Imperfections in Recycled Aluminium-Silicon Cast...

School of Engineering Jönköping University

Licentiate thesis Dissertation Series No. 8 • 2015

Imperfections in RecycledAluminium-Silicon Cast Alloys

ANTON BJURENSTEDT

LICENTIATE THESIS

Imperfections in Recycled Aluminium-Silicon Cast Alloys

ANTON BJURENSTEDT

Department of Materials and Manufacturing

SCHOOL OF ENGINEERING, JÖNKÖPING UNIVERSITY Jönköping, Sweden 2015

Imperfections in Recycled Aluminium-Silicon Cast Alloys Anton Bjurenstedt Department of Materials and Manufacturing School of Engineering, Jönköping University SE-551 11 Jönköping, Sweden [email protected] Copyright © Anton Bjurenstedt Research Series from the School of Engineering, Jönköping University Department of Materials and Manufacturing Dissertation Series No. 8 ISBN 978-91-87289-09-5

Published and Distributed by School of Engineering, Jönköping University Department of Materials and Manufacturing SE-551 11 Jönköping, Sweden Printed in Sweden by Ineko AB Kållered, 2015

Imperfections in recycled aluminium-silicon cast alloys


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ABSTRACT

In striving to produce high quality cast components from recycled aluminium alloys, imperfections have to be considered, because recycled aluminium usually contains more of it. However, there are great energy savings to be made by using recycled aluminium; as little as 5% of the energy needed for primary aluminium production may be required. High quality castings are dependent on, besides alloy chemistry, both melt quality and the casting process; the focus of this work is related to the melt quality.

This thesis aims to increase knowledge about imperfections, foremost about Fe-rich particles, oxides/bifilms, and porosity. Experiments were performed at industrial foundry facilities and in a laboratory environment. Melt quality was evaluated by producing samples with the reduced pressure test (RPT), from which both density index (DI) and bifilm index (BI) could be measured, results that were related to tensile test properties. Data from tensile test samples were analysed, and fracture surfaces and cross sections were studied in both light microscope and in scanning electron microscope (SEM). For the purpose of investigating nucleation of primary Fe-rich particles (sludge) differential scanning calorimetry (DSC) was used.

In the analysis of results, a correlation between the morphology of particles and tensile properties were found. And elongated Fe-rich β-particles were seen to fracture through cleavage towards the centre. However, DI and BI have not been possible to relate to tensile properties.

The nucleation temperature of primary Fe-rich particles were found to increase with increased Fe, Mn, and Cr contents, i.e. the sludge factor (SF), regardless of cooling rate. For a set SF, an increase of cooling rate will decrease the nucleation temperature.

Keywords: Imperfections, Recycled cast Al-Si alloy, Fe-rich particles, Melt quality, Fractography, Differential scanning calorimetry.


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ACKNOWLEDGEMENTS

I would like to express my sincere gratitude to some people who have helped me a little extra along the way:

My supervisors, Professor Anders E. W. Jarfors and Docent Salem Seifeddine, for their support, inputs, and fruitful discussions.

The technicians, Toni Bogdanoff, Esbjörn Ollas and Lars Johansson for their assistance when I was preparing samples.

I am grateful to the people at Stena Aluminium and Ljunghäll, for their help during long hours of experiments.

PhD student Stefano Ferro for some fun times and great collaboration.

Master students Marcello Gobbi and Francesco Bettonte for their excellent experimental work.

Friends and colleagues at the School of Engineering, Jönköping University, thanks for creating such a nice work environment.

Thanks to my family for being there, and for not talking too much work and keeping me fed.


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SUPPLEMENTS

The following supplements constitute the basis of this thesis.

Supplement I A. Bjurenstedt, S. Seifeddine, A. E. W. Jarfors: On the complexity of the relationship between microstructure and tensile properties in cast aluminium. Proceeding at APCMP 2014 and in International Journal of Modern Physics B 29(10&11), 2015, 1540011.

A. Bjurenstedt was the main author, S. Seifeddine and A. E. W. Jarfors contributed with advice throughout the work.

Supplement II A. Bjurenstedt, S. Seifeddine, A. E. W. Jarfors: Effect of Fe-rich particles on tensile properties in recycled cast aluminium. Submitted to Metallurgical and Materials Transactions A.

A. Bjurenstedt was the main author, S. Seifeddine and A. E. W. Jarfors contributed with advice throughout the work.

Supplement III S. Ferraro, A. Bjurenstedt, S. Seifeddine: On the formation of sludge intermetallic particles in secondary aluminum alloys. Accepted for publication in Metallurgical and Materials Transactions A.

A. Bjurenstedt was co-author and S. Ferraro was the main author, A. Bjurenstedt and S. Ferraro performed the experimental together, S. Seifeddine contributed with advice throughout the work.

Supplement IV A. Bjurenstedt, S. Seifeddine, T. Liljenfors: Assessment of Quality when Delivering Molten Aluminium Alloys Instead of Ingots. Materials Science Forum Volume 765, 2013, pages 266-270.

A. Bjurenstedt was the main author, S. Seifeddine and T. Liljenfors contributed with advice throughout the work.

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TABLE OF CONTENTS

CHAPTER 1. INTRODUCTION ................................................................................................................................... 1 1.1 BACKGROUND ...................................................................................................................................................................................... 1 1.2 ALUMINIUM ........................................................................................................................................................................................... 2 1.3 SUSTAINABLE PRODUCTION OF CAST COMPONENTS .................................................................................................... 6 1.4 IMPERFECTIONS IN CAST ALUMINIUM .................................................................................................................................. 7

CHAPTER 2. RESEARCH APPROACH .................................................................................................................. 13 2.1 PURPOSE AND AIM ......................................................................................................................................................................... 13 2.2 RESEARCH DESIGN ......................................................................................................................................................................... 13 2.3 MATERIAL AND EXPERIMENTAL PROCEDURES ............................................................................................................. 15

CHAPTER 3. SUMMARY OF RESULTS AND DISCUSSION .............................................................................. 17 3.1 IMPERFECTIONS IN RECYLED CAST ALUMINIUM .......................................................................................................... 17 3.2 FE-RICH PARTICLES ....................................................................................................................................................................... 22 3.3 LIQUID ALUMINIUM TRANSPORT ........................................................................................................................................... 28

CHAPTER 4. CONCLUDING REMARKS ............................................................................................................... 31

CHAPTER 5. FUTURE WORK ................................................................................................................................. 33 REFERENCES...… .......................................................................................................................... 35

APPENDED PAPERS .................................................................................................................... 39


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CHAPTER 1

INTRODUCTION

CHAPTER INTRODUCTION

This chapter first describes the background to the work, and provides an introduction to cast aluminium and imperfections.

1.1 BACKGROUND

A key issue in producing reliable high strength aluminium castings is to understand and minimise the amount of imperfections. There is a desire for reliable higher strength alloys, primary from the automotive industry in its pursuit of making vehicles lighter. The automotive industry is the largest user of cast aluminium, and the increase in motor vehicle sales on the global market (world car sales 2004-2014 [1]) has produced an increasing demand for aluminium alloys. As a consequence, the production of high quality castings from recycled aluminium becomes more and more important. However, each remelting of scrap to produce new alloys brings with it the chance of an increase in imperfections. On the other hand, the production of recycled aluminium alloys uses substantially less energy compared to the production of primary aluminium alloys produced from mined bauxite; the European Aluminium Association (EAA) has stated that up to 95% of the energy is saved when producing recycled aluminium alloys.

Imperfections include defects such as porosity and bifilms as well as dimensions of microstructural features such as secondary dendrite arm spacing (SDAS) and Si-particle length; features that may affecting tensile properties without normally being considered as defects. Variations in mechanical properties can often be correlated to imperfections in the castings [2]. Imperfections in the final casting can originate from the casting process [3, 4] or the melt itself [5, 6].

Recycled aluminium melts usually contain more imperfections than primary aluminium [7]. This causes degradation in mechanical properties and could possibly also affect physical properties, if not diluted with more pure aluminium. Samples cast with a gradient solidification technique where imperfections are pushed ahead of the solidification front, generate castings with significantly lower amount of imperfections. Tensile test samples cast with this technique show improved strength and elongation compared to samples cast through conventional casting procedures (see Figure 1). This shows that recycled aluminium castings are generally not yet reaching their fullest potential.

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Figure 1. Graph showing the potential of cast aluminium with lower amount of

defects [8].

Besides higher strength and better elongation to fracture (ef), scattering of mechanical properties is an issue. Data in a report from the Aluminum Association (AA), cited by Sigworth [9], show the variation in tensile properties of a heat treated alloy, cast in a standardised permanent mould by a number of different foundries. The mould had five different areas with variation in section thickness producing variation in solidification speed; Table 1 shows the range of values reported for ultimate tensile strength (UTS), yield strength (YS) and ef. As can be seen in the table, the properties had quite large variations.

Table 1. The range of values for the same alloy cast in the same mould by different

foundries [9].

Area UTS

(MPa) YS

(MPa) ef

(%) 1 235-276 166-242 1.8- 4 2 231-283 166-242 1.5- 4.5 3 252-297 173-162 3- 7.7 4 259-314 166-269 3.5- 9.5 5 248-293 173-162 3- 7.5

1.2 ALUMINIUM

1.2.1 Introduction to aluminium

There are two main types of aluminium:


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Wrought aluminium – originally cast as billets or ingots and then hot or cold formed into shape by, for example, rolling, extrusion, or forging.

Cast aluminium – directly cast into shape in a mould made from primarily sand or steel.

The main difference in chemical composition between the two types of aluminium is the silicon (Si) content. Cast aluminium typically has a higher Si content in order to increase castability; that is, to generate a sound casting with good mechanical properties.

Figure 2 shows the Al-Si phase diagram with the most frequently used Si contents. Compositions above the eutectic composition at 11.7% Si are referred to as hypereutectic, and compositions below are referred to as hypoeutectic.

Figure 2. The Al-Si phase diagram, showing the most frequently used Si contents

[10].

1.2.1.1 Cast aluminium

Casting is an economical way of producing near net shaped products with complicated geometries, since the as cast product normally requires only a minimum of expensive machining. The Si in the cast aluminium alloys improves the castability by increasing fluidity, resistance to hot cracking, and feeding [11].

Aluminium alloys are cast in two significantly different types of mould: expandable moulds and permanent moulds. Patterns can also be of expandable or permanent type. The dominant form of expandable mould is the sand mould with various types of binder materials, but one can also cast in, for example, plaster moulds. Permanent moulds are usually made of steel. They should have a good resistance to thermal fatigue, making the following properties desirable: high thermal conductivity, high strength at elevated temperature, low thermal expansion, and low modulus of elasticity [12]. The main advantage of gravity casting in a permanent mould is the faster solidification rate, which leads to finer structure and hence improved mechanical properties.

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In permanent moulds, pressurised molten metal may be used to fill the mould cavity, and with process automation, production rates can be high. However, pressure die casting requires a relatively high economic investment in the die, making it only suitable for large series, >20000 castings/year [13].

1.2.1.2 Main alloying elements

Alloying elements are used in order to improve properties such as; casting characteristics and strength. Commercially pure aluminium (< 99% Al) has low tensile strength and good elongation, with a UTS value of about 70MPa and elongation of about 43% [14].

Silicon (Si) increases the castability; that is, the ability to readily fill dies and to solidify as a component without hot cracking. For casting alloys, Si contributes to a higher degree of isothermal solidification which improves fluidity. Si has about five times greater heat of fusion than aluminium, contributing to high fluidity of hypereutectic alloys. In a well fed casting, the eutectic liquid lowers the risk of hot cracking [15].

Copper (Cu) improves strength and hardness in both as-cast and heat treated condition. Copper also improves machinability by increasing the matrix hardness, which makes it easier to generate small cutting chips and high surface finish. However, Cu reduces corrosion resistance [16].

Magnesium (Mg) improves the strength and hardness after heat treatment by forming Si2Mg precipitates which efficiently increase strength by precipitation hardening. Aluminium alloys containing Mg < 2% will form a surface oxide called spinel, MgAl2O4, and for Mg >2% the spinel converts to MgO; if entrapped in the melt, these oxides may be harmful to the final casting [17].

Iron (Fe) decreases the possibility of die soldering or die sticking [18], and improves resistance to hot tearing by increasing the high temperature strength [19]. However, Fe can also reduce mechanical properties (see section 1.4.2).

Manganese (Mn) is used to modify the morphology of harmful Fe platelets into α-Al15(Fe,Mn,Cr)3Si2, which has a more compact, shape and could therefore be less harmful. A maximum Fe:Mn ratio of 2:1 has been the accepted rule in the industry to suppress the formation of β-Al5FeSi [20].

Chromium (Cr), in AlSi9Cu3 alloys chromium promotes α-Al15(Fe,Mn,Cr)3Si2, and increases the size of the Fe-rich particles [21].

Strontium (Sr), antimony (Sb), sodium (Na), and calcium (Ca) modify the aluminium-silicon eutectic. Modifying changes the eutectic silicon from coarse continuous networks of thin platelets into finer fibrous or lamellar structures. When modifying the eutectic silicon, strength and ductility is increased [16].

Titanium (Ti) is used as a grain refiner together with Boron (B), commonly in a Ti:B ratio of 5:1 [16].

1.2.2 Solidification

Solidification can be monitored by thermal analysis (TA), in which the temperature of the metal is recorded over time. An example of a cooling curve of an A380 alloy is shown in Figure 3. From the peak temperature at the left hand side, the molten alloy


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is cooled down towards the nucleation temperature of the α-Al dendrites; the curve shows the undercooling needed to form the α-Al dendrites. If grain refiners are added, the required undercooling becomes less, and so the cooling curve can give information about the grain size in the casting.

The next region, between α-Al nucleation and eutectic nucleation, is where the α-Al dendrites grow and fill the casting. Growth after the dendrites have filled the casting will only occur laterally. This region is thus related to the SDAS, which is a common way to evaluate the local solidification time in a casting. Sometimes, especially in high pressure die castings, the SDAS can be hard to distinguish; in this case another measurement, called cell size or cell count, can be used. Cell size or cell count is the number of rounded Al-phase features in a measured length.

The second undercooling is related to the eutectic nucleation, and will give information about the modification level of the Si particles in the eutectic; a smaller undercooling indicates a higher level of modification. The temperature of eutectic nucleation also becomes lower for a modified alloy [22, 23].

The last reaction is related to the precipitation of intermetallics such as, Al2Cu and Al5Mg8Si2Cu2.

The nucleation of Fe-rich intermetallics is not detected in the cooling curve shown. They nucleate at temperatures before α-Al dendrites down to temperatures of the eutectic nucleation, depending on chemical composition and cooling rate.

Figure 3. Cooling curve of an A380 alloy, temperature versus time, showing a

cooling rate of 0.6°C/s (after solidification). After reference [22].

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Figure 4. Grain size, primary dendrite arm spacing (DAS,) and SDAS. After

reference [17].

1.3 SUSTAINABLE PRODUCTION OF CAST COMPONENTS

1.3.1 Recycling

The lifecycle of aluminium is shown in Figure 5. The top row illustrates the high energy consuming process of producing primary aluminium, and the bottom row shows the far less energy consuming process of recycled aluminium, also referred to as secondary aluminium.

There are great energy savings in producing aluminium alloys by recycling scrap. However, the demand for aluminium is high; and when used in vehicles, for example, the service time is long. This have led to that the amount of recycled aluminium from post-consumer scrap in new ingots were 27% in Europe 2004 [24].

Figure 5. The aluminium life cycle. Image curtesy of Constellium.


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Product design plays an important role in the economic viability of recycling the final product. Separation of material at the end of use should be taken into consideration during the design process, as even trace elements can affect the tensile properties [25, 26] and impact properties [27] of castings.

When it comes to recycling, the ideal situation is a closed loop like that shown in the lower part of Figure 5. However, there is limited availability of scrap metal; and, for example, a car manufacturer wanting to use recycled aluminium might end up using beverage can scrap, because of shortage of aluminium scrap from cars. This will lead to reduction of scrap in the beverage can recycling cycle, and more primary aluminium will be added into the aluminium can production instead. The goal is therefore to produce new engines out of old engines, and so forth, in order to add as little primary aluminium as possible to the cycle.

In order for refiners to produce recycled aluminium with a minimum of additional alloying elements, scrap is sorted in to various classes. Set amounts of scrap from different classes are melted in order to produce a melt as close as possible to the final chemical composition. The chemical composition is tested and alloying elements are added, if needed.

1.3.2 Transportation of aluminium melt

The traditional remelting method involves casting the aluminium alloys into ingots which can be transported to the foundries with standard lorries. An alternative is to pour the aluminium into large thermoses, and transport it to the foundries in its liquid state. This method requires one fewer remelting, thus saving energy. One drawback is that less aluminium can be transported by each truck: 22 tons for liquid versus 30 tons for ingots (Stena Aluminium). Another is that it requires greater cooperation between remelters and foundries, since delivery of the liquid metal is often time-critical. Small foundries may find the high volume in each delivery hard to handle, since the liquid metal does not store as well as ingots. For larger foundries, this means that a combination of ingots and liquid metal is the best practice from both an economic and an environmental standpoint [28].

It has been seen in small-scale experiments that precipitation and sedimentation could be a way of purifying the melt, thereby lowering the amount of imperfections in the casting [29-31]. On the other hand, there are concerns whether liquid metal delivery will introduce bifilms, hence degrading the final casting [17].

1.4 IMPERFECTIONS IN CAST ALUMINIUM

1.4.1 Introduction to imperfections

Imperfections include defects such as porosity and bifilms as well as the type and dimensions of microstructural features such as SDAS and Si-particle length; these features may affect tensile properties without being considered as defects. The consequence of the presence of imperfections is premature failure [2].

The three main imperfections in casting alloys are, in no particular order, intermetallics, oxides, and porosity. All three types of imperfections are related to both the melt and the casting process.

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1.4.2 Fe-rich intermetallics

Intermetallics form due to low solubility in the solid. Fe, which generates the most common intermetallics in recycled cast alloys, has a solubility in solid aluminium of about 0.03-0.05% at 655°C and even lower at room temperature [19]. Together with Al and Si, Fe forms β-particles (β-Al5FeSi) which are complex 3D plate structures (see Figure 6 a). On a polished 2D specimen, the β-particles have the shape of needles and are therefore at times incorrectly called “iron needles”. β-particles have a detrimental effect on mechanical properties; tensile testing has shown that they primarily reduce ductility, which is accompanied by a reduction in UTS [32, 33]. The explanation for the harmful effect is that the β-particles offer a preferred path way for crack growth (see Figure 7). There are two ways of altering the shape of β-particles: changing chemical composition and/or changing cooling rate. Addition of Mn converts the platelets into more compacted shapes, α-particles (α-Al15(FeMnCr)3Si2), which are considered as less detrimental to mechanical properties [34].

(a) (b)

Figure 6. Three-dimensional reconstructions of (a) β-Al5FeSi platelets (b) an α-Al15(Fe,Mn)3Si2 particle [35].

Figure 7. β-particles causing a faceted fracture surface. Secondary cracks are

shown in close connection to the β-particles.


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The commonly accepted Mn content needed to form α-particles is that it should be enough to keep the Fe:Mn ratio below 2:1. For a given Fe:Mn ratio, a slow solidification rate will form primary polyhedral α-particles called sludge (Figure 8 a). When the solidification rate it high the α-particles may grow as coupled eutectic and form Chinese script (see Figure 8 b) [36]. For intermediate cooling rates the particles may end up like in Figure 6 b, having a polyhedral crystal at the centre and convoluted arms connected to it [35].

(a) (b)

Figure 8. α-particles with different morphologies: (a) as primary polyhedral sludge particles and (b) as Chinese script in a quenched sample.

Problems with sludge is experienced in foundries, especially high pressure die casting foundries. The high specific gravity and makes them settle at the bottom of furnaces which reduces their capacity and sludge can restrict metal flow during casting [37, 38]. Sludge is hard and brittle [39], and when machining, tool life will be reduced [40]. The effect on tensile properties are more unclear, researcher have referred to sludge as “detrimental to mechanical properties” [34, 41], but when looking at results by Ji et al. [42], there is no significant reduction of UTS or ef for an Al-Mg-Si alloy with Fe:Mn ratio <2.

The sludge factor (SF) takes the Fe, Mn, and Cr contents into account in order to determine the whether an alloy sitting in the furnace at a given temperature will form sludge. The equation for the SF expressed by Dunn [43] is:

SF = wt.%Fe + 2×wt.%Mn + 3×wt.%Cr (1)

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A graph with temperature versus SF then show what temperature is sufficient to avoid precipitation of sludge. Figure 9 presents the results from two different researchers.

Figure 9. Formation of sludge relating to temperature and SF according to Jorstad

[37] and Gobrecht [44] (the dotted line is an extrapolation.)

Fe-rich particles have been suggested to increase porosity, and there are two major theories. One is that the Fe-rich particles act as nucleation sites for porosity, as porosity has been seen to be in contact with β-particles [45]. The other theory is that β-particles block feeding to the interdendritic regions, causing shrinkage porosity [45]. However, in observations of porosity in a polished sample it is hard to distinguish between cause and effect.

1.4.3 Oxides in cast aluminium

Aluminium alloys oxidize readily, forming oxide layers on surfaces in contact with the oxygen in the air. The growth of surface oxides increases with the temperature of the alloy [46]. Oxides grow from the surface, and as the layer becomes thicker the growth rate becomes slower, because the oxygen has to diffuse through the oxide layer. On an aluminium melt, the oxides will have a good bonding to the melt. If waves form at the surface, the oxides may become entrained in the melt and retain their bonding to the melt, but between the two dry surfaces there is no bonding and so the entrained oxide will act as a crack in the melt [17]. A sketch of this is shown in Figure 10, where the thicker oxide has been broken up and parts of it are becoming entrained into the melt while the new surfaces are quickly growing new oxide layers. The growth is initially fast; it takes only about 0.01 seconds to form an oxide layer [17]. This means that oxides form and become entrained in virtually all casting situations during mould filling. Once entrained, the oxides tend to stay suspended in the melt, due to the similar specific gravity of the oxides and the molten aluminium [47]. Oxides inside the melt or casting are referred to as bifilm due to their fold-over nature. Bifilms have been reported to nucleate porosity [48, 49], Fe-rich particles [50, 51], and Si particles in the eutectic [51].


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Figure 10. Entrainment of surface oxides [17].

In the filling process of a casting, the bifilms will experience bulk turbulence causing them to be ravelled into a relatively compact formation. When the filling is completed, the turbulence in the casting slows down into more gentle movements. In this turbulence free environment the bifilm may unravel into more crack-like formations, weakening the casting to a greater extent compared to the ravelled compact formations. A faster cooling rate will hinder this unravelling, thus enhancing mechanical properties [17].

In order to keep the amount of bifilms as low as possible in the final casting, a clean melt should be delivered to the mould and should fill up the mould cavity in a calm manner. This is achieved by proper design of the gating system and mould cavity [52].

1.4.4 Porosity

Porosity in cast aluminium has two major causes: hydrogen (H) and shrinkage porosity [53].

H is the only gas that has any significant solubility in aluminium. Its solubility is much lower in the solid aluminium compared to the melt, leading to rejected H forming porosity unless it escapes during solidification [54]. Figure 11 shows the dramatic solubility drop when the alloy solidifies; this drop is unique for aluminium in comparison to, for example, Fe, Mg, and Cu. As the aluminium solidifies, the H is rejected into the remaining liquid until the concentration is sufficient for the H to nucleate. The gas pores normally end up as more or less spherical pores, because the growth has been without restrictions in the melt (see Figure 12 a).

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Figure 11. Hydrogen solubility in pure aluminium and two of its alloys [54].

When cast aluminium alloys solidify the density increases and shrinkage occurs; an alloy with 8.9% Si will shrink 5.2% [55]. As regions of the casting start to solidify, the shrinkage has to be fed by liquid from neighbouring areas through the interdendritic channels. If areas are not fed, the built up tension due to shrinkage will be released by forming porosity. Shrinkage pores on a polished 2D sample are normally seen as empty interdendritic regions, (see Figure 12 b).

(a) (b)

Figure 12. Images showing (a) gas porosity and (b) shrinkage porosity. Etched with H2SO4.


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CHAPTER 2

RESEARCH APPROACH


This chapter starts with the purpose and aim of this thesis. Then is the research design described, followed by a description of the material used and experimental procedures.

2.1 PURPOSE AND AIM

The purpose of this work is to study the effects of imperfections in recycled aluminium-silicon cast alloys in general and Fe-rich particles in particular. Fe-rich particles in recycled aluminium alloys are a growing problem, as they increase with each recycling and today the only viable way of lowering Fe content is by diluting.

The aim is to increase knowledge in order to be able to improve strength and reliability (decrease scattering), which will facilitate greater usage of the relatively environmentally friendly recycled aluminium alloys, taking advantage of their long service life and the low energy consumption needed for recycling. As a result, industries such as the automotive industry, currently the largest user of cast aluminium components, can reduce weight and lower fuel consumption even further, which is beneficial for both the environment and the customer.

2.2 RESEARCH DESIGN

2.2.1 Research perspective

Two popular research perspectives whereby research can be conducted are, positivism and interpretivism, each having their way of reasoning. While positivism is related to deductive reasoning, interpretivism is related to inductive reasoning. Deductive reasoning is linked to a hypothesis that is either rejected or accepted after analysis of experimental results. While deductive reasoning starts with a hypothesis, inductive reasoning generates a hypothesis from analysis of experimental results without imposing pre-existing expectations [56].

The research perspective in this work combines these two ways of conducting research. Hypotheses were used in the design of the experiment, but the results were analysed without imposing pre-existing expectations.

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2.2.2 Research methodology

An overview of the research method is shown in Figure 13. This began with the topic of interest, which was established in cooperation with the industry in order to perform research for industrial purposes as well as for the scientific community. Building on the topic of interest, a literature study and gathering of information were performed to increase knowledge and uncover the latest research in the field. Narrowing of the topic of interest was then necessary in order to design feasible experiments. Gathering information made it possible to narrow the topic of interest and pinpoint the unexplored, any contradictions, and any areas where limited research had been performed. Design of the experiment began when the research issue had been pinpointed, in order to establish cause and effect between variables. Performing of a pre-experiment might sometimes be necessary in order to verify the experimental proceeding, making sure it will be possible to measure what is intended. When the pre-experiment is satisfactory the next steps are to perform the full experiment and do the analysis of the results, and to draw conclusions.

Figure 13. Overview of the research methodology.

2.2.3 Research questions

Below is an overview of the research questions, and the supplements in which each of them are addressed.

• Research question 1: Supplements I and II

What relations exist between imperfections and tensile properties?

• Research question 2: Supplement II

How do Fe-rich particles influence fracture?

• Research question 3: Supplement III

How does cooling rate influence nucleation temperature of sludge?

• Research question 4: Supplement IV

How does liquid aluminium transport affect melt quality?


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2.3 MATERIAL AND EXPERIMENTAL PROCEDURES

2.3.1 Alloys

The alloys used throughout this work were AlSi9Cu3(Fe) type, equivalent to A380, to which elements were added as master alloys to produce variation in chemical compositions for the experiments. The base alloy was EN AB-46000 (ingots); see Table 2 for chemical composition according to SS-EN 1676:2010. The major elements added were Fe, Mn, Cr, and Sr.

Table 2. Chemical composition (wt%) of EN AB-46000 according to SS-EN 1676:2010

[57].

Si Fe Cu Mn Mg Cr Ni Zn Pb Sn Ti EN AB- 46000

min 8.0 0.6 2.0 - 0.15 - - - - - - max 11.0 1.1 4.0 0.55 0.55 0.15 0.55 1.2 0.35 0.15 0.20

2.3.2 Sample preparation

Samples for Supplements I and IV were collected in Y-shaped moulds at a foundry from liquid transported alloy. Three tensile test specimens were subsequently machined from the lower part of the casting.

Samples for Supplements II and III were melted in an electric resistance furnace capable to melt about 7 kg of aluminium. Samples for tensile testing and differential scanning calorimetry (DSC) were cast into rods in a preheated copper mould. These rods were subsequently remelted in a gradient solidification furnace, with melting taking place in an argon (Ar) gas atmosphere. Next, the furnace was pulled along the stationary samples in a constant speed, and water cooling was applied, solidifying the samples from one end to the other, thereby well fed castings were produced with a low amount of casting defects. By varying pulling speed of the furnace the cooling speed of the samples can be controlled, thereby producing variation in SDAS.

Alloys for quench experiments were poured into a brine.

An underaged T6 heat treatment process was performed on tensile test samples in Supplement II in order to gain a dissolution and homogenous microstructure in order to study the sole effect of Fe-rich particles without the impact of Cu- and Mg- rich particles.

2.3.3 Sample and alloy testing

In Supplements I, II, and III, optical emission spectrometry (OES) was used to measure chemical composition. Certified standards were measured in connection to the unknown sample, allowing the results to be verified.

In Supplements I and IV, a reduced pressure test (RPT) apparatus was used to evaluate melt quality by means of density index (DI) and bifilm index (BI). Where DI, relates to how easy the pore formation is in the melt and BI (introduced by Dispinar [58]) relates to the amount of bifilms [58]. The RPT apparatus eases pore formation by solidifying a sample (about 70g of melt) at rescued pressure. DI is the per cent density decrease

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of a sample solidified in the RPT (density ρRPT) compared to the density of a sample solidified at normal atmospheric pressure (ρatm) and is calculated according to:

DI = (ρatm – ρRPT) / ρatm ×100. (2)

BI is measured as the sum of the largest Feret-diameter of each pore in a cross section of a RPT sample. This measure is assumed to relate to the amount of bifilms (oxides), and hence is related to mechanical properties [4, 59].

Tensile testing was performed in Supplements I, II, and IV. The round tensile test specimens and strain rate, controlled by constant crosshead separation rate, fulfilled the requirements of ISO 6892-1:2009 and ASTM B557M -07. Both a laser extensometer and a clip on extensometer with a gauge length of 20mm were used.

Thermal analysis (TA) in Supplement III was performed on alloys cast into preheated stainless steel tubes placed inside three different cooling media to attain different cooling rates. Temperature during solidification was measured with the two-thermocouple method developed by Bäckerud et al. [22], with one thermocouple in the centre and one close to the wall in order to record the “heat wave” produced when new phases are formed.

DSC analysis was performed in Supplement III in order to obtain higher resolution of the exothermic reaction related to sludge nucleation. In the DSC, the heat flow between a reference (an empty crucible) and the alloy sample (40mg) was determined by measuring the temperature difference during various cooling rates. This temperature difference is caused by the release of heat in the sample related to formation of phases. Three of more tests were performed for each alloy and cooling rate, and a new sample was used in each run.

2.3.4 Material characterization and evaluation

A stereo microscope was used to evaluate tensile test fracture surfaces in Supplements I, II, and IV, in order to determine which samples should be examined in the secondary electron microscope (SEM) and in the optical microscope (OM).

SEM and energy dispersive spectroscopy (EDS) were used in Supplements I, II, and IV in order to examine fracture surfaces and for identification of phases.

The optical microscope was used in all supplements. The samples were polished, and 10%-H2SO4 at 70°C [60] was used to etch Fe-rich particles black when extra contrast was needed for analysis and measurements.


17

CHAPTER 3

SUMMARY OF RESULTS

AND DISCUSSION


This chapter summarizes and discusses the main results of the supplemented papers.

3.1 IMPERFECTIONS IN RECYLED CAST ALUMINIUM (SUPPLEMENTS 1 AND II)

In Supplement I, melt process quality measures and microstructural features have been linked to tensile properties of the final casting to establish any possible correlations and evaluate inconclusive results. The measured features and tensile properties are shown in Table 3, and the chemical composition of the alloy used is presented in Table 4.

Table 3. Table showing the features and parameters measured.

Process quality measures Min Max Average Density index, DI 3.9 8.3 6.3 Bifilm index, BI (mm/cm2) 6.5 17.1 10.8 Sludge factor, SF 1.29 1.78 1.51 Microstructural features Porosity (area %) 0.04 1.19 0.29 Secondary dendrite arm spacing, SDAS (µm) 15.6 41.7 25 Si-length (µm) 7.5 75.6 42.1 Fe-rich particles (area %) 1.06 3.30 2.51 Output parameters Ultimate tensile strength, UTS (MPa) 140 281 213 Elongation to fracture, ef (%) 0.42 4.19 1.42

18

Table 4. Chemical composition (wt%) of the alloy in Supplement I. Average values are shown.

When plotting UTS versus ef a positive linear relationship was observed: i.e. graphs showing UTS versus imperfections showed the same trends as for ef versus imperfections. Therefore, in the following, only the results of the analysis of UTS versus imperfections are presented.

RPT samples were collected to examine correlation between BI, DI, and SF and tensile properties. The RPT samples (for BI and DI) and samples for chemical analysis were collected at the same time, from the same melt, as samples for tensile testing. Figure 14 shows fairly large variations in tensile strength for each BI and SF. Similar results were found for the DI; hence, no correlations were found between the process quality measures and tensile properties in our samples.

Our results where BI is correlated to UTS and ef are assessed to be in line with results shown by Dispinar and Campbell [61], who show graphs with very limited variation in UTS and elongation for different BI values. However, they draw the conclusion that the BI appears to define elongation but does not necessarily predict mechanical strength.

(a) b)

Figure 14. Graphs relating UTS to (a) bifilm index (BI) and (b) sludge factor (SF).

The microstructure of the tensile tested samples was examined, and their features measured. Porosity was measured and related to UTS; however, no correlation could be seen. This is in agreement with results by Herrera and Kondic [62].

A weak correlation was found between SDAS and UTS. A significantly stronger correlation was found between the length of Si particles (Si-length) and UTS (see Figure 15). The relationship between UTS and Si-length was supported by the observed crack propagation path, which was primarily along the Si particles (see Figure 16). The SDAS is believed not to have a direct influence on tensile properties,

Si Fe Cu Mn Mg Cr Ni Zn Pb Sn Ti Al SF Alloy 9.14 0.76 2.55 0.33 0.24 0.03 0.02 0.75 0.05 0.03 0.02 Bal. 1.51


19

but relates to local solidification time which affects the time available for the Si particles to grow.

(a) (b)

Figure 15. Correlation between UTS and (a) SDAS and (b) Si-length.

Figure 16. Fracture profile. The thin arrow indicates decohesion while thick arrows

indicate Fe-rich, Chinese script shaped particles. Etched with H2SO4.

In the samples, Fe-rich particles with two types of morphologies were found, Chinese scripts (see Figure 16) and polyhedral sludge (see Figure 17). The proposed stoichiometry for these two compacted morphologies is Al15(Fe,Mn,Cr)3Si2; this gives the theoretical atom per cent (at%) of the constituents as shown in Table 5. The table also shows results from the EDS analysis which confirms the proposed stoichiometry of the Fe-rich particles.

The surface area of Fe-rich particles was measured and related to the UTS and ef, but no correlation between Fe-rich particles in the measured area and tensile properties could be observed.

20

Figure 17. Polyhedral Fe-rich particles (sludge). Etched with H2SO4.

Table 5. EDS analysis of Fe-rich particles. A comparison between theoretical results

and EDS results is shown.

Al Fe + Mn + Cr Si Theoretical results (at%) 75 15 10 Average EDS (at%) 72 17 11

Four alloys were examined in Supplement II. Their chemical composition and denotations are shown in Table 6. The eutectic Si, seen to affect tensile properties, was modified with Sr. For alloys A(1.3) and B(1.8), the Fe:Mn ratios were higher, compared with the alloy in Supplement I. Higher Fe:Mn ratios caused the Fe-rich particles to predominantly grow as β-particles. The fracture profiles of A(1.3) and B(1.8) were examined (see Figure 18), and similar fracture profiles were observed as in samples from Supplement I, in terms of the fracture propagating in connection to elongated particles. Consequently the modification of Si, and promotion of β-particles, shifted the propagation path of the fracture, from Si-particles to β-particles.

Table 6. Chemical composition (wt%) of alloys examined in Supplement II.

Si Fe Cu Mn Mg Cr Ni Zn Pb Sn Sr Al SF Fe/ Mn

A(1.3) 9.30 0.74 2.79 0.25 0.05 0.03 0.04 0.96 0.06 0.03 0.022 Bal. 1.3 3.0 B(1.8) 9.40 1.17 2.77 0.25 0.05 0.03 0.04 0.91 0.05 0.03 0.027 Bal. 1.8 4.7 C(2.8) 9.23 1.29 2.65 0.53 0.04 0.15 0.04 0.86 0.05 0.03 0.025 Bal. 2.8 2.4 D(3.7) 9.30 1.59 2.64 0.80 0.04 0.18 0.04 0.80 0.05 0.03 0.029 Bal. 3.7 2.0


21

(a) (b)

Figure 18. Microstructure and fracture surface (towards the right) of a sample from (a) alloy A(1.3) SDAS 25 and (b) alloy B(1.8) SDAS 25. Etched with H2SO4.

Measurements of the β-particle length in A(1.3) and B(1.8) for two different SDASs was performed. The longest β-particle in 20-30 examined areas were measured. The results and the corresponding UTS and ef are shown in Table 7. The table shows that the slower cooling rate in alloy A(1.3) SDAS 25 resulted in a similar β-particle length as for alloy B(1.8) SDAS 10, having a faster cooling rate but higher Fe content. These two alloys with similar β-particle length were having comparable UTS and ef. Hence, there is a correlation between elongated particles, unmodified Si or β-particles, and tensile properties in cast aluminium alloys. The β-particles, and Fe-rich particles in general, are generally considered the more elusive of the two; hence, the further studies were conducted on the Fe-rich particles.

Table 7. Measurements of β-particle length and standard deviations with

corresponding UTS and ef.

A(1.3)

SDAS 10 A(1.3)

SDAS 25 B(1.8)

SDAS 10 B(1.8)

SDAS 25 Average max β-particle (µm) 40 ± 12 79 ± 26 74 ± 19 160 ± 45

UTS (MPa) 297 ± 14 258 ± 24 262 ± 50 228 ± 9 ef (%) 7.5 ± 2.6 3.4 ± 1.1 3.5 ± 2.4 1.1 ± 0.1

22

3.2 FE-RICH PARTICLES

3.2.1 Effect of Fe-rich particles on fracture (Supplement II)

In Supplement II, the effect on fracture of alloys containing β-particles and α-particles (in the form of sludge) was investigated.

The UTS and ef were related to the SF and as Figure 19 shows. Alloys with the highest SF have the lowest UTS and ef, these two alloys, C(2.8) and D(3.7), also have the lowest Fe:Mn ratio (see Table 6), which could have rendered an increase in tensile properties, owing to more compacted sludge particles rather than β-particles. However, in the samples with sludge, the sludge particles caused an increase in porosity, which degraded tensile properties. The primary sludge particles were believed to have nucleated and/or retained porosity in the gradient solidified samples.

(a) (b)

Figure 19. UTS and ef in relation to the sludge factor.

It has been observed that β-particles provide a preferred crack path, and it has been suggested that the crack propagates along the particle-matrix interphase [63]. In addition, it has also been proposed that the β-particles separate towards the centre, because β-particles grow on the outer surfaces of bifilms, and the centre becomes a “pre-cracked” surface. To study the fracture of β-particles, fracture surfaces were examined in the SEM. Figure 20 a and Figure 20 c show the opposite sides of fractured β-particles, as well as dendrite arms surrounded by β-particles. The β-particles and dendrite arms were identified using EDS (acceleration voltage 20keV) on both fracture surfaces. There were significantly higher Fe readings from the β-particles compared with the dendrite arms (see Figure 21). In higher magnification (Figure 20 b), river marks, indicating cleavage, were observed. On polished samples (Figure 20 d), secondary cracks were propagating towards the centre of the β-particles. These observations show that the β-particles, predominantly, fracture through cleavage, not by particle-matrix decohesion or separation of folded bifilms.


23

(a) (b)

(c) (d)

Figure 20. Opposite fracture surfaces of β-particles ((a) and (c)). A fracture surface of a β-particle with river marks can be seen in (b) and cleavage of β-particles in (d).

24

(a)

(b)

(c)

Figure 21. In EDS analysis of marked areas on the fracture surface (a), Fe was detected in the (b) β-particle and mainly Al and Cu was detected in the (c)

aluminum dendrite.

Figure 22 shows the influence of orientation on the β-particles. The particle oriented parallel to the tensile direction has multiple fractures, while particles oriented transverse do not have similar fractures. The fractures of the parallel β-particle confirmation strain in the area, and an existing “pre-cracked” surface (as in the case of precipitation of β-particles on bifilms) within the transverse β-particle would be expected to have caused a visible separation, which is not observed.

Figure 22. Influence of the orientation of β-particles away from the fracture.


25

Alloys C(2.8) and D(3.7), both with a lower Fe:Mn ratio, had Fe-rich particles mostly in the form of sludge. The fracture of sludge particles was examined in the SEM and optical microscope. Opposite sides of typical appearance of fracture involving sludge are shown in Figure 23. The sludge particles were having multiple fracture paths, as also seen in the optical microscope image (see Figure 23 b). The image shows that cracks in the sludge are mainly oriented parallel to the fracture surface, i.e. cracks appear to be related to the load direction, more than to a predetermined “pre-cracked” surface.

(a) (b)

(c)

Figure 23. Opposite fracture surfaces of α-particles ((a) and (c)). A fracture profile (towards the right) of alloy D(3.8) SDAS25, observed in the optical microscope, is

shown in (b). Etched with H2SO4.

26

3.2.2 Nucleation of Fe-rich particles (Supplement III)

It is of interest to know the nucleation temperature of sludge, foremost in the high pressure die casting industry, so as to avoid sludge formation. Four alloys, with different SFs (see Table 8), were investigated to evaluate the effect of cooling rate on sludge nucleation.

Table 8. Chemical composition (wt%) of alloys in Supplement III.

Si Fe Cu Mn Mg Cr Ni Zn Pb Sn Al SF

Fe/ Mn

A 8.39 0.75 2.60 0.23 0.26 0.02 0.03 1.07 0.05 0.03 Bal. 1.27 3.26 B 8.19 1.25 2.40 0.27 0.25 0.02 0.03 1.02 0.05 0.03 Bal. 1.85 4.63 C 7.72 1.32 2.28 0.58 0.24 0.14 0.03 0.96 0.05 0.03 Bal. 2.90 2.28 D 7.88 1.62 2.15 0.79 0.22 0.16 0.03 0.85 0.04 0.03 Bal. 3.68 2.05

The alloys were tested in a DSC. Typical heat flow curves for alloy D(3.7) in various cooling conditions are shown in Figure 24. (Five cooling rates were tested, three of them which are shown here.) The numbers refer to reactions in Table 9, which also shows that sludge nucleation was not observed by Bäckerud et al. [22] in their TA. Sludge nucleation was detected in the DSC, and the release of heat was seen to increase with SF and cooling rate. This is correlated to an increase in the volume fraction of sludge [64]. Furthermore, the DSC curves show that an increase in cooling rate lowers the sludge nucleation temperature. From the analysis of the four alloys it could be concluded that sludge formation is related to chemical composition; an increase of SF shifting the sludge nucleation towards a higher temperature, regardless of the cooling rate.

Figure 24. Solidification curves of alloy D(3.7), run at different cooling rates.


27

Table 9. Solidification reactions of the B380.1 type alloy, according to Bäckerud et al., cooling rate: 36°C/min [22].

No. Reactions Temp. (°C)

1 Precipitation of primary Fe-rich particles (sludge) - 2 Development of α-Al dendritic network 575 3 Precipitation of pro-eutectic Fe-rich particles 573 4 Main eutectic reaction involving precipitation of Si 564 5 Precipitation of Mg2Si - 6 Precipitation of Al2Cu 505

7 Formation of complex eutectics, containing Al2Cu and Al5Mg8Si2Cu2 492

Nucleation temperatures of the four alloys, at different cooling rates, could thereafter be correlated to sludge formation temperatures reported by Jorstad [37], Gobrech [44] and Warmuzek [65] (see Figure 25). As can be seen, the slope is the same for the three cooling rates shown, i.e. a decrease in the SF will lower the nucleation temperature by the same number for a higher cooling rate as for a lower cooling rate.

Figure 25. Sludge precipitation as a function of the SF by predecessors, and cooling

rate for the alloys studied in the present work. (The dotted line for Gobrechts results is an extrapolation.)

28

The sludge nucleation was confirmed by TA and quenching experiments. The alloys were quenched at four different temperatures and the samples examined in the optical microscope, allowing the growth of the sludge particles to be followed.

The findings can be useful in the high pressure die casting process when setting the holding furnace temperature to avoid sludge, since the melt is cooled during melt transfer from the furnace to the mould.

3.3 LIQUID ALUMINIUM TRANSPORT (SUPPLEMENT IV)

Liquid aluminium transport could be a way of both reducing energy consumption and cleaning the melt. Samples were taken from a thermos before transport at the supplier, and at the foundry (8 hours later) for analysis of chemical composition (see Table 10) and evaluation of RPT samples, by DI and BI (see Figure 26). These measurements did not show any significant difference between samples collected at the supplier and samples collected at the foundry.

The amount of Fe, Mn, and Cr was not lower at the foundry, indicating that there had been no significant nucleation and sedimentation of Fe-rich particles. Possibly this was due to a relatively high temperature, measured at the centre of the thermos, which was 708°C at the foundry, together with an anticipated natural convection.

Table 10. Chemical composition (wt%) of the alloy examined in Supplement IV.

Si Fe Cu Mn Mg Cr Ni Zn Pb Sn Ti Al SF Sup.1 9.40 0.81 2.87 0.36 0.23 0.03 0.02 0.80 0.05 0.04 0.02 Bal. 1.61 Fou.2 9.42 0.83 2.82 0.37 0.22 0.03 0.02 0.80 0.05 0.04 0.02 Bal. 1.64

1Supplier, 2Foundry

(a) (b)

Figure 26. Results from analysing RPT tests: (a) shows the density index (DI) and (b) shows the bifilm index (BI).


29

Materials for tensile testing were collected in Y-shaped moulds at the same time as the RPT samples. Analysis of tensile test results was made using a quality index, Q [66], to consider parameters combined, as follows:

Q = UTS + 0.4 × K × log10(ep × 100) (3)

where K is the strength coefficient and ep is the plastic elongation; both these parameters were extracted from the tensile test data.

The quality index results are given in Figure 27, which shows a tendency of increased quality index for the foundry samples; however, the results are not conclusive.

There was no conclusive decrease or increase in the melt quality when transporting liquid aluminium.

Figure 27. Average quality indices for the samples from the supplier and the

foundry. Error bars show 95% confidence intervals.


31

CHAPTER 4

CONCLUDING REMARKS


This chapter summarizes the work carried out in this thesis and some of the conclusions that have been drawn.

Melt quality measurements in the form of RPT, which have been proposed to relate to the pore formation through the DI, and to bifilms through the BI, have not been able to explain variations in tensile properties in our experiments. A decrease of the length of Si-particles and β-particles has been seen to improve tensile properties, where shorter particles foremost increase the elongation to fracture. SDAS can be an indirect measurement relating to tensile properties, because it is linked to the local solidification time and is hence related to the time available for Si-particles and β-particles to grow in size. The size of β-particles has also been seen to increase with higher Fe content.

Fe-rich particles, predominantly, fracture through cleavage. Fracture patterns show that there has been a physical bonding between surfaces of the Fe-rich particles. The crack orientation in particles were mainly oriented parallel to the fracture; this indicates that there is no preferred crack path within the particles.

The sludge nucleation temperature is closely related to the Fe, Mn, and Cr content, i.e. the SF. An increase in SF will increase the sludge nucleation temperature by the same amount, regardless of the cooling rate. For a set SF, an increase in cooling rate will decrease the sludge nucleation temperature. Moreover, the release of latent heat is increased, which is related to a larger volume fraction sludge, believed to be due mainly to an increase in the number of particles.

Transportation and holding a melt for an extended period of time, 8 hours, did not lead to any conclusive improvement or degradation of tensile properties, neither did the indices extracted from RPT samples, DI and BI, show significant effects of the transport. This suggests that transport of liquid aluminium, which is beneficial for the environment, does not considerably affect the melt quality.


33

CHAPTER 5

FUTURE WORK


In this chapter, ideas are presented on how to increase knowledge in the field of imperfections in recycled aluminium cast alloys.

It is known that properties in the final casting depends on both melt quality and the casting process, and reliable melt quality measurements would be of great help for the casting industry. It has been shown in this thesis that correlation between melt quality, in terms of RPT, and tensile properties is not so straight forward. Therefore future studies should involve analysis of available methods of analysing melt quality, such as; spiral test, PoDFA (Porous disc filtration apparatus)/Prefil test, and Alspek H by Foseco.

A deeper understanding for how Fe-rich particles effect tensile properties, possibly sludge particles in particular, since they have been referred to as detrimental to mechanical properties, however it has not yet been shown. In-situ tensile testing of samples in an SEM equipped with electron backscatter diffraction (EBSD) detector, can be used to study the details about how the microstructural constituents in general and Fe-rich particles in particular affect the tensile properties.

From the DSC experiments the nucleation temperature of sludge was identified for various cooling rates; next step would be to study the growth rate of sludge. The growth rate will be of interest when planning melt handling. Studying growth rate of sludge could be done in-situ, in an x-ray microradiograph.


35

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APPENDED PAPERS

Supplement I A. Bjurenstedt, S. Seifeddine, A. E. W. Jarfors: On the complexity of the relationship between microstructure and tensile properties in cast aluminium. Proceeding at APCMP 2014 and in International Journal of Modern Physics B 29(10&11), 2015, 1540011.

Supplement II A. Bjurenstedt, S. Seifeddine, A. E. W. Jarfors: Effect of Fe-rich particles on tensile properties in recycled cast aluminium. Submitted to Metallurgical and Materials Transactions A.

Supplement III S. Ferraro, A. Bjurenstedt, S. Seifeddine: On the formation of sludge intermetallic particles in secondary aluminum alloys. Accepted for publication in Metallurgical and Materials Transactions A.

Supplement IV A. Bjurenstedt, S. Seifeddine, T. Liljenfors: Assessment of Quality when Delivering Molten Aluminium Alloys Instead of Ingots. Materials Science Forum Volume 765, 2013, pages 266-270.

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