An aluminium alloy and a method of producing an aluminium alloy

Abstract

An aluminium-based alloy having the following composition: Albal. Mna Sin Xc Ze TMd, where X is one or more elements selected from the group consisting of Cu, Ni, Zn, Fe and Co; Z is one or more elements selected from the group consisting of Li, Ca, Sr, Ga, REE, mischmetals, Ce, Bi, Sn, Sb and Mg; TM is one or more elements selected from the group consisting of Cr, Ti, W, Mo, Zr and V; wherein a, n, c, e and d satisfy, by weight percent: 3≤a≤10, 0.1≤n≤3, 0.1≤c≤10, 0.1≤e≤8, and 0.1≤d≤3, and wherein 5≤a+c+e≤20; the balance being Al and inevitable impurities up to 0.5 wt.%, and comprising a metastable quasicrystal phase that is dispersed in a matrix composed of aluminium.

Description

[0001] An aluminium alloy and a method of producing an aluminium alloy TECHNICAL FIELD The present invention relates to a non-ferrous aluminium-based alloy having high strength and high hardness by constituting a structure in which, as a result of quenching of a melt having the composition defined by the present invention, at least a part is composed of a metastable quasicrystal phase that is dispersed in a matrix composed of aluminium. The aforesaid quasicrystals consist of an icosahedral phase (iQc phase) alone or a mixed phase of an iQc and a decagonal phase (dQC phase). The above structure is preferably such that the quasicrystals, various intermetallic compounds formed from aluminium and added elements are homogeneously and finely dispersed in the matrix composed of aluminium. BACKGROUND AND PRIOR ART Aluminium-based alloys with high strength (high-strength and light-weight) have attracted rapidly increasing interest because of the increase in the energy and environmental problems on the Earth. Aluminium parts can be produced economically using versatile manufacturing technologies. Since the discovery of an age-hardening phenomenon [1], various kinds of high- strength / high-performance Al-alloys have been produced. It is known that these Al-alloys have been developed by use of one or more of the following strengthening mechanisms: grain size reduction (modification, inoculation, deformation); introducing a hardening phase (intermetallic in a form of particles, flakes, whiskers etc.); strain hardening or the introduction of defects (dislocations, kinks, jogs, Lomer-Cottrell locks, point defects); impairing grain- boundary slide or creep; precipitation hardening (continuous, discontinuous); solution hardening; phase transformation and massive transformation. However, by using standard strengthening mechanisms limiting values of tensile strengths between 550 MPa and 600 MPa can be achieved in wrought alloys, but with reduced elongation (around 10 %). In Al-Si casting alloys alloyed with different alloying additions (e.g. AlSi12CuNiMg) tensile strengths up to 400 MPa can be attained, and the elongation is around 5 %. [2] Various attempts have been made recently to use a completely different strengthening mechanism, achieved by rapidly solidification of a molten aluminium base alloy, to impose the metastable solidification of the aluminium alloy to improve the mechanical and chemical properties of both wrought and cast aluminium alloys. In the last decade, alloys strengthened with a metastable phase (amorphous or quasicrystalline phases) have achieved a tensile strength exceeding 1000 MPa. However, these high strength values have been achieved in melt-spun strip alloys and extruded atomized powders, where cooling rates of about 106K / s are applied.[3- 5] For practical applications it would be preferred to produce an Al-based alloy with high strength in bulk form by conventional casting technique. In the present patent application we are presenting an alloy composition, which, when produced by use of economically versatile casting technologies enabling cooling rates from 500 K / s, or even as low as 100 K / s, to 100000 K / s, preferably to 10000 K / s, or even more preferably to 1000 K / s, results in an alloy reinforced with a metastable quasicrystalline (Qc) phase. First discovered quasicrystals in Al–Mn alloys system were thermodynamically metastable and formed only when the melt was rapidly cooled (~106K / s).[6] A few years later, it was shown that less rapidly solidified Al–Mn alloys forms a metastable decagonal Qc. Only in the case of even slower cooling rates a stable phase forms, which is periodic then [7, 8]. Figure 1 in the appended drawing summarizes the temperature-time-transformation (TTT) diagrams of Al88Mn12alloy for the Al6Mn, dQc and iQc phase, respectively. Higher cooling rates leads to the fraction of dQc growing at the expense of Al6Mn until no crystalline phase can occur at a critical cooling rate of 6.4 x 105K / s. At such cooling rates a phase mixture of the quasicrystalline dQc phase and iQc phase should occur. The faster crystallization of the iQc, compared to the periodic and dQc phase, was explained by its lower nucleation barrier, because atoms in the melt already existing in the icosahedral short-range order. However, the easy nucleation of structural subunits (clusters) with icosahedral symmetry does not necessarily imply quasiperiodic growth. The existence of Qc approximants clearly demonstrates that the same kind of cluster can basically constitute both quasiperiodic and periodic structures. The dominating factors determining whether a structure gets periodic or quasiperiodic are the chemical composition and cluster stacking sequence, which depends on the solidification conditions [9, 10]. Solidification conditions for versatile casting technologies do not vary much with process parameters and usually allow cooling rates from few 10 K / s to few 1000 K / s.. Therefore, to obtain bulk quasicrystal reinforced alloys, the chemical composition must be tailored. The most promising process to produce aluminium alloys reinforced with Qc phase in a bulk form is to alloying the aluminium melt with elements that reduce the cooling rate needed for primary metastable Qc formation. Different alloys were disclosed in reports by Zupanič et al.[11-13], Bončina et al.

[0014] , Markoli et al.[15, 16], Naglič et al.

[0017] , and Stan-Głowińska et al.[18, 19] the effect of different alloying elements on the formation of the primary metastable iQc phase in Al–Mn–X alloys (X = B, Be, C, Ca, Cu, Cr, Co, Ni, Ti, Sr, Fe, Si and Mg) was evaluated. They discovered that additions of boron, carbon, calcium, chromium, cobalt, titanium and strontium could not be proven to be effective in promoting the formation of iQcs in cast Al–Mn-based alloys. In contrast, beryllium, iron, copper, nickel, magnesium and silicon proved to be highly effective in promoting iQc formation at cooling rates between 500 and 1350 K / s. WO 2022 / 080319 discloses aluminium alloys that form metastable quasicrystals provided that the cooling rate is sufficiently high. Also Journal Of Material Science, Kluwer academic Publisher, Dortrecht, vol. 52, no. 23, 20 August 2017 (2017-08-20), pages 13657-13668), discloses aluminium alloys that form metastable quasicrystals provided that they are subjected to sufficiently high cooling rates. WO 2022 / 122670 is another example of a document that discloses aluminium alloys that form metastable quasicrystals provided that they are subjected to sufficiently high cooling rates. THE OBJECT OF THE INVENTION It is an object of the present invention to provide an aluminium-based alloy, that, upon solidifying and cooling with low cooling rates, for example from 500 K / s and higher, results in a material having superior strength and hardness. It is an object to present an alloy that, upon being produced by a casting technique using relatively low cooling rates (or at least not extremely high cooling rates), typically less than a 1000 K / s, results in a material that comprises: elements which form quasicrystals; an element which stabilize an icosahedral quasicrystalline phase; an element which promotes the formation of the quasicrystals; an element that reduces the surface tension between aluminium-based melt and the growing quasicrystalline phase; and an element with low diffusivity in the aluminium- based melt that stabilizes the supercooled state of the molten metal and further delays the crystallization of a crystal phase. It is also an object to present a method for the production of a material as defined hereinabove. SUMMARY The object of the invention is achieved by means of an aluminium-based alloy having the following composition: Albal. Mna Sin Xc Ze TMd, where X is one or more elements selected from the group consisting of Cu, Ni, Zn, Fe and Co; Z is one or more elements selected from the group consisting of Li, Ca, Sr, Ga, REE, mischmetals, Ce, Bi, Sn, Sb and Mg; TM is one or more elements selected from the group consisting of Cr, Ti, W, Mo, Zr and V; wherein a, n, c, e and d satisfy, by weight percent: 3≤a≤10, 0.1≤n≤3, 0.1≤c≤10, 0.1≤e≤8, and 0.1≤d≤3, and wherein 5≤a+c+e≤20; the balance being Al and inevitable impurities up to 0.5 wt.%, and comprising a metastable quasicrystal phase that is dispersed in a matrix composed of aluminium. According to one embodiment, 0.4≤n+d≤5, by wt.%. According to one embodiment, 5≤a. According to one embodiment, a≤8. According to one embodiment, 0.2≤n. According to one embodiment, n≤1.7. According to one embodiment, 0.5≤c, or, preferably, 1.0≤c. According to one embodiment, c≤8, or, preferably, c≤5. According to one embodiment, X is one or more elements selected from the group consisting of Cu, Ni and Zn. According to one embodiment, 0.3≤e. According to one embodiment, e≤6, or, preferably, e≤4. According to one embodiment, Z is one or more elements selected from the group consisting of Mg, Li, REE, and mischmetal. According to one embodiment, 0.4≤d. According to one embodiment, d≤2, or, preferably, d≤1.2. According to one embodiment, TM is selected from one or more elements selected from the group consisting of Ti, W, Mo, Zr and V. The object of the invention is also achieved by means of a method of producing an alloy comprising the steps of forming a melt of an alloy having the composition as disclosed hereinabove or hereinafter, and cooling the melt with a cooling rate of 500 K / s or more. According to one embodiment, the cooling rate is less than 5000 K / s. According to yet another embodiment, the cooling rate is less than 1000 K / s. Provided that the cooling rate is at least 500 the result will be a microstructure having a quasicrystalline phase with a super- saturated solid solution phase of aluminium, and various intermetallic compounds, which is a preferred microstructure resulting in higher strength and hardness. In the achieved structure the quasicrystals, various intermetallic compounds formed from aluminium and added elements are preferably homogeneously and finely dispersed in the matrix composed of aluminium as a result of the selected composition and cooling rate for the selected composition. DETAILED DESCRIPTION OF EMBODIMENTS In the following, the reasons for limiting the composition ratio of each component in the alloy according to the present invention are explained. Al (aluminium):Al≥75 wt.%, preferably Al≥80 wt.%. Al≤94,6 wt.%, preferably Al≤93 wt.%. Al weight percentage less than 75% leads to alloy embrittlement. On the other hand, an Al weight percentage greater than 94.6% leads to a reduction in the strength and hardness of the alloy. The weight percentage of Mn (manganese), which is an essential element for the formation of the quasicrystalline phase, is 3≤Mn, preferably 5≤Mn, and Mn≤10, preferably Mn≤8. Too low content of Mn leads to the formation of a crystalline Al6Mn phase rather than a highly symmetric Qc phase. On the other hand, a Mn content greater than 10% leads to the formation of a metastable crystalline µ, λ, or L phase. The weight percentage of Si (silicon), which is an essential element for stabilizing the icosahedral quasicrystalline cluster, is 0.1≤Si, preferably 0.2≤Si, and Si≤3, preferably Si≤1.7. A weight percentage for Si of less than 0.1% does not sufficiently affect the stability of the iQc phase due to the lack of Si atoms in the icosahedral (Mackay) cluster. On the other hand, a Si weight percentage higher than 3% leads to the formation of a primary crystalline β-AlMnSi or α-AlMnSi phase instead of the Qc phase. The amount of at least one metal element (X) selected from the group consisting of Cu (copper), Ni (nickel), Zn (zinc), Fe (iron) and Co (cobalt) by weight is at least 0.1%, preferably at least 0.5%, or even more at least 1%. The content of elements X should not exceed 10%, preferably not exceed 8%; and more preferably not exceed 5%. The elements from this group support the formation of the quasicrystals. They are incorporated into the Qc phase and reduce the critical cooling rate required for Qc formation. If the amount is less than 0.1%, a quasicrystalline phase cannot be achieved at moderate cooling rates, and the strength of the alloy is reduced. If, on the other hand, the amount exceeds 10%, intermetallic phases occur in addition to the quasicrystalline phase, which leads to embrittlement and a reduction in the formability of the alloy. The more preferred are Cu (copper), Ni (nickel), and Zn (zinc); and can range from 0.1 to 10%. Co (cobalt) content can range from 0.1 to 5%. A higher Co content increases the liquidus temperature of the alloy and intermetallic phases may form. The upper limit for the Fe content is defined by the equation XFe = 0.3 x XMn in wt% (upper limit 3 wt% (10x0.3=3)). A higher Fe content leads to the formation of intermetallic phases (AlxFe, Al6Fe) instead of Qc. Characteristic of this group of elements is that they all form binary eutectics with the aluminium-rich angle of the binary phase diagrams. The amount of at least one element (Z) selected from the group consisting of Li (lithium), Ca (calcium), Sr (strontium), Ga (gallium), REE (rare-earth elements), mischmetal, Bi (bismuth), Sn (tin), Sb (antimony) and Mg (magnesium) in percent by weight percentage is at least 0.1%, and even more preferably at least 0.3%. The content of element Z should not exceed 8%, preferably not exceed 6%; and even more preferably not exceed 4%. The elements from this group reduce the surface tension between the aluminium-based melt and the growing quasicrystalline phase, thus assisting the formation of the quasicrystalline phase. If the proportion is less than 0.01%, the surface tension between the aluminium-based melt and the growing quasicrystalline phase is not reduced sufficiently to promote Qc growth. If, on the other hand, the proportion exceeds 8%, intermetallic phases based on selected elements occur, which leads to embrittlement and a reduction in the ductility of the alloy. The preferred elements are Mg (magnesium), Li (lithium), REE (rare-earth elements) and mischmetal, whose content can range from 0.01 to 10%. Higher contents of these elements can form a strong oxide layer, so casting in a protective atmosphere (or vacuum) is desirable. Ca, Sr, Sb and Ba have low solubility in the aluminium matrix. The preferred range for these elements is limited by the amount where the binary eutectic point with the aluminium-rich angle of the binary phase diagram is located. These are: 1 wt% for Sr, 1 wt% for Sb, 3.5 wt% for Ba, and 7.6 wt% for Ca. Ga (gallium), Bi (bismuth), and Sn (tin) have low melting points and do not form intermetallic phases with aluminium. Their preferred content is between 0.01 and 1 wt.%. The amount of at least one metal element (TM=transition metal) selected from the group consisting of Cr (chromium), Ti (titanium), W (tungsten), Mo (molybdenum), Zr (zirconium) and V (vanadium) by weight is at least 0.1%, and even more preferably at least 0.4%. The content of element TM should not exceed 3%, preferably not exceed 2%; and more preferably not exceed 1.2%. Elements from this group have low diffusivity in the aluminium-based melt and help stabilize the supercooled state of the molten metal and further retard crystallization of a crystal phase. If the proportion is less than 0.01%, the selected elements do not have much influence on the solidification path. If, on the other hand, the proportion exceeds 3%, high- temperature intermetallic phases based on selected elements occur, which leads to embrittlement and a reduction in the formability of the alloy. The preferred elements are Ti (titanium), W (tungsten), Mo (molybdenum), Zr (zirconium) and V (vanadium), which can range from 0.1 to 3%. A higher content of these elements leads to the formation of high- temperature stable intermetallic phases, which may be present in the melt during alloy production and casting. The preferred content of Cr (chromium) is between 0.01 and 0.1 wt%. A higher Cr content leads to the formation of needle-like shape intermetallic phases, which cause embrittlement and a reduction in the ductility of the alloy. Characteristic of this group of elements is that they all form binary peritectic reactions with the aluminium-rich angle of the binary phase diagrams. BRIEF DESCRIPTION OF THE DRAWING Examples of the invention will be presented below with reference to the drawing, on which: Figure 1 is a temperature-time-transformation diagrams for Al6Mn, dQc and iQc phase calculated on the basis of transformation theory; modified from Gillessen et al. [8], Figure 2 is a microstructure of the alloy I (DIN 226) obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 3 is a microstructure of the alloy II obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 4 is a microstructure of the alloy III obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 5 is a microstructure of the alloy IV obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 6 is a microstructure of the alloy V obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figures 7 is a microstructure of the alloy VI obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 8 is a microstructure of the alloy VII obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 9 is a microstructure of the alloy VIII obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 10 is a microstructure of the alloy IX obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 11 is a microstructure of the alloy X obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, Figure 12 is a microstructure of the alloy XI obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification, and Figure 13 is a microstructure of the alloy XII obtained at 500 K / s (a) and 100 K / s (b). In the insets, phases with typical morphology for quasicrystalline phases are shown at higher magnification. EXAMPLES Aluminium alloys were produced in a chamber furnace. We used technical aluminium (99.8 wt%), manganese (99.9 wt%), silicon (99.8 wt%), titanium (99.7 wt%), magnesium (99.7 wt%), chromium (99.9 wt%), nickel (99.9 wt%), copper (99.9 wt%), iron (99,5 wt%), cobalt (99.8 wt%), tin (99.9 wt%), samarium (99.9 wt%), molybdenum (99.9 wt%), ferrovanadium (Fe (20 wt%), V (80 wt%)), and master alloys AlSr10, AlZr10, AlLi10, and AlCa6. The alloys were melted at 880°C and cast in copper moulds. The moulds consisted of a copper block with a cylindrical casting cavity of 5.5 mm diameter and a length of 70 mm, and a block casting cavity with a cross-section of 24 mm x 12 mm and a length of 130 mm. The microstructure was characterised using an optical microscope (LOM), ZEISS Axio Imager A1 m, equipped with an AxioCam ERc 5s digital camera (5 million pixels) and AxioVision software. X-ray fluorescence analyser NITON, model XL3t+ GOLDD 900S-He was used to confirm the chemical composition of alloys. A field-emission-gun scanning electron microscope (SEM) JEOL JSM-7600F equipped with electron backscatter diffraction (EBSD) was used for the phase characterization. The analyses were made using CHANNEL5 EBSD software with a Nordlys detector. The quasicrystal and crystal structure of the samples was confirmed by EBSD at 20 kV and specimen tilt of 70°. The tensile specimens were made from the cast alloys according to the standard DIN 50125. The cast rods with a diameter of 5.5 mm were used to produce type A tensile specimens with a test piece diameter of 4 mm, a diameter of gripped ends of 5 mm and a total length of the test pieces of 60 mm. The cast blocks with dimensions of 24 mm x 12 mm x 130 mm were cut in the middle of the width and length, resulting in 4 smaller blocks with dimensions of 11 mm x 12 mm x 60 mm. Type B tensile test specimens with a specimen diameter of 6 mm, a total length of 60 mm and a diameter of ISO metric thread M10 were made from these blocks. The samples for the microstructure and hardness measurements were cut from the centre of the castings under study. For the cast rods with a diameter of 5.5 mm, one-dimensional transient-heat-conduction analyses were used to evaluate the cooling rate at the beginning of the solidification process, as explained by Naglič et. al.

[0020] . The calculated cooling rates are in the range of 500 K / s. For the cast blocks with dimensions of 24 mm x 12 mm x 130 mm, cooling rates were measured using a National Instruments thermocouple input module NI-9211, which provides 14 measurements per second. Temperatures were measured using a type K thermocouple. The obtained cooling rates at the beginning of the solidification process are in the range of 100 K / s. Under these manufacturing conditions, the reference alloy DIN 226 and 11 alloys within the composition range for formation of Qc phase where produced. The chemical composition of alloys in wt.% are presented in Table 1. Table 1. Chemical composition of alloys in wt.%. Alloy Al Mn Si X Z TM I (DIN 226) bal 0.25 9 2.4 (Cu) / 0.1 (Cr) reference 0.7 (Fe) 0.9 (Zn) II bal 7 1.5 3 (Cu) 3 (Mg) 0.25 (Zr) 1 (Li) III bal 6.5 1.1 3 (Cu) 7 (Mg) 0.2 (Zr) 0.9 (Co) IV bal 6 0.3 0.5 (Ni) 0.5 (Sm) 0.25 (Ti) V bal 4.5 0.8 3 (Cu) 3 (Mg) 1 (V) 1.2 (Fe) 0.15 (Sr) VI bal 5.8 0.4 3 (Cu) 2 (Mg) 0.4 (V ) 0.2 (Fe) 0.3 (Ti) VII bal 4.5 0.3 0.5 (Cu) 0.1 (Sn) 0.25 (Ti) VIII bal 5.5 0.4 0.5 (Cu) 0.05 (Sr) 0.1 (Mo) 0.1 (Cr) 0.2 (Ti) IX bal 5.5 1.3 3.2 (Cu) 3 (Mg) 1.5 (V) 0.4 (Fe) X bal 6.5 2.6 3.2 (Cu) 8 (Mg) 0.1 (Zr) XI bal 9.5 1.8 3 (Cu) 6 (Mg) 0.05 (Ti) XII bal 5.8 0.3 1 (Cu) 0.3 (Mg) 0.3 (Ti) Microstructure The microstructures of the cast alloys are shown in Figures 2 to 13. Each figure shows the characteristic microstructure obtained at 500 K / s (a) and 100 K / s (b). In the inserts are phases with typical morphology for the quasicrystalline structures. Electron backscatter diffraction (EBSD) analyses of all castings with the proposed composition made at a cooling rate of 500 K / s confirm the presence of quasicrystalline phases. On the other hand, when analysing castings made at a cooling rate of 100 K / s, only the alloy (alloy VI) of the most preferred composition partially consists of the quasicrystalline phase. In other alloys, phases with a similar morphology to the iQc and dQc phases occur, but they are not quasicrystalline. This characteristic morphology, seen in the insets (Fig.3b - Fig.13b), could be due to the nucleation of metastable quasicrystalline phases, followed by an insufficient solidification rate, causing the Qc phase to transform into a crystalline phase and retain the morphology. The hardness (HV) and tensile strength (MPa) of each alloy type sample were measured. These results are shown in Table 2. Vickers hardness measurements were performed in air at room temperature under a constant loading condition using a load of 1 kg. The Kason Micro Vickers HTMV-1000 AD instrument was used for the measurements. Tensile strength measurements were performed in air at room temperature according to the test standard SIST EN ISO 6892- 1:2017. The ZwickRoell tensile testing machine Z 2500W with a test load of up to 50 kN was used for the measurements. Table 2: Tensile strength (MPa) and hardness (HV) of the reference alloy DIN 226 and 11 alloys within the composition range for the formation of the Qc phase at cooling rates of 500 K / s and 100 K / s. Cooling rate of 500 K / s Cooling rate of 100 K / s Alloy Strength Hardness Phases Strength Hardness Phases (MPa) (HV1) (MPa) (HV1) I (DIN 226) 105 104 crystalline 245 105 crystalline reference II 238 160 iQc, 138 130 crystalline crystalline III 222 147 iQc, dQc, 168 142 crystalline crystalline IV 175 80 iQc, 138 63 crystalline crystalline V 110 146 iQc, 191 130 crystalline crystalline VI 230 130 iQc, 251 137 iQc, crystalline crystalline VII 212 75 iQc, 159 60 crystalline crystalline VIII 179 80 iQc, 151 59 crystalline crystalline IX / 165 iQc, 186 153 crystalline crystalline X 274 155 iQc, 170 149 crystalline crystalline XI 134 154 iQc, dQc 129 150 crystalline crystalline XII 196 81 iQc, 174 80 crystalline crystalline The reference alloy is the commercial aluminium alloy DIN 226, which is mainly used for high- pressure die-cast parts. The tensile strength and hardness of cast DIN 226 are well known and are at least 240 MPa and 80 HB (85 HV), respectively. From Table 2, it can be seen that the tensile strength of the reference alloy at a cooling rate of 500 K / s is much lower at only 105 MPa. The reason for the lower tensile strength are casting defects, which cannot be avoided when gravity casting into a simple copper mould. Shrinkage porosity, which can be locally as large as almost 1 mm, has the greatest effect on tensile strength, followed by non-metallic inclusions (oxide layer) and other casting defects. Although the alloy DIN 226 has a high silicon (Si) content and the solidification shrinkage should be lower that of the alloys under consideration for the application, higher tensile strengths are obtained. As can be seen form the result shown in Table 2 and microstructures, any of the alloys according to the present invention has properties which are superior in strength or hardness at room temperature in comparison to reference alloy DIN 226, while cooling rate of 500 K / s is applied. At cooling rates of 500 K / s, any of the alloy structures had a multiphase composed of a quasi- crystalline phase (iQ, dQC or both of them) and aluminium or a super-saturated solid solution phase of aluminium, and various intermetallic compound were present in particular kinds of alloys. At cooling rates of 100 K / s, only alloy VI within the most optimal composition exhibit a multiphase microstructure consisting of a quasicrystalline phase (iQc) and a crystalline phase of aluminium or a supersaturated solid solution of aluminium and various crystalline intermetallic compounds. The content of the iQc phase was reduced compared to the alloy prepared with a cooling rate of 500 K / s. The difference was expressed in a higher content of crystalline intermetallic phases. Casting defects in type B tensile test specimens, such as shrinkage porosity and nonmetallic inclusions (oxide layer), were reduced compared to the 5.5 mm diameter roods, mainly due to the manufacturing process of the tensile specimens. As explained above, the cast blocks with dimensions 24 mm x 12 mm x 130 mm were cut in the middle of the width and length, resulting in 4 smaller blocks with dimensions 11 mm x 12 mm x 60 mm. In this way, the shrinkage porosity (located in the centre) was cut out of the cast blocks. This can also be seen from Table 2, where the alloy DIN 226 has a higher tensile strength value at a cooling rate of 100 K / s, although the microstructure of the alloy DIN 226 in Figure 2 clearly has a finer eutectic with a smaller size of α-Al dendrites. The much lower strength of the alloys of the present invention produced at cooling rates of 100 K / s is due to the formation of crystalline intermetallic phases with strong plate-like morphology, which enhances the notch effect of these alloys. It can also be seen from Table 2 that the hardness of the alloys of the present invention depends on the content of elements of the individual groups. The higher the content, the higher the hardness, which is due to the higher content of Qc or crystalline intermetallic phases. Conclusions From the experiments performed, it can be concluded that all alloys according to the invention consist of a metastable quasicrystalline phase (iQc, dQc, or both), a crystalline phase of aluminium or a supersaturated solid solution of aluminium and various crystalline intermetallic compounds, where the alloys have been prepared by casting process that cooling rates of 500 K / s are achieved. While the alloys have been prepared by a casting process that achieved cooling rates of 100 K / s, only one alloy (alloy VI) in the most preferred composition consists in part of the quasicrystalline phase (iQc), a crystalline phase of aluminium or a supersaturated solid solution of aluminium and various crystalline intermetallic compounds. References [1] Wilm A 1911 Metallurgie: Zeitschrift für de gesamte Hüttenkunde 8225-7 [2] ASM Handbook 1992 Properties and Selection: Nonferrous Alloys and Special-Purpose Materials (vol 02) ASM International Handbook Committee. pp 137-90 [3] Inoue A, Kong F, Zhu S, Liu C T and Al-Marzouki F 2015 Materials Research 181414- 25 [4] Inoue A and Kimura H 2000 Materials Science and Engineering: A 2861-10 [5] Inoue A, Kimura H and Horio Y 1999-01-12 High strength and high rigidity aluminum- based alloy and production method therefor. United States Patent Office, 5,858,131) [6] Shechtman D, Blech I, Gratias D and Cahn J W 1984 Physical Review Letters 531951- 3 [7] Battezzati L, Antonione C and Marino F 1989 Journal of materials science 242324-30 [8] Gillessen F and Herlach D 1991 Materials Science and Engineering: A 1341220-4 [9] Steurer W and Deloudi S 2008 Acta Crystallographica Section A: Foundations of Crystallography 641-11

[0010] Steurer W 2012 Chemical Society Reviews 416719-29

[0011] Zupanic F, Boncina T, Rozman N, Anzel I, Grogger W, Gspan C, Hofer F and Markoli B 2008 Development of an Al–Mn–Be–Cu alloy with improved quasicrystalline forming ability. In: Zeitschrift für Kristallographie International journal for structural, physical, and chemical aspects of crystalline materials, p 735

[0012] Zupanič F, Wang D, Gspan C and Bončina T 2015 Materials Characterization 10693- 9

[0013] Zupanič F, Lojen G, Barba L and Bončina T 2012 Materials Characterization 7048- 54

[0014] Bončina T, Markoli B and Zupanič F 2009 Journal of microscopy 233364-71

[0015] Markoli B, Delijic K, Strekelj N and Naglic I 2014 Contemporary materials 530-6

[0016] Markoli B, Boncina T and Zupanic F 2015 Recent patents on materials science 108109- 18

[0017] Naglič I, Samardžija Z, Delijić K, Kobe S, Dubois J-M, Leskovar B and Markoli B 2017 Journal of Materials Science 5213657-68

[0018] Stan-Głowińska K, Rogal Ł, Góral A, Wierzbicka-Miernik A, Wojewoda-Budka J, Schell N and Lityńska-Dobrzyńska L 2017 Journal of Materials Science 527794-807

[0019] Stan-Głowińska K and Lityńska-Dobrzyńska L 2017 Materials Characterization 128 203-8

[0020] Naglič I, Samardžija Z, Delijić K, Kobe S, Leskovar B and Markoli B 2018 JOM 70 2698-703

Claims

CLAIMS 1. An aluminium-based alloy having the following composition: Albal.MnaSinXcZeTMd, where X is one or more elements selected from the group consisting of Cu, Ni, Zn, Fe and Co; Z is one or more elements selected from the group consisting of Li, Ca, Sr, Ga, REE, mischmetals, Ce, Bi, Sn, Sb and Mg; TM is one or more elements selected from the group consisting of Cr, Ti, W, Mo, Zr and V; wherein a, n, c, e and d satisfy, by weight percent: 3≤a≤10, 0.1≤n≤3, 0.1≤c≤10, 0.1≤e≤8, and 0.1≤d≤3, and wherein 5≤a+c+e≤20; the balance being Al and inevitable impurities up to 0.5 wt.%, and comprising a metastable quasicrystal phase that is dispersed in a matrix composed of aluminium.

2. An alloy according to claim 1, wherein 0.4≤n+d≤5, by wt.%.

3. An alloy according to any preceding claim, wherein 5≤a.

4. An alloy according to any preceding claim, wherein a≤8.

5. An alloy according to any preceding claim, wherein 0.2≤n.

6. An alloy according to any preceding claim wherein n≤1.

7.

7. An alloy according to any preceding claim, wherein 0.5≤c, or, preferably, 1.0≤c.

8. An alloy according to any preceding claim, wherein c≤8, or, preferably, c≤5.

9. An alloy according to any preceding claim, wherein X is one or more elements selected from the group consisting of Cu, Ni and Zn.

10. An alloy according to any preceding claim, wherein 0.3≤e.

11. An alloy according to any preceding claim wherein e≤6, or, preferably, e≤4.

12. An alloy according to any preceding claim, wherein Z is one or more elements selected from the group consisting of Mg, Li, REE, and mischmetal.

13. An alloy according to any preceding claim, wherein 0.4≤d.

14. An alloy according to any preceding claim, wherein d≤2, or, preferably, d≤1.2.

15. An alloy according to any preceding claim, wherein TM is selected from one or more elements selected from the group consisting of Ti, W, Mo, Zr and V.

16. A method of producing an alloy comprising the steps of -forming a melt of an alloy having the composition according to any of claims 1-15, and -cooling the melt with a cooling rate of 500 K / s or more.

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