquidus boundary line, it is discernible that they exhibit an even courser eutectic microstructure. In FIG. 8 and FIG. 9, microscope images of the exemplary alloys 6 and 7 are also shown which have alloy compositions in a distant region of the pseudoeutectic point pE or at a great distance therefrom. It can be seen that a eutectic structure is already present, but with a relatively course structure and being notably less dominant and at a lower amount. In addition, high amounts of residual solidifications are also evident, identifiable in FIG. 8 and FIG. 9 in the form of light channels.

FIG. 13 and FIG. 14 show yield stress diagrams of the exemplary alloys 8 and 9, which have alloy compositions in proximity to or in a field around the pseudoeutectic point pE. It can be seen that both exemplary alloy 8 and also exemplary alloy 9 have high strength, in particular compressive strength, and pronounced deformability with yield stresses between 300 MPa and 400 MPa, wherein exemplary alloy 8 in particular, illustrated in FIG. 13, exhibits yield stresses of up to 400 MPa. By comparison, yield stress diagrams of the exemplary alloys 4 and 5 can be viewed in FIG. 15 and FIG. 16, which alloys have alloy compositions in proximity to the eutectic point E. The exemplary alloys 4 and 5 also exhibit high strength and, at least conditional on individual specimens, a high deformability, wherein yield stresses lie below those of the exemplary alloys 8 and 9 at approximately 300 MPa or, in relation to exemplary alloy 5, illustrated in FIG. 16, consistently below that. This result correlates with the finding that exemplary alloys with an alloy composition in the field of the pseudoeutectic point pE exhibit a particularly high fine structuring of the eutectic structure thereof, in particular also compared to the eutectic structure of exemplary alloys with alloy compositions in the field of a eutectic point E, which also explains the higher strength and pronounced elasticity of alloys in the field of the pseudoeutectic point.

In FIG. 20, a yield stress diagram of the exemplary alloy 10 is shown, the alloy composition of which is arranged at a somewhat greater distance from the pseudoeutectic point pE. Evident are slightly lower yield stress values and, in particular, a higher variance between the individual measurement results. In FIG. 17 and FIG. 18, it is furthermore shown that, by comparison, exemplary alloy 1 and exemplary alloy 2 with alloy compositions in the region of a liquidus boundary line, but at a distance from both the alloy composition of the pseudoeutectic point pE and also the eutectic point E, have notably poorer strength and deformability properties. In FIG. 19, a yield stress diagram corresponding to the alloy composition of the exemplary alloys 6 and 7 is additionally shown, the alloy composition of which is positioned at a relatively large distance from that of the pseudoeutectic point pE in the phase diagram. The corresponding yield stress curves show clearly reduced yield stresses compared to yield stresses of an alloy composition closer to the pseudoeutectic point pE, such as that of those shown in FIG. 13 for the exemplary alloy 8.

It is evident that an alloy composition in a field around a pseudoeutectic point pE corresponds to a finely structured eutectic microstructure and an accordingly high strength and pronounced deformability.

In a detailed view, it can be seen that, relative to the monovariant line or liquidus boundary line in the direction of Al3Mg2, the exemplary alloy 8 in the phase diagram from FIG. 2 lies above said line in the Mg2Si region, which is why a solidification begins with an undesirable formation of Mg2Si in particular, or a primary solidification is formed with an intermetallic Mg2Si phase, when the alloy is cooled from the liquid phase. It has been shown that a primary solidification formed with an intermetallic phase has negative effects for an embodiment of both high strength and also deformability. To achieve particularly advantageous strength and deformability, one therefore generally strives to keep a primary solidification having or being made of intermetallic phase as minor as possible, or to prevent it. However, the primary solidification for exemplary alloy 8 is so slightly pronounced that it entails virtually no restraint on mechanical properties. The microscope images of the exemplary alloy 8 in FIG. 10 show extensive regions with a fine eutectic structure, in this case formed with Al mixed crystal phase and Mg2Si. Advantageously, a residual solidification of Al mixed crystal phase is also only very slightly pronounced or hardly present. To keep from undermining the advantageous strength and deformability properties attainable with the eutectic structure, one strives to keep a residual solidification as small as possible or prevent it. In particular, the residual solidification is not bonded in a network-like manner, or is embodied in the form of units separated from one another, which likewise promotes an advantageous embodiment of high strength and pronounced deformability. The exemplary alloy 8 thus proves to be well suited, both with regard to low residual solidification and also low primary solidification, to controlling strength properties and deformability on the basis of the fine eutectic structure. This can be optimized even further if the alloy composition is chosen such that the primary solidification is formed having or being made of a mixed crystal phase and not with an intermetallic compound or phase, that is, if the primary solidification is located in the Al mixed crystal phase region in the case of the exemplary alloy 8.

This view of the exemplary alloy 8 and also the accompanying explanations apply analogously to the exemplary alloy 9. The exemplary alloy 9 has an alloy composition lying virtually at the pseudoeutectic point pE. The exemplary alloy 9, as can be seen in FIG. 11, also shows a fine eutectic structure with little residual solidification and little primary solidification. The somewhat lower strength in comparison with the exemplary alloy 8 is explained by the lower dissolved amount of Mg in the Al mixed crystal phase. A strength can be advantageously achieved by varying an amount of dissolved elements in the mixed crystal phase, with the primary solidification preferably lying, however, in the mixed crystal region and not in the region of an intermetallic phase, as stated above.

By comparison, the exemplary alloy 10, as can be seen in FIG. 12, also shows a fine eutectic structure, but with a greater amount of residual solidification, in the form of Al mixed crystal phase and Si, which residual solidification is also shaped in a network-like manner. Due to the low Mg content, most of the Mg is bonded in the form of Mg2Si so that a mixed crystal hardening of the Al mixed crystal phase is very slightly pronounced. This corresponds to lower yield stresses in the yield stress diagram from FIG. 20.

In a further detailed view of the exemplary alloys arranged at a distance from the pseudoeutectic point pE in relation to an alloy composition, it can be seen that the exemplary alloys 4 and 5, which lie in the field of the eutectic point E, illustrated in FIG. 6 and FIG. 7, comprise a low amount of primary solidification, around which a relatively course eutectic structure, formed with two phases, is arranged. A remaining predominant amount of eutectic structure is embodied as a ternary eutectic, formed with mixed crystal phase, Al2Si and Si. The mechanical properties, in particular strength and deformability, are negatively influenced by the course binary eutectic structure or phase in particular. A fine eutectic ternary structure is locally present to some extent, which structure transitions into markedly coarsened structures in some locations. The differences between the microstructures of exemplary alloys with alloy compositions at, or in the field of, the pseudoeutectic point pE compared to those at, or in the field of, the eutectic point E correlate with the finding of accordingly improved strength and deformability properties of alloy compositions at, or in the field around, the pseudoeutectic point pE.

It can furthermore be seen that the exemplary alloys 6 and 7 comprise course, polygon-shaped primary solidifications with the related microscope images shown in FIG. 8 and FIG. 9. This is explained by the positioning of the related alloy compositions in the Mg2Si region of the phase diagram, as a result of which a pronounced Mg2Si primary solidification forms. A course eutectic structure is identifiable therebetween, as well as a high amount of residual solidification, which is evident from the light regions or channels in FIG. 8 and FIG. 9. Due to this structural morphology, the exemplary alloys 6 and 7 exhibit markedly reduced strengths and yield stresses, which are associated in particular with crack initiation and brittle fracture.

In FIG. 2, a particularly advantageous region for the embodiment of an Al—Mg—Si alloy is drawn as a gray, planar region. This essentially designates or corresponds to an aforementioned alloy composition of the exemplary alloys 8 and 9, but with a variation of the alloy composition such that a mixed crystal phase is embodied as primary solidification and, in particular, no intermetallic phase is embodied. This enables an embodiment of particularly high strengths with pronounced deformability. A particularly advantageous implementation range for an Al—Mg—Si alloy of this type is thus ensured if the Al—Mg—Si alloy is arranged in a field around the pseudoeutectic point in the Al—Mg—Si phase diagram, wherein the alloy composition in the phase diagram is arranged starting from the aforementioned pseudoeutectic point of the phase diagram in FIG. 2 on a side of the corresponding monovariant line facing an increasing Al amount.

Al—Cu—Mg system:

FIG. 21 shows an illustration of a ternary phase diagram of an Al—Cu—Mg system. An exemplary alloy from the Al—Cu—Mg system was produced and examined. The related alloy composition is indicated in percentage by weight and atomic percent as exemplary alloy 13 in Table 2 and corresponds to reference numeral 13, which in particular denotes the alloy composition in the phase diagram from FIG. 21.

TABLE 2 Exemplary alloy from the Al—Cu—Mg alloy system. Al Cu Mg Exemplary alloy 13 wt % 66.00 24.00 10.00 at % 75.61 11.67 12.72

As can be seen in the phase diagram from FIG. 21, the exemplary alloy 13 has an alloy composition which is arranged in a field around a pseudoeutectic point pE. A related microstructure is illustrated in FIG. 22 with the aid of optical microscope images. Evident is a very fine-scale eutectic microstructure and a low amount of primary solidification formed with mixed crystal phase. In FIG. 23, a yield stress diagram is shown as the result of dilatometric test series of the Al—Cu—Mg exemplary alloy 13, wherein a yield stress, in MPa, is once again illustrated as a function of the degree of deformation. It is evident that very high strengths and yield stresses are achieved.

The elongation at break also lies in the technologically relevant range for this alloy system. The strength and deformability correspond to the fine eutectic microstructure and, in particular, to the low amount of primary solidification.

Mg—Al—Li system:

FIG. 24 shows an illustration of a ternary phase diagram of an Mg—Al—Li system. Three exemplary alloys from the Mg—Al—Li system were produced and examined. The alloy compositions of the exemplary alloys from the Mg—Al—Li system are respectively indicated in percentage by weight and atomic percent as exemplary alloy 14, 15, and 16 in Table 3 and correspond to the reference numerals 14, 15, and 16, which in particular denote the respective alloy composition in the phase diagram from FIG. 24.

TABLE 3 Three exemplary alloys from the Mg—Al—Li alloy system. Mg Al Li Exemplary alloy 14 wt % 55.00 29.00 16.00 at % 40.10 19.05 40.85 Exemplary alloy 15 wt % 56.00 24.00 20.00 at % 37.90 14.60 47.40 Exemplary alloy 16 wt % 65.00 15.00 20.00 at % 43.80 9.10 47.10

As can be seen in the phase diagram from FIG. 24, the exemplary alloys 14 through 16 respectively have an alloy composition which is arranged in a field around a pseudoeutectic point pE. The pseudoeutectic point pE is illustrated in FIG. 24 by a drawn reference line, wherein the pseudoeutectic point pE is located at the intersection of the monovariant line, or liquidus boundary line, and the reference line. With additions of CaY, in particular approximately 1 wt % Ca and approximately 0.5 wt % Y, oxidation properties of the exemplary alloys from the Mg—Al—Li system can feasibly be stabilized without negatively influencing how pronounced the structure is. In the phase diagram, the exemplary alloys 14 and 15 lie at a somewhat closer distance in a vicinity of the pseudoeutectic point and the exemplary alloy 16 somewhat farther away, wherein the alloy composition of the exemplary alloy 14 is positioned more or less at the pseudoeutectic point pE. According to currently available data, the exemplary alloys 14 through 16 are present in a mixed crystal region, in particular such that they form a body-centered cubic lattice, bec.

In FIG. 25 through FIG. 27, microstructures are respectively rendered visible with the aid of microscope images. The structural morphology from FIG. 25 and FIG. 26 indicates an embodiment of an extremely fine-scale structure which can no longer be resolved in the light microscope used. The grain boundaries which can thereby be recognized are attributable to oxidic impurities. The microstructure of exemplary alloy 16 was examined by means of scanning electron microscopy, illustrated in FIG. 27. Evident in FIG. 27 are, on the one hand, light grain boundary phases (in whitish-gray) that were identified as Al—Ca and, on the other hand, pronounced fine crystalline structures or morphologies in a region surrounded by grain boundary phases, in particular in a center section of said region, or in the interior of the mixed crystal phase, clearly visible in particular in the right-hand image from FIG. 27. In the phase diagram from FIG. 24, the alloy composition of the exemplary alloy 16 appears to lie at a relatively far distance from the monovariant line and the pseudoeutectic point pE. In this case, however, it should be noted that, according to established technical knowledge, the slope in the region of the body-centered cubic lattice, bec,—in which the exemplary alloy 16 is also arranged—in the phase diagram is very flat, and that the three elements Mg, Al. and Li also exhibit a high solubility in one another. It its thus possible to explain why such an expansive field around the pseudoeutectic point results, in which field an advantageous fine-scale eutectic microstructure can be embodied in a high amount.

FIG. 28 and FIG. 29 show yield stress diagrams of the exemplary alloys 15 and 16 as the results of dilatometric test series, wherein a yield stress, in MPa, is once again illustrated as a function of the degree of deformation, with FIG. 28 showing yield stress curves relating to the exemplary alloy 15 and FIG. 29 showing yield stress curves relating to the exemplary alloy 16. It is evident that both exemplary alloys have high strengths and yield stresses, as well as pronounced deformabilities, corresponding to the fine eutectic microstructures identified. In FIG. 29, which relates to the exemplary alloy 16, a possibility of a further property optimization by means of heat treatment is also illustrated.

FIG. 29 shows yield curves of alloy specimens immediately after a production of the exemplary alloy 16 (as cast), depicted in FIG. 29 as solid lines, denoted by reference numeral 16-1, and additionally yield curves of exemplary alloy specimens after a conducted heat treatment (aged) of the exemplary alloy 16, depicted in FIG. 29 as dashed lines, denoted by reference numeral 16-2. For this purpose, specimens of the exemplary alloy 16 were subjected to a heat treatment at 330° C. for 3 hours, and yield curves were then calculated by means of compression tests. A clear influence of the heat treatment on strength, in particular compressive strength, and deformability is evident, as a result of which there is the potential to set compressive strength and deformability in an optimized manner using heat treatment, in particular for an eventual intended application.

As previously explained above in the document, it has proven beneficial to the realization of an alloy with high application suitability if the alloy is a magnesium-based alloy comprising, in particular being made of, (in at %)

15% to 70.0% lithium,
greater than 0.0%, in particular greater than 0.01%, preferably greater than 0.05%, aluminum, magnesium and production-related impurities as a remainder,
wherein a ratio of aluminum to magnesium (in at %) is 1:6 to 4:6. The exemplary alloy 16 can be viewed as a representative example of this alloy definition, as is shown within the scope of European Patent Application Number 19184999.1 and also within the scope of International Application Number PCT. EP2020058280, both of which were filed in the European Patent Office. Here, reference is once again made in particular to FIG. 1 from each of these applications. In FIG. 24, a corresponding aluminum-to-magnesium ratio (in at %) of 1:6 is drawn as a dashed line. The aforementioned aluminum-to-magnesium ratio range (in at % or mol %) of 1:6 to 4:6 is thereby located to the left of this line in the phase diagram from FIG. 24 and, in particular, constitutes a specific embodiment in the field around the pseudoeutectic point pE.

A particularly advantageous implementation range for an Mg—Li—Al alloy that is usable as an application alloy, in particular for a structural part, is ensured if the Mg—Li—Al alloy is arranged in the Mg—Li—Al phase diagram in a region between the line indicating an aluminum-to-magnesium ratio (in at %) of 1:6 and the monovariant line or liquidus boundary line, in particular with an aforementioned Li content range. A range of this type is denoted in the phase diagram from FIG. 24 as a gray, planar region.

It becomes apparent, as was already the case previously within the scope of the exemplary alloys from the Al—Si—Mg system, that an alloy composition is preferably chosen such that the alloy composition lies in the field of the pseudoeutectic point pE and, moreover, preferably comprises a primary solidification having or being made of mixed crystal phase; that is, that the corresponding alloy composition is positioned in a mixed crystal region in the phase diagram.

Mg—Cu—Zn:

FIG. 30 shows an illustration of a ternary phase diagram of an Mg—Cu—Zn system. An exemplary alloy from the Mg—Cu—Zn system was produced and examined. The related alloy composition is indicated in percentage by weight and atomic percent as exemplary alloy 17 in Table 4 and corresponds to reference numeral 17, which in particular denotes the alloy composition in the phase diagram from FIG. 30.

TABLE 4 Exemplary alloy from the Mg—Cu—Zn alloy system. Al Cu Zn Exemplary alloy 17 wt % 58.00 16.5 25.5

As can be seen in the phase diagram from FIG. 30, the exemplary alloy 17 has an alloy composition which is arranged in a field around a pseudoeutectic point pE. A related microstructure is illustrated in FIG. 31 and FIG. 32 with the aid of optical microscope images. Evident is a very fine-scale eutectic microstructure that is at a limit of resolution of a light microscope. Here, a relatively large amount of primary solidification can be seen. It is therefore advantageous for a high strength and deformability if an alloy composition is selected even closer to the pseudoeutectic point pE or closer to the monovariant line or liquidus boundary line.

FIG. 33 shows a yield stress diagram as the results of dilatometric test series of the exemplary alloy 17, wherein a yield stress, in MPa, is once again illustrated as a function of the degree of deformation. It is evident that high strengths and yield stresses are achieved which, based on the pronounced amount of primary solidification apparent in the microscope images, can be further improved, however, by choosing an alloy composition even closer to the pseudoeutectic point pE.

In FIG. 33, yield curves of the exemplary alloy 17 immediately following a production of the exemplary alloy 17 (as cast) are thereby shown, denoted by reference numeral 17-1, and also yield curves of the exemplary alloy 17 after a conducted heat treatment, denoted by reference numeral 17-2. For this purpose, specimens of the exemplary alloy 17 were subjected to a heat treatment at 350° C. for 4 hours, and yield curves were then calculated by means of compression tests. A clear influence of the heat treatment on strength and deformability is evident, as a result of which there is the potential to further optimize strength and deformability by means of heat treatment.

Examinations of quaternary alloy systems and quaternary eutectics were then also carried out. The alloy systems Al—Cu—Mg—Zn and Al—Mg—Si—Zn in particular were considered for this purpose.

Al—Cu—Mg—Zn:

In regard to the alloy system Al—Cu—Mg—Zn, an exemplary alloy that lies in the field of a pseudoeutectic point pE was produced and examined. The alloy composition is indicated in percentage by weight and atomic percent as exemplary alloy 18 in Table 5 and corresponds to reference numeral 18.

TABLE 5 Exemplary alloy from the Al—Cu—Mg—Zn alloy system. Al Cu Mg Zn Exemplary alloy 18 wt % 6.20 75.40 5.40 13.00 at % 12.51 65.00 12.09 10.82

In order to examine the eutectic microstructure, electron microscope images of the exemplary alloy 18 were recorded, shown in FIG. 34. Evident is a finely structured eutectic structure, in particular with structural dimensions in the nanometer range, clearly visible in the right-hand image from FIG. 35 as an expansive grainy region in the center of the picture.

This is a binary eutectic structure in a system with four components or elements and thus an increase in the thermodynamic degree of freedom f, explained at the outset, from 1 to 3 (quaternary eutectic).

In FIG. 34, substructures are identifiable in the primary regions (in gray), wherein these are artifacts of an isostoichiometric structural transformation (bec to fec) in the solid state. In terms of a direct influence on strength and deformability, they are insignificant. Also visible is a relatively large amount of primary solidification in the form of a mixed crystal phase (in light gray to whitish), as well as an intermetallic secondary phase (in black), in particular in the form of a Laves phase.

FIG. 35 shows a yield stress diagram as the result of dilatometric test series with the exemplary alloy 18. Depicted are yield curves prior to a completed heat treatment, denoted by reference numeral 18-1, and yield curves following a completed heat treatment, denoted by reference numeral 18-2, wherein a yield stress, in MPa, is once again illustrated as a function of the degree of deformation. It is evident that the exemplary alloy 18 exhibits a very high strength with a simultaneously present elongation at break, wherein a deformability can be varied by means of heat treatment.

The pronounced primary solidification present as well as the secondary phase are to be regarded as brittleness-increasing factors, which is why it would be advantageous to further reduce these amounts in order to further improve strength and deformability, for example by reducing the distance of the alloy composition from or bringing it even closer to the pseudoeutectic point pE in the phase diagram.

Al—Mg—Si—Zn:

In regard to the alloy system Al—Mg—Si—Zn, an exemplary alloy that lies in the field of a pseudoeutectic point pE was examined by means of simulation. The alloy composition is indicated as exemplary alloy 19 in Table 6 and corresponds to reference numeral 19.

TABLE 6 Exemplary alloy from the Al—Mg—Si—Zn alloy system. Al Mg Si Zn Exemplary alloy 19 wt % 83.3 9.2 4.5 3.0

As a result of the simulation, phase amounts are illustrated in FIG. 36 as a function of the temperature of the exemplary alloy 19. Evident is a direct transition from the solid to the liquid phase, corresponding to an embodiment of a eutectic structure. In FIG. 35, corresponding thereto, an illustration of the solid amount as a function of the temperature is shown, determined by means of a Scheil-Gulliver solidification calculation. The equilibrium and Scheil-Gulliver solidification curves shown depict an alloy system which exhibits a binary eutectic solidification with four components or elements. Accordingly, there is therefore once again an increase in the thermodynamic degree of freedom from 1 to 3. In FIG. 37, the Scheil-Gulliver calculation shows a very small amount of primary solidification in the form of a mixed crystal phase at an amount of less than 3 mol % or at/o and in addition a virtually non-existent residual solidification.

It is thus analogously apparent that, in addition to a positioning of the alloy composition in the field of the pseudoeutectic point pE, an amount of primary solidification and/or residual solidification can advantageously also be minimized in order to further increase or improve strength properties and deformability properties.

An alloy according to the invention with more than three components having a eutectic structure created by a cooling from the liquid state to the solid state can thus advantageously be embodied with a finely structured eutectic structure, in particular with a fine structure in the nanometer range, which constitutes a dominant or principal phase amount or structure amount in the alloy if an alloy composition of the alloy is arranged in the field of or around a pseudoeutectic point in the phase diagram. The alloy can thus be embodied with advantageously high strength and pronounced deformability. This holds especially true if a primary solidification and/or residual solidification is embodied to be very small. Specifically, it is beneficial thereto if the primary solidification is formed having or being made of a mixed crystal phase, in particular not having or being made of an intermetallic phase, or if the alloy composition is chosen in a corresponding region in the phase diagram. An alloy formed in this manner thus offers the potential to realize, preferably depending on a specific purpose, robust and resilient components, especially structural components, in particular for an intended application in the automotive industry, aircraft industry, and/or space industry.

Claims

1. An alloy, in particular a light metal alloy, having an alloy composition with at least three components and a eutectic structure that is obtained by cooling the alloy from a liquid state to a solid state, under the condition that a composition of the alloy lies in a field around a pseudoeutectic point (pE) of a phase diagram of the alloy, so that at least 85 mol % eutectic structure is present in the alloy.

2. The alloy according to claim 1, wherein the eutectic structure has an average spacing of the phase amounts thereof of less than 3 μm, preferably less than 1 μm.

3. The alloy according to claim 1, wherein the alloy comprises a residual solidification at an amount of maximally 5 mol %, preferably maximally 3 mol %.

4. The alloy according to claim 1, wherein the alloy comprises a primary solidification at an amount less than 10 mol %, in particular less than 5 mol %.

5. The alloy according to claim 4, wherein the primary solidification is formed having a mixed crystal phase.

6. The alloy according to claim 1, wherein the alloy has a density less than 8 g/cm3.

7. The alloy according to claim 1, wherein the alloy is a magnesium-based alloy, aluminum-based alloy, lithium-based alloy, or titanium-based alloy.

8. (canceled)

9. The alloy according to claim 1, wherein the alloy is an Al—Mg—Si alloy.

10. The alloy according to claim 9, wherein the Al—Mg—Si alloy comprises between 0.01 wt % and 5.0 wt %, in particular approximately 3.0 wt %, zinc.

11. A method for producing an alloy, in particular an alloy according to claim 1, having a eutectic structure, wherein the alloy has an alloy composition with at least three components and wherein the alloy is cooled, starting from a liquid state, to a solid state of the alloy in order to form the eutectic structure, under the condition that the alloy composition is provided such that it lies in a field around a pseudoeutectic point (pE) of a phase diagram of the alloy, so that the eutectic structure is embodied at an amount of at least 85 mol % during the cooling to the solid phase.

12. A feedstock, semi-finished product, or construction material having an alloy according to claim 1.

13. A feedstock, semi-finished product, or construction material obtainable using the method according to claim 11.

Patent History
Publication number: 20220267881
Type: Application
Filed: Jul 7, 2020
Publication Date: Aug 25, 2022
Applicant: LKR LEICHTMETALLKOMPETENZZENTRUM RANSHOFEN GMBH (Braunau am Inn - Ranshofen)
Inventors: Stefan GNEIGER (Weistrach), Clemens SIMSON (Mattighofen), Alexander GROßALBER (Mattighofen), Simon FRANK (Neuhofen im Innkreis), Andreas BETZ (Pfarrkirchen)
Application Number: 17/625,360
Classifications
International Classification: C22C 21/02 (20060101); C22C 21/08 (20060101); C22F 1/043 (20060101); C22F 1/047 (20060101);