A HIGH STRENGTH ALUMINIUM ALLOY FOR HIGH PRESSURE DIE CASTING

Shouxun Ji, Feng Yan, Zhongyun FanBCAST, Brunel University London, Uxbridge, Middlesex, UB8 3PH, UK

Keywords: Aluminium alloys; High pressure die casting; Mechanical property; Microstructure

Abstract

In order to achieve the mechanical properties of thin-wall components made by high pressure die-casting at a level of 300MPa of yield strength, 420MPa of UTS and 3% of elongation, the composition of the new aluminium alloy and casting parameters have been optimized to reduce casting defects in final components. In the present paper, we describe the alloy development in an industrial scale HPDC system.

Introduction

Aluminium alloys are being increasingly used as lightweight materials in transport since weight reduction has been proved as one of the most effective approaches to improve fuel economy and reduce CO2 emissions [1,2]. High pressure die casting (HPDC) is a widely used process to make thin-wall aluminium components for structural applications. The solution and ageing heat treatment have been used to enhance the mechanical properties of castings [3]. However, the existing aluminium alloys for HPDC process are mainly based on Al-Si, Al-Si-Cu, and Al-Mg-Si systems. Most of these alloys are not particularly developed for property enhancement by heat treatment, especially not suitable for solution heat treatment because of the unsatisfied mechanical properties for industrial application [4,5]. Therefore, the development of high strength aluminium alloys becomes significant to promote the application of aluminium alloys where an improved strength and lightweight are essential.

In the present paper, we describe the alloy development including the effect of Mg, Mn and Zn in the Al-Mg2Si system on the mechanical properties and microstructure of the cast alloys. The effect of solution and ageing heat treatment on the mechanical properties and microstructure was also studied

Experimental

Commercial purity aluminium, magnesium, zinc, Al-20wt.%Mn, and Al-50wt.%Si master alloys were used as base materials to make the aluminium alloys. 10kg melt were made each time using a clay-graphite crucible in an electrical resistance furnace. During manufacturing, Al, Si and Mn were melted first and 15 ppm Be was added to the melt 10 minutes prior to the addition of Mg ingots. The melt was degassed at 750ºC using a rotatory degassing equipment with a rotating speed of 500 rpm for 5 minutes. After degassing, the melt was maintained at a temperature of 700ºC. A mushroom casting with a bottom flat of 60×10mm was made by a steel mould for composition analysis. The casting was cut 5mm off from the bottom and ground down to 800 grid sand SiC paper. The composition was obtained from a well-calibrated optical mass spectroscope, in which at least five spark analyses were tested and the average value was taken as the chemical composition of the alloy.

The castings were produced using a 4500 kN cold chamber high pressure die casting machine, in which most of the operation parameters including injection speed and intensification pressure are monitored. One set of die was used to cast standard tensile samples, during which three 6.35mm round samples and three flat tensile samples at 2mm, 3mm, and 5mm were made in one shot. The mechanical properties were obtained from 6.35mm round tensile samples. All parameters had been optimized for the HPDC machine during casting.

The tensile tests were conducted following ASTM standard B557, using an Instron 5500 Universal Electromechanical Testing System equipped with Bluehill software. All tensile tests were performed at ambient temperature. The gauge length of the extensometer was 25 mm and the ramp rate for extension was 2 mm/min. The microstructure of each alloy was examined using a Zeiss optical microscope with quantitative metallography, and a Zeiss SUPRA 35VP SEM.

Results

1. Effect of extra Mg in Al-Mg2Si alloys

Fig. 1 shows the equilibrium phase diagrams at Al-rich side for Al-Mg2Si system. The pseudo-binary Al-Mg2Si system behaves like a normal binary eutectic system with a eutectic point at 13.9 wt.% Mg2Si. It shows that a narrow three-phase region for L+α-Al+Mg2Si co-existence between 591 °C and 578 °C. With extra Mg in the Al-Mg2Si system, the shape of equilibrium phase diagram exhibits some changes. The eutectic point is shifted towards a lower Mg2Si concentration and the area of co-existence for the three phases is enlarged. This means that, for a given Mg2Si content, an extra Mg addition results in a decrease of the volume fraction of the primary α-Al phase and an increase of the volume fraction of Al-Mg2Si eutectic phase.

To confirm the calculated results, a series of experiments were performed with different amounts of extra Mg in the Al-Mg2Si alloys. The representative microstructures are shown in Fig. 2. The primary phase presented in the as-cast microstructure was different. With 1.9wt.% extra Mg, the alloy showed a typical hypoeutectic microstructure, consisting of dendrites or fragmented dendrites of primary α-Al phase surrounded by Al-Mg2Si eutectic phase. However, the microstructure was much different when the extra Mg was 5.8 wt.%. The coarse primary α-Al phase almost disappeared and the microstructure was dominated by fine primary α-Al phase and Al-Mg2Si eutectics. In general, the experimental results were consistent with the calculation results, although the Al-Mg2Si alloys without extra Mg cannot be obtained in high pressure die casting because of the severe die soldering problem.

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Fig. 1 Equilibrium phase diagrams of Al-Mg2Si pseudo-binary system and Al-Mg2Si-Mg ternary system with 6wt.%Mg.

The yield strength, UTS and elongation of the as-cast Al-Mg2Si-Mg alloys can be varied by the content of extra Mg. The yield strength was increased with increasing Mg content, but showed no significant difference with increasing the Mg2Si phase. Conversely, the elongation was decreased with the increase of Mg and Mg2Si in the alloys. The UTS were different for the alloys with different extra Mg contents. The decrease of the UTS was significant for the Al-9.9wt.%Mg2Si and Al-12.7wt.% Mg2Si alloys with the increase of extra Mg in the alloys. However, it is noted that the yield strength is much lower than the target of 300MPa.

2. Effect of Mn in Al-Mg2Si-Mg alloys

Fig. 3 shows the typical microstructure of the intermetallics in the Al-7.8wt.%Mg2Si- 5.8wt.%Mg-0.59wt.%Mn alloy with 0.12 wt.% Fe. Intermetallics with two different sizes were observed and the volume fraction of the intermetallics increased with the increase of Mn addition. The quantitative analysis showed that the average size of the fine intermetallics were about 0.8μm. An obvious increase of the volume fraction of the Mn-rich intermetallics could be seen for the increasing Mn content. The composition of the Mn-rich intermetallic phase was measured with different amounts of Mn additions and the results confirmed that the Mn-rich intermetallics were α-AlFeMnSi phase. The measured constituent of the intermetallics could be described by Al42-

46(Fe,Mn)3.5-Si7. The Mn-rich intermetallics are essentially associated with Fe in the die cast alloys.

The increase of Mn content in the alloys resulted in a slight increase of the yield strength and UTS, but a slight decrease of the elongation for the Al-7.8wt.%Mg2Si-5.8wt.%Mg alloy under as-cast condition. The elongation was decreased from 6.9% to 6.5% when the Mn content was increased from 0.19wt.% to 0.59wt.%,

which is the level of the commonly used in the die-cast Al-Mg-Si alloy.

Fig. 2 Optical micrographs showing the effect of excess Mg on the as-cast microstructure of Al-7.8 wt.%Mg2Si alloy with (a) 1.9wt.%Mg, (b) 5.8wt.%Mg.

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Mg2Si (wt.%)

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α-Al

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Fig. 3 The morphology of Mn-rich intermetallic phase and AlMg intermetallic phase in the Al-7.8 wt.%Mg2Si-5.8 wt.%Mg-0.59wt.%Mn alloy with 0.12wt.%Fe.

3. Effect of Zn in Al-Mg2Si-Mg-Mn alloys

The typical microstructure of Al-Mg2Si-Mg-Zn-Mn alloy is shown in Fig. 4. An AlMgZn phase was observed, which were characterised by irregular shapes. The Zn-containing intermetallic compounds were located at the boundaries of the primary α-Al phase or between the eutectic cells and the primary α-Al phase. The measured volume fractions of the AlMgZn phase were increased significantly with increasing the Zn addition. The quantitative measurement confirmed that the intermetallic compounds contained 25±1 at.% Mg, 10±0.7 at.% Zn, which corresponded to a formula of Al13Mg5Zn2. Therefore, the Zn-rich intermetallic phase was essentially the same compound in the experimental alloys.

Fig. 4 Backscattered SEM images showing (a) the microstructure and (b) the morphology of Zn- containing intermetallics in the Al-7.8wt.%Mg2Si-5.8wt.%Mg-0.59wt.%Mn alloy with 3.51wt.% Zn.

The yield strength was increased but the elongation was decreased significantly with increasing the Zn content in the Al-7.8wt.% Mg2Si-5.8wt.%Mg-0.59wt.%Mn alloy under as-cast condition.

When Zn content was 3.51wt.%, the elongation was still higher than 2% and the yield strength was at a level of 250MPa and the UTS was at a level of 350MPa.

4. Effect of solution and ageing in Al-Mg2Si-Mg-Zn-Mn alloys

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Fig. 5 (a) Mechanical properties with different Zn contents in the Al-7.8wt.%Mg2Si-5.8wt.%Mg-0.59wt.%Mn and (b) backscattered SEM microstructure of the alloy with 3.51wt.%Zn alloy under a condition of solution treated at 490oC for 30 minutes and subsequent aged at 180oC for 90 minutes.

Solution and ageing treatments were performed at 490oC and 180oC, respectively. Fig. 5(a) shows the effect of Zn on the yield strength, UTS and elongation of the alloys after solution and ageing heat treatment. Significant increases in yield strength and UTS and a decrease in elongation were observed in the alloy with increasing the Zn content. The tensile properties can be easily achieved at a level of 340MPa of yield strength, 430MPa of UTS and 3.5% of elongation when Zn is 3.51wt.%.

The typical microstructures after solution at 490oC for 30 minutes and subsequent aged at 180oC for 90 minutes are shown in Fig. 5b. During solutionising, the AlMgZn intermetallics were gradually dissolved into the matrix. However, no obvious change

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(a)

(b)

(a)

(b)

was observed for the AlFeMnSi intermetallics and eutectic Mg2Si phase after the solution treatment. The TEM examination revealed that a large numbers of fine MgZn precipitates were formed in the α-Al phase after solutioning and ageing, which provide effective precipitates strengthening.

Discussion

In the Al–Mg2Si–Mg–Zn alloy, the as-cast microstructure contains several phases formed during solidification, including primary α-Al, eutectic Al–Mg2Si, β-AlMgZn and α-AlFeMnSi phases.

The microstructural characteristics of the alloy can be understood from the solidification and phase formation process. According to the equilibrium phase diagram of the Al–Mg2Si–Mg–Zn–Mn system on the cross section of Al–8 wt%Mg2Si–6 wt%Mg–0.6 wt%Mn calculated by Pandat software [5]. It is seen that the solidification sequence of the alloy is L-α-Al-Mg2Si- AlMgZn-AlFeMnSi-AlMg when Zn is with in the calculation range. The prior phase is α-AlFeMnSi, followed by the formation of other phases with temperatures being reduced. However, the calculation results are slightly different to the experimental results where no AlMg phase was actually found in the experimental results. The disappearance of the AlMg intermetallics is mainly because of the existence of Zn in the alloy. Although the exact mechanism is still not clear, the formation temperatures of both AlMg and AlMgZn intermetallics is close to each other according to the equilibrium shown in given non-equilibrium solidification, the AlMgZn phase is formed before AlMg phase, resulting in suppression of the AlMg phase.

As the microstructure in the diecastings compromises the primary α-Al phase, Al–Mg2Si eutectics and intermetallic phases. Mg,Si, Zn and Mn can provide solution strengthening to the α-Al phase. The precipitates of intermetallics under as-cast condition can strengthen the alloy by acting as pins to prevent dislocations from sliding under stress. Therefore, the fine and compact α-AlFeMnSi phase can reduce the detrimental effect of needle-shaped intermetallics. However, the intermetallics themselves are brittle and can initialize cracks, resulting in the reduction in ductility. The further strengthening by solution and ageing heat treatment results from the precipitation of MgZn phase in the α-Al phase. The AlMgZn compounds decompose during solution and subsequently precipitate during ageing. The fine precipitates provide much more effective strengthening effect in the alloy, during which the detrimental effect from irregular AlMgZn phase is minimized and the elongation is thus improved. On the other hand, during heat treatment, no significant change is observed for the α-AlFeMnSi andMg2Si phase because of their high thermal stability, indicating that the enhancement in mechanical properties from the α-AlFeMnSi andMg2Si intermetallics is limited after the heat treatment.

Summary

A high strength aluminium alloy for high pressure die casting has been developed based on Al-Mg2Si-Mg-Zn system. The typical alloying elements comprises of 8.0wt.%Mg2Si, 6.0wt.%Mg, 3.5wt.%Zn, 0.6wt.%Mn. The alloy consists of primary α-Al, Al-Mg2Si eutectics, β-AlMgZn, and α-AlFeMnSi phases in the as-cast microstructure. The typical tensile properties of the die-cast

tensile samples are 300MPa of yield strength, 420MPa of ultimate tensile strength, and 3% of elongation under solution and aged condition. A quick solution and ageing is an effective way to enhance the mechanical properties, by which the AlMgZn intermetallics are decomposed into Al matrix and fine Zn-rich intermetallics are precipitated during the subsequent ageing treatment.

Acknowledgement

The authors acknowledge the Engineering and Physical Sciences Research Council (EPSRC), Technology Strategy Board (TSB) and Jaguar Land Rover (JLR) in United Kingdom for financial support.

References

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