WO2009132388A1

WO2009132388A1 - Improved aluminium based casting alloy

Info

Publication number: WO2009132388A1
Application number: PCT/AU2009/000532
Authority: WO
Inventors: Roger Neil Lumley; Dayalan Romesh Gunasegaram; Maya Gershenzon; Andrew Colin Yob
Original assignee: Commonwealth Scientific and Industrial Research Organization CSIRO
Current assignee: Commonwealth Scientific and Industrial Research Organization CSIRO
Priority date: 2008-04-30
Filing date: 2009-04-30
Publication date: 2009-11-05
Anticipated expiration: 2010-10-30
Also published as: US20110100515A1; CN102016092A; EP2285995A1; AU2009242962A1

Abstract

An aluminium based alloy has a weight percentage composition of from 5 to 15% silicon, 0 to 0.25% magnesium, 0 to 0.25% titanium, 0.2 to 0.65% manganese, 0.1 to 0.6% iron, 1 to 4% copper, 0 to 3% zinc, less than 0.01% in total of silicon modifiers (with less than 0.007% strontium), less than 0.05% tin, less than 0.2% in total of other transition or rare earth metals (with less than 0.05% chromium), less than 0.5% in total other elements, and a balance of aluminium. The limits for iron and manganese are constrained such that the amount of iron present in the alloy is 0.4 to 1.6 times the manganese content and the alloy has a sludge factor (SF), calculated as SF = (1 x wt% Fe)+(2 x wt% Mn), of from 0.8 to 1.6. A casting made from the alloy has enhanced fracture resistance relative to casting of the same product made of a conventional HPDC alloy when compared in the as cast or same heat treated state.

Description

Improved Aluminium Based Casting Alloy

Field of the invention

This invention relates to an aluminium based alloy for manufacture of cast parts which display enhanced fracture resistance.

Background to the Invention

Typical alloys for high pressure diecasting (HPDC) are based around the Al-Si alloying system, with a range of additive elements present. Some common HPDC alloys from different regions of the world and their compositions are shown in Table 1. More recently however, there has been a range of alloy compositions developed for high ductility such as the five specialty alloys Magsimal, Silafont, Aural 2, 367.0 and 368.0, and each of these has found application in components used in crash sensitive applications where high energy absorption is required. The compositions of these alloys are shown in Table 2. These specialty alloys generally display levels of ductility between 8 and 20% in the as-cast conditions. Some common features of these three alloys are as follows:

1. All have a specification for Cu which is very low, typically less than 0.25%.

2. All have very low levels of Fe, typically less than 0.25%.

3. All have Mn present in the range of 0.25 to 0.8% to help prevent die sticking.

4. All are generally used as primary alloys to avoid contamination or degradation. Table 1 - Composition Limits of Some HPDC Alloys From Different Regions

K)

"may also contain up to 0.1 Cr

Table 1 (continued)

Table 2 - Some Specialty Primary High Ductility Alloys

Max Sr in casting 0.015

Silafont, Aural 2, 367 and 368 are similar and based around the Al-Si system, whereas Magsimal is based on the Al-Mg alloying system. However, several of the alloys of Table 2 also are required to have low contents of Zn, which may also assist in avoiding the formation of Zn vapour porosity in the components. Generally, these three alloys also preferably are produced using vacuum diecasting, to reduce porosity that can be considered detrimental to ductility.

All of the above features 1 to 4 are known to assist in improving ductility of other HPDC alloys. Furthermore it should be noted that iron and manganese are present in HPDC alloys to minimize die soldering (sticking) so that die life is increased and productivity is improved. In general, the Fe content needs to be above 0.5% to avoid soldering, but can be substituted by other elements such as Mn, as in the case of the specialty alloys mentioned. In other cases, Sr may be used for the prevention of die soldering. However, Fe allowances are preferred because a large proportion (estimated at greater than 95%) of HPDC alloys are made of secondary metal. Additionally, Fe tends to accumulate during recycling operations, and, as a result, is an important factor to be considered in the development of alloys. Fe, Mn and Cr tend to form intermetallics that will be present in differing amounts and morphologies, depending on the so-called "sludge factor" of the alloy. In high pressure diecasting, a proportion of sludge particles form in the melt, some in the shot sleeve, and some during solidification in the die. Generally, a simple rule to be followed is that:

Sludge factor (SF)=(I x wt%Fe)+(2 x wt%Mn)+(3 x wt%Cr),

where SF less than 1.7 is generally considered to be most acceptable in practice.

Cr is purposefully omitted in most compositions worldwide, due to its toxicity. It can however sometimes be present as a minor impurity. When no Cr is present, the relationship is then: SF=(I x wt%Fe)+(2 x wt%Mn) The morphology of the sludge particles is strongly influenced by the Fe:Mn:Cr ratio of the alloy, and because of this, Mn and Cr are sometimes referred to as Fe correctors since they can eliminate a proportion of the undesirable, needle-like Fe-Al intermetallic phases within the alloy, causing the Fe-Al intermetallics to form more innocuous particle morphologies such as star-like rosettes, blocky shaped, Chinese script or polyhedral particles. In general, it is the Chinese script morphology that is considered to be most desirable. However, the Chinese script form is generally associated with additions of Cr, which is now purposefully omitted from most alloy specifications due to toxicity concerns. Although other transition metal elements (for example Co, Ni, V) may also serve as Fe correctors, Mn is the most common in commercial alloy products since it often appears in alloys produced by recycling. When the ratio of Fe:(Mn+Cr) is high, the sludge particles tend to be needle-like. When the ratio of Fe:(Mn+Cr) is low, the particles tend to exhibit more Chinese script, blocky, star-like or polyhedral particles. In general, the recommended amounts of Fe correctors are of the order of half the concentration of Fe. However, adding Fe correctors to an alloy also increases the total fraction of hard, brittle particles that may be detrimental to machining, fracture resistance and ductility of the alloy. Some sludge particles may also be more detrimental than needle-like particles to fracture resistance (e.g. star like or rosette shaped) when they are present in larger amounts because they may behave as large brittle plates. The size, shape and distribution of these particles are therefore very important.

Although ductility and fracture resistance may both be affected by similar microstructural features, they are not necessarily directly related, and high ductility does not necessarily mean high fracture resistance. A survey of the literature concerning the fracture properties of HPDC products has revealed that little information exists, with the exception of a limited amount of data on Charpy impact energy. It has also become a widely accepted practice in the design of HPDC products to rely solely on tensile elongation as a measure of fracture resistance. However, neither Charpy impact energy nor tensile ductility is considered in scientific literature to be a valid indicator of the comparative fracture toughness properties of different aluminium alloys. Although alloys which have high ductility may often show high fracture resistance, similar alloys with lower tensile ductility may also display high fracture resistance. Similarly, for a specific ductility value, fracture resistance may vary from alloy to alloy. Fundamental to solving this dilemma therefore is the acquisition of valid fracture toughness data.

HPDC products typically approximate to either a single projected plate or a series of interconnected plates. Such castings have a nearly constant wall thickness, often 2 to 6mm. There are, however, exceptions when locally increased section thickness is required to increase stiffness, to improve metal feeding during casting or to allow for geometric features such as bolt holes and most parts display varying thickness as well as geometric stress raisers. Such changes in thickness or geometric stress raisers may approximate to a notch within the material, which then influences the overall performance of the product in service. As may be appreciated, it is rare to have applications that fail by a purely tensile mode.

ASTM standard B871 describes a method for comparative evaluation of fracture resistance using a technique developed specifically for testing wrought aluminium sheet and plate materials, which is also an acceptable representation of fracture resistance for cast aluminium alloys. The technique also provides an estimate of the notch sensitivity of the material. In general, the tear test method of ASTM standard B871 provides load displacement curves relating to crack initiation and propagation. The area of initiation up to the point of maximum loading when cracking begins, as well as the area of propagation under the curve describes the entire fracture process as it relates to thin plates. The maximum load obtained during the testing of ASTM B871 may be used to determine the tear strength of the material. The ratio of the tear strength to the yield strength (0.2% proof stress) or TYR (Tear to Yield ratio) value gives a measure of notch sensitivity. The area under the whole curve provides the unit total energy, and the area from the peak load to the end of cracking gives the unit propagation energy. Results are presented along with tensile data (0.2% proof stress, tensile strength, ductility) taken from the same material batch. The inherent reliability of the tear test method has also been shown to provide reasonable correlations with other fracture toughness data such as the critical strain energy release rate G₀.

As might be expected, the fracture resistance of HPDC products is related to a range of failure types. For example, when a fatigue crack reaches a critical length during service, conditions for rapid and uncontrollable fracture may result. As a result, it has also been shown that, at least in some wrought alloys, there is a relationship between the unit propagation energy derived from the tear test, and the rate of crack growth in fatigue. This is not altogether unexpected, since increased fracture toughness may also correspond to improved fatigue resistance.

Optimally, a HPDC component is utilized in the as-cast state. However, there is a range of thermal processes that may be applied to a HPDC alloy. For example, in patent application WO2006/066314, methods are shown by which conventionally produced HPDC alloys and components may be successfully heat treated without displaying surface blistering or dimensional instability, and a range of heat treatments has been demonstrated that take advantage of the procedures disclosed in application WO2006/066314.

Summary of the Invention

The present invention provides an aluminium based casting alloy which provides high fracture resistance mostly irrespective of the ductility of the alloy, although high ductility in the alloy is an additional advantage in the alloy when casting quality is improved. The alloy is applicable to castings in which porosity may be present. Castings of the alloy may be produced by what can be regarded as a conventional or usual die casting technique, such as with a cold-chamber die casting machine. The castings may be produced with or without either an applied vacuum or use of a reactive gas. The castings may alternately be produced by high integrity casting processes to achieve minimum levels of porosity.

An aluminium based alloy according to the invention has a weight percentage composition of:

silicon 5 to 15% magnesium 0 to 0.25% titanium 0 to 0.25% manganese 0.2 to 0.65% iron 0.1 to 0.6% copper 1 to 4% zinc 0 to 3% silicon modifiers less than 0.01 % in total (with less than

0.007% strontium) tin less than 0.05% other transition or rare earth metals less than 0.2% in total (with less than

0.05% chromium) other elements less than 0.5% in total,

and a balance of aluminium,

wherein the limits for iron and manganese are constrained such that the amount of iron present in the alloy is 0.4 to 1.6 times the manganese content and the alloy has a sludge factor (SF), calculated as SF = (1 x wt% Fe)+(2 x wt% Mn), of from 0.8 to 1.6. The alloy preferably is free of beryllium, rare earth elements and transition metal elements other than those individually identified (that is, other than Ti, Mn, Fe, Cu and Zn).

The invention also provides a casting having enhanced fracture toughness relative to casting of the same product made of a conventional HPDC alloy, such as an A380 alloy, when compared in the as cast or same heat treated state, wherein the casting having enhanced fracture toughness is cast from an aluminium based alloy having a weight percentage composition of:

0.007% strontium) tin less than 0.05% other transition or rare earth metals - less than 0.2% in total (with less than

0.05% chromium) other elements less than 0.5% in total,

and a balance of aluminium,

wherein the limits for iron and manganese are constrained such that the amount of iron present in the alloy is 0.4 to 1.6 times the manganese content and the alloy has a sludge factor (SF), calculated as SF = (1 x wt% Fe)+(2 x wt% Mn), of from 0.8 to 1.6; and wherein the casting having enhanced fracture toughness has a microstructure exhibiting silicon formed from solidified eutectic which is modified and substantially free of acicular silicon particles, and exhibiting iron-bearing phases which are substantially fine polyhedral particles.

The casting preferably is free of beryllium, rare earth elements and transition metal elements other than those individually identified (that is, other than Ti, Mn, Fe, Cu and Zn).

The role of each of the elements of both the alloy and the manufacture of the casting of the invention now will be discussed in turn.

Silicon is required in the alloy to depress the melting temperature, aid fluidity and increase strength. Compositions ranging from hypoeutectic through to hypereutectic are applicable within the limits of 5 to 15%, but all require good fluidity to aid casting. The Si level preferably is from 6.5 to 10.5% and more preferably from 6.5 to 8.5%. This corresponds to the optimal casting conditions for the majority of instances. Below the lower limits of Si content, castability may be adversely affected. Above the upper limits of Si content, the high proportion of the Si phase produces proportionately higher embrittling effects and reduced resistance to crack propagation.

Copper is present to also aid fluidity and to provide strengthening to the alloy, optionally by heat treatment, where required. In general, Cu levels around 1.5 to 3 wt% are optimal in the present invention, but levels as low as 1 % and as high as 4% may also be considered suitable in some applications. Below 1 %Cu, any heat treatment response will be limited, and above 4%Cu, embrittling effects arising from residual Cu-based intermetallics may be evident. Additionally, these Cu-based intermetallics may adversely influence corrosion resistance. Higher levels of Cu may also incur a cost penalty. Iron and manganese concentrations are interrelated. Although some Cr, rare earth elements or other transition metals may be as functional in practice as Mn, their presence should be restricted or reduced to trace element levels wherever possible to promote the formation of (AI-Is(Mn, Fe)sSi2 in a fine polyhedral form. Due to toxicity concerns regarding Cr, it is preferable to limit Cr content to a minimum. Because of this, the correct proportions and amounts of both Fe and Mn are required. Following the general rule for sludge factor (SF) of the alloy, for the purpose of the present invention, SF= 0.8 to 1.6 = (1xFe)+(2xMn), and preferably SF=LO to 1.3. This ensures good castability, no die sticking and the presence of the correct forms of intermetallic phases. As a result of this sludge factor limitation and that of the alloy composition, the relative amounts of Fe and Mn for the invention are readily determined. If Fe content is proportionately high compared to Mn, the alloy tends to form distributed particles of FeSiAIs .which are present as needles in the microstructure. If the Mn content is proportionately high compared to Fe, the alloy may more readily exceed the sludge factor limits or not form the appropriate morphology of hard particles.

Ideally, the Fe content should be 0.4 to 1.6 times the Mn content by weight, and most preferably 0.5 to 0.9 times, within the scope of the compositional limitations and sludge factor limits. For example, if 0.4% Mn is present in the alloy, then 0.4 times the Mn content gives minimum Fe content of 0.16% with an alloy sludge factor of 0.96. 1.6 times the Mn content gives 0.64% Fe, but there is a maximum Fe allowance of 0.6% to give an alloy sludge factor of 1.4. If the alloy contains 0.6%Fe, then the minimum Mn allowable is 0.375% and the maximum content of Mn permitted is 0.5%, in order to stay within the limits of the sludge factor maximum of 1.6. In the present invention, these ratios and limitations positively influence the microstructure that develops both in the as-cast and heat treated conditions by promoting the formation of innocuous transition metal containing intermetallic particles that are polyhedral rather than plate or needle like in shape. The Fe content preferably ranges from 0.2 to 0.4%, while the Mn content preferably is from 0.3 to 0.6%.

Rare earth metals are reported to provide both Fe modification and Si modification, but they are preferably omitted from the present invention, due to toxicity, cost and availability. Trace amounts of rare earth elements may however, be present in secondary aluminium alloys where contamination has arisen from other sources. The rare earth elements, combined with other transition metals, should be kept below 0.2% in total.

Strontium is known as a modifier to silicon in cast aluminium alloys. However, it's presence is not necessary in the present invention to promote modification and eliminating it from the alloy has various benefits, such as avoiding the potential formation of Sr-Fe intermetallic compounds which will increase relative sludge content. Sr also is known to increase porosity in aluminium castings. Similarly, Na, Ca and P are omitted or kept to very low levels. Use of any of these four elements as purposeful additions provides a cost penalty associated with the use of master alloys typically utilized, but Sr in particular may appear as trace amounts if alloy is manufactured from recycled material. In this case, the combined levels of silicon modifying elements should be kept at a level below 0.01 wt%, with strontium below 0.007 wt%.

Titanium may be present in small quantities, of up to 0.25%, as an optional element but it is not necessary for the efficacy of the present invention. Above this limit, Ti will form coarse sludge particles which may be detrimental to fracture resistance.

Zinc may be present at levels up to 3%, but its presence is not necessary for the functionality of the invention. Zinc may be present to improve castability, machinability or corrosion resistance in the alloy and optimally, may be 0.3 to

1.0%. Elevated Zn contents such as up to 3% may arise due to the presence of Zn diecast material in recycling operations, but higher levels should be avoided where possible due to the high vapour pressure of Zn resulting in additional gas phase porosity in castings.

Tin should be omitted within the alloy of the invention, or restricted to trace element levels as specified. Tin may have a detrimental effect on the shape of Fe-containing intermetallic particles causing them to form rosettes comprised of brittle phases. Tin may also cause severe problems with die sticking and hot tearing, even when present at very low levels.

Magnesium is often present simply as an impurity in high pressure diecasting alloys which contain Cu. It may however increase strength and modify the precipitation occurring in the alloy during cooling following casting or during heat treatment. At the lower limits, ~0.05%Mg is required to initiate an age hardening response in HPDC alloys. Higher Mg levels (e.g. greater than 0.2wt%) lead to progressive reductions in fracture resistance.

Beryllium is known to provide various advantages to aluminium alloys. It is however highly toxic and expensive and should therefore not be permitted or included in the alloy other than trace, incidental amounts which may arise due to the use of secondary metal.

By the manufacture of the abovementioned alloy, it may be appreciated that due to the generous Fe allowance, that the alloy may be manufactured by blends or proportions of secondary alloys. Similarly, as the alloy is similar to the broad specifications of many alloys utilized worldwide, it may be readily recycled with other secondary material without segregation. As may be appreciated, it may also be recycled into alloy within the limits of the alloy composition of the present invention.

The sludge factor SF, as detailed above, generally is taken as: SF=(I x wt%Fe)+(2 x wt%Mn)+(3 x wt%Cr), although Cr usually is not present as more than a minor impurity, such that effectively:

SF=(I x wt%Fe)+(2 x wt%Mn).

In the present invention, Cr is constrained by the requirement that transition metal elements other than Ti, Mn, Fe, Cu and Zn, plus rare earth metals, are present at less than 0.2 wt% in total. As a practical matter, that limit of less than 0.2 wt% constrains the Cr content to a sufficiently low level at which its influence in raising the sludge factor can generally be disregarded. That is, given that any Cr present will be accompanied by other metals subject to the limit of less than 0.2 wt% in total, the Cr content will be sufficiently low. However, it generally is desirable that the content of Cr alone is less than 0.05 wt%, such as below 0.02 wt%.

The alloy of the present invention is most highly suited to the process of high pressure diecasting, but may also be utilized on other casting techniques such as squeeze casting or thixo-casting. There is also the possibility that the current alloy may also be suitable for permanent mold casting or sand casting.

Brief Description of the Figures

Figure 1 shows typical tear test fracture curves for either low tear resistance or high tear resistance fracture, as determined by the standard tear test method in accordance with ASTM B871 ;

Figure 2(a) shows the configuration of the plate shaped castings used for tear test development purposes with the alloy of the present invention;

Figure 2(b) shows a tear test fracture sample and a tensile sample taken from castings as in Figure 2(a); Figure 3 shows typical tear test fracture curves for the as-cast condition of a alloy within the specification of A380 (also representative of CA313 (Australian designation), ADC10 (JIS designation), AISi8Cu3(Fe) (European designation)) compared to a similar alloy according to the present invention;

Figure 4 shows typical tear test fracture curves for the T4 condition of the same alloys as for Figure 3;

Figure 5 shows typical tear test fracture curves for the T6 condition of an alloy of A380 (also representative of CA313, ADC10, AISi8Cu3(Fe)), compared to a similar alloy according to the present invention;

Figure 6 shows microstructures at the centre of high pressure diecastings in the as-cast condition, comparing a conventional A380 alloy (also representative of CA313, ADC10, AISi8Cu3(Fe)) with a similar alloy according to the present invention;

Figure 7 shows microstructures at the centre of high pressure diecastings in the as-cast condition, similar to Figure 6 but now at higher magnification;

Figure 8 shows microstructures at the centre of high pressure diecastings in the T6 condition, comparing a conventional A380 alloy (also representative of CA313, ADC10, AISi8Cu3(Fe)) with a similar alloy according to the present invention;

Figure 9 shows different castings prepared similarly to those shown by Figure 7, now comparing the same A380 alloy (also representative of CA313, ADC10, AISi8Cu3(Fe)) with a different alloy of the present invention;

Figure 10 shows the evolution of properties with duration of ageing for an alloy according to the invention; Figure 11 is similar to Figure 10, but relates to a different casting of the same alloy;

Figure 12 is similar to Figure 11 , but unit total energy and unit propagation energy; and

Figure 13 provides a comparison of the 0.2% proof stress and unit propagation energy for four conventional HPDC alloys with a range of the newly developed alloys according to the current invention in as-cast and heat treated (T4 or T6) conditions.

Detailed Description of the Figures

Figure 1 shows a respective typical tear test fracture curve for each of low tear resistance and high tear resistance fracture, as determined by the standard tear test method in accordance with ASTM B871. The curves are characterized by three main features. The energy of fracture initiation is given as the area under the curve up to the point of fracture initiation (i.e. the maximum load). Secondly, the maximum load P from which the tear strength is derived as 4P/bt, where b is the width from the notch root to the back edge, and t is the thickness. Finally, the energy of fracture propagation is given as the area under the curve after fracture initiation until final failure when the load approaches zero. As may be appreciated, the total energy of fracture is the sum of fracture initiation and fracture propagation.

The curves of Figure 1 for fracture according to ASTM B871 provide information regarding the initiation and propagation of cracks in a sharply notch plate shaped specimen. Of great significance in the standard is the root radius of the notch, which is specified as being 25±12μm. For the current results, this was achieved on samples using a single tooth cutter lapped to a very sharp edge at the 60 degrees specified angle. Each sample tested according to ASTM B871 was individually examined prior to testing to ensure compliance with the standard. Figure 1 shows the most important features of the tear test load displacement curve, being the maximum force P, from which the tear strength may be derived; the energy associated with crack initiation, and the energy associated with crack propagation. To derive the unit total energy (UTE), the total area under the curve is divided by the cross section; to derive the unit propagation energy (UPE) the area past the peak load is divided by the cross section from the notch root to the back of the sample (see ASTM B871 for further details). Unit initiation energy is simply UTE - UPE.

Figure 2(a) shows the configuration of the plate shaped castings used for tear test development purposes with the alloy of the present invention, as well as the runner system, overflows and ejector pin positions. The dimensions of the plates are 70mm wide, 60mm long and approximately 2mm thick. From these and in accordance with the standard, one tear test fracture sample and one tensile sample were taken from each plate (Figure 2b). For each condition tested, 5 tear and 5 tensile samples were tested to generate data.

The samples were oriented either with the propagating crack running parallel to the direction of metal flow (i.e. stress axis perpendicular to flow direction), or with the propagating crack running perpendicular to the direction of metal flow (i.e. stress axis parallel to the direction of metal flow). According to convention as it relates to rolled products, the fracture samples are hereafter designated as TF where the propagating crack runs in the direction of metal flow, and FT where the propagating crack runs perpendicular to the direction of metal flow. This is necessary because the current work has revealed that some alloy compositions prepared from the aforementioned HPDC plates display directionality of properties. Corresponding tensile samples are hereafter referred to as the T direction where the propagating crack in the tensile sample runs in the direction of metal flow, (i.e. corresponding to the TF fracture orientation) or the F direction where the propagating crack is perpendicular to the direction of metal flow. (i.e. corresponding to the FT fracture direction). As an example, the machined samples shown in Figure 2(b) are representative of the TF and T orientations. For each condition tested, 5 tear and 5 tensile samples were tested to generate data.

Figure 3 shows typical tear test fracture curves for the TF orientation for the as- cast condition of an A380 specification alloy (also representative of CA313, ADC10, AISi8Cu3(Fe)) having a composition of AI-9Si-0.86Fe-3.1 Cu-0.16Mn- 0.10Mg-0.11 Ni-0.53Zn-(less than 0.2 other), compared to an alloy (designated alloy X for the purposes of this comparison) according to the present invention and having a composition of AI-7.6Si-0.27Fe-2.74Cu-0.48Mn-0.12Mg-0.43Zn- (less than 0.2 other), for which SF=1.23. The curve marked "A" on the tear test plot corresponds to the standard A380 alloy, and the curve marked "B" is the alloy X. As is readily observed, the maximum load is higher and the energy absorbed (as determined by the area under the respective initiation and propagation portions of the curves) is increased for the alloy X of the present invention compared to the conventional A380 alloy composition. All samples were produced by high pressure diecasting with a metal velocity at the gate of 56 m/s. The average tear and tensile properties for the samples tested are shown in Table 3 (five tear and five tensile samples were tested for each comparative data set) and it will be seen that the values of fracture resistance are higher for the alloy X of the present invention in the as-cast condition. The unit propagation energy is 6.46 KJ/m² for the A380 alloy, while the unit propagation energy is 16.88 KJ/m² for the alloy X. The initiation energy (and hence UTE) is also increased for the alloy of the present invention. Additionally, the notch sensitivity index (TYR), as given by the ratio between tear strength and yield stress (0.2% offset) shows reduced sensitivity for the alloy X compared to the conventional A380 alloy. That is, the higher is the number of the index, the lower is the sensitivity to the presence of the sharp notch. Figure 4 shows typical tear test fracture curves for the T4 condition in the TF orientation of an A380 alloy having a composition as for Figure 3, compared to the same alloy X of the invention as for Figure 3. The curve marked "A" on the tear test plot is the standard A380 alloy, and the curve marked "B" is the alloy X according to the present invention. The alloy X of the present invention displays superior fracture properties when compared to the conventional A380 alloy, as noted in Table 3. The yield stress is similar for each alloy, although the ductility is higher for the alloy X. The average data set for the samples are shown in Table 3 and it will be seen that values describing fracture resistance are higher for the alloy X. For example, compared to a unit propagation energy (UPE) of 12.39 KJ/m² for the A380 alloy composition, a UPE value of 33.47 KJ/m² for the alloy X represents an increase in energy absorbed during crack propagation of 170%. Additionally, the notch sensitivity index, as given by the ratio between tear strength and yield stress (0.2% proof stress) shows reduced sensitivity (i.e. a higher value) for the alloy X. Interestingly as well, the comparative curves A and B in Figure 4 suggest that a degree of ductile yielding is occurring in the alloy X (curve B) prior to cracking despite the presence of the sharp notch. This is also reflected in the comparison of the initiation energy values for the two conditions provided in Table 3, where the initiation energy from curve A is 1.05 N.m, and the initiation energy from curve B is 1.48 N.m. The notch sensitivity index (TYR), as given by the ratio between tear strength and yield stress (0.2% proof stress) shows beneficial reduced sensitivity for the alloy of the present invention.

Figure 5 shows typical tear test fracture curves for the T6 condition of an A380 alloy of the same composition as for Figure 3, compared to the same alloy X according to the present invention. The curve marked "A" on the tear test plot is the normal A380 alloy, and the curve marked "B" is the alloy X according to the present invention. Consistent with Figures 3 and 4, the alloy X of the present invention displays superior fracture properties when compared to the A380 alloy of similar composition. The yield stress is slightly reduced for the alloy X although the average ductility is higher (Table 3). The data set for the samples are shown in Table 3 and it will be seen that the values are consistently and significantly higher for the alloy of the present invention. Additionally, the notch sensitivity index, as given by the ratio between tear strength and yield stress (0.2% offset) shows beneficial reduced notch sensitivity for the alloy of the present invention.

Table 4 shows tear test and tensile data of the same alloy X composition of the present invention, as in Figure 3, tested in the as-cast, T4 , T6 and T7 conditions, but now with the crack running perpendicular to the direction of metal flow (FT) compared to parallel to the direction of metal flow (TF) as shown in Table 3. In general, the fracture properties are moderately higher in the FT orientation, and the tensile ductility is also slightly higher in the F orientation shown in Table 4 compared to the T orientation shown in Table 3. Five tear test samples and five tensile samples were tested in each condition.

Table 5 shows tear test and tensile data of the same alloy X of the present invention as in Figure 3, tested in as-cast, T4 and T6 conditions, tested in the FT and F orientations. Five tear test samples and five tensile samples were tested in each condition. Whereas the data generated for Tables 3 and 4 was done with samples produced at a melt velocity at the gate of 56 m/s, the data for Table 5 was generated from samples produced at a melt velocity at the gate of 26 m/s. The difference in general between the two manifests in the part quality, which is often superior in the samples produced at 56 m/s than for those produced at 26 m/s. As can be seen by comparing Tables 4 and 5, the tensile properties, in particular the ductility, is lower in the examples of Table 5. However, the fracture properties determined from tear testing remain relatively high.

From Tables 3, 4 and 5, it can be surmised that the two main factors influencing the fracture properties of the alloy are the composition and the choice of temper applied to the alloy. The levels of ductility present do not appear to strongly relate to the fracture properties, but do show some relationship where comparisons are made for the same temper condition. Table 3 - Comparison of Alloys for Figures 3 to 5 - Tear Test (TF) and Tensile Data (T)

K) K)

Table 4 - Tear Test (FT) and Tensile Data (F) for Alloy X

K)

Lk)

Table 5 - Tear Test (FT) and Tensile Data (F) for Alloy X

K)

Because of the differences between the two alloys examined in Table 3, further investigations of the microstructure were made as in Figure 6.

Figure 6 shows microstructures at the centre of high pressure diecastings in the same as-cast condition, comparing the respective A380 alloy and alloy X compositions as used for Figures 3 to 5. The microstructures were taken from as close as possible to the same position in the respective castings, and show similar aluminium grain sizes. Both micrographs were taken at the same magnification. Figure 6(a) shows the microstructure of the A380 alloy casting and Figure 6b shows the microstructure of alloy X. Both alloys were produced identically by cold chamber high pressure diecasting with a metal velocity at the gate of 82 m/s. What is particularly surprising is that significant morphological changes to the Si phase forming from the solidified eutectic have occurred in the alloy X of the present invention. This is despite the fact that no known Si modifiers were present, added or detected in chemical analysis.

Figure 7 shows microstructures at the centre of high pressure diecastings in the as-cast condition, similar to the compositions represented in Figure 6 but now at higher magnification. The conventional A380 alloy is again compared with a similar alloy X according to the present invention. As for Figure 6, the micrographs were taken from as close as possible to the same location in the respective castings, and were taken at the same magnification. Figure 7 shows similar information to Figure 6, but the higher magnification emphasizes the differences in the conventional A380 microstructure compared to the microstructure of the alloy X.

Figure 8 shows microstructures of high pressure diecastings in the T6 condition, made from the same A380 alloy (Figure 8a) and alloy X (Figure 8b) that were used to obtain the data presented in Figure 3. Microstructures were taken from the same positions in each of the cast test pieces. In the T6 heat treated condition, of Figure 8b, it will be observed that the Si structure shown by Figure 7b has been changed to smaller Si particles in the microstructure of the T6 heat treated alloy. Additionally, the Fe-bearing phases are now more readily resolved. In the conventional alloy A380, Fe-bearing phases are predominantly needle like (some are arrowed). In contrast, the alloy of the present invention Fe-bearing phases are predominantly non-acicular and mostly small polyhedra, of a size similar to the spheroidized silicon particles. These fine polyhedra are arrowed in Figure 8 (b). Importantly, there are substantially no Fe bearing rosettes, plates or needles present in the microstructure of the alloy X for the present invention.

The change to the morphology of the Si plates in the as-cast condition as shown by Figures 6 and 7, and their differing sizes as shown in Figure 8 are similar to some known effects provided by additions of elements such as Sr, Na, P, Sb, Be or some rare earth elements, but in the current alloy the effect is observed without the presence of these elements. The combined effect of correcting the

Fe-bearing particles, transforming them from needle like to small polyhedra, as well as changing the morphology of the Si phases present, has been reported for some rare earth bearing casting alloys, but in the present invention no rare earth elements are present.

Figure 9 shows microstructures of castings prepared in a similar manner to those shown by Figure 7, but now comparing the same A380 alloy with a different alloy according to the present invention which has a composition of AI-8.8Si-0.56Fe-

3.3Cu-0.36Mn-0.19Mg-0.07Ni-1.4Zn-(<0.2 other), consistent with the upper limits of Fe content in the present invention and having SF=1.28. The micrographs were taken from as close to the same position as possible from the respective castings.

The microstructure with the alloy according to the present invention shown in Figure 9(b) is again changed, by the modifications to the alloy chemistry, relative to the A380 alloy shown in Figure 9(a). However, although the alloy of the present invention does not form Fe-bearing needles in the microstructure, some larger, less detrimental, blocky Fe-Mn bearing phases appear in the microstructure due to the higher Fe content. Since tensile ductility is not an optimal representation of fracture properties as indicated by the above examples, it is also desirable to develop alloys displaying both elevated levels of ductility and fracture toughness. Also, as indicated above, for the same temper condition, there is a rough correlation evident. In general, provided all conditions are constant and samples are the same, ductility of the alloy does allow for a rapid approximate determination of compositional ranges or limitations that may not be suited to high fracture resistant applications. In particular, limitations that arise due to alloy chemistry are important.

Figure 10 presents results for alloy X in which tensile properties are plotted against ageing times between 4 and 24h at 15O⁰C. In this case, cylindrical test bars of gauge length 25mm and diameter 5.55(±0.1 mm) were prepared from the same alloy as used to obtain the data shown in Figure 3. These bars were first solution treated at 48O⁰C for 15 minutes and cold water quenched before ageing was commenced. In the figure, open circles represent yield stress (0.2% proof stress), open diamonds represent tensile strength and solid triangles show ductility as measured by elongation at failure. Five tensile samples were tested for each comparative data point.

Figure 11 is similar to Figure 10, and is representative of the same alloy X of Figure 10, but now tensile samples were taken from the plate castings as shown in Figure 2, as part of testing for tear resistance. The markers for the experimental points are the same as for Figure 10. Five tensile samples were tested for each comparative data point.

Figure 12 plots the same 0.2% proof stress and tensile stress values as Figure 11 , as open circles and open diamonds respectively, plotted with values for the unit total energy (UTE) and unit propagation energy (UPE) (solid triangles and solid squares respectively). Comparison of Figures 10 to 12 suggests that for this alloy X, there is a rudimentary trend relating tensile ductility and fracture resistance. Five tensile samples and five tear test samples were tested for each comparative data point.

It can be seen from Figure 11 that for the same alloy composition, tensile samples machined from the plates display only a poor correlation to the fracture properties of the alloy, due to the elevated levels of scatter present in the data. However, the results of Figure 10 (cylindrical tensile bars) do seem to show a moderately better correlation to the fracture properties (Figure 12) for the same alloy composition.

Tables 6 to 8 show the chemistry of a range of alloys and their corresponding tensile properties in the as-cast, T4 and T6 conditions, respectively. For T4 conditions of Table 7 and T6 conditions of Table 8, solution treatment was conducted at 490⁰C for 15 minutes followed by cold water quenching. T4 conditions were prepared by holding for 14 days at 25°C, and the T6 temper involved ageing for 24h at 150⁰C. For alloy 3, in Tables 7 and 8, a solution treatment temperature of 480⁰C was also tested, as noted in the table.

Table 6 reveals little difference in tensile ductility of the different alloys examined. A comparison of Alloy 1 and Alloy 3 shows only a minor change in ductility, which is raised from 4.1 to 6.1 %, yet displays some reduction in 0.2% proof stress. Comparisons of Alloy 1 with Alloy 2, or Alloy 3 and Alloys 4, 5 and 6, shows that increasing the Mg content from -0.1 to ~0.3wt% tends to increase 0.2% proof strength and slightly decrease ductility in the as-cast condition. In relation to Alloys 1 and 3, these correspond to the two alloys compared in Table 3.

A comparison of alloy 3 with alloy 7, shows that increasing Cu content also tends to decrease ductility. As for Alloy 1 and Alloy 2, or Alloy 3 and alloy 4, comparison of alloy 7 with alloy 8 also shows a reduced ductility with change in Mg content. Table 6 - Tensile Properties for Alloy Compositions in As Cast Condition

K)

¹O

Table 7 - Tensile Properties for Alloy Compositions in T4 Condition

O

Table 8 - Tensile Properties for Alloy Compositions in T6 Condition

Table 7 shows that comparison of alloy 3 with alloy 1 shows a greater difference in ductility for the T4 condition compared to that present in Table 6. This is also consistent with the tear test results presented for the same two alloys in Table 3. Here the 0.2% proof stress and UTS is similar for both alloys. Similar trends to the as-cast results of Table 6 are noted in terms of Cu and Mg contents.

Table 8 shows that there is little difference in the respective ductility of alloy 3 and alloy 1 in the T6 condition. This result is consistent with that shown for the same two alloys in Table 7, and the tear test results presented in Table 3. Results shown in Table 8 confirm trends apparent in Table 6 in that increased Mg contents above 0.25% are preferably avoided, and that increased Cu content also produces some reduction in ductility.

From Tables 6 to 8, the upper limits on some elements may therefore be determined. For example, results from Alloy 6 suggest that Sr should preferably be omitted or minimized wherever possible. Mg contents are optimal when around 0.1 %. Cu contents are optimal when below 3.5%, and are better as Cu levels reduce to below 3%. To ensure some heat treatment response is possible, Cu levels can be kept at above 1 %. In the T4 and T6 treated conditions alloys 3 and 7 are preferred compared to the other alloys examined.

Although morphological alterations to the Si phase as shown in Figures 7 to 9 is expected for several of these alloy compositions included in Tables 6 to 8, it is apparent that it is not just changes to the Si and optimizing the Fe and Mn concentrations that produce optimum properties.

In relation to the results of Tables 6 to 8, it is important to note the relationship of the alloy chemistries in reference to the present invention. Alloy 1 has a sludge factor of 1.18, but is outside of the limits of the current invention with regards to the Fe and Mn contents. Alloy 2 is similar to Alloy 1 , having a sludge factor of 1.14, but is outside of the limits of the current invention in regards both to Fe and Mn contents and also the Mg content. Alloy 3 is inside the limits of the present invention, and has a sludge factor of 1.22. Alloy 4 has a sludge factor of 1.26 but is outside the limits of the current invention with regards to Mg content. Alloy 5 is similar to alloy 4, has a sludge factor of 1.22 and is outside the limits of the current invention with regard to Mg content. Alloy 6 is similar to alloy 5, has a sludge factor of 1.28 and is outside the limits of the current invention in regards to both the Mg and Sr contents. Alloy 7 is within the limits of the current invention and has a sludge factor of 1.6. Alloy 8 is outside the limits of the current invention in regards to Mg content and has a sludge factor of 1.46. Alloy 9 is similar to alloy 8, has a sludge factor of 1.23 and is outside of the limits of the current invention with regard to Mg content.

The preferred lower limits of Mg would ideally be 0.0% for alloys being heat treated, but some surprising results have also been found in this regard. Tables 9 to 11 show tensile results for a range of alloy compositions wherein the Mg, Cu, Zn and Ti contents were varied. All alloys in Tables 9 to 11 are within the scope of the current invention. For these alloys, it is surprising that for very low Mg contents of around 0.02%, no improvement to tensile properties are observed by heat treatment despite the presence of age hardening elements such as Cu. Additions of Zn appear to slightly decrease ductility, whereas additions of Ti appear to counteract this effect, slightly raising ductility. In the T4 tempers for all alloys, ductility is highest. The as-cast conditions for these alloys also display relatively high ductility. Although these alloys would be expected to exhibit excellent fracture resistance, the reduced strength would limit their potential applications.

Results shown for alloys 13 and 14 highlight the importance of optimizing the Mg content so as to gain a good hardening response, while maintaining adequate ductility. It is worth noting however, that even the alloys showing no age hardening response may find a range of applications where strength is less important than a high level of energy absorption. Table 9 - Alloy Tensile Properties in As Cast Condition with Variation in Mg, Cu, Zn and Ti

4*.

Table 10 - Alloy Tensile Properties in T4 Condition with Variation in Mg, Cu, Zn and Ti

LU

Table 11 - Alloy Tensile Properties in T6 Condition with Variation in Mg, Cu, Zn and Ti

Os

To examine this further, samples of a different casting shape (tensile fatigue test bars) having a composition of AI-7.2Si-0.23Fe-1.8Cu-0.49Mn-0.07Mg-0.43Zn- 0.05Ti-(<0.2 other total) were tested and were found to also display a good age hardening response by heat treatment, displaying elevated T6 tensile properties and high ductility.

Therefore, Mg contents for heat treated conditions of the alloy of the invention are ideally above 0.03% and less than 0.26%, and preferably in the range of 0.05 to 0.15 to ensure a good response of the alloy to heat treatment without causing any significant decrease in ductility.

Tables 12 to 14 show results for a range of 11 alloys in which the Mg content was raised progressively from 0.005 Mg to 0.22 in progressive increments. All alloys were within the scope of the present invention. As before, five tensile tests were conducted for each alloy in each condition examined. Table 12 shows results for the as-cast condition, Table 13 results for the T4 condition and Table 14 results for the T6 condition. Table 12 shows there is little variation in the tensile properties in the as-cast condition for the 11 alloys examined. Table 13 shows that for the T4 temper, the results are quite different. Alloys 15 and 16 do not show any advantage to tensile yield stress (0.2% proof stress) compared to the as-cast condition, but the tensile ductility is approximately doubled. Alloys 17 to 25 all show improvement to the yield stress (0.2% proof stress) compared to the as-cast condition as well as high levels of ductility, where the Mg content is above 0.05wt% Mg. Alloys 18 to 23 all display similar levels of ductility, and alloys 24 to 25 show lowered levels of ductility.

Table 14 shows that as for Table 13, there is effectively no response to heat treatment for alloys containing very low Mg levels (e.g. alloys 15 and 16.). Alloys 17 to 25 all display good levels of T6 treated tensile properties. Alloys 17 to 23 all display optimum tensile ductility while maintaining a good response to heat treatment. Table 12 - Alloy Tensile Properties in As Cast Condition with Principal Variation in Mg Content

LU

OO

Table 13 - Alloy Tensile Properties in T4 Condition with Principal Variation in Mg Content

LU

¹O

Table 14 - Alloy Tensile Properties in T6 Condition with Principal Variation in Mg Content

O

Additionally, comparison of as-cast alloys as shown in Tables 9 and 12 gives an indication of the preferred limits of Si content. For Table 9, ductility ranges from 7.4 to 9.8%, where the Si levels were between 7 and 8%. For Table 12, the Si content is higher, of the order of 10 to 10.5 wt%, and the ductility values are lower at 4.2 to 5.3%.

Table 15 shows compositions for a range of seven alloys (alloys 26 to 32) within the scope of the present invention. Alloys 26 to 32 were used for determination of tear and tensile properties, (shown in Tables 16 to 19). The major changes were to the Cu content, which ranges from 1.82%Cu to 3.12%Cu in progressive increments. Table 16 shows results for the as-cast condition. Tables 17 to 19 respectively, show results for the T4 condition, the T6 condition, and for an underaged T6 condition (UA) and, in each case, the alloys were aged 6h at 15O⁰C following solution treatment and quenching. All the T4, T6 and UA samples were solution treated at 48O⁰C before quenching. As before, five tensile test specimens and five tear test specimens were prepared for each alloy in each condition examined.

Tensile results shown in Table 16 for samples machined from plates and tested in the T direction for alloys 26 to 32 show that the 0.2% proof stress increases moderately with Cu content in the as-cast condition, while tensile elongation does not change appreciably across the compositions examined. The tear strength from plates tested in the TF direction also does not change appreciably across the seven alloys tested, although the Unit Propagation Energy does decrease as Cu content is raised. Similarly, the values of the Tear-to Yield-Stress ratio decrease from 1.73 for Alloys 26 and 27, down to 1.54 for Alloy 32.

Table 17 shows results as for Table 16, but for material treated to a T4 temper for samples machined from plates and tested in the T direction (tensile) or TF direction (tear) for alloys 26 to 32. In this case, the 0.2% proof stress increases more significantly with Cu content, from a lower value of 150 MPa for Alloy 26 up to a higher level of 196 MPa for Alloy 32. The tear strength increases from 305 MPa to 343 MPa with increasing Cu content for alloys 27 to 30, before then decreasing to a lower level of 317 MPa for Alloy 31. It is then little changed with a further addition of Cu, increasing only to 321 MPa for Alloy 32. The values for UPE change more significantly with the differences in Cu level, changing from 54.63 KJ/m² for Alloy 26, and down to 34.5 KJ/m² for Alloy 32. The tear-to-yield- ratio decreases from an upper level of 2.05 for Alloy 26 down to a lower level of 1.64 for Al by 32.

Table 18 shows results as for Table 16 and 17, but now for material treated to a T6 temper for samples machined from plates and tested in the T direction (tensile) or TF direction (tear) for alloys 26 to 32. Here, the 0.2% proof stress rises with increases in Cu content from 217 MPa for Alloy 26 up to 304 MPa for Alloy 32. The tear strength is above 300 MPa for Alloys 26 to 30, and then falls below 300 MPa for Alloys 31 and 32. The values of UPE decrease monotonically with increasing Cu content, from 29.45 KJ/m² for alloy 26 down to 9.72 KJ/m² for Alloy 32. Similarly, the tear-to-yield-ratio is again at its highest for Alloy 26 at 1.45, before then falling to below a value of 1 for Alloys 31 and 32.

Table 19 shows results as for Tables 16 to 18, but now for material treated to an underaged T6 temper for samples machined from plates and tested in the T direction (tensile) or TF direction (tear) for alloys 26 to 32. Here, tensile results are rather similar to the results for the T4 temper shown in Table 17 despite the difference in the heat treatment procedure, in that the T4 treated material is held at 25⁰C for 14 days, whereas the UA treated material is held 6h at 15O⁰C. Tear strength values are all above 300 MPa for the UA treated material. The values of UPE decrease monotonically with increases in Cu content, similar to the T4 treated alloy shown in Table 17. The tear-to-yield ratio is again highest for alloy 26 at 1.95, and decreases to a lower value of 1.49 for alloy 32. Table 15 - Alloy compositions examined for Tear Test (TF) and Tensile (T) properties with Principal Variation in Cu Content

Alloy Composition

26 (1 ) AI-7.4Si-0.24Fe-1.82Cu-0.45Mn-0.1 Mg-0.25Zn- <0.1 (other total)

27 (2) AI-7.2Si-0.23Fe-2.0Cu-0.45Mn-0.1 Mg-0.26Zn- <0.1 (other total)

28 (3) AI-7.4Si-0.24Fe-2.35Cu-0.45Mn-0.11 Mg-0.26Zn-<0.1 (other total)

29 (4) AI-7.4Si-0.24Fe-2.59Cu-0.45Mn-0.11 Mg-0.27Zn-<0.1 (other total)

30 (5) AI-7.5Si-0.24Fe-2.82Cu-0.46Mn-0.11 Mg-0.26Zn-<0.1 (other total)

31 (6) AI-7.2Si-0.24Fe-2.91 Cu-0.45Mn-0.1 Mg-0.26Zn- <0.1 (other total)

32 (7) AI-7.2Si-0.24Fe-3.21 Cu-0.44Mn-0.1 Mg-0.26Zn- <0.1 (other total)

Table 16 - Tear Test (TF) and Tensile Data (T) for As Cast Alloys 26-32 with principal variation in Cu Content

t

Table 17 - Tear Test (TF) and Tensile Data (T) for T4 Treated Alloys 26-32 with principal variation in Cu Content

Table 18 - Tear Test (TF) and Tensile Data (T) for T6 Treated Alloys 26-32 with principal variation in Cu Content

ON

Table 19 - Tear Test (TF) and Tensile Data (T) for UA Treated Alloys 26-32 with principal variation in Cu Content

Table 20 shows compositions for a range of additional different alloys (alloys 33 to 36) within the scope of the present invention, which were used for determination of tear and tensile properties.

Table 21 shows tear and tensile test results for alloy 33 and corresponds similarly to the composition used for tensile results shown for alloy 10 in Tables 9 to 11 , being made from the same base ingot material. Five tensile samples and five tear test samples were tested in each condition. Tensile testing was in the T direction and tear testing in the TF direction. As is the case for this composition as shown in Tables 9 to 11 , the alloy displays little or no heat treatment response in regards to its yield stress. However, heat treatment displays significant and substantial advantages to the fracture resistance of this alloy. The tear strength of this alloy is little changed with heat treatment, but the energy absorbed during fracture, particularly in the UPE, is greatly improved. Additionally, there is less difference between the UPE of the T4, T6 and UA tempers, which results because the alloy does not display a noticeable heat treatment response. The initiation energy of the heat treated conditions is also substantially better than the as-cast condition, because the material yields more before crack initiation, despite the presence of the sharp notch present for the tear test samples. The tear-to-yield ratio is very high for all conditions, but particularly so for the T4 and UA material where this value is over 2.5.

Table 22 shows tear and tensile test results for alloy 34, which is similar to alloy 26 from Table 15. Five tensile samples and five tear test samples were tested in each condition. Tensile testing was in the T direction and tear testing in the TF direction. In the T4 temper the alloy display similar levels of 0.2% proof stress and UTS as for the as-cast condition, but the elongation at failure is substantially increased from 4.9% to 9%. For the T6 temper, the 0.2% proof stress is raised to Table 20 - Alloy compositions examined for Tear Test (TF) and Tensile (T) properties with Principal Variation in Mg Content

Alloy Composition

33 (10) AI-7.1Si-0.21 Fe-1 6Cu-0.5Mn-0.01 Mg-(O. 1 other total) (alloy 10, table 9)

34 (A) AI-8Si-0.28Fe-1.84Cu-0.5Mn-0.12Mg-0.2Zn-(<0.1 other total)

35 (B) AI-7.8Si-0.29Fe-1 78Cu-0.52Mn-0.18Mg-0 2Zn-(<0.1 other total) 4*.

¹O

36 (C) AI-7.5Si-0.27Fe-1 74Cu-0.48Mn-0.23Mg-0 2Zn-(0.1 other total)

Table 21 - Tear Test (TF) and Tensile Data (T) for Alloy 33 in different tempers

LtX

O

Table 22 - Tear Test (TF) and Tensile Data (T) for Alloy 34 in different tempers

216 MPa, while the UTS changes only marginally to 309 MPa. The elongation at failure is reduced compared to the T4 temper at 4.5%. Values of tear strength for alloy 34 are relatively high in all tempers examined, being consistent with the data for alloy 26 in Tables 16 to 18. The tear strength is however superior for alloy 34 in the T6 temper, compared to alloy 26 in the T6 temper. The tear-to- yield ratio is also similar between alloys 26 and 34 for the respective equivalent tempers. Values of UPE are also superior in all tempers examined for alloy 34 compared to alloy 26.

Table 23 shows tear and tensile test results for alloy 35. Alloy 35 is similar to alloy 34 except that the level of Mg was slightly increased, and the level of Cu slightly decreased. Five tensile samples and five tear test samples were tested in each condition. Similar to alloy 34, there is only a small change in tensile 0.2% proof stress and UTS for alloy 35 in the T4 temper when compared to the as-cast condition. The elongation at failure is however improved compared to both of the other conditions examined. Tensile testing was in the T direction and tear testing in the TF direction. The heat treatment response in the T6 temper for alloy 35 is again good. There is little difference between the tear strength of alloys 34 and

35 in their respective tempers, however in the T4 and T6 tempers the tear-to- yield ratio is reduced for alloy 35. Similarly, the values of UPE are also slightly reduced for alloy 35 compared to alloy 34. The overall fracture resistance of alloy 35 is slightly less than that of alloy 34.

Table 24 shows tear and tensile test results for alloy 36. Alloy 36 is similar to alloy 35 except that the level of Mg was again slightly increased, and the level of

Cu again slightly decreased. Five tensile samples and five tear test samples were tested in each condition. The levels of tensile properties developed in alloy

36 are again similar to those for alloys 34 and 35 in as-cast and T4 tempers, although there is a slight increase across the three alloys as the levels Table 23 - Tear Test (TF) and Tensile Data (T) for Alloy 35 in different tempers

Table 24 - Tear Test (TF) and Tensile Data (T) for Alloy 36 in different tempers

of Mg are raised. Similarly, the levels of tensile 0.2% proof stress are raised slightly across the three alloy compositions, in line with the increasing Mg content. Levels of TYR are again slightly decreased at this higher Mg content. The major difference for alloy 36 appears in the UPE values for the T6 temper, where the energy absorbed during crack propagation is more significantly reduced, despite the similar levels of tensile properties. This would tend to suggest that at a level of 0.23%Mg, that the fracture resistance is approaching its upper preferred limit.

Figure 13 compares the values of yield stress (0.2% proof stress) and unit propagation energy for four conventional HPDC alloys with the new alloys of the current invention in as-cast and heat treated (T4 or T6) conditions. The data for the new alloys of the current invention correspond to the values for alloys 26 to 36. The conventional HPDC alloys are Al-Si-Cu alloys A380 (US specification), C380 (US specification), ADC12 (JIS specification) and AK9M2 (CIS specification). As cast, the alloys display yield stress (0.2% proof stress) and unit propagation energies shown by the line A. Heat treated to T4 or T6 tempers, the four conventional alloys display yield stress (0.2% proof stress) and unit propagation energies shown by line B. For the range of newly developed alloys as cast, the yield stress (0.2% proof stress) and unit propagation energy is shown by line C. For the range of newly developed alloys heat treated to a T4 or T6 temper, the yield stress (0.2% proof stress) and unit propagation energy is shown by line D. For given levels of 0.2% proof stress, the energy absorbed during crack propagation is substantially higher for the alloys of the present invention than for the conventional alloys. The differences are however minimized at the upper end of yield strength (0.2% proof stress) for the alloys of the present invention, which also corresponds to the T6 condition for Alloy 32, containing 3.2%Cu, suggesting there is little benefit to be gained in a T6 temper above this Cu concentration. In a T4 temper however, there is still benefit to be gained to the energy absorption during fracture at this level of Cu, above comparative conventional alloys. In the examples above, upper and lower limits on key alloying elements are detailed for the alloys of the invention.

Finally, it is to be understood that various alterations, modifications and/or additions may be introduced into the constructions and arrangements of parts previously described without departing from the spirit or ambit of the invention.

Claims

1. An aluminium based alloy according to the invention has a weight percentage composition of:

Silicon 5 to 15% magnesium 0 to 0.25% titanium 0 to 0.25% manganese 0.2 to 0.85% iron 0.1 to 0.8% copper 1 to 4% zinc 0 to 3% silicon modifiers less than 0.01 % in total (with less than

0.05% chromium) other elements less than 0.5% in total,

and a balance of aluminium,

wherein the limits for iron and manganese are constrained such that the amount of iron present in the alloy is 0.4 to 1.6 times the manganese content and the alloy has a sludge factor (SF), calculated as SF = (1 x wt% Fe)+(2 x wt% Mn), of from 0.8 to 1.8.

2. The alloy of claim 1 , wherein the chromium level is less than 0.02%

3. The alloy of claim 1 or claim 2. wherein the composition is free of beryllium, rare earth elements, and free of chromium and other transition metal elements not including Ti, Mn, Fe, Cu and Zn.

4. The alloy of any one of claims 1 to 3, wherein the silicon ievei is from 8.5 to 10.5% such as from 6.5 to 8.5%.

5. The alloy of any one of claims 1 to 4, wherein copper is present at from 1 to 3.5% such as from 1.5 to 3%.

6. The alloy of any one of claims 1 to 5, wherein SF=LO to 1.3.

7. The alloy of any one of claims 1 to 6, wherein the iron content is from 0.5 to 0.9 times the manganese content by weight.

8. The alloy of any one of claims 1 to 7. wherein the iron content is from 0.2 to 0.4% and the Mn content is from 0.3 to 0.6%.

9. The alloy of any one of claims 1 to 8. wherein zinc is present at from 0.3 to 1.0%.

10. The alloy of any one of claims 1 to 9, wherein magnesium is present at from 0.05 to 0.15%.

11. A casting having enhanced fracture resistance relative to casting of the same product made of a conventional HPDC alloy when compared in the as cast or same heat treated state, wherein the casting having enhanced fracture resistance is cast from an aluminium based alloy having a weight percentage composition of:

silicon - 5 to 15% magnesium - 0 to 0.25% titanium - 0 to 0.25% manganese - 0.2 to 0.85% iron - 0,1 to 0.8% copper - 1 to 4% zinc - 0 to 3% silicon modifiers - less than 0.01 % in total (with less than 0.007% strontium) tin - less than 0.05% other transition or rare earth metals - less than 0.2% in total (with less than

0.05% chromium) other elements - less than 0.5% in total,

and a balance of aluminium,

wherein the limits for iron and manganese are constrained such [hat the amount of iron present in the alloy is 0.4 to 1.6 times the manganese content and the alloy has a sludge factor (SF), calculated as SF = (1 x wt% Fe)+(2 x wt% Mn), of from 0.8 to 1.6; and

wherein the casting having enhanced fracture toughness has a microstructure exhibiting silicon formed from solidified eutectic which also contains iron-bearing phases which are substantially fine polyhedral particles.

12. The casting having enhanced fracture resistance according to claim 11 , wherein the casting is cast from the alloy of any one of claims 2 to 10.

13. The casting having enhanced fracture resistance according to claim 11 , wherein the casting is in the as cast condition.

14. The casting having enhanced fracture resistance according to claim 11 , wherein the casting has its fracture resistance and/or microstructure enhanced relative to the as cast condition by heat treatment.