the role of alloy composition in the heat treatment of aluminium high

2 Metallurgical Science and Technology Vol. 26-2 - Ed. 2008

The role of alloy composition in the heat treatment of aluminium high pressure die castings

RR.N. Lumley, R.G. O’Donnell, D.R. Gunasegaram, T. Kittel-Sherri, M. Gershenzon, A.C Yob, I.J. PolmearLight Metals Flagship, CSIRO Materials Science and Engineering, Private Bag 33, Clayton South MDC, Australia

ABSTRACT

High pressure die-cast (HPDC) aluminiumcomponents that respond to age hardeningcannot normally be solution treated at hightemperatures because the presence of internalporosity and entrapped gases leads to theformation of surface blisters. Parts may alsobecome dimensionally unstable due toswelling. These factors that prevent heattreatment present significant limitations to theutilisation of HPDC components. Now it hasbeen found that blistering and dimensionalchange can be avoided by using modifiedshorter solution treatment procedures whichstill allow strong responses to age hardeningto be achieved with a wide range of Al-Si-(Cu/Mg) alloys. In the present paper, the rolesof critical alloying elements are considered inboth current commercial and experimentalalloy compositions in this series. It is shownthat values of 0.2% proof stress exceeding400 MPa may be readily achieved by heattreating conventionally produced HPDCcomponents.

RIASSUNTO

I componenti in lega di alluminio ottenuti per pressocolata da rafforzare tramite invecchiamento non possono subirenormalmente un trattamento disolubilizzazione ad alta temperatura inquanto la presenza di porosità interne o digas intrappolati conduce alla formazione dirigonfiamenti superficiali. Alcune partipossono inoltre diventare dimensionalmenteinstabili a causa di dilatazioni. Questi fattoriche contrastano il trattamento termicocomportano una limitazione importantenell’utilizzo di componenti pressocolati. Si èora riscontrato che rigonfiamenti e variazionidimensionali possono essere evitati riducendola durata del trattamento di solubilizzazione econsentendo così a numerose leghe Al-Si(Cu/Mg) di essere notevolmente rafforzateper invecchiamento. Nel presente lavoro èstato preso in considerazione il ruolo deglielementi leganti presenti sia nelle usuali leghecommerciali sia in leghe sperimentali. È statoevidenziato che trattamenti termici di

KEYWORDS

Heat Treatment, Precipitation, Aluminium,HPDC

INTRODUCTIONAs shown in Fig. 1, high pre s s u re diecastings (HPDC’s) normally retain a highgas content which arises from the presenceof entrapped air, hydrogen and vaporizeddie lubricant. Because of this, HPDCc o mponents cannot be given ac o nventional solution treatment (e.g. at520°C for 8h) as the pores containinggaseous elements may expand resulting inthe unacceptable surface blisters, distortionand lower mechanical pro p e rties. Thedimensions of the die cast parts may alsochange due to swelling with a consequentadverse effect on mechanical properties.Recent work within the CSIRO Light MetalsFlagship [1-4] has revealed a heattreatment cycle for HPDC aluminium alloysthat avoids these problems. The modified

GAS CONTENT IN ALUMINIUM CASTINGS, NORMALISED

0.01 0.1 1.0 10 100 cm3/100g

VACUUM DIECASTINGS

SAND AND PERMANENT MOULD CASTINGS

LOW PRESSURE DIECASTINGS

SQUEEZE CASTINGS AND SEMISOLID

DIECASTINGS

Fig. 1: Comparative gas contents arising from different casting processes [5].

componenti ottenuti con pressocolatetradizionali possono dare facilmente luogo acarichi di snervamento (0,2%) superiori a400 MPa.

3Metallurgical Science and TechnologyVol. 26-2 - Ed. 2008

cycle involves a severely truncated solutiont reatment sta ge at lower than norm a ltemperatures to avoid blistering, which,due in part to the unique microstructuregenerated by the high pressure die-castingprocess, is sufficient to attain at least apartial solid solution of the potent ageh a rdening elements Cu, Mg and Si.Comparisons between the surface finishand inte rnal micro st ru c t u res of highpressure die-cast test components in the as-c a st, conve n t i o n a l ly solution tre a ted, orsolution tre a ted according to the newprocedures are shown in Fig. 2. Followingsolution treatment and qu e n ching, th eHPDC alloys may be naturally aged to a T4te mper at ambient te mp e ra t u re, ora rt i fi c i a l ly aged at an inte rm e d i a tetemperature (e.g. 150-180°C). As a result,large increases in tensile properties may berealized and a summary of hardness andstrength values of heat treated HPDC’scompared with those for the as-cast alloysis shown in Fig. 3. A T4 temper normallyraises the 0.2% proof stress by an averageof 30% above the as-cast condition, oftenwith elevated levels of ductility. Ageing to aT6 te mper allows peak st re n g th to beachieved and values of 0.2% proof stressmay be raised by 80 to 100% or more.Wi th respect to their potential fo rstrengthening by heat treatment, normal Al-Si-(Cu/Mg) HPDC alloys contain amountsof the critical elements Si, Mg and Cu thatcontribute to age hardening. Although upto 3% Zn is permissible in some HPDC alloyva riants, th e re is often insufficient Mgpresent to form substantial quantities ofMgZn2 pre c i p i ta tes that could alsoc o n t ri b u te to st re n g thening. The curre n tpaper is concerned with the roles of thevarious alloying elements in contributing toage hardening in a wide range of HPDCalloys.

Fig. 2: Comparisons of surfaces and internal microstructures for a HPDC alloy in the as-cast, conventionallyheat treated and heat treated according to the modified heat treatment conditions.

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AC: 0.2% P.S; y = 1.9x - 18.8

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HT: 0.2% P.S; y = 2.7x - 64.2

HT: UTS y = 2.1x + 96.3

0.2% ProofStress

Fig. 3: Relationships between hardness and tensile properties in as-cast (solid symbols) or heat treated(open symbols) HPDC’s. Each data point represents 5 or more tensile tests. All heat treated samples wereartificially aged to peak strength and hardness at different temperatures. Ranges for 0.2% proof stress andtensile strength are shown as dotted lines. Equations approximating the line of best fit are also provided.

EXPERIMENTAL METHODSHPDC alloy specimens for tensile testingwere produced using a Toshiba horizontalcold chamber die-casting machine with a250 tonne locking force, a shot sleeve withan internal diameter of 50mm and a strokeof 280mm. The die provided two cylindricaltensile specimens and one flat te n s i l especimen from each shot, and eachconformed to specification AS1391. The

cylindrical tensile test bars used for thecurrent work had a total length of 100 mm,with a central parallel gauge section 33mm long and a diameter of 5.55±0.1mm.The runner and samples produced areshown in Fig. 4. Five tensile samples werete sted in each condition and ten alloyc o mpositions we re examined which arelisted in Table 1. Levels of Cu, Mg, and Zn

we re va ried within the specifi e dcomposition ranges to assess their effectson mechanical pro p e rties of the heattreated alloys.Tensile properties were determined for allof the alloys in each of the three conditionsexamined (as-cast, T4 and T6 tempers).Solution treatment for all heat treatmentswas standardized at 490°C for 15 minutes


Table 1. Compositions of ten alloys examined for heat treatment

Alloy Al Si Cu Mg Zn Fe Mn other

1 balance 9 3.1 0.1 0.53 0.86 0.16 <0.2

2 balance 9.1 3.18 0.29 0.6 0.86 0.14 <0.2

3 balance 8.6 3.6 0.1 0.53 0.93 0.18 <0.2

4 balance 8.6 3.6 0.3 0.53 1.0 0.2 <0.2

5 balance 8.5 4.9 0.1 0.51 0.97 0.2 <0.2

6 balance 8.7 4.9 0.27 0.51 1.0 0.21 <0.2

7 balance 8.8 4.0 0.7 0.56 1.1 0.19 <0.2

8 balance 9.1 2.04 0.26 0.53 0.79 0.36 <0.2

9 balance 9.2 3.11 0.09 2.9 0.9 0.16 <0.2

10 balance 9.1 4.2 0.22 1.2 1.3 0.2 <0.25

Table 2. Mechanical properties of 10 alloys examined tested in the as-cast, T4 or T6 temper

Alloy 1 2 3 4 5 6 7 8 9 10

As Cast 172, 189, 176, 200, 193, 206, 237, 157, 165, 198,

354, 358, 358, 362, 363, 369, 377, 311, 347, 351,

4.1% 3.2% 3.7% 3.0% 3.3% 2.9% 2.2% 3.8% 3.8% 2.5%

T4 217, 246, 234, 258, 244, 275, 261, 192, 224, 267,

387, 411, 397, 411, 412, 418, 413, 306, 405, 405,

5.8% 5.2% 4.8% 4.3% 5.5% 3.8% 3.9% 3.3% 7% 3.4%

T6 352, 374, 379, 419, 381, 440, 413, 291, 370, 437,

441, 458, 457, 474, 446, 490, 474, 368, 442, 484,

3.4% 2.4% 2.7% 2.2% 1.9% 1.6% 1.6% 2.6% 3% 1.6%Results are shown as 0.2% proof stress (MPa), tensile strength (MPa) and elongation at failure (%).

total immersion time in a circulating airfurnace, followed by water quenching. Thisprocedure has been shown elsewhere toavoid blistering and still produce a highresponse to age hardening in these testspecimens [1-4]. Samples for the T6tempers were subsequently aged in oil at150°C for 24h whereas those for T4tempers were aged at 25°C in air for astandard time of 14 days.

Fig. 4: The runner and test specimens produced forthe current evaluation.

Transmission electron micro s c o py was conducted on selected alloys to observeprecipitation occurring within the aluminium grains. TEM discs were prepared by standardpolishing techniques, and examined in a Jeol 2000EX microscope operating at 200 kV.

RESULTS AND DISCUSSIONTENSILE PROPERTIESS

The tensile pro p e rties of the 10 alloysexamined in the as-cast, as solution treated,T4 and T6 tempers are presented in Table2, with re fe rence to numbere dc o mpositions listed in Table 1. As-castvalues of 0.2% proof stress ranged from157 MPa to 237 MPa. Typically, ageing tothe T4 temper raised the 0.2% proof stressby about 30% above the as-cast condition,often with elevated levels of ductility. T40.2% proof stress values generally rangedbetween 190 MPa for the low Cu Alloy 8,up to 275 MPa for the high-Cu Alloy 6.Ageing to a T6 temper increased the as-cast values of 0.2% proof stress by 75 to120%, with values that ranged from 291MPa for the low Cu Alloy 8 up to 440 MPafor the high-Cu Alloy 6. Raising Mg levelsto higher amounts (e.g. 0.7% Mg in Alloy7) proved to be disadvantageous whencompared to similar alloys with lower Mgcontents (e.g. 0.22% Mg in Alloy 10).

ROLE OF ALLOYING ELEMENTS

SSiliconSilicon is present in casting alloys largely tofacilitate high fluidity in the molten state, tofo rm a eutectic, and to pro m ote th efo rmation of st re n g thening pre c i p i ta te s

Fig. 5: Microstructures of as-cast HPDC Alloy 1 (a) at the sample edge and (b) at the sample centre. Lightgrey needles and blocky phases are Fe-based intermetallics.


through the expected reaction of Si and Mgpresent in solid solution. Silicon also formsintermetallic sludge particles with Fe andMn in the melt (e.g. ‚-AlFeSi5). As the Al-Sib i n a ry system displays an abnorm a le u tectic, solidification invo lves th eformation of Si particles in the solidifiedmelt. For HPDC alloys, the Si plate sizevaries greatly between the edge and centreof a part, because the rapid cooling rate atthe surfaces of components produces anextremely fine aluminium grain size and Siparticle size, such as shown in Fig. 5(a).Towards the centre of a component, the Algrain size is larger and the Si particles arepresent in fewer numbers and larger sizes(Fig. 5(b)).D u ring heating of the HPDC alloy to490°C, the Si present in the solidifi e d

Fig. 6: The refinement of silicon particles during heating to 490°C. Note surface blisters form at times of?20 minutes. See also Fig. 7 for quantification.

Fig. 7: Quantification of the changes to silicon particle size (diamonds) and number (asterisks), for (a) edge and (b) centre regions of the die-castings. Total areaexamined for each data point is ~122000Ìm2. The silicon particles fragment and spheroidize, then undergo Ostwald ripening, decreasing in number.

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e u tectic ra p i d ly undergoes signifi c a n tmorphological changes; the Si particlesf ragment and then qu i ck ly sphero i d i z e .Examples of this are shown in Figs. 6 and 7.Fig. 6 shows the microstructural changes tothe silicon particles at the edge of samples,and Fig. 7 quantifies these changes forboth edge and centre regions. It may benoted that most of the changes to the Sip a rticles occurred during an immers i o ntime of less than 20 minutes, whichc o rresponds closely to the appro p ri a teheat treatment time used to avoid blisteringon the surfaces of samples. Fig. 8 shows themicrostructure of a sub-surface region that

Fig. 8: A laminar defect leading to blistering in an A380 alloy, solution treated 20 minutes at 490°C. Asimilar, unexpanded laminar defect is also shown arrowed and at higher magnification. Note thedistribution of pinning silicon particles (arrowed at right) along the defect length.

A) B)


c o n tains a blister under which th e rea p p e a rs to be a thin laminar defe c td e c o ra ted with small Si part i c l e s(arrowed). It may be that these particleshave so far served to pin this defect therebyp reventing its expansion into an actualblister. Such features have been observedto be discontinuous defects displaying onlypartial metallurgical bonding across theirinterfaces. These appear to be responsiblefor a significant proportion of blisters thatare generated when samples are held forlonger times or at higher temperatures. Asimple explanation is that small pockets ofgas entrapped at high pressure in theselaminar defects expand to fo rm smallcracks. These cracks can then extend andcoalesce with similar cra cks fo rming inother gas pockets so that they producelarge surface blisters like the one shown atan early stage of formation in Fig. 8.The role of Si on precipitation processesduring the age hardening of many HPDC’sis qu i te comp l ex, as it is invo lved inprecipitation associated with the complexQ? phase, Al5Cu2Mg8Si6 or the precursorL phase and will be discussed later.

CCopperThe presence of Cu in HPDC Al-Si-Cu alloysmay also assist fluidity by depressing thesolidus temperature of the alloy and die-c a sting ex p e riments have revealed th a tAlloys 5 and 6 are particularly well suitedto production of thin walled cast i n g s .Copper also allows as-cast HPDC’s todevelop reasonable strengths due to thel i m i ted pre c i p i tation that occurs duri n gcooling from the casting te mp e ra t u re .D u ring solution treatment, Cu dissolve srapidly into the aluminium matrix despitethe short duration employed (Fig. 9) andthis element is critical in facilitating agehardening, particularly when Mg and Siare also present.The strengthening that Cu additions provideto the Al matrix is caused by precipitationof solute clusters and fine Guinier-Preston(GP) zones in the T4 condition, or as otherphases such as � ( A l2Cu) or Q’( A l5C u2M g8S i6) that fo rm in the T6condition (e.g. Fig. 10). In Fig. 10 it isevident that the fine dispersion ofp re c i p i ta te plates present within th ealuminium grains is similar to that observedin binary Al-Cu wrought alloys. There is,however, a large difference in the strength

Fig. 9: (a) As cast and (b) T6 treated HPDC Alloy 1. Despite the sort duration of the solution treatment step,the Cu-containing phases are largely dissolved leaving remnant Fe and Mn-bearing phases.

Fig. 10: TEM micrographs of HPDC A380 Alloy 1 in the T6 temper. (a) shows the � -aluminium grainsbounded by the silicon phase and (b) finely dispersed �‘ precipitates within the �-aluminium grains in an[001]� orientation.

that develops when binary wrought andHPDC alloys containing similar amounts ofCu are age hardened (Table 3). In eachcase, the alloys were aged to peak strengthat 150°C (T6 temper) and the 0.2% proofstress of the more complex HPDC alloy is143 MPa higher. This difference can beattributed mainly to the presence of thespheroidized particles of Si that providesubstantial dispersion hardening.The tensile properties of Alloys 2,5,7&9,which contain similar levels of Mg andvarying levels of Cu, were determined forthe as-cast, T4 and T6 tempers and theresults are summarized in Fig. 11. Here itmay be seen that increasing Cu contentimproves the strength properties of the die-castings in all three conditions. For the as-cast alloys, the 0.2% proof stress increasedlinearly between 2 wt% and 3.6 wt%, withno further imp rovements being ev i d e n twhen the Cu was raised further to 4.9 wt%.The tensile st re n g th of these alloys

i n c reased pro p o rt i o n a te ly to the 0.2%proof stress. Ductility was little changedwith only a slight decrease occurring as theCu content was raised from 2 to 4.9%. Forthe T4 te mp e r, the 0.2% proof st re s sincreased significantly as the Cu contentwas raised from 2 to 3.2wt% and then at aslower rate for higher levels of Cu. In allalloys except Alloy 8, ductility in the T4temper was in general higher than those forthe as-cast condition. Another observationis that, for Alloy 6 (Cu 4.9%) the 0.2%proof stress for the T4 temper was 275MPa which is similar to that for the Al-Si-Mg(permanent mold or sand cast) alloy A357aged to a T6 temper [e.g. 6].Alloys aged to a T6 temper showed thehighest strength levels. The 0.2% proofstress for Alloy 8 (2%Cu) was 291 MPaand rose linearly with Cu content to 419MPa for Alloy 4 (3.6%). Further smallincreases occurred as Cu level was raisedto 4.9%.


Levels of tensile strength followed the samege n e ral trends. As with the as-castcondition, ductility for alloys aged to a T6te mper decreased gra d u a l ly as the Cucontent was raised.In summary, the differences in levels of0.2% proof st ress bet ween alloyscontaining 2% and 3.6% Cu was 43 MPafor the as-cast condition (increase of 27%),66 MPa for the T4 temper (34% increase)

Table 3. A comparison of the properties developed from either binaryAl-Cu alloy or the more complex HPDC alloy, displaying similar Al:Cu ratio

Atomic TensileProduct Alloy Alloy Ratio 0.2% Proof strength Elongation

form (wt%) (at%) Al:Cu stress (MPa) (MPa) (%)

Wrought T6 Al-4Cu Al-1.74Cu 56.5:1 236 325 5%

HPDC T6 Al-8.6Si-3.6Cu- Al-8.54Si-1.58Cu- 56.3:1 379 457 2.7%(Alloy 3) 0.93Fe-0.1Mg- 0.46Fe-0.11Mg-

0.53Zn-0.18Mn 0.23Zn-0.09Mn

Fig. 11: Effect of increasing Cu content on mechanical properties of Al-Si-Cu HPDC’s containing ˜0.3Mg content and tested in (a) as-cast, (b) T4 or (c) T6conditions. Squares show 0.2% proof stress, diamonds tensile strength and triangles elongation at failure. One standard deviation is shown on each data point.See text for details.

1Copper Content (wt%)

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and 128 MPa for the T6 temper (44%i n c rease). Ageing most alloys to a T6temper more than doubled their as-castlevels of 0.2% proof stress.

MMagnesiumWhen present in combination with Cuand/or Si, Mg is also a very efficienta l l oying addition for st re n g th e n i n galuminium alloys, although in 380 typeHPDC alloys it is often restricted to less than0.3%, and even below 0.1% for somevariants. In the related alloy AlSi9Cu3(Fe)(European designation), up to 0.55% Mgmay be present whereas in alloy SC84R(Canadian designation), up to 0.75% ispermissible. Even small amounts of Mg canhave a profound effect on age hardening.The hardness-time curves for Alloys 1 and 2aged at 150°C following solution treatmentat 490°C for 15 minutes are shown in Fig.12. Each has a similar shape but Alloy 2,which has the higher Mg content of 0.3%,shows rapid ear ly hardening at ageingtimes up to 1hr. In general, the greatestbenefit to tensile properties arises when thelevels of Cu and Mg are raised together(see Tables 1&2) although increases in Mgabove 0.3% appear to have little effect ontensile properties (at least when Cu level ishigh). In fact, comparison of results forAlloys 7 & 10 suggests that higher Mgc o n tents may be det ri m e n tal to th edevelopment of improved tensile propertiesbecause increasing the Mg content from0.22 to 0.7% has led to a decrease of5.5% in 0.2% proof stress.The role of Mg on precipitation processes isdiscussed further below, in relation to thecombined effects of Si, Cu and Mg.

Combined Roles of Silicon, Copper andMagnesium In relation to the precise roles of Cu, Mgand Si during strengthening of HPDC’s bya ge hardening, it is also imp o rtant toconsider the solute contents of th ealuminium grains separately from the bulka l l oy composition. Fig. 13 shows th equaternary Al-Si-Cu-Mg phase diagram at460°C for a Si content of 1.2% [7]. If it isassumed that all Cu, ~1.2Si, and all Mg,are present in the solid solution then it ispossible to predict the phases that mayform during heat treatment of HPDC’s. Thevast majority of HPDC alloy compositions

fall within regions of the diagram in whichit can be predicted that the complex Qphase, or its precursors Q’ or L will forminstead of �’ , �” or S which are betterknown as strengthening precipitate phases.For example, Fig. 14 shows examples ofAlloys 3, 4 and 8, heat treated to a peaka ged condition at 150°C. Each of th ea l l oys displays a diffe rent pre c i p i ta testructure despite the only minor differences

Fig. 12: The effect of Mg on the age hardening response of Al-Si-Cu-(Fe) alloys. Alloys were solution treatedfor 15 minutes at 490°C, quenched, then aged at 150°C. The major difference between the ageing curvesoccurs within the first hour, where rapid early hardening occurs for the alloy containing 0.3% Mg.

1

Alloy 2 (0.3 Mg)

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Time (Hours)

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Fig. 13: The quaternary Al-Si-Cu-Mg phase diagram (isothermal section at 460°C) for a constant 1.2wt%Si[7].

1

Magnesium Weight %

20 3 4 5 6 7

7

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0

�+Si+�

� Al�: Mg2Si�: Al2CuQ: Al5Cu2Mg8Si6S: Al2CuMg

�+�+S

Isotherm 460°C1.2Si (wt%)

�+�

�+�+Si+Q�+Si+�

in composition. For the alloy containingonly 0.1%Mg, (Alloy 3) the majority of thep re c i p i ta tes present within th emicrostructure can be identified as �’. Aspredicted by the phase diagram, for thesimilar alloy containing 0.3Mg (Alloy 4),most of the precipitates present within themicrostructure are still the �‘ phase, butn ow a substantial amount of fi n eprecipitates of the Q’ phase, or it’s pre-


c u rsor L, can be dete c ted in th eb a ckground (viewed as fine, ofte nrectangular dots). Alloy 8, which has lessCu present, forms fewer �’ precipitatesto gether with the ve ry fine, densedispersion of the L or Q’ phases throughoutthe matrix. As expected from Fig. 13, themode of decomposition of th esupersaturated solid solution varies withalloy composition. For example, where Cucontent is reduced further and Mg contenti n c reased, the alloys are ex p e c ted toprecipitate a combination of the Q’ phaseand �” phase (including its GP zone andcluster precursors). As Cu is reduced stillfurther, the relative proportions continue toch a n ge, and the dominant pre c i p i ta tephase becomes �”. The most impor tantobservation, however, is that the Q’ phaseor its pre c u rsor L are the dominantp re c i p i tating phases in most HPDCcompositions. This also helps to explain thep ri m a ry diffe rence in (early sta ge )hardening between Alloy 1 and 2, (Fig.12). This follows because, whereas �’ is thedominant pre c i p i tating phase in th ealuminium grains of Alloy 1, Alloy 2 haspresent a combination of L/Q’ laths thatprecipitate first, together with �’ that formslater in the ageing cycle.

ZZincZinc is normally added to HPDC alloys toi mp rove casta b i l i t y, mach i n a b i l i t y, andcorrosion resistance [8] and up to 3% ispermitted in some compositions. WhereasZn and Mg together can cause significantage hardening through the precipitation of

Fig. 14: Peak aged T6 precipitation within the aluminium grains. (a) is of the low- Mg Alloy 3, (b) is of the higher-Mg Alloy 4, and (c) is of the higher-Mg, lower-Cu Alloy 8. TEM images [001]�.

MgZn2 in many aluminium alloys (e.g.7000 series wrought alloys), the levels ofMg in the present alloys are generally lowand the effect on tensile properties is verysmall. The increase in tensile properties thatis facilitated by increasing Zn content in lowMg alloys is therefore more likely to arisef rom solution st re n g thening ra ther th a nf rom pre c i p i tation. High additions (i.e.

~3%) of Zn do appear to incre a s esoftening during overageing (Fig. 15).

TTinTin may be present in HPDC alloys due toresidual contamination arising from the use

Fig. 15: Comparison of hardness-time curves for Alloy 1 (standard) and Alloy 9 (high zinc) showing theeffect of raised Zn content on overageing behavior. The alloy containing 2.9 wt% Zn overages more duringprolonged exposure at 150°C.

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Alloy 1 (0.5 Zn)

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Alloy 9 (2.9 Mg)

of secondary metal. In general, Sn mayi n c rease fluidity during casting andi mp rove mach i n a b i l i t y, so th a ts p e c i fications for HPDC comp o s i t i o n sworldwide include varying allowances forSn, usually up to 0.35%. Sn however mayd i s p l ay delete rious effects during die-c a sting and it is recommended that itscontent is kept to below 0.05%. This followsbecause additions as low as 0.08% havebeen observed to cause severe flashing,die st i cking and hot te a ring. Thesec o mplications may result in re d u c e dp ro d u c t i v i t y, gre a ter die we a r, anincreased reject rate and variable partquality in the as-cast state.


Corrosion may also be adversely affectedwhen Sn is present [8]. Sn may also act asan Fe modifier, producing large rosetteshaped Fe bearing sludge particles (Fig.16), which act as preferred sites for crackpropagation. It may be noted, however,that trace additions of Sn (e.g. 0.05%) areknown to stimulate age hardening in Al-Cualloys [e.g. 9]. It was decided therefore tomake a preliminary study of the effects onheat treatment of adding a small amount(0.08%) of Sn to one HPDC alloy (similarto Alloy 1).Tensile properties were determined aftereither a standard T4 treatment at 25°C, ora T6 ageing treatment at 15 0 ° C .Generally, addition of Sn in binary Al-Cualloys reduces the T4 ageing response andi n c reases the response to T6 age i n g .However, as shown in Table 4, neithereffect was observed in the HPDC alloy.Therefore, since Sn was found to provideno benefit in stimulating an incre a s e dresponse to age hardening, it was decidedthat levels should be kept as low as possiblefor the reasons explained above.Transition Metal Elements; (Fe, Mn, Cr )Iron and manganese are present in HPDCalloys to minimize die soldering so that diel i fe is increased and productivity isimproved. In general, the Fe content needsto be above 0.5% which is fo rt u n a tebecause >95% of HPDC alloys are made ofsecondary metal in which Fe contents maybe relatively high. Fe, Mn and Cr tend toform intermetallics that will be present ind i ffe ring amounts and morp h o l o g i e s ,depending on the so-called “sludge factor”of the alloy [10]. A proportion of sludgeparticles form in the melt, some in the shotsleeve, and some during solidification in thedie. Genera l ly, a simple rule that isfollowed by diecasters is that:Sludge factor (SF)=1x wt%Fe + 2x wt%Mn+ 3x wt%Cr ,where SF<1.7 is generallyc o n s i d e red to be most acceptable inpractice (e.g. Table 1).The morphology of the sludge particles isstrongly influenced by the Fe:Mn:Cr ratio ofthe alloy. Because of this, Mn and Cr ares o m etimes re fe rred to as Fe corre c to rssince they can eliminate a proportion ofu n d e s i rable, needle-like, Fe - c o n ta i n i n gintermetallic phases from forming within thealloy, instead causing them to form as moreinnocuous particles. Other transition metalelements (e.g. Co, Ni, V) may also serve asFe corre c to rs and contri b u te to sludge

Fig. 16: Fractography of a HPDC Al-Si-Cu-(Fe) alloy containing 0.08wt%Sn. Sn even in very low amountscauses severe die sticking, hot tearing (example in (a) shown at higher magnification in (b)) as well asacting as an Fe modifier. Sn causes the formation of rosette shaped Fe-bearing particles (c), which arebrittle and therefore preferred crack propagation sites. In (a), all the bright phases seen are therosetteshaped Fe-bearing particles. Backscattered SEM.

Table 4. The effect on Sn on the ageing response of Al-Si-Cu HPDC’s

As Cast properties T4 properties T6 properties

Alloy 0.2% P.S UTS %El. 0.2% P.S UTS %El. 0.2% P.S UTS %El.

380 alloy+ 0.01Sn 176 MPa 329 MPa 2.9% 239 MPa 368 MPa 3.2% 386 MPa 448 MPa 1.7%

380 alloy+ 0.08Sn 182 MPa 323 MPa 2.6% 242 MPa 363 MPa 3.0% 382 MPa 428 MPa 1.4%

formation, but Mn is the most commonlyfound in commercial alloy products since itis often introduced when secondary metal(e.g. canstock) is recycled. When the ratioof Fe:(Mn+Cr) is high, the sludge particlestend to be needle-like whereas when theratio is low, the particles tend to exhibitm o re Chinese scri pt, block y, sta r- l i kerosettes or polyhedral particles. Examplesof different Fe-containing particles present

CONCLUSIONS1. High pre s s u re diecast aluminium

alloys that respond to age hardeningm ay be successfully heat tre a te dwithout causing blistering, by solutiontreating for much shorter times and atlower temperatures.

2. When heat treated to a T4 temper,

in die-cast microstructures are shown inFigs. 5,9&16.The transition metal elements have nodistinct role in heat treatment. Rather theycause some increase in as-cast and hightemperature strength as well as improvingwear and creep resistance. Because of apropensity to form brittle phases, they mayhave adverse effects on fracture propertiesand fatigue resistance.

HPDC’s generally develop increasesin 0.2% proof stress of around 30%c o mp a red to the as-cast condition,to gether with simulta n e o u simprovements to ductility.

3. In a T6 temper, HPDC’s generallydevelop increases in 0.2% proof stress


of bet ween 75 and 120% whenc o mp a red to the as-cast condition,usually with little change to ductility.The high pro p e rties deve l o p e dcompare very fa vorably with thoseobserved for heat treated aluminiumpermanent mold castings and manywrought products.

4. During solution treatment, Si rapidlyundergoes significant morphologicalchanges. Initially it fragments beforespheroidizing and then grows due toOstwald ripening. When the numberof Si particles is high, internal gas-containing laminar defects tend to bepinned which appears to prevent ordelay their expansion that can lead tos u rface bliste ring. Ad d i t i o n a l ly, th ech a n ges to the silicon morp h o l o g yappear, in part, to be responsible forthe imp roved st re n g th / d u c t i l i t yrelationships in the heat tre a te dconditions.

5. S p h e roidized Si particles prov i d esignificant dispersion strengthening to

HPDC aluminium alloys.6. Cu is the pri m a ry element fo r

promoting a strong response to heattreatment in HPDC aluminium alloys.For the compositions studied th a tcontain ~0.3% Mg, the propertiescontinue to increase on addition of upto 4.9%Cu, whereas for lower levelsof Mg (~0.1%), the properties arelittle changed beyond about 3.6%Cu.

7. Magnesium has a st rong effect ini n c reasing the tensile pro p e rties ofheat treated HPDC Al-Si-Cu alloys. Inparticular, increasing the Mg contentfrom 0.1% to above 0.2% appears toi n i t i a te rapid early hard e n i n g .However, increasing the Mg contentabove 0.3% and up to as much as0.7% causes no further improvementin tensile pro p e rties. Where Cucontent is 3.6% or above and Mgcontent is >0.2%, 0.2% proof strengthvalues may exceed 400 MPa.

8. Zinc additions have little effect oni n c reasing the tensile pro p e rties of

a ged HPDC aluminium alloys ,p roviding only minor increases inst re n g th via solution st re n g th e n i n g ,because the Mg content is low. Theelement has the disadva n ta ge ofa c c e l e rating softening andoverageing at elevated temperatures.

9. Tin is re ga rded as having adeleterious effect because it causesdie soldering, flashing and hot tearingin HPDC aluminium alloys, as well aspromoting the formation of large, Fe-b e a ring sludge particles in th emicrostructure. Since it has been foundthat this element has no effect on theresponse of the Al-Si-Cu alloys to agehardening, it is recommended that itscontent be kept as low as possible.

10. Transition metals present in HPDCalloys have no significant effect on theresponse to ageing. These elementsform brittle intermetallic compoundsand as a result may have an adversee ffect on fra c t u re and fa t i g u eproperties.

ACKNOWLEDGEMENTSThe auth o rs would like to thank SamTa rta glia (Sam.Ta rta gl i a @ c s i ro.au) fo rbusiness development.

REFERENCES[1] Lu m l ey, R.N, R.G. O’Donnell, D.R.

Gunasegaram, M. Givord, InternationalPatent Application PCT/2005/001909,(2005).

[2] Lu m l ey, R.N., R.G. O’Donnell, D.R.G u n a s e ga ram, M.Givo rd: Mat. Sci.Forum, vols. 519-522, (2006), pp. 351-359.

[3] Lu m l ey, R.N., R.G. O’Donnell, D.R.G u n a s e ga ram, M.Givo rd: Proc. 13 thDie Casting Confe rence of th eAu st ralian Die Casting Association,M e l b o u rne, Au st ralia, (2006), Pa p e r25.

[4] Lu m l ey, R.N., R.G. O’Donnell, D.R.G u n a s e ga ram, M.Givo rd: Met. andMat. Trans A., 38A, (2007), pp. 2564-2574.

[5] Bonollo, F., J. Urban, B. Bonatto, M.B ot te r, Proc. 13 th Die Cast i n g

C o n fe rence of the Au st ralian DieC a sting Association, Melbourn e ,Australia, (2006), Paper 10.

[6] Kaufman, J.G. “Fracture Resistance ofAluminum Alloys: Not ch To u g h n e s s ,Tear Re s i stance and Fra c t u reToughness”, ASM international, (2001),p. 71.

[7] Axon, H.J., J. Inst. Metals, 81, (1952-53), pp. 209-213.

[8] Wang, L., M. Makhlouf & D. Apelian,Int. Mat. Rev., vol.40, #6, (1995) pp.221-238.

[9] Po l m e a r, I.J., “Light Alloys, Fro mTraditional Alloys to Nanocrystals”, 4thed., Butte rwo rth-Heinemann, (2006),pp. 49-51.

[10] Jorstad, J.L. Die Casting Eng., Nov/Dec.(1986), pp. 30-36.

the role of alloy composition in the heat treatment of aluminium high

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