TECHNICAL FIELD The present disclosure relates to a casting alloy and particularly, but not exclusively, to aluminium-silicon casting alloys for use in castings of structural components. Aspects of the invention relate to an alloy composition, and to a method of making and using such alloys. BACKGROUND

Aluminium alloys are used in the production of castings for engine components used in the manufacture of vehicles. Such components include engine blocks, cylinder heads, or heat shields.

Aluminium is favoured for the production of such components due to its ease of handling, its low cost, its ease of machinability and castability, and because it is relatively lightweight. Castings typically comprise complex geometries including parts that are exposed to elevated temperatures and pressures. In order to meet the stringent requirements of strength, ductility and resistance to metal fatigue whilst retaining good castability, aluminium-silicon eutectic system alloys (silumin) are favoured for the purposes of castings.

The continuing drive to reduce fuel consumption and C0 ₂ emissions means that there is an increasing focus towards increasing the operating efficiency of vehicle engines. The performance of motor vehicle engines can be significantly improved by increasing the operating pressure and temperature conditions.

Conventionally, aluminium-silicon alloys used in vehicle castings consist of a major component of aluminium, between 7wt% and 17wt% of silicon as well as minor amounts (between 0.1 wt% and 1wt%) of additional grain refining and strengthening additive elements such as copper, zinc, manganese, magnesium, titanium, silver and other elements. A significant problem associated with the use of these alloys is that there is a significant degradation in their performance at elevated temperatures, such as those associated with the operating conditions of a vehicle engine with improved performances.

In particular, the strengthening phases which are formed during casting and particularly during heat treatment of the alloy, such as AI ₂ Cu, Mg ₂ Si and AI ₂ CuMg, and which maintain the alloy's strength at room temperature, are known to coarsen or dissolve into the solid solution at elevated temperatures. The dissolution of these strengthening phases results in a considerable reduction in the alloy's performance at high temperatures. This high temperature strength reduction limits the suitability of the alloy for use in engine components, particularly in engine components that offer improved performance in the form of increased fuel efficiency and higher output torques of vehicle engines.

Various alloying elements such as, for example, nickel have been previously added to Al-Si alloys in order to enhance the mechanical properties at elevated temperatures. However, it has been shown that such alloying additions significantly decrease the castability of the alloy as well as causing a reduction in the mechanical properties at ambient temperatures and pressures.

Hence, there is a dependence in the art towards using alloying elements that exhibit both limited solid-solubility at typical working temperatures and low diffusivity in the a-AI matrix. This has led to the development of complex Al-Si-Fe based systems, which rely on the addition of alloying elements such as silver, hafnium, nickel, caesium, cobalt, molybdenum and vanadium, which have been shown to exhibit high- strength and thermally-stable properties. These alloys, however, derive their high- temperature strength from a large volume fraction of stable precipitates that form directly from the melt during rapid solidification as a result of supersaturation of the alloying elements. Such precipitates are practically impossible to reach in sand casting and gravity casting, where solidification rate is too low to achieve the required supersaturation. Consequently, such requirements make it difficult to cast Al-Si alloys that exhibit suitable mechanical properties at elevated temperatures. It would be a considerable advantage if a means of fabricating a high performance, castable, light and affordable alloy could be formulated, which exhibited acceptable material and mechanical characteristics at room temperature but which could also withstand the elevated temperatures associated with the operating conditions of a vehicle engine. It would be particularly advantageous if such results could be achieved for methods such as sand casting and gravity casting that involve comparatively slow solidification times.

The present invention provides alloy compositions that meet the aforementioned objectives. These and other uses, features and advantages of the invention should be apparent to those skilled in the art from the teachings provided herein.

SUMMARY OF THE INVENTION Against this background, the invention resides in an aluminium-silicon casting alloy composition comprising: at least 7 wt% and at most 1 1 wt% silicon; at least 0.5 wt% and at most 4.5 wt.% copper; at least 0.1 wt%, and at most 1wt.% magnesium; and at least 0.1 wt% and at most 1wt.% zirconium. The balance of the composition comprises aluminium. The sum of the copper and magnesium compositions is at least 1.0 wt% and at most 4.6 wt%.

When the alloy is cast to form a finished product, by virtue of the balance between the zirconium content, the balance of copper and magnesium, and the silicon content, the finished alloy comprises a combination of phases that provide high strength at both room temperature and high temperature, even if the alloy is cast at low solidification rates.

In particular, the inventors have found that the selected zirconium and silicon contents lead to the formation of a Zr ₂ Si phase in the microstructure of the cast alloy. This phase increases the strength of the alloy and is thermally stable, such that it provides improved strength at both low and high temperatures. Additional zirconium-containing phases including Al-Si-Zr phases also provide high-strength phases that are thermally- stable. Phases such as AI ₂ Cu (Θ), Al-Si-Cu, and Al-Si-Cu-Mg (Q), tend to provide high strength at low temperatures, but easily dissolve at high temperatures. However, the inventors have found that under the high zirconium content of the invention, the AI ₂ Cu (Θ), Al-Si-Cu and Al-Si-Cu-Mg (Q) phases of the alloy all contain zirconium. The presence of zirconium in these phases advantageously increases the stability of these phases at high temperature due to the large size of the zirconium ion, thereby further improving the high temperature strength of the alloy.

Thus, the careful balance of zirconium, copper, magnesium and silicon all give rise to a combination of precipitates that provide high strength at both low and high temperatures. This is particularly beneficial for application in vehicle engines, in which high strength is required both at normal engine operation where temperatures are comparatively low, and during engine start-up, where the engine components are subject to high temperatures.

The composition may comprise at least 1 wt% and at most 4.5 wt% copper.

The sum of the copper and magnesium compositions may be at least 2.0 wt% and at most 4.6 wt%.

To achieve a particularly desirable ductility and/or castability that is particularly suitable for a casting alloy, the content of silicon may be about 6 wt% more than the content of copper. In embodiments the aluminium-silicon alloy may further comprise at least 2 wt% and at most 4 wt% copper.

The aluminium-silicon alloy may further comprise at least 0.2 wt% and at most 0.6 wt% magnesium.

The aluminium-silicon alloy may further comprise at most 10 wt% silicon, and optionally from 7.5 to 9 wt% silicon. The zirconium composition may be set at a level to ensure that, in the cast alloy, a zirconium-containing intermetallic phase is formed that assumes an elliptical and/or spherical morphology having a semi-coherent or coherent interface with the Al-matrix microstructure of the alloy.

The zirconium composition is set at a level to ensure that, in the cast alloy, the zirconium-containing intermetallic phase forms as micro and/or nano-sized particles within an aluminium matrix of the cast alloy. Herein, the term micro-sized particles relates to particles which have a diameter that is substantially equal to or greater than 1 μηι and the term nano-sized particles refers to particles that have a diameter that is less than 1 μιη. In embodiments, the nano-sized particles refers to particles that have a diameter that is substantially less than 1 μιη, alternatively less than approximately 500nm. The alloy may further comprise at least 0.1 wt% and at most 0.3 wt% titanium. The alloy may further comprise at most 0.2 wt% manganese. The alloy may further comprise at most 0.15% iron.

The iron composition may be set at a level to ensure that, in the cast alloy, iron- containing intermetallic phase particles form within an interdendritic region of the cast alloy. The manganese to iron ratio may be set at a level to ensure that, in the cast alloy, the majority of iron-containing intermetallic phase particles assume a substantially elliptical morphology having a semi-coherent interface with an aluminium matrix of the cast alloy. The semi-coherent interface with the aluminium matrix enhances the microstructure's resistance to dislocation movement, which thereby increases the mechanical strength of the alloy, particularly at elevated temperatures. Advantageously, the morphological modification of the iron-containing phase caused by the iron/manganese ratio not only avoids the particularly undesirable 'script' morphology, but additionally the elliptical morphology of the iron-containing phases acts to retard Ostwald ripening of these phases and thus improves the coarsening resistance of the cast alloy at elevated temperatures.

The alloy may further comprise at least 0.01 wt% and at most 0.05 wt% strontium.

The alloy may further comprise less than 0.05% nickel, and/or less than 0.05% chromium.

According to a further aspect of the invention there is provided a cast component comprising an aluminium-silicon alloy having a composition comprising:

at least 7 wt% and at most 1 1 wt% silicon;

at least 0.5 wt% and at most 4.5 wt% copper;

at least 0.1 wt% at most 1 wt% magnesium; and

at least 0.1 wt% and at most 1 wt% zirconium,

the balance of the composition comprising aluminium;

wherein the sum of the copper and magnesium compositions is at least 1 .0 wt% and at most 4.6 wt%.

According to a further aspect of the invention there is provided a cast component comprising an aluminium-silicon alloy having a composition comprising:

at least 7 wt% and at most 1 1 wt% silicon;

at least 0.5 wt% and at most 4.5 wt% copper;

at least 0.1 wt% at most 1 wt% magnesium;

at least 0.1 wt% and at most 1 wt% zirconium;

the balance of the composition comprising aluminium ;

wherein the sum of the copper and magnesium compositions is at least 1 .0 wt% and at most 4.6 wt%;

and wherein the cast component comprises an Al-Si-Zr phase within the microstructure of the alloy.

The component may comprise a Zr ₂ Si phase within the microstructure of the alloy.

The Zr ₂ Si phase may have a semi-coherent or coherent interface with an aluminium matrix of the alloy. The Zr ₂ Si phase may take the form of micro- or nano-sized particles.

The Zr ₂ Si phase may take the form of particles having an elliptical and/or spherical morphology.

The component may further comprise an Al-Si-Zr phase within the microstructure of the alloy. The component may comprise a zirconium-containing Al-Cu phase within the microstructure of the alloy.

The component may comprise a zirconium-containing AI ₂ Cu (Θ) phase within the microstructure of the alloy.

The component may comprise a zirconium-containing Al-Si-Cu-Mg (Q) phase within the microstructure of the alloy.

The component may comprise a zirconium-containing Al-Si-Cu phase within the microstructure of the alloy.

The composition may further comprise at most 0.2 wt% manganese.

The composition may further comprise at most 0.15% iron, and the alloy comprises iron-containing intermetallic phase particles.

The iron-containing intermetallic phase may be contained within an interdendritic region of the alloy. The majority of the iron-containing intermetallic phase particles may have a substantially elliptical and/or spherical morphology and have a semi-coherent or coherent interface with an aluminium matrix of the alloy.

The composition may comprise at least 1 wt% and at most 4.5 wt% copper. The sum of the copper and magnesium compositions may be at least 2.0 wt% and at most 4.6 wt%. The aluminium-silicon alloy may be the aluminium-silicon casting alloy according to the above described invention.

The component may be for use within a powertrain of a motor vehicle, optionally the engine.

According to a further aspect of the invention, there is provided a vehicle engine comprising the cast component according to the above described invention.

According to a further aspect of the invention, there is provided a powertrain for a vehicle, the powertrain comprising the cast component according to the above described invention.

According to a further aspect of the invention, there is provided a vehicle including the cast component according to the above described invention.

According to a further aspect of the invention, there is provided a method of casting a component, the method comprising: providing a mould that defines the configuration of the component, pouring the molten aluminium-silicon casting alloy according to the above described invention into the mould and allowing the alloy to cool to form a solid cast component.

According to a further aspect of the invention, there is provided a method of casting a component, the method comprising: providing a mould that defines the configuration of the component; providing an aluminium-silicon casting alloy comprising: at least 7 wt% and at most 1 1 wt% silicon; at least 0.5 wt% and at most 4.5 wt% copper, at least 0.1 wt%, and at most 1 wt% magnesium; and at least 0.1 wt% and at most 1 wt% zirconium; and at least 10 wt% and at most 90.9% aluminium; wherein the sum of the copper and magnesium compositions is at least 1 .0 wt% and at most 4.6 wt%. The method also comprises heating the casting alloy to form a molten casting alloy and pouring the molten casting alloy into the mould and allowing the alloy to cool to form the above-described cast component. The method may further comprise heating the casting alloy to form a molten casting alloy at a temperature of at least 700°C, optionally 750°C before pouring the casting alloy into the mould.

The method may further comprise degassing the molten casting alloy with an inert gas source for at least 1 minute, optionally at least 3 minutes, optionally up to 30 minutes.

The method may further comprise heat treating the solid cast component after the casting alloy has solidified. The method may further comprise holding the solid cast component at a solutionising heat treatment temperature of at least 400°C and at most 550°C, optionally at least 535°C and at most 540°C.

A solutionising heat treatment period may be at least 5 hours, optionally at least 6 hours, optionally up to 10 hours.

The method may further comprise following the solutionising heat treatment, quenching the casting alloy to a quenching temperature below 50°C, optionally 25°C, suitably down to 23°C.

The method may further comprise holding the solid cast component at an ageing heat treatment temperature of at least 150 °C and at most 220 °C.

The method may further comprise holding the solid cast component at the ageing heat treatment temperature for an ageing time of at least 0.5 hours and at most 40 hours.

In embodiments, the alloy may further comprise at most 0.15% iron, in which case the method may further comprise conducting the heat treatment such that the majority of iron-containing intermetallic phase particles assume a substantially elliptical morphology having a semi-coherent interface with the microstructure of the alloy. The method may further comprise adding the zirconium alloying element to the composition after the casting alloy has been heated to form a melt at a temperature of at least 800°C.

In embodiments, the alloy may further comprise at least 0.1 wt% titanium and at most 0.3 wt% titanium.

In embodiments, the alloy may comprise at least 1 wt% and at most 4.5 wt% copper. Optionally, the sum of the copper and magnesium compositions may be at least 2.0 wt% and at most 4.6 wt%.

In embodiments, the method of adding the zirconium alloying element to the composition may comprise adding an Al-Zr master alloy to the melt. The Al-Zr 'master' alloy may be first heated up to 450°C before it is added to the melt In order to maximise the dissolution of Zr in the melt.

In embodiments, the alloy of the molten aluminium-silicon casting alloy according to the above described invention may be formed using the above described method wherein the method may further comprise adding the zirconium and titanium alloying elements to the composition after the casting alloy has been heated to form a melt at a temperature of at least 800°C.

The melt may be allowed to fill the mould under the action of gravity alone.

According to a further aspect of the invention, there is provided a use of the aluminium-silicon die casting alloy according to the above described invention for manufacturing a cast part of a motor vehicle. Within the scope of this application it is expressly intended that the various aspects, embodiments, examples and alternatives set out in the preceding paragraphs, in the claims and/or in the following description and drawings, and in particular the individual features thereof, may be taken independently or in any combination. That is, all embodiments and/or features of any embodiment can be combined in any way and/or combination, unless such features are incompatible. The applicant reserves the right to change any originally filed claim or file any new claim accordingly, including the right to amend any originally filed claim to depend from and/or incorporate any feature of any other claim although not originally claimed in that manner.

BRIEF DESCRIPTION OF THE DRAWINGS

One or more embodiments of the invention will now be described, by way of example only, with reference to the accompanying drawings, in which:

Figure 1 is a micrograph showing a section through an as cast Al-Si alloy of an embodiment of the invention, which comprises a eutectic Si phase and AI ₂ Cu precipitates within an aluminium matrix. Figure 2 is a micrograph of a section of the alloy of Figure 1 which has been heat treated according to an embodiment of the invention, which comprises a eutectic Si phase together with AI ₂ Cu and Mg ₂ Si precipitates within the aluminium matrix.

Figure 3 is an electron back scattered diffraction (EBSD) micrograph of a section of a comparative conventional Al-Si alloy after it has been heat treated according to an embodiment of the invention.

Figure 4 is an EBSD micrograph of a section of the heat treated alloy of Figure 2, which has been heat treated according to an embodiment of the invention.

Figures 5a and 5b are back scattered electron (BSE) micrographs of different regions of the as-cast alloy of Figure 1 , taken using a scanning electron microscope (SEM).

Figures 6a, 6b and 6c are BSE micrographs of different regions of the heat treated alloy of Figure 2, taken using a SEM.

Figure 7 is a transmission electron microscope (TEM) micrograph of the as cast Al-Si alloy of Figure 1 , showing an Fe containing phase distributed within the aluminium matrix. Figure 8 is a TEM micrograph of the as-cast alloy of Figure 1 , showing a Zr containing phase distributed within the aluminium matrix. Figure 9 is an energy dispersive x-ray spectroscopy (EDS) spectra taken from the region shown in Figure 7, showing the constituent elements of the Fe containing phase.

Figure 10 is an EDS spectra taken from the region shown in Figure 8, showing the constituent elements of the Zr containing phase.

Figure 1 1 is a graph of the results of a tensile strength test on multiple samples of parts manufactured from the as-cast and heat treated alloy of an embodiment of the invention, showing the yield strength (YS) of the as-cast and heat treated samples at room and 200 °C.

Figure 12 is a graph of the results of a tensile strength test on multiple samples of parts manufactured from the as-cast and heat treated alloy of an embodiment of the invention, showing the ultimate tensile strength (UTS) of the as-cast and heat treated samples at room and 200 °C.

Figure 13 is a schematic drawing of a vehicle comprising an Al-Si cast component according to an embodiment of the invention. DETAILED DESCRIPTION

Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this invention belongs.

Prior to setting forth the invention, a number of definitions are provided that will assist in the understanding of the invention. The term "alloy" is used herein to denote a metallic substance comprising a mixture of a predominating metallic element and other elements, including impurities. The basic aluminium-silicon (Al-Si) alloy forms a simple eutectic system, with around 12 weight% (wt%) silicon being the eutectic composition at 577°C. Conventional Al-Si alloys suitable for use in casting applications typically meet industry standards such as European Standard EN AC-42000 (AISi ₇ Mg), which defines the requirements for various grades of alloyed aluminium ingots intended for re-melting. It is believed that the alloys of the present invention show high levels of equivalence in terms of suitability for use in casting, especially for automotive and aerospace parts.

"Casting" is a manufacturing process that can produce metal components through the use of moulds. In some instances the mould may be sacrificial, as in the case of gravity casting where a sand-mould is used, or the mould may take the form of a reusable die, such as may be used in high pressure die casting processes. The casting process involves a furnace, metal and the mould. Alternatively, a casting machine may also be used to apply pressure to the mould or die during the casting process. The metal, typically a non-ferrous alloy comprising aluminium, is first melted in the furnace and then poured or injected, optionally under pressure, into the mould. Once the casting has cooled and has set the part can be subjected to additional tooling or trimmed and finished. Casting processes can produce large and small component parts, with geometrically complex shapes. The cast parts are typically of high strength and can be subjected to considerable loads when in use. The process supports a reasonably high rate of production and is favoured producing consistent parts with good surface finish.

As mentioned, components manufactured via die casting processes of the present invention are also referred to as 'parts' and broadly refer to any metal object manufactured via the casting process. The parts are often comprised of geometrically complex shapes that perform a defined function. Components manufactured according to the present invention are particularly suitable for use in motor vehicles and aircraft.

As used herein, the term 'balance', when used in reference to a particular element, is used to describe that the remainder of the composition (in wt%) excluding any alloying additions is comprised of the designated element. Hence, the total composition including the 'balance' element in combination with other stated alloying elements is equal to 100 wt% of the alloy composition. The term "impurity" refers to a metallic or non-metallic element that is present in an alloy but which is not added intentionally. Hence, no lower limit is specified for the presence of the impurity.

The microstructure of a solid alloy takes form when a metal cools from molten liquid phase to a solid metal. This solidification forms a polycrystalline metal with different orientations of the grains. These grains will form the microstructure of the alloy and will contribute to its material properties including strength, resistance to fatigue and ductility. The form of this microstructure depends not only upon the heat transfer but also upon the alloy composition. The two major growth types of microstructures include the dendritic and the eutectic. Both types of growth are present in almost every aluminium alloy casting because of the improved castability of a near-eutectic or eutectic alloy than that of other compositions. Al-Si alloys typically used in casting applications will usually comprise a two-phase microstructure. This means that the alloy comprises larger grains (known as dendrites) that are predominantly comprised of aluminium with a small silicon component. The atomic structure of the dendritic phase can be referred to as an 'aluminium matrix'. During cooling, at the point that the eutectic temperature is reached the eutectic phase will begin to solidify between the dendrite arms filling this volume out. The eutectic phase consists of a primary alpha-aluminium (a-AI) phase and a silicon phase. This will give a two-phase microstructure where the primary a-AI crystals (dendrites) are distributed within the Al- Si eutectic phase. The ratio between the two phases will depend on the content of silicon and also the solidification time.

Turning now to the alloy of the invention, an aluminium-silicon casting alloy has a composition that comprises at least 7 wt% and at most 1 1 wt% silicon; at least 0.5 wt% and at most 4.5 wt.% copper; at least 0.1 wt% and at most 1wt.% magnesium; and at least 0.1 wt% and at most 1wt.% zirconium. The sum of the copper and magnesium compositions is at least 1 .0 wt% and at most 4.6 wt%.

The careful balance between the elements gives rise to certain precipitates in the microstructure of the cast alloy that provide an advantageous achievement of high strength at both low and high temperatures. In particular, as will now be described, the high proportion of zirconium (Zr) in the alloy firstly causes the formation of certain strengthening phases that are stable at high temperatures, and secondly modifies some of the strengthening phases that are stable at low temperatures so as to increase their high-temperature stability.

The high proportion of Zr in the alloy leads to the formation of Zr-AI-Si phase precipitates which have been identified through experimentation as Zr ₂ Si phases in which some Si atoms are substituted by Al atoms. Increasing the Zr composition leads to a reduction in the proportion of Al atoms in the Zr ₂ Si phase, which leads to the formation of an intermediate Zr ₂ (Si ₁ - _x ,Al _x ) phase. For ease of reference, this phase is referred to as a Zr ₂ Si phases, though it will be appreciated that term 'Zr ₂ Si phase' encompasses a phase in which some of the silicon ions are substituted with aluminium ions, as well as an unsubstituted Zr ₂ Si phase.

These Zr ₂ Si phases demonstrate higher thermal stability than the Al-Cu phases, and thereby act to inhibit dislocation gliding at elevated temperatures, which results in a significant increase in the mechanical properties of the cast alloy, as will be described in more detail below.

The Zr ₂ Si phases take the form of micro- or nano- sized particles that are uniformly dispersed within the microstructure of the alloy. The Zr ₂ Si phases have also been shown to exhibit a semi-coherent or coherent interface with a-AI matrix which results in an elliptical or spherical microstructure and enhances the microstructure's resistance to dislocation movement. The Zr ₂ Si phases therefore account for a significant contribution to the mechanical strength of the alloy, particularly at elevated temperatures.

The addition of Zr to the alloy composition also leads to the formation of AI ₃ Zr phases. These phases are shown to be present within the microstructure even at elevated temperatures (up to 400 °C). They have also been shown to exhibit a coherent interface with a-AI matrix which enhances the microstructure's resistance to dislocation movement, and accounts for a significant contribution to the mechanical strength of the alloy, particularly at elevated temperatures. The presence of the copper (Cu) in the alloy composition leads to the formation of Al- Cu phases, exemplified as AI ₂ Cu phase (θ-phase) precipitates, within the aluminium (a-AI) matrix, as is known in Al-Cu alloys. The θ-phase precipitates are considered to be the most effective strengthening phase for enhancing the alloy's mechanical properties at ambient temperatures. Such precipitates would usually be susceptible to dissolution at high temperatures due to Ostwald ripening, and so are known to be ineffective at providing strength at high temperatures. However, the high proportion of Zr in the alloy causes the presence of Zr in the AI ₂ Cu θ-phase precipitates. In this way, although the θ-phase precipitates are referred to as AI ₂ Cu precipitates for ease of reference, the AI ₂ Cu precipitates will in practice contain substituted Zr. The presence of Zr in these phases advantageously increases the stability of these phases at high temperature due to the large size of the Zr ion, thereby further improving the high temperature strength of the alloy.

The presence of Mg in the alloy composition in the specified content leads to the formation of magnesium-containing precipitate phases that act to strengthen the alloy, namely Mg ₂ Si and AI ₂ CuMg. In particular, the relative composition of Cu and Mg is provided such that the sum of the Cu and Mg compositions is at least 1 wt% and at most 4.6 wt% (1 wt%<(Cu+Mg)≤4.6wt%), which leads to the formation Al-Si-Cu-Mg containing phases (Q-phase) within the a-AI matrix. The formation of Q-phase precipitates is particularly enhanced due to the tuning of the Cu/Mg composition when casting at low solidification rates.

A further consequence of the Zr composition is the effect it has on the other intermetallic phases contained within the alloy microstructure such as the Q-phase precipitates, and the Al-Si-Cu intermetallic phase. The Zr atoms interact with, and/or are partitioned within, the Q-phase and Al-Si-Cu phase such that both the Q and Al-Si- Cu phases also contain Zr atoms.

In this way, although certain intermetallic phases are referred to for convenience as Al- Si-Cu and Al-Si-Cu-Mg intermetallic phases, these phases in practice contain substituted Zr. The presence of Zr in these phases causes the precipitates to exhibit enhanced thermal stability at elevated temperatures. This is caused, primarily, due to the very low diffusivity of Zr which acts to inhibit the Ostwald ripening of the host precipitates at elevated temperatures. The resulting increase in the alloys coarsening resistance leads to enhanced mechanical properties of the cast alloy.

It is considered that the presence of silicon in the above described ranges leads to an increase in the hardness of the resulting cast alloy. The presence of silicon also decreases the elongation due to the formation of silicon pools in the microstructure. Moreover, the addition of silicon enhances the alloy modulus and expansion coefficient, thereby making it an essential element for increasing the castability and fluidity of the alloy. This is particularly necessary for gravity casting applications where the alloy is poured into a mould without any external forces being applied to help fill the mould.

The relative composition between Si and Cu is maintained according to the relationship Si = 6 + Cu. In this way the composition of the Cu is tuned relative to the composition of Si, in order to maintain the alloy's castability. For example, if the Si composition were tuned according to the abovementioned ratio, for a baseline Cu composition of 4.5wt% the Si composition would be 10.5wt%. Furthermore, it is considered that by the combination of controlling the relative composition of Cu/Mg and Si/Cu leads to an improvement in the castability of the resulting alloy, thereby reducing the cost of manufacture of the cast components. By maintaining the relative Si/Cu composition within the above described parameters, it has been determined that the alloys exhibit beneficial mechanical properties at elevated temperatures whilst maintaining the copper composition below levels which would otherwise result in a significant increase in the risk of hot tearing and galvanic corrosion. Therefore maintaining the Si/Cu composition at these levels significantly improves the castability of the resulting alloy, whilst retaining the desired mechanical properties.

The alloy includes further optional alloying elements, as will now be described. In particular, the alloy includes the following elements: The alloy further comprises manganese in a proportion of at most 0.09 wt%. The alloy also further comprises iron in a proportion of at most 0.15% iron. Iron (Fe) is considered an inevitable alloying element in the formation of Al-Si alloys. Usually, its content is minimised, since its presence results in an iron-containing intermetallic phase that forms within an inter-dendritic region of the microstructure and that can lead to weakening, increased vulnerability to fatigue and reduced ductility. In particular, iron-containing intermetallic phase particles tend to adopt the 'script' morphology, which is detrimental to the properties of the alloy because it contributes to the fracture mechanism, reducing strength and also bringing about an increase in porosity.

However, an advantageous effect of the alloy described is that because of the proportions of elements present in the composition, when the alloy is cast, the majority of iron-containing intermetallic phase particles assume a substantially elliptical morphology having a semi-coherent interface with the Al-matrix of the cast alloy. This not only avoids the particularly undesirable 'script' morphology, but additionally the elliptical morphology of the Fe-rich precipitate phases form semi-coherent interfaces with the aluminium matrix, which have lower energy and thus cause these phases to exhibit higher phase stability at elevated temperatures. In this way, the advantageous Fe-Mn ratio of the invention produces Fe-rich phases whose morphology acts to retard Ostwald ripening of these phases and thus improves the coarsening resistance of the cast alloy at elevated temperatures.

The alloy also comprises at least 0.1 wt% and at most 0.3 wt% titanium. The presence of titanium leads to the formation of Al-Si-Zr-Ti phases, which are also stable at high temperatures, and which contribute still further to the high-temperature strength of the alloy.

The alloy of the present invention, once cast, therefore comprises many different precipitate phases that give rise to high strength at both ambient temperatures and at high temperatures (for example, temperatures above 200°C). These beneficial mechanical properties at both high and ambient temperatures make the cast component ideally suited for use in motor vehicle engine applications.

It is particularly beneficial that all the precipitates described above form without the need to induce supersaturation during casting of the alloy. In this way, the precipitates described above form even at low solidification rates of the sort experienced during simple casting methods such as sand casting and gravity casting.

Furthermore, owing to the balance of elements, and particularly the silicon content, the alloys of the present invention exhibit good castability allowing them to be used in the casting of components with complex geometries and dimensions.

An example of an alloy is provided as set out in Table 1 , with the balance made up from aluminium:

Table 1

In addition to the elements shown in Table 1 , the alloy composition may further comprise 0.01 -0.02 wt% Ni, 0.005-0.01 wt% Yb, and/or 0.005-0.01 wt% Gd. Further suitable alloying elements may be added if desired.

The inventors have found that the Al-Si alloy composition described forms an alloy structure with well dispersed micro- and nano-sized precipitates, which are thermally stable and resistant to deformation, and which result in an a 0.2% yield strength (YS) of up to 300 MPa, and an ultimate tensile strength (UTS) of up to 340 MPa. Furthermore, it is envisaged that the alloys of the invention would not require the presence of additional grain refining elements such as vanadium and molybdenum, which are otherwise generally required. Hence, the alloys of the invention provide a saving in terms of cost and ease of manufacture.

A method of casting a cast component 12 of a vehicle 10 (as shown in Figure 13) using the casting alloy will now be described. The cast component 12 is a part of a powertrain of the vehiclel O. In embodiments the cast component is the engine of the vehicle.

First, a casting alloy having the components described above is provided. To this end, the alloy is fabricated by melting a primary Al-Si alloy having an appropriate Al-Si ratio in a suitable furnace, such as an electrical resistance furnace of the type typically known in the art. The furnace operates at a suitable temperature, for example approximately 740°C, optionally 750°C, or approximately 700 °C. Next, the primary alloying elements including Cu and Mg are added.

The furnace temperature is then increased up to 800 °C and the Zr and Ti alloying elements are added. The furnace temperature is increased up to 800 °C at this stage to ensure the dissolution of the Zr and Ti additions in the melt. The Zr alloying element is added to the composition through the addition of an Al-Zr master alloy. The term master alloy would be understood by a person skilled in the art to describe a base, in this case aluminium which is combined with a relatively high percentage of one or two other elements, in this case Zr. In order to maximise the dissolution of Zr in the melt, the Al-Zr master alloy is first heated up to 450°C before being added to the melt. Following the addition of the Zr and Ti, further alloying elements are added to the melt including Sr and Mn.

The melt is then subjected to degassing, during which nitrogen gas (N ₂ ), alternatively argon gas (Ar), is blown into the melt by a commercial rotatory degasser at 400 rpm for 3 min. Alternatively, the degassing may proceed for as little as 1 minute or as many as 10 minutes, optionally as many as 30 minutes. The alloy is then processed by applying various suitable casting techniques, for example sand casting, investment casting, lost foam casting, gravity permanent mould casting, low pressure casting, high pressure die casting and squeeze casting etc. However, the best casting properties may be achieved using permanent mould gravity casting and low pressure die casting whereas sand casting provides the greatest control of the cooling rate of the alloy and therefore allows for greater control over the microstructure and mechanical properties of the alloy.

Improved mechanical properties are then achieved through retarding of precipitate coarsening at elevated temperatures. Further refinement, modification and dispersement of the abovementioned advantageous intermetallic precipitates is achieved through applying a number of different heat treatment steps (hereafter referred to as the T6 heat treatment) in which solutionising and ageing are used. The T6 heat treatment comprises two principle steps: a) homogenizing and b) age hardening. During the homogenizing step, the as-cast parts are held at high temperatures in the range from about 400°C to 600°C for a period of time which varies from 10 minutes to over 10 hours. In particular the temperature may be about 540 °C, optionally 550°C and the time may be about 6 hours, optionally 5 hours.

Once the microstructure is homogenized, the part is immediately quenched to a temperature below 50 °C in order to avoid the selective precipitation of unwanted precipitate phases. The quenching time may suitably be less than 10 seconds. After quenching, the quenched parts are subjected to another heat treatment which is an ageing treatment. This age-hardening step is carried out at low temperatures, preferably in the range of about 150 to 220 °C for a time of 0.5 hours to 40 hours.

The invention is further illustrated by the following non-limiting examples.

Example 1 - Compositional Analysis of Developed Alloy Cast components made from alloys having compositions listed in Table 2 were formed by gravity casting through a procedure of melting primary alloys and master alloys.

The alloys were melted in an electrical resistance furnace at 740°C. The melt was subjected to degassing, during which Ar gas was blown into the melt by a commercial rotatory degasser at 400 rpm for 3 min.

Table 2

After casting, the alloys of Table 1 were subjected to the T6 heat treatment, in which the alloy was homogenised at 540°C for 6 hours and then immediately quenched in water before being aged at 155°C for 24 hours.

Each of the samples were mechanically tested at 23°C, 200°C and 300°C. Mechanical testing undertaken at room temperature (23°C) was carried out according to the ASTM-E8 standard whereas the testing undertaken at elevated temperatures (200°C and 300°C) was carried out according to the ASTM-E21 standard. Each of the tests used a soaring time of 30 minutes for each testing temperature. The results of the mechanical testing are shown in Table 3.

Table 3 Ultimate Tensile

Alloy Yield strength (MPa)

Strength (MPa)

No. 1 261 181 18 282 210 21

No. 2 267 180 20 270 197 16

No. 3 298 226 88 337±16 244 95

No. 4 280 208 61 320 223 82

No. 5 197 167 15 231 205 73

Test Temp. (°C) 23 200 300 23 200 300

The results in Table 3 show that the yield strength (YS) and ultimate tensile strength (UTS) of alloy No. 3 is greater than that of alloy No. 1 across all three testing temperatures. In particular, Alloy No. 3 exhibits YS and UTS values which are, respectively, 25% and 17% higher than that of Alloy No. 1 at 200°C. The difference in the properties of alloy No. 1 and No. 3 become much more pronounced at even high temperatures. The YS and UTS of alloy No. 3 are several times greater than that of alloy No. 1 when tested at 300°C. Thus, according to the Table 2 and 3, it is demonstrated that the addition of Cu and Zr to the composition of Alloy No. 3 leads to a significant improvement in the alloy's mechanical properties at elevated temperatures (i.e. greater than 200°C) whilst maintaining its mechanical properties at ambient temperatures (i.e. 23°C).

The relative castability of the tested alloys was also investigated, the results of which are shown in Table 4. The castability of each alloy sample was determined by analysing a spiral casting of each alloy, each casting having a trapezoid cross section with a top edge, bottom edge and a height of 10mm, 6mm and 6mm, respectively. The results in Table 4 show that the each of the casting alloys, Nos. 1 to 5, exhibit excellent castability ratios (i.e. greater than 80%). Moreover, the results also show that the addition of Cu and Zr to the composition of Alloy No. 3 does not result in a significant reduction of the castability of the resulting alloy.

Table 4 Alloys The Ratio of Castability (%)

No. 1 95

No. 2 90

No. 3 80

No. 4 85

No. 5 100

Example 2 Microstructural Analysis of the Developed Alloy.

A cast component was made using a casting alloy as described above and the microstructure of the alloy was examined. Figure 1 shows the two-phase microstructure of the as-cast Al-Si alloy. The alloy formed with the composition according to an embodiment of the invention includes a number of distinct phases (including a number of Cu-rich phases, such as the labelled AI ₂ Cu phase, and an Si-rich eutectic phase) within an a-AI matrix. In particular, it is considered by the inventors that the Cu-rich phases form preferentially due to the relatively high percentage of Cu content in the alloy.

As can be clearly seen in Figure 1 , the Si-rich eutectic phases are concentrated in inter-dendritic areas while the presence of the Cu-rich phases are dispersed within a-AI matrix. As a result, the uniform distribution of these strengthening phases enhance the strengthening mechanism that operate within the alloy microstructure when it is subjected to stress.

Figure 2 is a micrograph of a section of the alloy shown in Figure 1 which has been heat treated. The heat treatment was performed at 540°C for 6-10 hours followed by an immersion of each sample in into water at ambient room temperature.

The microstructure shown in Figure 2 comprises eutectic Si and AI ₂ Cu phases. Additionally, a Mg ₂ Si precipitates phase was identified within the aluminium matrix. From Figure 2 it is clear that the heat-treated microstructure differs from the as-cast microstructure. In particular, the eutectic phases are transformed into separate, finer and more dispersed phases. It is also noted that the microstructure of the heat treated alloy (Figure 2) contains precipitate phases with predominantly uniform shape and morphology, which is considered a significant factor in the activation of the microstructural strengthening mechanisms that determine the enhanced mechanical properties of the resulting alloy. The refinement in the shape and size of the precipitate phases, as observed in the heat treated alloy of Figure 2, is considered to be enhanced due to the advantageous Zr alloying composition.

With further reference to Figure 2, the eutectic structure is clearly finer and more dispersed following the application of the heat treatment, when compared with the connected and coarse eutectic phases observed in the as-cast specimen shown in Figure 1 . It is also considered that the eutectic phases shown in Figure 2 are rounder and more distinct than those observed in an unmodified Al-Si alloy (not shown). It is considered that Sr plays a significant role in the modification of the eutectic phase. However, the presence of Zr in the alloys of Figures 1 to 4 is considered to further enhance the eutectic phase modification.

Figures 3 and 4 show the grain structure of the EN-AC-42000 alloy and the alloy of the invention, respectively, after they have each been subjected to the T6 heat treatment. The grain size of the heat treated alloy of the invention was observed to be considerably lower than that of the unmodified EN-AC-42000 alloy (Figure 3), which exhibited an average grain size of approximately 400 μιη in the heat treated form compared to the average grain size of the heat treated alloy of the invention (Figure 4) which exhibited a grain size of 201 μιη. Thus a 50 % reduction in grain size was achieved due to the advantageous composition of the alloy of the invention.

The result of an SEM/EDS phase composition analysis obtained from the regions of the as-cast and heat treated alloys described is shown in Figures 5a, 5b, 6a, 6b and 6c and listed in Table 5.

Table 5

The analysis of the EDS data indicates that the as-cast alloy consists of Al-Si phases known as eutectic silicon (#1 ), AI ₂ Cu θ-phase (#2), Al-Si-Zr (#9), Al-Si-Zr-Cu (#10) and Al-Si-Zr-Ti (#12) phases. It is considered that, during formation of the Al-Zr phase precipitates, Si occupies the Al sublattice sites of the L12 ordered phases. Additionally, it was determined that Titanium (Ti) atoms also incorporate in the formation of Al-Si-Zr phase (see Figure 6c and Table 4). Also observed in the alloy's microstructure were Fe-containing phases including Al-Si-Fe (#13), described as β-phase, Al-Si-Fe-Cu (#14) and Al-Si-Fe-Mn-Cu (#15) which are found predominantly in the interdendritic region of the microstructure.

Typically, Fe precipitates form flake-, needle- and script-shaped phases that induce detrimental effects to the mechanical properties of an alloy. However, the Fe phases shown in Figure 6 can be seen to form fine and round shapes, which are considered to contribute to the thermal stability of the microstructure at elevated temperatures. This observed modification of Fe-containing phases is attributed to the effects of the T6 heat treatment and also due to the addition of Mn in this alloy.

In addition to the micro-sized phases present in the respective as-cast and heat treated microstructures of Figures 1 , 2, 5 and 6, TEM micrographs revealed that nano- size phases are distributed homogenously throughout the matrix of these microstructures. Figures 7 and 8 show, respectively, TEM images of Fe- and Zr- containing particles which are each distributed within the aluminium matrix of the as- cast alloy of Figure 1 . The TEM micrographs show that both the Zr and Fe nano-sized precipitate phases form in an elliptical shape.

EDX spectra of the Fe and Zr nano-sized particles are shown in Figures 9 and 10, respectively. The Fe-containing phase also contains Al, Si, Mn and Cu elements along with Fe. It is considered that the Mn modifies the Al-Si-Fe-Cu phase to form AI ₁₅ (FeMn) ₃ (CuSi) ₂ , according to the atomic ratio derived from the EDX spectral analysis shown in Figure 9. It is considered that the Zr ions in the Al-Si-Cu-Mg alloys are partitioned among other phases and/or elements such as Al, Si, Cu resulting in the formation of Zr ₂ Si phases (as shown Figure 8). It is noted that the Zr-containing phases form finer and more dispersed precipitates than those containing Fe phases. The Zr ₂ Si phase exhibits a hexagonal-close-packed (HCP) crystal structure which determines the formation of the elliptical shape of the Zr ₂ Si particle (at equilibrium). Consequently, the Zr ₂ Si is formed of a series of large coherent/semi-coherent facets connected by shorter incoherent curved edges. A similar equilibrium structure is observed for the Fe-containing phases.

Mechanical Testing

Mechanical testing was carried out on samples of the cast alloy. The following results, graphically represented in Figures 1 1 and 12, were obtained from ambient and high temperature tensile testing, which was performed according to the ASTM-E8 and ASTM-E21 standard, respectively. For comparison, the results for the base EN-AC- 42000 alloy are also shown.

According to the Figures 1 1 and 12, modification of the Cu and Zr composition is shown to cause significant improvement in both the YS and UTS of alloys according to the invention as compared with that of EN-AC-42000 alloy.

The improvement is observed in the as-cast alloy, indicating that the composition selection alone gives rise to beneficial strength effects. The improvement is also seen in the heat-treated alloy, indicating that the heat treatment also plays a role in the improved mechanical properties.

The maximum yield strength was observed after ageing the castings at 155-165°C for 24 hours. However, the lower ageing time resulted in a decline in the alloy's mechanical properties at both room and high temperatures. It is therefore considered that optimization of intermetallic strengthening phases is strongly influenced by ageing time. This is most likely due to the significant increase in the micro- and nano-size phases which are dispersed uniformly throughout the matrix of heat treated alloy (Figure 2) compared with the as-cast alloy (Figure 1 ).

With reference to Figure 1 1 and 12, an improvement in the YS and UTS is observed in the samples which underwent the heat treatment according to an embodiment of the invention. In particular, the respective YS and UTS values for the as-cast developed alloy are 135 MPa and 206 MPa, respectively. The YS and UTS for the heat treated samples is 291 MPa and 335 MPa, respectively. The heat treatment therefore results in the observed enhanced mechanical properties by encouraging phase formation, refinement and morphological modification of the alloying phases that form in the microstructure during the casting of the alloy.

In addition to the phases, grain structure has a substantial influence on the yield strength. Figure 4 demonstrates that the alloy according to the invention consists of fine grains in the heat treated form. As is mentioned above, this fact can be attributed to the addition of Zr-forming phases which are stable in the Al liquid phase. Moreover, grain refinement is also brought about due to the Fe-rich phases which improve heterogeneous distribution of Cu in the a-AI matrix due to refinement of Cu-containing phases. Therefore, a higher number of nucleation sites in the alloy of the invention is thought to have led to the enhanced refinement of the grain structure that is observed in Figure 4.

The enhancement of YS in finer grain structure can be described through Hall-Petch equation representing the relationship between the grain size and YS. Mechanical strengthening by grain refining has an advantage over phase and particle dispersion strengthening. In spite of the particle strengthening mechanism decreasing elongation, grain refinement results in simultaneous increment of strength and toughness.

The YS and UTS of the developed alloy are 30% and 25% higher than that of EN-AC- 42000 at room temperature. Surprisingly, higher improvement of mechanical properties was observed at elevated temperatures as compared to the room temperature for the developed alloy, as both YS and UTS of the developed alloy were around 45% higher than that of EN-AC-42000 alloy at 200 °C. The significant reduction of YS and UTS at 200 °C is attributed to the softening of alloy on account of dissolution and coarsening of strengthening phases. The degradation of strengthening phases due to coarsening is considered to be the dominant mechanism, at elevated temperatures, causing the observed reduction of alloy mechanical properties.

It is noted that the coarsening rate of each phase is strongly controlled by the bulk diffusion of alloying elements. Hence, a rise in the diffusivity of the alloying elements leads to an acceleration in the coarsening rate. The Zr ₂ Si phase precipitates consist of Zr, which has a relatively high diffusion activation energy, 242 KJ/mol, and a relatively low diffusivity, 1 .20x10 ^"20 m ² /s at 400°C, in a-AI.