EP3997251A1

EP3997251A1 - Alloy having fine-scale eutectic, in particular nanoeutectic, structure and production of such an alloy

Info

Publication number: EP3997251A1
Application number: EP20735621.3A
Authority: EP
Inventors: Stefan Gneiger; Clemens Simson; Alexander GROSSALBER; Simon Frank; Andreas Betz
Original assignee: LKR Leichtmetallkompetenzzentrum Ranshofen GmbH
Current assignee: LKR Leichtmetallkompetenzzentrum Ranshofen GmbH
Priority date: 2019-07-08
Filing date: 2020-07-07
Publication date: 2022-05-18
Also published as: CN114026260A; WO2021005062A1; US20220267881A1; WO2021004662A1; EP3763845B1; EP3763845A1; JP2022540544A; CA3138658A1; CA3137604A1; CN114096690A; US20220259705A1; KR20220030243A; CN114026260B; KR20220030244A; JP2022540542A

Abstract

The invention relates to an alloy, in particular a light metal alloy, which has an alloy composition with at least three components and a eutectic structure which is obtained by cooling the alloy from a liquid state to a solid state, with the proviso that a composition of the alloy lies in an area 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 invention further relates to a method for producing such an alloy.

Description

Alloy with fine-scale eutectic, in particular nano-eutectic, structure and manufacture of the same

The invention relates to an alloy, in particular a light metal alloy, comprising an alloy composition with at least three components and a

eutectic structure, which is obtained by cooling the alloy from a liquid state to a solid state.

The invention also relates to a method for producing an alloy, in particular a light metal alloy, with a eutectic structure, the alloy having an alloy composition with at least three components and the alloy for forming the eutectic structure starting from a liquid state into a solid state of the alloy is cooled.

It is known that it can be advantageous if a portion of a structure of an alloy is formed with a eutectic structure in order to influence the casting properties or strength properties of an alloy. As a technical

Application alloys are often binary cast alloys, that is, alloys with two components that have eutectic structures. These are usually characterized by a eutectic point in their phase diagram, at which a liquid phase of the alloy and two solid phases of the alloy are in thermodynamic equilibrium with one another or, when the alloy cools from the liquid phase, a direct transition from a liquid state to a solid state takes place, whereby a eutectic structure is formed. According to Gibbs' phase rule for solids with 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

Degree of freedom f = 0. The direct transition from the liquid phase to the solid phase often leads to the formation of a fine and lamellar structure.

With regard to ternary alloy systems, too, attempts have become known in an analogous manner to form alloys with compositions close to a ternary eutectic point in order to form a eutectic structure

To improve strength properties. According to Gibbs' phase rule f = N - P +1 this corresponds in an analogous manner with three components and four phases also to a degree of freedom of f = 0. However, high cooling rates are usually required to form such alloys in order to produce a eutectic structure with pronounced fineness with usable amounts of alloys and is often a coordinated additional Combination with other elements required for precipitation hardening in order to significantly increase the strength of the alloy. For this purpose, cooling rates in a range from 50 K / s to 200 K / s are usually used. In particular, however, the requirement for high cooling rates limits the technical usability of such alloys to small-scale components.

This is where the invention comes in. The object of the invention is to specify an alloy with at least three components which has high strength and good formability.

It is also an object of the invention to provide a method for producing such an alloy.

The object is achieved according to the invention if, in an alloy of the type mentioned at the outset, the eutectic structure is obtained with the proviso that a

Composition of the alloy lies in a region around a pseudo-eutectic point of a phase diagram of the alloy, so that at least 85 mol% or at% eutectic structure are present in the alloy.

The invention is based on the knowledge that when an alloy is composed with at least three components or elements, which in the phase diagram of the alloy is at or near a pseudo-eutectic point, a particularly fine-scale or finely structured eutectic structure can be formed, which, in particular, can have a finer eutectic structure than an alloy with a selected composition that is at the “usual” eutectic point in the phase diagram. In particular, in this way characteristic structural distances of the eutectic structure in the low micrometer range and in particular nanometer range can be implemented, also referred to as nanoeutectic structure. In addition, it has been shown that the eutectic structure formed thereby generally represents an essential or dominant structure and in particular in the vicinity or in an area around, in particular on, a eutectic Point often only a small or negligibly small or no primary solidification phase and / or residual solidification phase occurs. This combination of the exceptionally fine structure of the eutectic structure and its

dominant presence in the alloy enables the alloy to be formed with both high strength, in particular compressive strength, and also pronounced

Formability. The pseudo-eutectic point is usually abbreviated as "e" or "pE" in representations and the eutectic point is abbreviated as "E".

In a ternary phase diagram, the liquidus line and solidus line known from the binary phase diagram usually correspond to curved surfaces and binary phase surfaces correspond to phase volumes. The intersection lines of liquidus surfaces in the ternary phase diagram usually form eutectic grooves, also referred to as liquidus boundary lines or monovariant lines, which open into a ternary eutectic point of the phase diagram. The pseudo-eutectic point represents a point on the liquidus boundary line, which forms a saddle point, i.e. a local extreme along the liquidus boundary line and a minimum perpendicular to it - in relation to the delimiting single-phase regions.

Sometimes two-component border systems or

Content cuts of, in particular ternary, phase diagrams, including binary eutectics, inconsistently referred to as pseudo-eutectic points. Such a conceptual one

However, designation is not in the sense of the present concept and explicitly not meant or designated by the nomenclature pseudo-eutectic point in this document or encompassed by this. In particular, the pseudo-eutectic point is characterized in that its existence requires an addition or presence of at least a third component or a third element.

With regard to Gibbs' phase rule, the pseudo-eutectic point pE in the ternary alloy system represents a local extremum along the liquidus boundary line which, compared to the ternary eutectic E, has a number of

Degrees of freedom and compared to a single-phase solidification MK has a number of degrees of freedom reduced by 1. With 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:

/ (M / O = 3-2 + 1 = 2

This higher degree of freedom of 1 at the pseudo-eutectic point pE compared to the degree of freedom of 0 at the ternary eutectic point E is the cause of the

Formation of the eutectic microstructure, often up to several orders of magnitude, finer in the area of the pseudo-eutectic point compared to the microstructure formed at the eutectic point.

In the case of an alloy with four components, the liquidus boundary line corresponds accordingly to a two-dimensional surface and the pseudo-eutectic point to a pseudo-eutectic line. For alloys with a corresponding higher component and more than four components, the dimensionality of the corresponding increases

Status areas in an analogous manner. In the context of this document, the designation pseudoeutectic point should therefore be understood in particular as a generic term, which includes both a pseudoeutectic point in a phase diagram of a ternary alloy and a corresponding pseudoeutectic line in a phase diagram of an alloy with four components or a corresponding pseudoeutectic line multidimensional area in one

Phase diagram of an alloy with more than four components. In this context, the designation pseudo-eutectic point and pseudo-eutectic area are used synonymously in particular. It goes without saying that a pseudo-eutectic point of a ternary alloy system represents a special variant.

In accordance with this explanation, in particular for the ternary alloy system, the following therefore applies to a pseudoeutectic point of a phase diagram of a ternary alloy or to a pseudoeutectic point, in particular one

pseudo-eutectic line or a pseudo-eutectic area, one

Phase diagram with more than three components according to Gibbs' phase rule: / (£) <f (pE) </ (M / 0,

so:

0 <f (pE) <N - 1.

The alloy composition at the pseudo-eutectic point or the

The pseudo-eutectic point of the phase diagram of the alloy with at least three components N is thus characterized in particular by the fact that, according to Gibbs' phase rule, the number of degrees of freedom f lies between 0 and N-1.

It has been shown to be both high strength and pronounced

The formability of the alloy is sufficient if the alloy composition is close to or in an area around the, in particular at, the pseudo-eutectic point or the saddle point represented by this, so that in the alloy at least 85 mol% or at% ( given in mol percent or atomic percent) eutectic structure is present. It is preferred 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 combined with good formability can be particularly pronounced.

Often up to 98 mol% or at% in particular can be achieved, so that the mechanical properties of the alloy are practically only determined by the eutectic microstructure. The eutectic structure usually forms during a liquid-solid phase transition or when the alloy solidifies.

High strength and pronounced formability of the alloy can be achieved both when the alloy is a ternary alloy and when the alloy has four components or at least five components. In particular, the alloy can have a large number of components, depending on the application objective, for example in the form of further added components for solid solution hardening and / or precipitation hardening. The alloy with high strength and formability can be formed particularly simply and practically if the alloy is ternary

Alloy or quaternary alloy.

The eutectic structure usually has average characteristics

Structural distances or an average distance between their phase components, especially lamellas, of less than 3 pm. A particularly pronounced strength and formability can be achieved if the average distance is less than 2 pm, in particular less than 1 pm. This can be achieved, for example, if the alloy composition of the alloy is selected in closer proximity to the stoichiometric composition of the pseudo-eutectic point. Particularly high strengths can be achieved if the average distance is less than

800 nm, in particular smaller than 600 nm. Additionally or alternatively, the average spacing of the phase components can be influenced by varying a cooling rate of the alloy when the alloy solidifies.

It is advantageous if the alloy has a residual solidification with a proportion of at most 5 mol% or at%, preferably at most 3 mol% or at%, particularly preferably at most 2 mol% or at%. -%, having. As a result, the aforementioned properties advantageously achieved by the eutectic structure are only insignificantly influenced or not influenced by the proportion of residual solidification. The proportion of residual solidification can be adjusted by choosing the alloy composition closer to the stoichiometric composition of the pseudo-eutectic point. Residual solidification usually denotes that part of the structure in which a residual part of the liquid phase solidifies after the formation of the eutectic structure in the form of a no longer eutectic structure or at the end of the eutectic solidification the number or type of the forming phase changes. The proportion of residual solidification usually represents a factor which adversely affects the properties brought about by the eutectic structure formed, which is why it is advantageous if the residual solidification is kept as small as possible. It is particularly favorable for this if the residual solidification is not configured in the manner of a network or in the form of a network structure, but rather, if present, in the form of islands or units that are separated from one another. As a rule, the residual solidification is formed in a proportion of at least 1 mol% or at%, but can preferably also be less.

It is useful for a pronounced strength and formability if the

Alloy has a primary solidification with a proportion of 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 formation of the eutectic structure or a formation of the eutectic structure with a high proportion of the structure advantageous to achieve the aforementioned properties. The primary solidification, which describes that part of the solidified microstructure which does not solidify in the form of a eutectic structure immediately before the formation of the eutectic structure, is less relevant than the aforementioned residual solidification with regard to an impairment of the properties to be achieved with the formation of the eutectic structure, but should are also preferably kept as small as possible. As a rule, the primary solidification is formed with a proportion of at least 1 mol% or at%, but can preferably also be less.

For a development of high strength and particularly pronounced formability, it is advantageous if the primary solidification is formed with or from a mixed crystal phase and in particular not with or not from an intermetallic phase. This appears to be an advantageous criterion for all alloy systems in order to achieve particularly user-friendly strength and formability properties.

The proportion of the aforementioned residual solidification and / or primary solidification can be in

can be checked or predetermined in the usual way with a thermodynamic calculation according to Scheil-Gulliver. The Scheil-Gulliver calculation or equation, sometimes just called Scheil calculation or equation, describes one

Distribution of an alloy component in an alloy during solidification, whereby a local equilibrium at an advancing solidification front and neglected diffusion in the solid phase is assumed. Such a calculation represents a customary tool or textbook knowledge in the field of metallurgy, which is assumed to be known to the person skilled in the art. For example illustrated in the textbook "Solidification" by J.A. Dantzig et al., (ISBN: 978-2-940222-17-9).

It is favorable if the alloy has a density of less than 8.0 g / cm ³ , in particular less than 7.5 g / cm ³ , preferably less than 6 g / cm ³ . As a result, the alloy can have a particularly advantageous strength-to-weight ratio in relation to an application, in particular as a structural component. It is particularly advantageous if the alloy is designed as a light metal alloy. This makes it possible to achieve a particularly high level of application suitability for the alloy. It is advantageous if the alloy for this purpose has less than 5.0 g / cm ³ , in particular less than 3.0 g / cm ³ . For practical use as application material, it is beneficial if the

Alloy a magnesium-based alloy, aluminum-based alloy,

Lithium-based alloy or titanium-based alloy is.

It is advantageous if the alloy is a cast alloy. This enables a particularly practicable production, in particular of structural components with in particular the aforementioned properties.

It has proven useful if the alloy is an Al-Mg alloy. Depending on the specific application, the alloy can have additional alloy components. In this way, application components are particularly relevant in practice, in particular

Structural components that can be produced with or from the alloy. It is particularly favorable here if the alloy is an Al-Mg-Si alloy. The alloy can also advantageously contain zinc (Zn), in particular with a proportion of more than 0.01% by weight, usually more than 1% by weight. This allows a compressive strength of the alloy to be optimized. The alloy usually has less than 15% by weight, in particular less than 10% by weight, preferably between 1.0% by weight and 5.0% by weight, particularly preferably about 3.0% by weight , Zinc on.

A high application suitability, for which both strength and pronounced formability 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 applicability, especially in the form of a

Structural component, which shows particularly high strength and formability, can be achieved if the alloy is a magnesium-based alloy, having, in particular consisting of, (in at .-%)

15.0% to 70.0% lithium,

more than 0.0%, in particular more than 0.01%, preferably more than 0.05%, aluminum,

Remainder magnesium and production-related impurities,

wherein a ratio of aluminum to magnesium (in at .-%) is 1: 6 to 4: 6

Such a Mg-Li-Al alloy has an alloy composition in an area around or in the vicinity of an alloy composition pseudoeutectic point in the Mg-Li-Al phase diagram, so that a

finely structured or microscale eutectic microstructure can be achieved. The fine-scale microstructure is associated with high strength, in particular high compressive strength, while at the same time good formability of the

Magnesium alloy is given with the corresponding aforementioned proportions of lithium in the magnesium alloy. Orientation composition or

The orientation line in the phase diagram is in particular a ratio of aluminum to magnesium (in atomic percent, abbreviated to atomic%) of approx. 3: 6, as this ratio results in a particularly homogeneous fine-scale or homogeneous fine lamellar microstructure or morphology . In a range around this ratio, especially with a ratio of aluminum to magnesium (in at .-%) of 1: 6 to 4: 6, the fine, especially fine lamellar, microstructure or morphology continues to be found with varying degrees which usually corresponds to different characteristics of a level of strength, in particular a level of a

Compressive strength, as well as formability or ductility of the magnesium alloy is connected. Because of this special morphological behavior in the specified

In the composition range, it is thus possible to form a magnesium alloy which has both high strength, in particular compressive strength, and good formability. This magnesium alloy and a process for its production as well as implementation as a starting material, semi-finished product or component as well as special design variants of these were included in the European patent application with the application number 19184999.1 and also in the international application with the international file number PCT / EP2020 / 058280

European Patent Office, the disclosures of which are hereby incorporated in their entirety into the disclosure of this document. This is particularly true if, as set out in these applications, the Mg-Al-Li alloy (in at .-%) 30.0% to 60.0%, in particular 40% to 50%, preferably 45% to 50%, particularly preferably 45% to 48%, lithium. It is also advantageous if the Mg-Al-Li alloy (in at.%) Has more than 0.05%, in particular more than 0.1%, generally more than 1% aluminum. It has been shown that the Mg-Al-Li alloy can be formed with a, in particular lamellar, microstructure with high fineness if the ratio of aluminum to magnesium (in at.%) 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 for a pronounced fineness or fine, especially lamellar, microstructure if the ratio of aluminum to magnesium (in At .-%) from 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 thereby be achieved. This applies particularly to a ratio of aluminum to magnesium (in at.%) Of 2.8: 6 to 3.3: 6, preferably about 3: 6, at which a very homogeneous fine

Morphology or microstructure is available. It is particularly advantageous if the magnesium alloy (in at.%) Is 30.0% to 60.0% lithium and a ratio of aluminum to magnesium (in at.%) Of 2.5: 6 to 3.5: 6, in particular 2.8: 6 to 3.3: 6, preferably about 3: 6. In this context, reference may be made in particular to FIG. 1 of the aforementioned application documents, in which a corresponding arrangement is shown schematically in a Mg-Li-Al phase diagram, and the disclosure and the associated description are likewise to be regarded as part of this document. A particularly pronounced homogeneity can also be achieved if the magnesium alloy (in at.%) Has 40.0% to 60.0% lithium. As set out in the aforementioned applications and accordingly included in the present disclosure, the properties of the Mg-Al-Li alloy can be further optimized if, in addition, proportions of calcium, rare earth metals, in particular yttrium, zinc and / or silicon according to the aforementioned applications with corresponding The salary ranges specified in the aforementioned registrations are available. For example, such an alloy can be used as Mg-20% Li-15% Al-1% Ca-0.5% Y (in% by weight) or Mg-20% Li-24% Al-1% Ca-0.5 % Y (in% by weight).

The further aim of the invention is achieved by a method of the type mentioned at the beginning with the proviso that the composition is provided lying in an area around a pseudo-eutectic point of a phase diagram of the alloy, so that when it cools into the solid phase or when it solidifies Alloy forms the eutectic structure with a proportion of at least 85 mol% or at%. As shown above, the alloy can thereby be formed with high strength and pronounced formability. By choosing the

Alloy composition in an area around the pseudo-eutectic point occurs when the alloy cools from the liquid to the solid state or during the liquid-solid transition, a eutectic phase reaction or phase transformation takes place, which the eutectic microstructure with particularly high fineness or fine structuring as an essential structural component of the Alloy forms. It goes without saying that the method according to the invention corresponding or analogously to the features, advantages, implementations and effects which are implemented within the framework of a

alloy according to the invention, in particular described above, can be formed. The same also applies to the alloy according to the invention with regard to a method according to the invention.

A starting material, semi-finished product or component is advantageously implemented with, in particular made of, an alloy according to the invention or obtainable by a method according to the invention for producing an alloy according to the invention. Corresponding to the above statements, features and effects of the invention

Alloy or an alloy produced using a method according to the invention also has a pre-material, semi-finished product or component formed with an alloy, advantageously high strength and good formability. Further features, advantages and effects result from the following

illustrated embodiments. In the drawings to which reference is made:

1 and 2 are phase diagram representations of an Al-Mg-Si system in which alloy compositions of alloy examples are given;

3 to 12 are optical microscopic photographs of alloy examples from FIGS. 1 and 2, respectively;

FIGS. 13 to 20 flow stress diagrams of alloy examples of FIGS. 1 to 12;

21 shows a phase diagram representation of an Al-Cu-Mg system with the alloy composition of an alloy example drawn in;

Fig. 22 is optical microscopic photographs of the alloy example of Fig. 21;

23 is a yield stress diagram of the alloy example of FIGS. 21 and 22, respectively; 24 is a phase diagram representation of a Mg-Al-Li system in which

Alloy compositions of alloy examples are given;

25 and 27 are optical microscopic photographs of alloy examples of FIG. 24; 28 and 29 flow stress diagrams of alloy examples of FIGS. 24 to 27; 30 shows a phase diagram representation of an Mg-Cu-Zn system with the alloy composition of an alloy example drawn in;

31 and 32 are optical microscopic photographs of the alloy example of FIG. 30; 33 is a yield stress diagram of the alloy example of FIGS. 30 to 32; Fig. 34 is an electron microscope photograph showing an alloy example of an Al-Cu-Mg-Zn system;

Fig. 35 is a yield stress diagram of the alloy example of Fig. 34;

Fig. 36 is a phase fraction diagram of an alloy example of an Al-Mg-Si-Zn system;

37 is a fixed component diagram of a Scheil-Gulliver calculation of the

Alloy example of FIG. 36.

As part of a development of the alloy according to the invention,

Experiments with different alloy compositions from

different alloy systems. In each case, alloys with an alloy composition in the area of or around one

Selected pseudo-eutectic point of a respective associated phase diagram and formed a eutectic structure by cooling the alloy from a liquid state to a solid state. The microstructure was then examined using microscopy. In addition, various dilatometric test series or

Compression tests, carried out as standard at room temperature, around 20 ° C., the result being that flow curves were determined which show a yield stress, in MPa, as a function of a degree of deformation, represented as a proportion of a change in length AL relative to a

Initial length L ₀ , ie -, and correspondingly dimensionless.

i-0

In the following, representative test results for alloy examples 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 order to clarify the aforementioned concept on a broad basis. Al-Mg-Si system:

1 and 2 show representations of a ternary phase diagram of an Al-Mg-Si system, FIG. 2 being a detail representation of the phase diagram in order to show the relevant alloy composition range in detail. There were ten Alloy examples of the Al-Mg-Si system produced and investigated. The alloy compositions of the alloy examples of the Al-Mg-Si system are given in Table 1 as alloy example 1 to alloy example 10, in each case in weight percent and atomic percent and correspond to the reference numerals 1 to 10, which in particular in the phase diagram of FIGS. 1 and 2 the respective

Designate alloy composition.

Table 1: Ten alloy examples from the Al-Mg-Si alloy system.

As can be seen in the phase diagram of FIG. 1 and FIG. 2, the alloy examples 8 to 10 each have compositions which, in an area around one

pseudoeutectic point pE are arranged, the alloy examples 8 and 9 very close to the pseudo-eutectic point and the alloy example 10 are positioned at a somewhat greater distance from the pseudo-eutectic point pE. The

Alloy composition of alloy example 9 is practically at the pseudo-eutectic point pE. The pseudoeutectic point pE is illustrated in FIG. 2 with a drawn reference line, the pseudoeutectic point pE being at the intersection of the monovariant line in the direction of AhMg2 and the reference line. It can also be seen in FIG. 2 that the alloy examples 3 to 5 are arranged in an area around a eutectic point E of the phase diagram. Alloy Examples 6 and 7 are also provided as comparisons, their

Compositions are located at a great distance from the pseudo-eutectic point pE, as can be seen in FIG. 2, as well as the alloy examples 1 and 2, which are positioned in the immediate vicinity of a liquidus boundary line, but at a greater distance from both the pseudo-eutectic point pE and the eutectic point E. in Fig. 1.

In FIGS. 3 to 12, optical microscope images of the alloy examples 1 to 10 are shown in order to illustrate a respective microstructure. In FIGS. 13 to 20, flow stress diagrams are shown as the results of dilatometric test series of the Al-Mg-Si alloy examples, which were carried out at room temperature, about 20 ° C. Yield stress curves are shown, with a yield stress, in MPa, being shown as a function of the degree of deformation. Each of the

Yield stress diagrams show several yield stress curves of alloy samples with an alloy composition corresponding to the alloy composition of one of alloy examples 1 to 10. Each flow stress diagram therefore represents an alloy composition of one of alloy examples 1 to 10.

As can be seen in FIG. 10 to FIG. 12, the microscopic images of FIG

Alloy examples 8 to 10, which have alloy compositions in the vicinity or an area around the pseudo-eutectic point pE, a dominant, finely structured or finely-scaled eutectic structure. In comparison, in FIGS. 6 and 7, microscopic photographs of alloy examples 4 and 5, which have an alloy composition in the vicinity of the eutectic point E, can be observed. These show an expression of a eutectic structure, which in comparison to the structure of the alloy examples 8 and 9 is coarse Has structure. If one compares this with those shown in FIGS. 3 and 4

Microscopic photographs of alloy examples 1 and 2, their

Alloy compositions are at a great distance, but in the area of a liquidus boundary line, it can be seen that these show an even coarser eutectic structure. 8 and 9 are also microscopic photographs of alloy examples 6 and 7, which alloy compositions in a

Far range of the pseudo-eutectic point pE or at a great distance to this have shown. It can be seen that a eutectic structure is already present, but with a relatively coarse structure and significantly less dominant or with a lower proportion. A high proportion of residual solidifications can also be seen, recognizable in FIGS. 8 and 9 in the form of light-colored channels.

13 and 14 show yield stress diagrams of Alloy Examples 8 and 9 which have alloy compositions in the vicinity and an area around the pseudo-eutectic point pE, respectively. It can be seen that both alloy example 8 and alloy example 9 have high strength, in particular compressive strength, and pronounced formability with yield stresses between 300 MPa and 400 MPa, with alloy example 8, shown in FIG. 13, in particular.

Shows yield stresses of up to 400 MPa. In comparison, in FIGS. 15 and 16, yield stress diagrams of alloy examples 4 and 5, which alloy compositions have in the vicinity of the eutectic point E, can be observed. Alloy examples 4 and 5 also show high strength and, at least depending on the individual sample, high formability, with yield stresses below those of alloy examples 8 and 9 at approximately 300 MPa or in relation to

Alloy Example 5 shown in Figure 16 is consistently below this. This result correlates with the finding that alloy examples with a

Alloy composition in the area of the pseudo-eutectic point pE has a particularly high fine structuring of its eutectic structure, especially in comparison to the eutectic structure of alloy examples

Have alloy compositions in the area of a eutectic point E, which also explains the higher strength or pronounced ductility of alloys in the area of the pseudoeutectic point. In addition, FIG. 20 shows a flow stress diagram of the alloy example 10, the alloy composition of which is arranged at a somewhat greater distance from the pseudo-eutectic point pE. Slightly lower yield stress values and, in particular, a higher scatter between the individual measurement results can be seen. It is further shown in FIGS. 17 and 18 that, in comparison therewith, alloy example 1 and alloy example 2 with alloy compositions in the range of a

Liquidus boundary line, however, both remote from the alloy composition of the pseudo-eutectic point pE and eutectic point E, have significantly poorer strength and deformation properties. In Fig. 19 is also a

Yield stress diagram corresponding to the alloy composition of alloy examples 6 and 7, the alloy composition of which is positioned at a relatively great distance from that of the pseudo-eutectic point pE in the phase diagram, is shown. The corresponding yield stress curves show clearly reduced yield stresses in comparison to the yield stresses of an alloy composition closer to the pseudo-eutectic point pE, such as that of the alloy example 8 shown in FIG. 13.

It can be seen that an alloy composition in an area around a pseudo-eutectic point pE corresponds to a finely structured eutectic microstructure or correspondingly high strength and pronounced formability.

In a detailed examination it can be seen that the alloy example 8 in

Phase diagram of FIG. 2 relative to the monovariant line or liquidus boundary line in the direction AhMg2 above this in the Mg2Si range, which is why solidification begins when the alloy cools from the liquid with in particular an undesired formation of Mg2Si or primary solidification with an intermetallic phase Mg2Si is formed. It has been shown that a primary solidification formed with an intermetallic phase has negative effects on the formation of both high strength and formability. In order to achieve particularly advantageous strength and formability, the aim is therefore mostly to keep primary solidification with or from an intermetallic phase as small as possible or to avoid it. However, the primary solidification in alloy example 8 is so low that it practically does not lead to any inhibition of mechanical properties. The

Microscopic photographs of alloy example 8 in FIG. 10 show large areas with a fine eutectic structure, in this case formed with Al mixed crystal and Mg Si. A residual solidification of Al mixed crystal is also advantageous only very slightly or hardly present. In order to avoid undermining the advantageous strength and deformability properties that can be achieved with the eutectic structure, the aim is to keep residual solidification as small as possible or to avoid it. In particular, the residual solidification is not connected in the manner of a network or is designed in the form of units separated from one another, which likewise promotes an advantageous design of high strength and pronounced formability. Alloy example 8 thus proves to be both low and low in terms of residual solidification

Primary solidification as well suited to control strength properties or formability on the basis of the fine eutectic structure. This can be further optimized if the alloy composition is chosen such that the primary solidification is formed with or from a mixed crystal phase and not with an intermetallic compound or phase, i.e. the primary solidification in the case of alloy example 8 is in the Al mixed crystal range .

This consideration of alloy example 8 and the associated explanations apply in an analogous manner to alloy example 9. Alloy example 9 has an alloy composition practically at the pseudo-eutectic point pE. As can be seen in FIG. 11, alloy example 9 also shows a fine eutectic structure with hardly any residual solidification and hardly any primary solidification. The somewhat lower strength in comparison with alloy example 8 is explained by the lower dissolved content of Mg in the Al mixed crystal phase. A strength can advantageously be achieved by varying a proportion of dissolved elements in the mixed crystal phase, with the primary solidification, however, as stated above, preferably in

Solid solution range and not in the range of an intermetallic phase.

In comparison with this, alloy example 10, which can be seen in FIG. 12, also shows a fine eutectic structure, but with a higher proportion of residual solidification, in the form of Al mixed crystal and Si, which is also characterized by a network. Due to the low Mg content, most of the Mg is bound in the form of Mg Si, so that solid solution hardening of the Al solid solution phase is very little pronounced. This corresponds to lower yield stresses in the yield stress diagram in FIG. 20. On further detailed consideration of the alloy examples arranged away from the pseudo-eutectic point pE with respect to an alloy composition, it can be seen that the alloy examples 4 and 5, which are in the region of the eutectic point E, shown in FIGS. 6 and 7, have a small proportion of primary solidification, around which a relatively coarse eutectic structure, formed with two phases, is arranged. The remaining, predominant part of the eutectic structure is designed as a ternary eutectic, formed with mixed crystal, Al2Si and Si. The mechanical properties, in particular strength and formability, are negatively influenced in particular by the coarse binary eutectic structure or phase. In some places there is a fine eutectic ternary structure, which in places merges into highly coarse structures. The differences between the microstructural structures of alloy examples with alloy compositions at or in the area of the pseudo-eutectic point pE compared to those at or in the area of the eutectic point E correlate with the finding of correspondingly improved strength and formability properties of alloy compositions at or in the area around the pseudo-eutectic point pE.

It can also be seen that the alloy examples 6 and 7 with the associated microscopy images shown in FIGS. 8 and 9 have coarse, polygonal primary solidifications. This is explained by the positioning of the associated

Alloy compositions in the Mg2Si range of the phase diagram, so that a pronounced Mg2Si primary solidification is formed. A coarse eutectic structure and a high proportion of residual solidification can be seen between these, which can be seen in the light areas or channels in FIG. 8 and FIG. 9

Alloy examples 6 and 7 have significantly reduced structural morphology

Strengths or yield stresses, which are associated in particular with crack initiation and brittle fracture.

In Fig. 2, as a gray, flat area is a particularly advantageous area for

Formation of an Al-Mg-Si alloy is shown. This characterizes or essentially corresponds to an aforementioned alloy composition of

Alloy Examples 8 and 9, but with a variation in the

Alloy composition such that a primary solidification

Mixed crystal phase and in particular no intermetallic phase forms. This enables the formation of particularly high strengths with pronounced

Formability. A particularly advantageous conversion range for such an Al-Mg-Si alloy is therefore given when the Al-Mg-Si alloy is arranged in the Al-Mg-Si phase diagram in an area around the pseudo-eutectic point, the

Alloy composition in the phase diagram based on the aforementioned pseudo-eutectic point of the phase diagram of FIG

increasing Al proportion facing side of the corresponding monovariant line is arranged.

Al-Cu-Mg system:

Fig. 21 shows a representation of a ternary phase diagram of an Al-Cu-Mg system. An alloy example of the Al-Cu-Mg system was prepared and examined. The associated alloy composition is shown in Table 2 as

Alloy example 13 indicated in weight percent and atomic percent and

corresponds to reference number 13, which in particular denotes the alloy composition in the phase diagram of FIG. 21.

Table 2: Alloy example from the Al-Cu-Mg alloy system.

As can be seen in the phase diagram of FIG. 21, the alloy example 13 has an alloy composition which, in an area around a

pseudo-eutectic point pE is arranged. Using optical

An associated microstructure is illustrated in FIG. 22 using microscopic images. A very fine-scale eutectic structure and a small amount of primary solidification formed with mixed crystal can be seen. In Fig. 23 is a

Yield stress diagram as the result of dilatometric test series of Al-Cu-Mg alloy example 13 is shown, again showing a yield stress, in MPa, as a function of the degree of deformation. It can be seen that very high

Strengths or yield stresses can be achieved. The elongation at break is also in the technologically relevant range for this alloy system. Strength or Formability corresponds to the fine eutectic microstructure and in particular the small amount of primary solidification.

Mg-Al-Li system:

24 shows a representation of a ternary phase diagram of a Mg-Al-Li system. Three alloy examples of the Mg-Al-Li system were prepared and examined. The alloy compositions of the alloy examples of the Mg-Al-Li system are in Table 3 as alloy examples 14, 15 and 16, respectively, in percent by weight and

Atomic percent and correspond to the reference numerals 14, 15 and 16, which in particular in the phase diagram of FIG

Designate alloy composition.

Table 3: Three alloy examples from the Mg-Al-Li alloy system.

As can be seen in the phase diagram in FIG. 24, the alloy examples 14 to 16 each have an alloy composition which is arranged in a region around a pseudo-eutectic point pE. The pseudo-eutectic point pE is illustrated in FIG. 24 with a reference line drawn in, the

pseudo-eutectic point pE at the intersection of the monovariant line or

Liquidus limit line and the reference line is located. With admixtures of CaY, in particular about 1 wt .-% Ca and about 0.5 wt .-% Y, can be practicable

Oxidation properties of the alloy examples of the Mg-Al-Li system are stabilized without adversely affecting a microstructure.

Alloy examples 14 and 15 lie somewhat closer in the phase diagram in a near area of the pseudo-eutectic point, alloy example 16 somewhat further away, the alloy composition of alloy example 14 being approximately at pseudo-eutectic point pE is positioned. The alloy examples 14 to 16 are, according to the current data situation, in a mixed crystal region, in particular forming a body-centered cubic lattice, bcc.

With the aid of microscopic recordings, structural structures can be seen in FIGS. 25 to 27. The microstructure morphology of FIGS. 25 and 26 indicates a

Formation of an extremely fine-scale structure which can no longer be resolved in the light microscope used. The grain boundaries that can be recognized are oxidic

Due to impurities. The microstructure of alloy example 16 was examined by means of scanning electron microscopy, shown in FIG. 27. Visible in FIG. 27 are, on the one hand, light grain boundary phases (in whitish-gray), which were identified as Al-Ca, and, on the other hand, pronounced fine crystal structures or

Morphologies in an area surrounded by grain boundary phases, in particular in a central section of this area, or in the interior of the mixed crystal, clearly visible in particular in the right image of FIG. 27. In the phase diagram of FIG. 24, the alloy composition of alloy example 16 appears to be relatively far from the mono-variant Line or the pseudo-eutectic point pE away. In this case, however, it should be noted that according to known specialist knowledge, the slope in the area of the body-centered cubic lattice, bcc, - in which also that

Alloy example 16 is arranged - is very flat in the phase diagram and also the three elements Mg, Al and Li have a high solubility in one another. This explains why there is a correspondingly extensive area around the

pseudo-eutectic point results in which an advantageous fine-scale eutectic microstructure can be formed with a high proportion.

28 and 29 show flow stress diagrams of alloy examples 15 and 16 as the results of dilatometric test series, again showing a flow stress, in MPa, as a function of the degree of deformation, FIG. 28

Yield stress curves relating to alloy example 15 and FIG. 29

Fig. 16 shows yield stress curves relating to alloy example 16. It can be seen that both alloy examples have high strengths or flow stresses as well as pronounced formability, corresponding to the fine eutectic microstructures determined. In FIG. 29, relating to alloy example 16, a possibility of further optimization of properties by means of heat treatment is also shown. 29 shows flow curves of alloy samples immediately after a production of the alloy example 16 (as-cast), shown in FIG. 29 as solid lines, identified by reference symbols 16-1, as well as flow curves from FIG

Alloy sample samples after heat treatment (aged) of the

Alloy example 16 shown in Fig. 29 as dashed lines denoted by reference numeral 16-2. For this purpose, samples of alloy example 16 were used

Subjected to heat treatment at 330 ° C for 3 hours and then flow curves determined by means of pressure tests. A clear influence of the

Heat treatment for strength, in particular compressive strength, and formability, whereby the potential is given to optimize compressive strength and formability, in particular for a later application, through heat treatment

adjust.

As already explained above in the document, it has proven to be advantageous for the implementation of an alloy with high application suitability if the alloy has a

Magnesium-based alloy is, having, in particular consisting of, (in at .-%)

15.0% to 70.0% lithium,

more than 0.0%, in particular more than 0.01%, preferably more than 0.05%, aluminum, the remainder magnesium and production-related impurities,

wherein a ratio of aluminum to magnesium (in at .-%) is 1: 6 to 4: 6. The alloy example 16 can be seen as a representative example for this alloy definition, as in the context of the European patent application with the

Application number 19184999.1 as well as in the context of the international

Application with the international file number PCT / EP2020 / 058280, which was filed with the European Patent Office, is shown. In particular, reference may again be made to the respective FIG. 1 of these applications. In FIG. 24, a corresponding ratio of aluminum to magnesium (in at.%) Of 1: 6 is shown as a dashed line. The aforementioned ratio range of aluminum to magnesium (in at.% Or mol%) from 1: 6 to 4: 6 is located in the phase diagram of FIG. 24 to the left of this line and in particular represents a special one

Variant in the area around the pseudo-eutectic point pE.

A particularly advantageous implementation range for a Mg-Li-Al alloy, which can be used as an application alloy, in particular for a structural component, is given, if the Mg-Li-Al alloy in the Mg-Li-Al phase diagram is in a range between the line indicating a ratio of aluminum to magnesium (in at .-%) of 1: 6 and the monovariant line or liquidus boundary line, in particular with an aforementioned Li content range. One such area is in

The phase diagram of FIG. 24 is identified as a gray, flat area.

As was already the case with the alloy examples of the Al-Si-Mg system, it is shown that an alloy composition is preferably selected such that the alloy composition is in the area of the pseudo-eutectic point pE and preferably also has primary solidification with or from a mixed crystal phase , i.e. the corresponding alloy composition is positioned in a mixed crystal area in the phase diagram.

Mg-Cu-Zn:

Fig. 30 shows a representation of a ternary phase diagram of a Mg-Cu-Zn system. An alloy example of the Mg-Cu-Zn system was prepared and examined. The associated alloy composition is shown in Table 4 as

Alloy example 17 indicated in percent by weight and percent by atom and

corresponds to reference symbol 17, which in particular denotes the alloy composition in the phase diagram of FIG. 30.

Table 4: Alloy example from the Mg-Cu-Zn alloy system.

Alloy compositions, which in an area around a

pseudo-eutectic point pE is arranged. Using optical

Microscopic recordings show an associated microstructure in FIGS. 31 and 32. A very fine-scale eutectic microstructure can be seen, which is at the limit of a light microscopic resolution. A relatively large proportion of primary solidification can be seen here. It is therefore advantageous for high strength and formability if an alloy composition is even closer to pseudoeutectic point pE or closer to the monovariant line or

Liquidus limit line is chosen.

Fig. 33 shows a yield stress diagram as results of dilatometric

Test series of alloy example 17, again showing a yield stress, in MPa, as a function of the degree of deformation. It can be seen that high strengths or yield stresses are achieved, which, however, in view of the pronounced proportion of primary solidification visible in the microscope images, can be further improved by choosing an alloy composition closer to the pseudo-eutectic point pE. In FIG. 33, flow curves of the alloy example 17 are shown immediately after a production of the alloy example 17 (as-cast), denoted by the reference symbol 17-1, as well as flow curves of the

Alloy example 17 after the heat treatment has been carried out, identified by reference symbol 17-2. For this purpose, samples of alloy example 17 were used

Heat treatment at 350 ° C. for 4 hours and then flow curves determined by means of pressure tests. A clear influence of the

Heat treatment for strength and formability, which gives the potential to further optimize strength and formability by means of heat treatment.

Subsequently, studies of quaternary

Alloy systems or quaternary eutectics carried out. For this were

in particular the alloy systems Al-Cu-Mg-Zn and Al-Mg-Si-Zn are considered.

Al-Cu-Mg-Zn:

With respect to the Al-Cu-Mg-Zn alloy system, an in the area of

Pseudoeutectic point pE located alloy example produced and investigated. The alloy composition is shown in Table 5 as alloy example 18 in

Weight percent and atomic percent indicated and corresponds to

Reference number 18.

Table 5: Alloy example from the Al-Cu-Mg-Zn alloy system.

To investigate the eutectic structure, electron microscopy images of alloy example 18 were taken, shown in FIG. 34. A finely structured eutectic structure, in particular with structural dimensions in the nanometer range, can be clearly seen in the right-hand image of FIG. 35 as an extensive granular area in the center of the Image.

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

In FIG. 34, substructures can be seen in the primary regions (in gray), which are artifacts of an isostoichiometric structural conversion (bcc to fcc) in the solid state. In terms of a direct influence on strength and

Formability are irrelevant. A relatively large proportion of primary solidification in the form of a mixed crystal phase (in light gray to whitish) and an intermetallic secondary phase (in black), in particular in the form of a Laves phase, can also be seen.

35 shows a flow stress diagram as a result of dilatometric

Experiments with alloy example 18. Shown are flow curves before heat treatment carried out, denoted by reference number 18-1, and flow curves after carried out heat treatment, denoted by reference number 18-2, again showing a yield stress, in MPa, as a function of the degree of deformation. It can be seen that the alloy example 18 has a very high strength with a simultaneous elongation at break, with deformability by means of

Heat treatment is variable.

The existing pronounced primary solidification and secondary phase are as

To assess factors that increase brittleness, which is why it would be advantageous to further reduce these proportions in order to further improve strength and formability,

for example by reducing the distance to or further approximating the alloy composition in the phase diagram to the pseudo-eutectic point pE. Al-Mg-Si-Zn:

With respect to the Al-Mg-Si-Zn alloy system, one in the area of one

Alloy example located at the pseudo-eutectic point pE was examined by means of simulation. The alloy composition is given in Table 6 as alloy example 19 and corresponds to reference symbol 19.

Table 6: Alloy example from the Al-Mg-Si-Zn alloy system.

As a result of the simulation, phase components as a function of the temperature of the alloy example 19 are shown in FIG. A direct transition from solid to liquid phase can be seen, corresponding to the formation of a eutectic structure. Corresponding to this, FIG. 35 shows a representation of the solid fraction determined by means of the Scheil-Gulliver solidification calculation as a function of the temperature. The equilibrium and Scheil-Gulliver solidification curves shown show an alloy system which shows binary eutectic solidification with four components or elements. Accordingly, there is again an increase in the thermodynamic degree of freedom from 1 to 3. The Scheil-Gulliver calculation in FIG. 37 shows a very small proportion of primary solidification in the form of a mixed crystal with a proportion of less than 3 mol% or at% and also a virtually non-existent residual solidification.

It is thus shown in an analogous manner that, in addition to positioning the alloy composition in the area of the pseudo-eutectic point pE, a proportion of primary solidification and / or residual solidification can also be kept small in order to further increase or improve strength properties or deformation properties.

An alloy according to the invention with more than three components with a eutectic structure produced by cooling from the liquid state to the solid state is therefore advantageous with a finely structured eutectic structure,

in particular with a fine structure in the nanometer range, can be formed which has a forms dominant or substantial phase portion or structural portion in the alloy when an alloy composition of the alloy is arranged in the phase diagram in the area of or around a pseudo-eutectic point. In this way, the alloy can be formed with advantageously high strength and pronounced formability. This is particularly true if there is very little primary solidification and / or residual solidification. In particular, it is favorable for this if the primary solidification is formed with or from a mixed crystal, in particular not with or from an intermetallic phase, or the alloy composition is selected in a corresponding area in the phase diagram. An alloy formed in this way thus offers the potential for robust and resistant components, in particular

Construction elements, in particular for a purpose in the

Automotive, aircraft and / or space industries are preferred

purpose-dependent to implement.

Claims

1. Alloy, in particular a light metal alloy, comprising a

Alloy composition with at least three components and a eutectic structure, which is obtained by cooling the alloy from a liquid state to a solid state, with the proviso that a composition of the alloy lies in a region around a pseudo-eutectic point (pE) of a phase diagram of the alloy so that there is at least 85 mol% eutectic structure in the alloy.

2. Alloy according to claim 1, characterized in that the eutectic structure has an average distance between their phase components of less than 3 pm, preferably less than 1 pm.

3. Alloy according to claim 1 or 2, characterized in that the alloy has a residual solidification with a proportion of a maximum of 5 mol%, preferably a maximum of 3 mol%.

4. Alloy according to one of claims 1 to 3, characterized in that the alloy has a primary solidification with a proportion of less than 10 mol%, in particular less than 5 mol%.

5. Alloy according to claim 4, characterized in that the primary solidification is formed with a mixed crystal phase.

6. Alloy according to one of claims 1 to 5, characterized in that the alloy has a density of less than 8 g / cm ³ .

7. Alloy according to one of claims 1 to 6, characterized in that the alloy is a magnesium-based alloy, aluminum-based alloy,

Lithium-based alloy or titanium-based alloy is.

8. Alloy according to one of claims 1 to 7, characterized in that the alloy is a cast alloy.

9. Alloy according to one of claims 1 to 8, characterized in that the alloy is an Al-Mg-Si alloy.

10. Alloy according to claim 9, characterized in that the Al-Mg-Si alloy has between 0.01% by weight and 5.0% by weight, in particular about 3.0% by weight, zinc.

11. A method for producing an alloy, in particular an alloy according to one of claims 1 to 10, with a eutectic structure, wherein the alloy has an alloy composition with at least three components and wherein the alloy for the formation of the eutectic structure starting from a liquid state in a solid state of the alloy is cooled, with the proviso that the alloy composition is provided lying in an area around a pseudo-eutectic point (pE) of a phase diagram of the alloy, so that when cooling into the solid phase, the eutectic structure with a proportion of at least 85 mol -% trains.

12. Pre-material, semi-finished product or building material with an alloy according to one of claims 1 to 10 or obtainable with a method according to claim 11.