Design of Al-Mg-Si-Mn alloys with Zn, Cr and Sc additions with unique strengthening response
Design von Al-Mg-Si-Mn-Legierungen mit Zn-, Cr- und Sc- Zusätzen mit besonderem Verfestigungsverhalten
Der Technischen Fakultät / der Friedrich-Alexander-Universität Erlangen-Nürnberg
zur Erlangung des Doktorgrades
DOKTOR‐INGENIEUR
vorgelegt von
Oleksandr Trudonoshyn, M.Sc.
aus Bila Tserkva, Ukraine
1
Als Dissertation genehmigt von der technischen Fakultät der Friedrich‐Alexander Universität Erlangen‐Nürnberg
Tag der mündlichen Prüfung: 15.06.2020
Vorsitzender des Promotionsorgans: Prof. Dr.-Ing. habil. Andreas Paul Fröba Gutachter/in: Prof. Dr.‐Ing. habil. Carolin Körner Prof. Dr.‐Ing. Karsten Durst
ii
ACKNOWLEDGMENTS
First of all I would like to thank Prof. Dr.-Ing. habil Carolin Körner for giving me an opportunity to provide current research under her supervision and for the scientific freedom. It is a great honor to work under her supervision. I would also like to thank Peter Randelzhofer for the comprehensive help in all areas of my life during the project period. I would like to express my gratitude to my colleagues from WTM FAU, for creating a great friendly working environment, their technical support and helps. I would also like to thank Nicklas Volz from the Department of Materials Science & Engineering (WW1) University Erlangen- Nürnberg for the help with the TEM investigations and Sebastian Rehm for the help with the partly automated image analysis. I am grateful to Prof. Dr.-Ing. Karsten Durst (Head of Physical Metallurgy Department, Institute of Materials Science, TU Darmstadt) for the willing and motivated second appraisal, as well as his expert discussion contributions to the dissertation. I would also like to thank Prof. Dr. rer. nat. Mathias Göken for the chairmanship of the examination committee. Many thanks to my first supervisor Prof. K. Mykhalenkov, who inspired me to plunge into the world of science as well as all detractors, who motivated me to work harder. Finally, I owe a deep sense of gratitude to my family (N. Trudonoshyna, I. Trudonoshyn, A. Trudonoshyn), all my friends (especially O. Kalashnikova, K. Pryhornytska, Plokhotniuk bros., the Nykonenko family, T. Nekrasov, O. Vergeles, the Vashchuk family, the Petryna-Druzhchenko family, captain Baranov, E. Gershevich, A. Ryabinin, I. BUzhanska) for their constant encouragement throughout my research period. Current research have been founded by German Academic Exchange Service (DAAD) and was done in cooperation with Olena Prach (PhD student from TU-Darmstadt, DAAD scholarship holder). I would like also to thank Olena Prach for the discussion and comprehensive help not only in case of the research project but throughout the entire period of our friendship.
iii
iv
ABSTRACT
An excellent combination of the properties of Al, high strength-to-weight ratio, good formability, good electrical mass conductivity, unique corrosion behavior, and recycling potential make it the essential material for different fields of application. The growing demand for more fuel-efficient and ecological vehicles to reduce energy consumption and air pollution is a challenge for the transport sector. Al is the second most used material of the total weight of the car. Each 1 kg of Al is able to replace about 2 kg of steel or cast iron, more and more types of car parts and components are produced from Al. The automotive industry in Europe has tripled the average amount of Al used in cars during the last three decades. A unique feature of Al alloys is that they can be cast by all known casting technologies. The high pressure die casting (HPDC) is the most useful casting technique; about 50% of the total amount of the castings from light alloys are produced by this method. The aluminium-intensive car body structures have a great demand of the thin-wall HPDC structural parts. However, the currently available HPDC aluminium alloys do not meet all the requirements of car bodies. First of all, the most commonly used aluminium alloys have insufficient levels of ductility that is essential for joining casting parts with sheets and extruded parts. Basically, the consumer properties of HPDC alloys are determined by alloy composition, defect levels and microstructure. Al-Mg-Si alloys are well known as alloys capable of providing an excellent combination of high strength and high ductility levels. The Al-Mg-Si system always attracts the attention of researchers since its solidification and precipitation processes are very complex and sensitive to chemical composition. However, the existing studies on the Al-Mg-Si alloys are mainly focused on the wrought alloys (with low Mg and Si content) and hyper-eutectic casting alloys. Therefore, the development of high-strength, high-ductility Al-Mg-Si based alloys for the HPDC process can be very valuable for improving the quality of automotive components. The present study was carried out to the alloy design for the HPDC process in order to satisfy the requirement of mechanical properties, in particular, ductility for the application in automotive body structure. The effects of Sc, Cr and Zn on the solidification and microstructural evolution and the mechanical properties of hypoeutectic Al-5.7Mg-2.6Si-0.6Mn base alloy have been investigated by the combination of thermodynamic calculations and the experimental validation. A comprehensive literature review of the features of the structure and mechanical properties of commercial Al-Mg-Si wrought and cast alloys is given in Chapter 1. In this chapter also an extensive review of the strengthening methods and especially of the precipitation strengthening as the main one for Al-Mg-Si alloys is represented. Fundamental differences
v between cast and wrought alloys of the Al-Mg-Si system that allow additional alloying of the cast Al-Mg-Si alloys to enhance the precipitation strengthening effect were established. In addition to a literature review on the structure and properties of the studied system, Chapter 2 provides an overview of the analysis of the Al-Mg-Si phase diagram, as well as the possible effects of various alloying elements and their concentrations on the studied system. To select alloy compositions for research, phase diagrams and also solidification curves of alloys were calculated by Thermo-Calc software with TCAl2:Al-alloys v2.1 database. Chapter 4 deals with the detailed characterization of the microstructure with scanning electron microscopy, the structure of α-Al dendrites with transmission electron microscopy and common (hardness, tensile test) and local (microhardness) mechanical properties of the studied alloys in as-cast state. Chapter 4 describes the effects of Sc, Cr and Zn on the solidification behavior, microstructural changes, the relationship between the mechanical properties and the microstructure. In addition to the formation of new intermetallics, the Zn addition leads to the formation of nanosized strengthening precipitates in the α-Al dendrites. Zn-containing precipitates can form even in the as-cast state that promote significant strengthening effects. Chapter 5 deals with the detailed characterization of the alloys after heat treatment. Two types of heat treatment were applied to the studied alloys: artificial aging from as-cast state (one- step heat treatment) and solution treatment with subsequent artificial aging (two-step heat treatment). Cr-containing alloys did not show any significant differences in the mechanical properties in comparision to the base alloy. Sc-containing alloys showed the most prominent increase in strength after one-step heat treatment. Zn-containing alloys showed the most interesting combination of the properties after two-step heat treatment. Differences in the character of changes in the mechanical properties were explained based on the TEM results and the strengthening mechanisms in the designed alloys. It has been shown that in the states in which alloys have the highest values of hardness and strength, the structure of α-Al dendrites contains a significant amount of nanoscale precipitates. Current research was founded by German Academic Exchange Service (DAAD) and was done in cooperation with TU-Darmstadt.
vi
ZUSAMMENFASSUNG
Die hervorragende Eigenschaftskombination von Al, hohe spezifische Festigkeit, gute Verformbarkeit, gute elektrische Leitfähigkeit, einzigartiges Korrosionsverhalten und das Recyclingpotenzial machen es zum unverzichtbaren Material für viele Anwendungsbereiche. Die wachsende Nachfrage nach effizienteren und umweltfreundlicheren Fahrzeugen zur Reduzierung des Energieverbrauchs und der Luftverschmutzung ist eine wesentliche Herausforderung für den Industriebereich Automotive. Al ist das am zweithäufigsten verwendete Material gemessen am Gesamtgewichts des Autosmobils. Je 1 kg Al können etwa 2 kg Stahl oder Gusseisen ersetzt werden, mehr und mehr Autoteile und -komponenten werden aus Al hergestellt. Ein einzigartiges Merkmal von Al-Legierungen ist, dass sie mit allen bekannten Gusstechnologien gegossen werden können. Druckgießen (HPDC) ist die gebräuchlichste Gießtechnik, mit diesem Verfahren werden ca. 50% der Gesamtmenge der Leichtmetall-Gussteile hergestellt. Für aluminiumintensive Karosseriestrukturen besteht ein hoher Bedarf an dünnwandigen Druckguss-Bauteilen. Die derzeit erhältlichen Aluminium-Druckgusslegierungen erfüllen jedoch nicht alle Anforderungen für Karosseriebauteile. Vor allem weisen die am häufigsten verwendeten Aluminiumlegierungen ein unzureichendes Maß an Duktilität auf, was aber für das Verbinden von Gussteilen mit Blechen und extrudierten Teilen wesentlich ist. Grundsätzlich werden Eigenschaften von Druckguss- legierungen wie etwa die Duktilität durch die Zusammensetzung, das Defektniveau und die Mikrostruktur bestimmt. Al-Mg-Si-Legierungen sind als Legierungen bekannt, die eine hervorragende Kombination aus hoher Festigkeit und hoher Duktilität bieten können. Das Al-Mg-Si-System ist interessant für die Forschung, da die Erstarrungs- und Auscheidungsprozesse sehr komplex sind und empfindlich auf die chemische Zusammensetzung reagieren. Die bisherigen Studien zu Al- Mg-Si-Legierungen konzentrierten sich jedoch hauptsächlich auf Knetlegierungen (mit niedrigem Mg- und Si-Gehalt) und hyper-eutektische Gusslegierungen. Daher kann die Entwicklung hochfester, duktiler Druckguss-Legierungen auf Al-Mg-Si-Basis für die Verbesserung von Kraftfahrzeugkomponenten von großem Wert sein. Die vorliegende Arbeit befasst sich mit dem Legierungsdesign von Druck- gusslegierungen, um die Anforderung an mechanische Eigenschaften, insbesondere Duktilität, für die Anwendung in der Karosseriestruktur zu erfüllen. Die Auswirkungen von Sc, Cr und Zn auf die Erstarrung und die Entwicklung der Mikrostruktur sowie die mechanischen Eigenschaften der hypoeutektischen Al-5.7Mg-2.6Si-0.6Mn-Basislegierung wurden durch die Kombination von thermodynamischen Berechnungen und experimentellen Validierungen untersucht. Der Stand des Wissens über Struktur und mechanische Eigenschaften handelsüblicher Al- Mg-Si-Knet- und Gusslegierungen bildet Kapitel 1. In diesem Kapitel wird auch eine umfassende
vii Übersicht über die Härtungsmechanismen und insbesondere über die Ausscheidungshärtung als Hauptmechanismus bei Al-Mg-Si-Legierungen gegeben. Grundlegende Unterschiede zwischen den Guss- und Knetlegierungen des Al-Mg-Si-Systems, die ein zusätzliches Legieren der Al-Mg- Si-Gusslegierungen zur Festigkeitssteigerung durch Ausscheidungshärtung ermöglichen, werden aufgezeigt. Nach der Literaturübersicht über Struktur und Eigenschaften des untersuchten Systems bietet Kapitel 2 einen Überblick über das Al-Mg-Si-Phasendiagramm sowie über die möglichen Auswirkungen verschiedener Legierungselemente und deren Konzentrationen auf das untersuchte System. Zur Auswahl der zu untersuchenden Legierungszusammensetzungen wurden Phasendiagramme sowie Erstarrungskurven mit Thermo-Calc unter Verwendung der Datenbank TCAl2: Al-alloys v2.1 berechnet. Kapitel 4 befasst sich mit der detaillierten Charakterisierung der Mikrostruktur mit Rasterelektronenmikroskopie, der Struktur von α-Al-Dendriten mit Transmissionselektronenmikroskopie und den globalen (Härte, Zugversuch) und lokalen (Mikrohärte) mechanischen Eigenschaften der untersuchten Legierungen im Gusszustand. Dieses Kapitel beschreibt den Einfluss von Sc, Cr und Zn auf die Erstarrung, auf das Gefüge und auf den Zusammenhang zwischen den mechanischen Eigenschaften und dem Gefüge. Neben der Bildung neuer intermetallischer Verbindungen führt die Zugabe von Zn zur Bildung nanoskaliger, Festigkeit steigernder Ausscheidungen in den α-Al-Dendriten. Zn-haltige Ausscheidungen können sich bereits im Gusszustand bilden und signifikante Verfestigungseffekte fördern. Kapitel 5 beschäftigt sich mit der detaillierten Charakterisierung der Legierungen nach der Wärmebehandlung. Bei den untersuchten Legierungen wurden zwei Arten der Wärmebehandlung angewendet: Warmauslagerung im gegossenen Zustand (einstufige Wärmebehandlung) und Lösungsglühung mit anschließender Warmauslagerung (zweistufige Wärmebehandlung). Cr-haltige Legierungen zeigten im Vergleich zur Basislegierung keine signifikanten Unterschiede in den mechanischen Eigenschaften. Sc-haltige Legierungen hatten nach einstufiger Wärmebehandlung die stärkste Zunahme der Festigkeit. Zn-haltige Legierungen wiesen nach zweistufiger Wärmebehandlung die interessanteste Kombination der Eigenschaften auf. Die Änderungen der mechanischen Eigenschaften wurden auf Grundlage der TEM- Untersuchungen und der Härtungsmechanismen in den entworfenen Legierungen erklärt. Es hat sich gezeigt, dass in den Zuständen, in denen Legierungen die höchsten Werte für Härte und Festigkeit aufweisen, die Matrix der α-Al-Dendriten eine signifikante Menge nanoskaliger Ausscheidungen enthält. Die vorliegende Arbeit wurde vom Deutschen Akademischen Austauschdienst gefördert und erfolgte in Kooperation mit der TU Darmstadt.
viii
TABLE OF CONTENTS
ACKNOWLEDGMENTS ...... iii
ABSTRACT...... v
ZUSAMMENFASSUNG ...... vii
TABLE OF CONTENTS ...... ix
INTRODUCTION AND MOTIVATION ...... 1
CHAPTER 1 Fundamentals and literature overview ...... 5
1.1. General characteristic and classification of aluminium alloys ...... 5 1.1.1. Wrought aluminium alloys ...... 5 1.1.2. Casting aluminium alloys ...... 6 1.1.3. High pressure die casting technique ...... 8
1.2. Strengthening mechanisms in aluminium alloys ...... 9 1.2.2. Solid solution strengthening ...... 10 1.2.3. Coarse two-phase strengthening ...... 12 1.2.4. Precipitation strengthening ...... 12
1.3. Features of Al-Mg-Si alloys ...... 16 1.3.1. Wrought alloys of 6XXX series ...... 17 1.3.2. Casting Al-Mg-Si alloys ...... 19 1.3.3. Characterization of strengthening phases in Al-Mg-Si alloys ...... 21
1.4. Summary...... 25
CHAPTER 2 Alloy design...... 27
2.1. Analysis of Al-Mg-Si (-Mn) system phase diagram ...... 27
2.2. Effect of alloying elements and their proportions...... 30 2.2.1. Basic alloying elements Mg, Si and their ratio ...... 30 2.2.2. Common alloying elements on Al-Mg-Si alloys ...... 33 2.2.3. Zn addition to Al-Mg-(Si) alloys ...... 37 2.2.4. Summary ...... 39
2.3. Thermodynamic calculations of Al-Mg-Si-Mn system ...... 40
ix 2.3.1. Al-Mg-Si-Mn base alloy ...... 40 2.3.2. Al-Mg-Si-Mn + Sc ...... 43 2.3.3. Al-Mg-Si-Mn + Cr ...... 44 2.3.4. Al-Mg-Si-Mn + Zn ...... 45 2.3.5. Discussion ...... 47 2.3.6. Summary ...... 49
2.4. Selection of heat treatment parameters for studied alloys ...... 50 2.4.2. Solution treatment ...... 52 2.4.3. Artificial aging ...... 53 2.4.4. Summary ...... 54
CHAPTER 3 Experimental Details ...... 55
3.1. High pressure die casting and alloys composition ...... 55
3.2. Differential scanning calorimetry ...... 56
3.3. Microstructural investigation ...... 57 3.3.1. Light microscopy and specimens preparation ...... 57 3.3.2. Scanning electron microscopy and electron microprobe analysis ...... 57 3.3.3. Transmission electron microscopy ...... 57
3.4. Partly automated image analysis for determination structure parameters ...... 58 3.4.1. Particle size and volume fraction ...... 58 3.4.2. Dendrite arm spacing (DAS) ...... 59 3.4.3. Interlamella spacing (ILS) ...... 59 3.4.4. Strengthening precipitate density ...... 60
3.5. Mechanical tests...... 60 3.5.1. Hardness measurements ...... 60 3.5.2. Tensile Tests ...... 61
3.1. Heat treatment...... 62
CHAPTER 4 As-cast state of Al-Mg-Si-Mn-X alloys ...... 63
4.1. Differential scanning calorimetry ...... 63 4.1.1. Cr and Sc containing alloys ...... 63
x
4.1.2. Zn containing alloys ...... 64
4.2. Microstructure and element distribution ...... 65 4.2.1. Al-Mg-Si-Mn base alloy ...... 65 4.2.2. Al-Mg-Si-Mn with Sc and Cr additions ...... 69 4.2.3. Al-Mg-Si-Mn-Zn alloys ...... 73
4.3. Structure of α-Al dendrites ...... 76
4.4. Mechanical properties ...... 80 4.4.1. Al-Mg-Si-Mn with Sc and Cr additions ...... 81 4.4.2. Al-Mg-Si-Mn-Zn alloys ...... 82
4.5. Discussion...... 83 4.5.1. Influence of the Sc- and Cr- additions to the base alloy in the as-cast state ..... 84 4.5.2. Influence of the Zn- addition to the base alloy in as-cast state ...... 86
4.6. Summary...... 90
CHAPTER 5 Heat treatment of Al-Mg-Si-Mn-X alloys ...... 93
5.1. One-step heat treatment ...... 93 5.1.1. Artificial aging of the base alloy ...... 93 5.1.2. Artificial aging of the Al-Mg-Si-Mn alloys with Sc and Cr additions ...... 95 5.1.3. Structural changes in Al-Mg-Si-Mn-Zn alloys during artificial aging ...... 98 5.1.4. Mechanical properties of Al-Mg-Si-Mn-Zn alloys after artificial aging...... 101
5.2. Two-step heat treatment ...... 104 5.2.1. Two-step heat treatment of the base alloy ...... 104 5.2.2. Two-step heat treatment of the Al-Mg-Si-Mn-Cr alloys ...... 108 5.2.3. Structural changes in Al-Mg-Si-Mn-Zn alloys ...... 110 5.2.4. Mechanical properties of Al-Mg-Si-Mn-Zn alloys ...... 112
5.3. Discussion...... 115 5.3.1. Heat treatment of the Al-Mg-Si-Mn base alloy and after Cr addition ...... 115 5.3.2. Heat treatment of the Al-Mg-Si-Mn-Sc alloys ...... 117 5.3.3. Heat treatment of the Al-Mg-Si-Mn-Zn alloys ...... 118
5.4. Summary...... 121
xi SUMMARY...... 123
REFERENCES...... 125
APPENDICES...... 137
Appendix A. List of acronyms ...... 137
Appendix B. Classification of Al-alloys...... 139
Appendix C. Common Al-alloys for automotive application ...... 141
Appendix D. Commercial Al-Mg-Si alloys...... 142
Appendix E. Short characteristic of the casting methods ...... 144
Appendix F. Designation system of the heat treatment tempers...... 145
Appendix G. Binary equilibrium diagrams ...... 146
Appendix H. Equilibrium phase diagrams of the studied system ...... 147
xii
INTRODUCTION AND MOTIVATION
Al is one of the most widespread metals in the earth crust [1–4]. An excellent combination of the properties of Al, high strength-to-weight ratio, good formability, good electrical mass conductivity, unique corrosion behavior, and recycling potential make it the essential material for different fields of application (Figure 1 a).
electrics, electronics transportation 14% 26% packaging, consumer Heat Exchangers market 9% 19% civil machinery, engineering equipment 26% 10%
Figure 1. Fields of the application of Al alloys [2] (a) and distribution of Al in cars [5] (b)
Figure 2. Al content in europian cars (adapted from [5])
1 Al is the second most used metal in the car design. Due to the fact that 1 kg of Al is able to replace about 2 kg of steel or cast iron, more and more types of car parts and components are produced from Al (Figure 1 b). The automotive industry in Europe has tripled the average amount of Al used in cars during the last three decades (from 50 kg in 1990 to 150 kg in 2016) (Figure 2) [5]. According to European Aluminium [5] the amount of Al used in cars can increase up to 200 kg in 2025. The growth of Al and the automotive sector is closely connected. The consecutive replacement of steel and cast iron parts and components by Al alloys increase production volumes of Al not only in Europe but also in the world (Figure 3) [6,7].
60 Africa Asia (ex. China, in. GCC) China Unreported 50 Oceania South America North America Europe 40
30
20
10 Million Million metrictonnes of Al
0
1975 1977 1979 1981 1983 1985 1987 1989 1991 1993 1995 1997 1999 2001 2003 2005 2007 2009 2011 2013 2015 2017 2019 Figure 3. The volume of Al production in the world [6]
The increasing demand for more fuel-efficient and eco-friendly vehicles in order to achieve efficiency of energy consumption and to reduce air pollution is a challenge for the transport sector. The main directions in the area of aluminum alloys design are [8–10]: Development of new and improvement of of commonly used casting methods. Thereupon last decades bring to life SSM-processes (thixocasting and rheocasting) that requires new appropriate alloys for these processes [8,9]; Improvement of methods of refinement and degassing of alloys during production process that would satisfy all requirements of final products quality and to comply with modern environmental protection trends; Design of new alloys that can provide the basis for new product design. This topic is particularly important for casting alloys because their final properties are only the result of the casting and heat treatment parameters. Notwithstanding the fact that the first two are of vital importance for the improvement of a quality and properties of the cast products, the third one is of great demand for the elaboration of new cost-efficient casting methods and development of new parts for the next generation of cars, trucks, aircrafts, trains or ships. Alloy development and characterization of physical and mechanical properties provides the basis for new product development.
2
In connection with the foregoing, the main objective of the current study is development new alloy composition of the Al-Mg-Si system using the specific features of commercial casting alloys of this system available on the market. The main feature of the commercial alloys (the most common alloy composition AlMg5Si2Mn) is a relatively high strength with simultaneous high ductility that makes these alloys in demand especially in the automotive industry. This combination of properties is caused by the main hardening phase, Mg2Si (CHAPTER 1 provides a detailed analysis). Another feature of the alloys is the presence of Mg in excess that avoids embrittlement of alloys (the absence of free Si prevents formation of brittle silicides). The lack of Si in the solid solution and excess Mg deprive the alloys of the possibility of age strengthening. On the other hand, the composition of the alloy solid solution (close to alloys of Al-Mg system) promotes the prerequisites for additional alloying with a series of elements (used for alloying Al-Mg alloys) and obtaining alloys with the advantages of both Al-Mg-Si and Al-Mg systems. The current study is devoted to the analysis of theoretical premises for the possible further strengthening of alloys, selection of suitable alloying elements, adaptation alloy composition (using thermodynamically calculations) and experimental testing (using conventional methods) to obtain new generation alloys of the Al-Mg-Si system. The current research was done in cooperation with Olena Prach (PhD student from TU- Darmstadt).
3
4
CHAPTER 1 Fundamentals and literature overview
1.1. General characteristic and classification of aluminium alloys
The production of Al products includes three different types of Al: super purity (99.99%), commercial purity (up to 1% of impurities), and alloys. Alloys are used for producing castings or fabricating wrought products. Pure Al doesn’t have high mechanical properties. However, the addition of alloying elements (Cu, Si, Mg, Zn, Mn) can increase the strength of Al and produce an alloy with properties adapted to specific applications. Al alloys are available in two forms: cast and wrought. Wrought products dominate, traditionally accounting for about 80% of all Al used [6,11]. On the other hand, cast Al alloys prevail in automobile production, hence the steady replacement of cast iron components by lighter Al alloy castings increases this cast-to-wrought alloy ratio.
1.1.1. Wrought aluminium alloys
The International Organisation for Standardization (ISO) in 1970 adopted an International Alloy Designation System (IADS) for wrought products which were created by the North American Aluminum Association in 1954 (Appendix B Table B.1) [4,11,12]. The microstructure of the wrought Al alloys consists of the Al solid solution grains and intermetallic phases. The phase composition also depends on the heat treatment temper of the alloy. The number of intermetallic phases in most wrought Al alloys is low (phase can be solvable between solvus and solidus temperatures), which allows them to be dissolved during solution treatment with the formation of a supersaturated solid solution (ssss) of Al (during rapid cooling to RT) and the subsequent formation of nanosized precipitates during aging [13–15]. Taking into account the need for further heat treatment of alloys after their production the wrought Al alloys can be divided into 2 groups: Heat-treatable alloys can be strengthened by special heat treatment modes; Non-heat-treatable alloys cannot be strengthened by heat treatment. Their strength is determined by the chemical composition and can be increased by work hardening. The alloys that constitute heat-treatable Al alloys include alloys from the 2XXX, 6XXX, 7XXX, and some of the 4XXX and 8XXX alloys. The alloys with the highest strength belong to 2XXX (Al-Cu-Mg-(Li)), 7XXX (Al-Zn-Mg-Cu) and some of the 8XXX (Al-(Cu)-Mg-Li) groups with the strength values close to 600 MPa. These alloys are used for the most demanding fields like aircraft industry, due to their highest strength/weight ratio [16–18]. All these alloys belong to
5 age strengthen types in which nanosized Cu-, Mg- and Li-containing strengthening precipitates are formed during heat treatment that involved artificial or natural aging [14–16]. Similarly to aircraft, requirements for structural automotive and marine components depend on the application. However, for automotive and marine industries mostly used alloys are from 5XXX and 6XXX series (Appendix C, Table C.1) [19–21]. Manufacturers provide different applications with specific properties for achieving of which a wide range of Al alloys is used. As already stated, the driving force of the development and implementation of a new Al alloys into automotive applications is the weight-to-strength ratio. Unlike the aircraft, alloys for the structural automotive components must be not only strong enough to carry the required loads but also ductile enough to absorb a significant amount of impact energy in case of an accident [19].
1.1.2. Casting aluminium alloys
The Aluminum Association has adopted the designation system for casting Al alloys similar in some respects to that for the wrought alloys (Appendix B, Table B.2) [11,22]. Casting Al as well as wrought alloys can be divided into 2 groups: heat-treatable and non-heat-treatable. The as-cast microstructure of Al alloys normally shows significant segregation of the alloying elements and consists of α-Al solid solution (grains or dendrites) and intermetallic phases. During alloy solidification and the structure formation, the elements distribute differently (Figure 1.1). A gradual change in the distribution of alloying elements in the dendrite scale (from the center to the edge) is a phenomena of microsegregation, which expresses by the partition coefficient k. The coefficient k is usually defined as the slope of the liquidus over the slope of the solidus lines using equilibrium phase diagrams (k=CS/CL) [23]. For most elements, the coefficient of k in Al is less than one that causes the elements to remain and accumulate in the liquid phase during solidification. With an increasing an element concentration in the liquid phase the liquid composition may reach the eutectic point and, thus, eutectic phases can be formed. Mg is an example of an element with the segregation coefficient k less than 1 (~0.5). Mg enriches a dendritic edge (Figure 1.1 b). In case of Al-Mg-Si alloys, Mg together with Si (~0.1) forms Mg2Si eutectic phase in the interdendritic region (Figure 1.1 b). The alloying element has a little tendency for segregation in the Al and so distributed homogeneously in the α-Al if the k is close to 1 (such as Mn) (Figure 1.1 c). Ti (as well as Zr) is good example of element that has a segregation coefficient in Al higher than 1 (~9). Thus, concentration of Ti is higher in the center of dendrite and lower near dendritic edge (Figure 1.1 d) [23,24].
6
Casting alloys are typically heat treated at T6-mode condition (i.e., solution heat-treated, quenched and aged, see Appendix F). But the casting alloys contain a greater amount of alloying additions than the wrought alloys. So, this leads to the fact, that intermetallics during solution treatment don’t dissolve completely, but can change their morphology [25,26].
a) b)
c) d) Figure 1.1. Element distribution across Al dendrites (Al-5.7Mg-2.6Si-Mn alloy) a) SEM microstructure, b) distribution of Mg (k<1), c) distribution of Mn (k~1), d) distribution of Ti (k>1),
The most demanded alloys for automotive applications stay alloys that are based on the Al- Si-Mg system due to their fluidity with near eutectic point, low coefficient of thermal expansion, high corrosion resistance and good weldability. In this regard, Al-Mg-Si alloys (compared to Al- Si-Mg*, the Mg/Si ratio is shifted to the Mg-rich concentrations) is underestimated and even do not have their own group. On the other hand, several Al-Mg-Si casting alloys (DIN EN 1706) are already existing and used in the industry. Alloys of this system possess high strength and ductility in the as-cast state and belong to age strengthened alloys. In order to maximize the benefits of Al-intensive car body structures, the castings should have equal mechanical properties with the wrought and extruded components. However, the
*Al-Si-Mg system of alloys (casting 5XX.X system) has high Si content and minor Mg addition (<1%). The main strengthening phase is Al-Si eutectic. Al-Si-Mg system has low ductility levels. Al-Mg-Si system of alloys has similar to Si (wrought 6XXX system) or higher (casting alloys) Mg content. The main strengthening phase is Mg2Si in form of strengthening precipitates (wrought alloys) or Al-Mg2Si eutectic (casting alloys). Al-Mg-Si system has high ductility levels. 7 mechanical properties of currently available die-cast alloys (Appendix C, Table C.2) [19,20,27] are not competitive and cannot satisfy the industrial requirement (Figure 1.2). In particular, ductility must be enhanced for use in automotive components. Low ductility levels of the casting alloys result in the decreasing their volume fraction in car bodies. The overall increase in Al in cars is due to the growing amount of wrought alloys. Figure 1.2 shows the percentage of cast and wrought alloys depending on tensile characteristics.
32%68% 12% 3% 49% 26% 74% 10% 2% 54% 22% 78% 9% 1% 55%
12% UTS>300 A<5 14% UTS<300 A<3 17% Other 7% UTS<300 A<15 UTS<280 A>15 6% UTS>300 A<15 6% 16% 14% Outer: 2016 Outer: 2016 13% Middle:2012 Middle:2012 Inner: 2001 Inner: 2001 a) b) Figure 1.2. The change over time in the volume of the cast and wrought parts in the structure of automobiles (a), and the volume of cast and wrought alloys in cars depending on their tensile characteristics (b) [5,28,29]
1.1.3. High pressure die casting technique
The main requirement in the design of the new alloys is their universality. A unique feature of Al alloys is the ability to be produced using all known casting technologies (see Appendix E). High pressure die casting (HPDC) is the most useful casting technique; about half of the total amount of the castings from the light alloys are produced by this method. Particularly used materials for HPDC are Zn, Mg, and Al alloys. The production of castings from Al alloys, in turn, represents up to half of the total production of high pressure die castings. HPDC technique is based on the forced filling a die cavity of a reusable metal mould with the molten melt and the solidification of the casting under the pressure. The HPDC moulds can possess complex shapes with a high accuracy level and repeatability. HPDC has some features, such as high turbulence nascent by the forced shot of the molten metal into the die cavity and the high solidification and cooling rates of the casting. This leads to the presence of internal pores (due to the presence of trapped gases) in the casting. The shrinkage of the molten metal during solidification and other defects, such as oxide skins or cold shuts, also lead to the porosity formation [30,31]. 8
The HPDC technique can provide castings from tens to thousands per hour with high mechanical properties, low surface roughness and dimensions of the finished part as close as possible. The wall thickness can be less than 1 mm, and the mass can range from a few grams to dozens of kilograms. Die casting occupies one of the leading position in the foundry industry. Castings with the closest quality to the finished parts are produced by this method. Two methods of the HPDC technique are usually applied: hot and cold chamber. For Al alloys, the cold chamber is used (Figure 1.3). Al alloys based on the Al-Si system with the addition of Cu, Mg and Mn that possess good fluidity and mechanical properties (such as AlSi9Cu3, AlSi7Mg) are the most common for HPDC technique. These alloys have comparable strength to wrought alloys but significantly lower ductility. The casting alloys with comparable to wrought alloys ductility level will have a great potential. The Al-Mg-Si system of alloys is a Figure 1.3. Cold-chamber die casting machine [32] prospective candidate for HPDC [29,31,33]. The features (see Appendix E) of the die casting processes have become a prerequisite for the further development of new technological processes such as liquid impact (metal) forging (LIF) and semi-solid metal casting (SSM) [30].
1.2. Strengthening mechanisms in aluminium alloys
Any impact that leads to an increase in the resistance to deformation of metals and alloys can be reduced to the mechanisms described in Table 1.1. All the mechanisms listed in the table are present in various Al alloys. Usually, Al alloys are strengthened by several mechanisms at once. Nevertheless, cast alloys cannot be strengthened by work hardening, substructure strengthening, texture strengthening and grain size strengthening (by recrystallization) mechanisms since cast alloys are not exposed to deformation. The final properties of cast Al alloys are only the result of the casting solidification parameters and the following heat treatment parameters. In this regard, the resulting strength properties are the sum of the primary phases strengthening (eutectic and two- phase) and strengthening due to changes in solid solution (solid solution and precipitation). The studied system of alloys can be strengthened by all three main mechanisms: two-phase strengthening, precipitation strengthening, solid-solution strengthening.
9 Table 1.1. Strengthening mechanism in Al alloys [34] Mechanism Obstacles Outline Designation Alloys
Work hardening, ∆R =R =α Gb ρ W Dislocations 1 d 1 √ d substructure strengthening Grain size strengthening, Grain boundary, k' W ∆R2=Rgs= Texture strengthening anisotropy √d ∆R =R =α Gсn W,C Solid solution strengthening Solute atom 3 ss 2 0,5≤n≤1 rm W,C Coherent ∆Ra =R =α γm Precipitation strengthening 4 c 3 eff precipitates l+2r (Fine two-phase m=1…1,5 strengthening) Incoherent Gb r W,C ∆Rb=R =α ln precipitates 4 l 4 l b k'' C Eutectic spheres, ∆Rс =R = Coarse two-phase 4 s particles, fibers √λ strengthening α C and lamellas ∆Rd=R = 5 4 l λ Multi-phase strengthening Two or more W,C ∆Rе =R =(R -R )V phases 4 tf А В В Where G is the shear modulus, b is the Burgers vector; d is the grain size; c is the foreign atom content in at. %, r is radius of spherical precipitates; l is free distance, γeff is the interface energy decisive for cutting; λ is the mean free particle spacing (ILS); VB is the volume fraction of the second phase, RB and RA are the deformation resistances of phases A and B; k', k'' - material-dependent constants, α1-α5 are constants.
1.2.2. Solid solution strengthening
The solid solution strengthening is one of the simplest strengthening mechanisms for metallic materials where solute atoms of the alloying element are dissolved into the lattice of matrix metal (solvent) to form a solid solution. Dissolved solute atoms lead to distortion of the lattice during solidification. There are two basic forms of solid solutions [35] (Figure 1.4): Substitutional solid solution - atoms take the place of a normal atom (1), smaller atoms tense the crystal lattice (2), larger atoms compress the crystal lattice (3); Interstitial solid solution - atoms take positions between the lattice points (4), in case when solute atoms are much Figure 1.4. Solute atoms in the matrix smaller. This type of solid solution strengthening is related to the interaction of the moving dislocations and the solute atoms. Solute atoms make distortions of the crystal lattice that act as "pinning points", which obstruct the moving of dislocations.
10
There are several mechanisms of substitutional strengthening of Al alloys, but the most weighty are two [35–37]: Atomic size mismatch (the sizes of the solute and matrix atoms are different that produced a strain field around the solute atom); Modulus mismatch (the binding forces between the solute and the matrix atoms are different what produces a "hard spot" or "soft spot" in the matrix); The dislocations and the solute atoms in the matrix interrelate in two ways [38]: Dislocation locking: the interaction of stationary dislocations with solutes; Dislocation friction: the interaction of gliding dislocations with solutes. In real conditions, both effects have a contribution to the common strengthening effect. Mg and Mn are the most widely used alloying elements for solid solution strengthening of Al alloys [36]. An interest in the Mg addition is in its high solubility in Al. It is possible to dissolve up to 2 wt.% during regular solidification (permanent mould casting (PM) or HPDC) and increase this value up to 6 wt.% and even more with fast cooling from eutectic temperature (using squeeze casting or quenching after PM or HPDC). Mn doesn’t have such high values of solubility (no more than 2 wt.%), nevertheless, a usual Mn addition to commercial Al alloys is not more than 1%. Table 1.2 represents the solubility of the alloying elements in Al in binary systems.
Table 1.2. Summary of the solubility of some alloying elements in Al at different temperatures [39–41] Temperature, [°C] Solubility, [wt.%] System (e - eutectic; p - peritectic) at e/p point at RT High solubility Al-Zn 382 (e) <70 <2.5 Al-Ag 566 (e) 55.6 <0,2 Al-Mg 450 (e) 17.4 <2 Al-Ge 425 (e) 7.2 <0.5 Al-Cu 548 (e) 5.7 <0.1 Al-Li 600 (e) 4.2 <1 Middle solubility Al-Mn 660 (p) 1.8 <0.1 Al-Si 577 (e) 1.6 <0.1 Al-Ti 665 (p) 1.2 <0.1 Al-Cr 660 (p) 0.8 <0.1 Low solubility A1-V 660 (p) 0.37 - Al-Zr 660 (p) 0.28 - Al-Sc 665 (e) 0.27 - Al-Ni 640 (e) 0.05 - Al-Fe 652 (e) 0.05 - Equilibrium diagrams of the selected binary systems calculated by Thermo-Calc are presented in Appendix G
11 Some elements (for example Mn, Fe, Si [36]) show ambiguous results on the effect on strength properties of alloys. It depends on the chemical composition of the alloys, since many elements are able to form dispersoids* (Mn, Fe, Cr, V) [42–44] and precipitates (Si, Zn, Cu) [45– 47] in the Al matrix that can significantly affect the final properties of the alloys.
1.2.3. Coarse two-phase strengthening
This is an important strengthening mechanism especially for casting alloys. The microstructure of casting alloys may consist of eutectic, eutectoid, peritectic, peritectoid and other coarse phases (in contrast with nanoscale precipitates formed in the matrix, see below) [48]. Coarse second-phase particles affect all bulk properties of alloys. Thus, if the particles have higher Young’s modulus than the matrix, the load will be transferred from the matrix to the particle if the material is subjected to elastic stress, and the strength increases. The distribution of the load has a strong dependency on the shape of the particles and their distribution in the matrix [49]. Coarse second-phase particles also can increase the strength of the alloy by their interaction with dislocations (dislocations bend and loop around the particles by the Orowan mechanism, increase the density of dislocations and lead to more work hardening) [49]. The strength of alloys in which particles form eutectic colonies is influenced partly by Orowan mechanism and also by Hall-Petch strengthening depending on the size of the fine eutectic colonies. Alloys with two-phase strengthening attract the attention of researchers because they have good ductility and energy- absorbing characteristics even if one of the phases is brittle [48]. Despite large particles have a multitude of ways of affecting properties of alloys, they do not affect as strong as fine particles (precipitation strengthening in heat-treated alloys).
1.2.4. Precipitation strengthening
Precipitation strengthening (age-hardening) is the most effective and a very important method of strengthening [35,50]. For precipitation strengthening, aging treatment is applied (at room temperature - natural aging, at elevated temperatures - artificial). During aging the supersaturated solid solution (ssss) decomposes to form finely dispersed precipitates in α-Al matrix. These finely dispersed precipitates have a significant effect on yield and tensile properties and can also influence other properties (e.g. dimensional stability, corrosion resistance) [51,52].
*Dispersoids – type of precipitates in Al-alloys with size 10-500nm in diameter that form during high temperature treatment (400-550ºC) in the presence of Mn, Fe, Cr and Si. The dispersoids is non-strengthening precipitates that increase the resistance to recrystallization and improve fracture toughness of Al-alloys [135,136,170]. It is necessary to differentiate precipitated dispersoids and reinforcing dispersoids (in the particle-reinforced metallic composites) [35] 12
It is also necessary to differentiate precipitation- from dispersion-strengthened metals (particle-reinforced metallic composites) [35]. In the last one strengthening particles are introduced by mixing hard, insoluble second phases in a soft metallic matrix. In precipitation- strengthened metals the strengthening precipitates form with time in the alloy matrix. A precondition for precipitation strengthening is that the second (strengthening) phase must be solvable between solvus and solidus temperatures. Following this condition precipitation can start both with heat treatment (solution treatment with rapid quenching for achieving ssss and following artificial aging) and even at room temperature with the time (natural aging)[53]. The essence of the precipitation strengthening is in obstructing the dislocation motion by those small second-phase particles distributed in the ductile matrix and thus increase the strength and hardness of alloys. The aging process of Al alloys is complicated and is composed of several stages (Table 1.3) [35,54]. The aging process starts with the decomposition of metastable supersaturated solid solution (ssss) to coherent with matrix clusters of the solute atoms of alloying elements (Figure 1.5 a, b), but these clusters still have the same crystal structure as the matrix. These clusters are usually called “zones” rather than “precipitates” in order to underline that these zones have not yet taken the form of precipitated particles. These zones are transition structures and are named Guinier-Preston zones (GP zones) [35,55]. Matsuda et al. [56] presented a comprehensive review of the structure of GP zones in the Al-Mg-Si alloys.
Table 1.3. Some precipitation-hardening systems and precipitation sequences [35,57,58] Precipitation Sequences System Supersaturated Transition structures Aged phases solid solution Al-Cu-(Mg) ssss Clusters, GP-I GP-II, Q′′, S′′ θ′ θ(CuAl2) S′S(Al2CuMg) Al-Zn-Mg ssss Clusters, GP-I GP-II, η′′, T′′ η′ η-MgZn2 T′T-Mg32(Al,Zn)49 Al-Mg-Si-(Cu) ssss Clusters, GP-I GP-II, β′′, Q′′ β′β-Mg2Si θ′ θ(CuAl2) Al-Mg-Ge-(Si) ssss Clusters, GP-I GP-II, β′′ β′β-Mg2Ge β-Mg2Si Al-Li-(Cu) ssss Clusters, GP-I GP-II, δ′′, θ′′ δ′δ(Al3Li) θ′θ(CuAl2) T(CuAl2Li) Al-Ag-(Mg) ssss Clusters, GP-I GP-II, γ′′, T′′ γ′γ(Ag,Al) T′T(Mg32(Ag,Al)49) β(Mg2Al3)
13
a) b) c) d) Figure 1.5. Different crystallographic relationships between matrix and second phase[35]. a) complete coherency; b) coherency with strain, but continuous lattice planes; c) semicoherent, partial continuity of lattice planes; d) incoherent equilibrium precipitate; no continuity of lattice planes across the interface.
The GP zones are very small and have a very small lattice mismatch with the matrix (they are coherent with the matrix). The lattice planes cross the interface continuously. Such coherent interfaces lead to the formation of small elastic strain fields in the matrix. The next stage of the aging process is formation of aged phase (semicoherent and incoherent precipitates). With the growth of the strain fields (formed by GP-zones), their elastic energy reduces by the formation of semicoherent precipitates (Figure 1.5c). Further growth of the precipitates results in the formation of an incoherent interface (Figure 1.5 d) between the precipitate and the matrix) [35]. The strain fields in the matrix surrounding the coherent particles are even more important than precipitates for increasing hardness and strength by inhibiting the dislocations movement in the matrix [35,55]. The dislocations move through the precipitation-strengthened matrix is more difficult because of their interaction with precipitates. There two main interactions between dislocations and precipitates (Figure 1.6): Friedel cutting, Orowan looping.
Figure 1.6. Illustration of interaction mechanisms between dislocations and precipitates: particle looping and particle cutting
14
The first mechanism (Friedel cutting) (Figure 1.6) occurs when the dislocations continuously slide in the matrix through the precipitates. This is possible when the stress of the dislocation motion in the precipitate is comparable to its motion in the matrix. Small coherent particles can be cut by dislocation. If there is an interface between the precipitates and the matrix or an abrupt change in the crystal structure (incoherent particles) cutting is no longer possible. In this case, dislocations have to bend and loop around the precipitates (Figure 1.6). This situation is known as the looping mechanism or the so-called Orowan mechanism. The stress of interaction between dislocations and precipitates depends on the size (radius, r) of the particles, volume fraction (Vf), their mean free spacing (free distance between precipitates, l) and by the nature of the particles themselves. The influence of the shape of precipitates on the alloy strength has not been confirmed so far. However, r, l, and Vf are not independent (if the volume fraction is constant, the spacing of the particles increases with increasing their radius) [48,49]. The relation between them can be expressed by the equation [35]: 4휋푟3 푉 = (2.1) 푓 3푙3 The last equation can be also presented according to [49]: 4휋푟3 푉 = 푁 (2.2) 푓 3푉 where N – the number of precipitates, V – volume (calculated as cube volume)
The stress (in addition to τm – critical shear stress of the matrix material), required to bend and loop a dislocation around the precipitates, is inversely proportional to the average free spacing l of the particles in the matrix or to their size [35,49]: 퐺푏 퐺푏 ∆휏 ≈ ; ∆휏 ≈ (2.3) 2푟 푙 The increase in the stress because of the precipitation strengthening (no particle shear) is [35,49]:
휏 = 휏푚 + ∆휏 (2.4) The stress, which is necessary for particle shear is less than stress, which is necessary for bending the dislocation. It can be written as [35]:
휏 < 휏푚 + ∆휏 (2.5) The original Orowan equation (2.3) was improved by Ashby [59] incorporating the radius of the precipitate: 퐺푏 ∙ ln(푟⁄ ) 2.38휋 2√(1 − 휐) 푏 (2.6) ∆휏 = 푙 where υ is Poisson's ratio of the matrix
15 An interesting feature of strength evolution in precipitation strengthening alloys is its strong dependency on the size of precipitates. Thus, small-size precipitates (coherent with the matrix) are cut by dislocations (Friedel cutting); large-size precipitates (incoherent with the matrix) are bypassed by dislocations (Orowan looping). Figure 1.7 represents of a typical one- peak precipitation-strengthening curve. The strength at first stage (underaging) increases; at the second stage (overaging) the strength gradually decreases with increasing size of precipitates, even though the volume fraction of precipitates and the interface energy remain
constant. With the increasing size of Figure 1.7. Schematic aging curve precipitates, the precipitate-dislocation interaction becomes stronger. Precipitates force the dislocations to adopt increasing curvatures before being cut. The number of obstacles increases, which in turn increases strength. Further growth of the precipitates needs solute atoms from the matrix. The growing of large precipitates occurs as a result of the dissolving small precipitates or to absorb smaller precipitates by larger. During coarsening of the precipitates (with fix volume fraction) the free spacing of the precipitates increases until the Orowan mechanism for bypassing the obstacles becomes predominant. Consequently, large-sized incoherent precipitates are disadvantageous because the strength lowers [48,49].
1.3. Features of Al-Mg-Si alloys
The commercial heat-treatable Al alloys are based on ternary or quaternary systems (with some additional alloying elements) according to the solutes involved in developing strengthening precipitates in the Al matrix. Most of the heat-treatable alloys contain Cu, Si, and Zn in combination with Mg. Even small amounts of Mg in the presence of the listed elements force the precipitation strengthening. Al alloys, that belong to the heat-treatable alloys, include 2XXX, 6XXX, 7XXX wrought series and 2XX.X, 3XX.X, 7XX.X casting series. [22,55] In the heat treatable wrought Al alloys (except 2024, 2219, 6066, 7178) the content of alloying elements doesn’t exceed the level of the solubility at temperatures near the eutectic (or peritectic) point. In contrast, most of the casting alloys have concentrations of the alloying elements that close to eutectic concentration. In these alloys, the phases formed during solidification can’t be dissolved during solution treatment [11,22,55].
16
The Al-Mg-Si system is the important group of Al-alloys that is widely used in both cast and wrought forms since their advantageous properties (high strength-weight ratio, excellent plasticity, good formability and corrosion resistance), but more frequently used as wrought 6XXX alloys system in the automotive industry.
1.3.1. Wrought alloys of 6XXX series
Most of Al alloys for forming processes (including 6XXX series), such as extrusion or rolling, are produced using the direct-chill casting. The as-cast material possesses low formability due to the several inhomogeneities [60,61]: [60,61] Microsegregation, Grain boundary segregation, Low melting point eutectics, Brittle intermetallic compounds; These effects can be solved by the solution treatment of the cast ingot [61]: Removing of microsegregation; non-equilibrium fusible eutectics and areas, brittle particles in order to avoid cracking or tearing during hot working; Shape control (spheroidization) of refractory needle-shaped particles with sharp edges (such as Fe-based intermetallics) that reduce ductility; Formation of dispersoids that improve grain size control during extrusion or rolling; Homogenization of the elements distribution along the grain. For these reasons, wrought 6XXX alloys are formed in the T4 temper (solution treatment without aging) and then strengthened (aged) after forming to full T6 properties. The plastic deformation of 6XXX is carried out by metalworking processes that can be divided into several types according to the applied force [60]: Rolling and forging (direct compression type); Extrusion and deep drawing (indirect compression type); Stretch forming (tension-type). The metalworking is also classified according to the applied temperature of plastic deformation: hot working; cold working [60]. Wrought Al–Mg–Si alloys usually contain Si and Mg in the proportions sufficient for formation β”, β’-Mg2Si precipitates, that defines them as heat treatable alloys. The wrought alloys usually have a Mg:Si ratio close to 1 and a concentration of Si less than the solubility at the eutectic point (c < 1.6 wt. %, but most of the alloys have concentration 0.5…1 wt. % Si).
17 Also, alloys in this series may contain significant levels of Cu: (Mg, Si) > Cu and Cu < 1.2 wt. % (6066, but most of the alloys have concentration 0.1…0.4 wt. % Cu). Many alloys in this class contain either Mn or Cr for increased strength and control of grain size (Appendix D, Table D.1) [43,62–65]. Depending on the Mg and Si content the desirable mechanical properties of these Figure 1.8. Composition vs. yield strength (T6) alloys may vary according to Figure 1.8. of the 6XXX alloys [4,61] Wrought Al–Mg–Si alloys are in general stronger than 5XXX Al-Mg (non-heat treatable) alloys, but not as strong as 2XXX and 7XXX alloys. Nevertheless, 6XXX series alloys have good castability, formability, machinability and weldability. Furthermore, alloys of this group belong to the most ductile alloys (see Appendix B, Table B.1) [65–67]. Al alloys of 6XXX series have good corrosion resistance compared to Cu or Zn containing high strength Al-alloys (2XXX, 7XXX) [66,67]. However, they can be subjected to intergranular corrosion caused by Si and Cu depletion, the formation of Si and Cu-containing strengthening precipitates or due to the anodic dissolution of the Mg2Si phase along grain boundaries [64,67–69]. The typical as-cast structure of 6XXX series of Al alloys (based on the chemical compositions) consists of a β-AlFeSi, α-AlFeMnSi intermetallics and small amounts of Mg2Si particles distributed along the cell boundaries. Some alloys also contain Cu and Cr containing particles [70–74]. During the hot working of 6XXX alloys, the coarse intermetallic phases reduce and align along the plastic flow direction, that leads to the band structure formation. During homogenization of alloys a coarse needle-shaped β-AlFeSi phase transforms in a stable α-
AlFeMnSi phase and the β-Mg2Si particles are dissolved. During decomposition of a supersaturated solid solution (in the aging process), the dissolved β-Mg2Si particles precipitate as fine-scale precipitates [60,70,74]. The typical microstructure changes during treatment of alloys are presented in Figure 1.9 by the example of the structure of 6082 alloy. There are works [75–77] on the possibility of natural aging of the alloys of the Al-Mg-Si system, however, the results are controversial because of many complex processes that occur at the same time, but affect properties differently [75–77]. Based on the typical chemical composition, production process and methods of heat treatment, 6XXX series alloys can be strengthened mainly by means of 3 strengthening
18 mechanisms [60,78,79]: grain-boundary strengthening; work hardening; precipitation strengthening.
a) b)
c) d) Figure 1.9. The microstructure of the 6082 alloy a) as-cast state; b) after hot extrusion c) after homogenization at 570ºC/6 h and cooling in water d) after aging process [70]
1.3.2. Casting Al-Mg-Si alloys
The ductility of Al cast components for car body structures requires thin wall die castings with at least 15% of elongation. To achieve such levels of ductility, several critical aspects have to be controlled during production [80]: Alloy composition, Level of gas and impurities in the melt, Level of defects, Casting parameters. It has been found [20,80,81] that HPDC alloys of Al-Mg-Si system are able to have high ductility levels with a good combination of other mechanical properties already in the as-cast state. For the last years several versions of this alloy group have been provided in Europe and North America: Aluminium Rheinfelden produces alloys under the brand name "Magsimal", Salzburger Aluminium Group (SAG) "Maxxalloy" and Rio Tinto Alcan "Aural". The nominal compositions for most of these alloys are almost the same - AlMg5Si2Mn, (Appendix D, Table D.2). Casting Al-Mg-Si alloys have principal differences in chemical composition compared to wrought alloys of 6XXX series (Appendix D). The wrought alloys usually have Mg:Si weight ratio close to 1, and concentration of Si less than the point of solubility in eutectic point. At such
19 concentrations, an insignificant amount of eutectic Mg2Si particles is formed in the structure. At elevated temperatures, the particles of Mg2Si dissolve in the solid solution but because of the decrease in solubility at lower temperatures form strengthening Mg2Si precipitates. On the contrary, commercial casting Al-Mg-Si alloys have compositions close to eutectic concentration
(Al-Mg2Si quasi-binary eutectic section), and have Mg:Si weight ratio higher than 1.73 (according to stoichiometric composition of the Mg2Si compound: 63.2 wt. % Mg and 36.8 wt. % Si) [40,61,82,83]. In other words, if the weight ratio Mg to Si is less than 1.73, the alloy belongs to an “excess-Si-type”. The high Si content leads to the formation of metastable phases and to the decreasing of whole set of the mechanical properties in the as-cast state [83,84], that requires the need for further heat treatment [84–86]. The most popular alloys on the market (Appendix D, Table D.2) have Mg:Si ratio at least 2.0 and belongs to “excess-Mg-type”. The Si content in the α-Al solid solution of the casting Al-Mg-Si alloys is insignificant for the formation of β’’ precipitates during aging treatment. In comparison to the most successful used and popular casting alloy AlSi7Mg0.3 (A356), alloys AlMg5Si2Mn have a good combination of strength-ductility in the as-cast state. Thus, the further heat treatment of these alloys is not necessary, however, some alloys also can be heat treated in order to achieve higher elongation (Appendix D, Table D.3) [87,88]. The advantages of the Al-Mg-Si casting alloys can be summarized as follows: Good fluidity that provides the ability to produce very thin-walls (e.g. HPDC); Good corrosion resistance and stress corrosion cracking; Good combination of the levels of strength and ductility in the as-cast condition (that satisfy the requirements of ductile alloys for automotive applications); The typical structure of the Al-Mg-Si casting alloys consists of α-Al dendrites, Al-
Mg2Si eutectic, Mn,Fe-containing phases (Figure 1.10). The Mg:Si ratio in casting alloys promotes around 40 vol. % of the
Mg2Si eutectic fraction. Such amount of
Mg2Si eutectic in the Al-Mg-Si casting alloys leads to the improvement in the corrosion Figure 1.10. Structure of Magsimal 59 [89] resistance and strength in the as-cast state (UTS up to 350 MPa) [89,90]. After analysing and comparing the chemical compositions with the reported mechanical properties, it can be summarized that the increasing Mg and Si content (and correspondingly amount of Mg2Si) increases hardness and strength (the alloy with the highest strength is
20
Thermodur72, containing 8 wt. % Mg, 3 wt. % Si, and the alloy with the lowest strength is Magsimal22 with 1 wt. % Mg and 0.15 wt. % Si). At the same time, the values of the elongation of these alloys are inversely related. Magsimal59, Maxxalloy59, and Aural11 have mean values of strength and elongation. The elongation as well as strength increase with decreasing wall thickness of the casting. Thus, for HPDC 3-mm-thickness plate of Magsimal59, the UTS reaches 360 MPa, the YS 220 MPa, and elongation to a fracture 18% in the as-cast condition, while for 4-mm- thickness plate these values are UTS 320 MPa, YS 170 MPa and elongation 14% [89,91]. Hu et. al. [88] investigated the influence of casting method on the structure and mechanical properties of AlMg5Si2Mn alloy. It was concluded that tensile properties of HPDC alloy are significantly higher than properties of alloy cast in the permanent mould. The difference of the properties between the PM and HPDC alloys can be attributed to the facts that HPDC leads to decreasing the size of the α-Al dendrites, changes the eutectic morphology from plate-like to fibrous, prevents the formation of large shrinkage pores, prevents the formation of metastable β-AlFeSi phase and promotes the formation of compact α-AlMnFeSi phase. That means that the preferred casting method for Al-Mg-Si alloys is HPDC. The continual interest in this system by researchers as well as by manufacturers confirm the high potential of the casting Al-Mg-Si alloys. Thus, alloys with enhanced composition (with adjusted proportions of the main elements) and additional alloying elements are continually appearing on the market: Maxxalloy-Ultra (SAG Aluminium Lend GmbH&Co) with 0.1-0.3 wt. % Cr, possesses one of the highest strength among all commercial Cr-containing casting alloys in as-cast state [92]. Magsimal-59Plus (Rheinfelden Alloys) containing 0.1-0.2 wt. % Zr that shows improved hardness, YS and UTS as compared to original Magsimal-59 alloy [89]. Al-Mg-Si-Mn-Sc alloys (Rheinfelden Alloys), several compositions were patented, however, not yet implemented in the production [93]
1.3.3. Characterization of strengthening phases in Al-Mg-Si alloys
From a scientific point of view Al-Mg-Si alloys are very interesting since the process of precipitation strengthening is very complex and sensitive to the chemical composition and parameters of the heat treatment. The most fundamental studies of dispersion hardening in Al-Mg- Si alloys were made by M. H. Jacobs [94], G. A. Edwards et. al. [51] and D. J. Chakrabarti et. al.
21 [95,96]. The simplified precipitation sequence in the Al-Mg-Si system with conventional heat- treatment [96–99] is: 푠푠푠푠 → 푐푙푢푠푡푒푟푠, 퐺푃퐼 → 퐺푃퐼퐼, 훽′′ → 훽′, 푈1, 푈2, 퐵′ → 훽, 푆푖 The most important and strengthening phases in Al-Mg-Si system of alloys are the very fine fully coherent Guinier-Preston (GP) zones (with diameters about 2.5 nm) and semicoherent needle-shaped β"-precipitates (with a typical size 4x4x50 nm). They are forming during aging at temperatures typically between 150 and 190 °C. The density of the β" in the matrix is about 104/μm3. This is approximately equal to 1% of the volume of material [97–102]. In the Table 1.4, the general characteristics of the precipitates that can be observed in Al-Mg-Si-(Cu) system were summarized (according to studies Chakrabarti et. al. [95,96], Andersen et. al. [97–99] and Matsuda et. al. [103,104]). Table 1.4 presents precipitation sequence of the formation of strengthening phases including the peak age, overaging, and the stable phases. Each of the systems has noticeably different precipitation sequences.
Table 1.4. Precipitation sequence and precipitate structure for different Al-Mg-Si-(Cu) alloy compositions [95,97,105,106] System Peak age Stages of overaging Equilibrium
Al-Mg2Si bal. GPI + β" β' + β" β' β' β Al-Mg2Si-Mg GPI + β" β" + β β" + β'+ β β' + β β Al-Mg2Si-Si GPI + β" β' + U2 U2 + U1 U1 + B' β + (Si) Al-Mg2Si-Cu GPI + β" + L L + β' + β" Q' + β' β' + Q' β + Q GP-zone (Needle) 퐴푙푀푔4푆푖6, 퐶2/푚 a=1480, b=405, c=648, β=105.3 GP-zone (Plate) 푆푖/푀푔 = 1, (푓푐푐 퐿10) a=405 β" (Needle) β' (Needle) β (Plate/cube) 푀푔5푆푖6, 퐶2/푚 푀푔1.8푆푖, 푃63 푀푔2푆푖, 퐹푚3̅푚 a=1516, b=405, c=674, β=105.3 a=b=715, c=405, γ=120 a=635.4 U2 (Needle) U1 (Needle) B' (Lath) Si (Plate) 푀푔퐴푙푆푖, 퐹푛푚푎 푀푔퐴푙2푆푖2, 퐹3̅푚1 푀푔9퐴푙3푆푖7, 푃6̅ 푆푖, 퐹푑3푚 a=675, b=405 a=b=405, c=674 a=1040, c=405 a=543.1 c=794 γ=120 γ=120 L (Lath) Q' (Lath) Q (Lath) 푀푔6퐴푙4푆푖6퐶푢1, 푃6̅ 푀푔8퐴푙4푆푖7퐶푢2, 푃6̅ a=800, c=170 a=910, c=1040 a=1039.32, c=4.0173 γ=120 Space group of the monoclinic crystal system: C2/m Space groups of the hexagonal crystal system: 푃63, 푃6̅ Space groups of the cubic crystal system: 퐹푛푚푎, 퐹3̅푚1, 퐹푚3̅푚, 퐹푑3푚
Atomic clusters. During exposure after heat treatment (e.g. solution treatment) or after hot working (e.g. extrusion) the supersaturated solid solution starts to decompose with formation of clusters from solute elements. For the first time, the formation of the clusters in Al-Mg-Si system was reported by Jacobs et al. [94,107] to explain the two-step aging behavior. As a next step,
22
Edwards et al. [51] described more detailed separate Si- and Mg- clusters, Mg-Si co-clusters and complete precipitation sequence in Al-Mg-Si system. It should be mentioned that the solubility of Si (see Table 1.2) in Al is lower than solubility of Mg. This leads to the formation of the first clusters from the Si atoms. However, clustering of Mg atoms is also possible [108]. The clusters nucleate at quenched-in defects (e.g. vacancies) immediately after quenching even at room temperature (when the vacancy movement becomes very low). Storing specimens at a temperature above 50ºC leads to diffusion of Mg atoms to the Si-containing clusters, and Mg-Si co-clusters and then phases are formed [101,109]. GP-zones and β" precipitates. The next stage is precipitation of the GP zones (or pre-β" precipitates) on the Mg/Si co-clusters. A model for these zones was proposed by Thomas [110]: Atoms of Mg and Si substitute Al atoms with a ratio when they take the same volume. Thus, one atom of Si (r = 0.11 nm) and two of Mg (r = 0.16 nm) can substitute 3 atoms of Al (r = 0.14 nm) with a simple sequence Mg-Si-Mg-Mg-Si-Mg along 110-direction. Usually, two types of GP zones are distinguished. The GP-I type is fully coherent with the size in a range of 1-3 nm. Particles with equal dimensions to the GP-zones apparently have Mg/Si ratios less or close to 1, that is therefore different from the Thomas model [100,101]. The GP-II zones form next after GP-I and are also named β" precipitates. These strengthening phases are needle-shaped with typical dimentions about 4x4x50 nm3. After the aging treatment, the density of the needles has typically reached a level of 104/μm3. The β" phase is fully coherent only along the “needle-axis”. The conventional unit cell is monoclinic, space group C2/m, with the composition Mg5Si6 that is also different from the Thomas model but closer to the stoichiometric Mg2Si [100,101]. However, later study by Andersen [111,112] showed that more energetically favourable compositions are Mg5Al2Si4 and Mg4Al3Si4 with high Al content of Mg sites along the {130} interface. The majority of the listed studies reported that the aged alloys with maximum hardness contain a high density of both GP-I and GP-II zones in the α-Al matrix. β' phase. During heat treatment the β" precipitates dissolve/transform into the β' phase and the last one remains stable during the long-time heat-treatment. The β' has a higher Mg/Si ratio than the β" (Mg1.8Si) but lower than equilibrium β phase. The β'-phase is coherent only along the c-axis (0.405 nm), coarsens fast and forms rods of 10x10x500 nm [102]. Typically, the process of the transformation from one phase to next (e.g. β"β'β or other) is related to a step-type process, where each of existing phases with reaching of the specific temperature firstly dissolves before following-stage phase precipitates [62,113]. However, Andersen et. al. [114,115] reported that the close relationship in atomic structure between these precipitates suggests consecutive structural transformation from one phase to the following.
23 Stable β-phase. The equilibrium β-phase in this system was firstly described by Jacobs [94]. It is the only phase up to now with a well-known structure. It has an equilibrium formula
Mg2Si with the structure of the three-atom merger (Mg-Si-Mg) on the corners and faces of a cube, directed along the diagonals. The β-phases precipitates on grain boundaries and can have morphology as cubic-shaped and platelet-shaped with dimensions from dozens of nm to several μm [101,105]. In the quasi-binary alloys and in the alloys with excess Mg, β-phase can form at low-temperature aging and can co-exist with β', β" precipitates [105,116,117]. The possibility of the co-existing β, β' and β" phases also was confirmed by Matsuda et al. [113] who reported that the β'-phase is not directly transformed into the equilibrium β-phase. The β-phase is nucleated independently at other sites. However, during the further growth of the β-phase, the β'-phase decomposes and dissolves into the matrix. The increase of excess Si content results in the decrease of the density of β-phase or completely inhibit its formation [105,116,117]. Ohmori et al. [116] reported that the cube-shaped phase has the same crystal structure as the platlet β-phase, but with different orientation relationships. Matsuda et al. [103] reported that the cube-shaped phase contains a higher Mg content than equilibrium β-phase and probably has a formula Mg3Si. Both authors assume that cube-shaped β-phase is an intermediate phase between the β'- and the equilibrium β-phase, and plays an important role in the precipitation process.
a) GP-I [108] b) β' and β" [118] c) Stable β-Mg2Si [103]
U2
B'
U1
B
d) U1 [97] f) U1, U2, B'[113] g) β', B' [117] Figure 1.11. TEM microstructure of the α-Al matrix of Al-Mg-Si-(Cu) wrought alloys
24
U1, U2, B' precipitates in Si-rich Al-Mg-Si alloys. In Al-Mg-Si alloys with an excess of Si, several extra phases are precipitated during aging. There are U1, U2, B’ (or Type A, B, and C respectively) and stable Si. U1 is trigonal and has composition MgAl2Si2. U2 is orthorhombic with composition Mg2Al4Si5. It is structurally related to both β’ and β". All phases except Si plates, have acicular morphology, with the longer dimension coherent to the <100> direction and are often found to coexist with β" and β’. The precipitation sequence in Si-rich alloys is quite complex and controversial [101,115]. Marioara et al. [114] concluded that the composition (Mg/Si ratio) of precipitates formed in the alloy correlates well with the alloy’s own Mg/Si ratio and the stability of a precipitate and also highly depends on the alloy composition. Moreover, in the study [98] they reported that precipitates neither GP zones nor any U1, U1, B' form in alloys with excess Mg content and the higher the amount of Mg in the alloy, the less β" is found and more β’, β precipitates are formed. This leads to the fewer values of hardness in heat treated state in comparison with Si-rich alloys (for both wrought and cast Al-Mg-Si alloys) [84,98,114]. L, Q' and Q phases in Cu-containing Al-Mg-Si-(Cu) alloys. L-phase has a lath-shaped morphology and occurs at peak age along with β" (it is known only the lattice parameter but no crystal structure data of L-phase). L phase is replaced by Q' and Q phases in the overaged states. With increasing Cu content the ratio L/β" phase at peak age increases [95,96]. Laughlin and Miao [77] reported that all the Cu-containing alloys of the 6XXX series contain the quaternary Q-phases. The Q' phase is a metastable phase that preceds the formation of equilibrium Q-phase and has the same crystal structure. The Q' phase has a lath-shaped morphology and a hexagonal structure, with the orientation of the long axis being parallel to the <100> direction of Al. The Q' phase has a similar size of crystal lattice as B' phase but contains Cu [106].
1.4. Summary
The following facts have governed the choice of the selection of the casting method: Gravity casting methods are little used in mass production. Moreover, the solidification and cooling rates are too low for formation of fine structures. HPDC is the most useful casting technique, about 50% of the total amount of the castings from the light alloys are produced by this method; The most weighty disadvantage of HPDC (gas porosity in castings) can be minimized with the optimization of the casting parameters; The composition of existing Al-Mg-Si casting alloys is not optimal for LIF or SSM; Despite the available information [119] on the application of casting Al-Mg-Si alloys for extrusion, such data is insufficient, and the preferred technology for these alloys is HPDC [89].
25 The alloys used in this study are based on AlMg5Si2Mn alloys and have been produced using high-pressure die casting machine. Despite the fact that alloys have a number of advantages (good fluidity, high corrosion resistance and stress corrosion cracking, perfect strength-ductility combination in as-cast state), their main disadvantage is the inability to achieve strengthening by heat treatment. In this connection, further improvement of the mechanical properties can be achieved by additional alloying.
26
CHAPTER 2 Alloy design
In the second chapter, the literature review on the studies of equilibrium phase diagrams of Al-Mg-Si-(Mn) system as well as the possible effects of various alloying elements on the system will be discussed. According to the obtained literature analysis, several possible compositions for new casting alloys are proposed, and analysis of multicomponent phase diagrams of their systems are presented. The diagrams represent graphical interpretation of the phase transformation in the alloys as a function of temperature, element concentrations and, less commonly, pressure. The thermodynamic calculations is an effective method to predict the possible set of phases, their compositions in alloying system and solidification behavior on the basis of available databases. The thermodynamic and phase diagram calculations for multicomponent systems were performed using Thermo-Calc (TC) software package with the TCAl2:Al-alloys v2.1 database.
2.1. Analysis of Al-Mg-Si (-Mn) system phase diagram
Due to the fact that wrought Al alloys of the 6XXX (Al-Mg-Si) system are among the most popular materials for the forming processes (e.g. extrusion or rolling) the analysis of the ternary phase diagram, especially its Al corner, is presented in a large number of publications. The most fundamental analysis of the Al-Mg-Si system diagram is presented in the classic work of Mondolfo [40]. Also, the analysis of the Al-reach side near quasi-binary eutectic Al-Mg2Si line was performed by Feufel et al. [120], Raghavan [121], Fan et al. [20,80,122–124] and some others [120,125–127].
The most specific characteristic of the Al-Mg-Si system is a quasi-binary Al-Mg2Si section around 13.9 wt. % Mg2Si (the stoichiometric composition: 63.2 wt. % Mg and 36.8 wt. % Si). In the works [125–127] was reported that the point with zero range of eutectic solidification does not lie on the section with the stoichiometric composition of Mg2Si. Thus, the real quasi- binary section with zero range is shifted to the Mg-side (Figure 2.1 line A-A') with an eutectic temperature of 594°C. Table 2.1 represents reactions in Al-Mg-Si system during solidification. An increasing of Si concentration changes the Mg and Si diffusion in Al. The increasing of the Si content in the Al-Mg-Si system over stoichiometric ratio increases crystallization range of the alloy. Due to the low solubility of Si in Al, even a small amount of excess Si leads to the formation of ternary Al-Mg2Si-Si eutectic (Figure 2.2 a). According to [123] an increase of the Mg content in the Al-Mg-Si system over stoichiometric ratio (wt. % Mg/Si > 1.73) shifts the eutectic reaction to the side with lower Mg2Si content (Al corner) (Figure 2.2 b). The increase of the Mg increase the volume fraction of the
Al+Mg2Si eutectic. Moreover, an excess of Mg rises sharply in the solidification range [123,124].
27
a) b) c) Figure 2.1. a) ternary Al-Mg-Si diagram; b) section Al-Mg2Si; c) section A-A' [127] e'-ternary eutectic reaction of Lα(AI)+Mg2Si+Si; e°- the temperature range of solidification equals zero in binary eutectic reaction, Lα(Al)+Mg2Si; e-аnу point on the binary eutectic line close to the ternary eutectic point, Lα(Al)+Mg2Si+Mg5Al8; A-A'-position of the section of the quasi-binary Al-Mg2Si.
Table 2.1. Invariant reactions in the Al-Mg-Si ternary system [61,125,126] № Reaction Temperature, [C] Phase Composition, [at%] Al Mg Si L 85.3 10.8 3.9 1. L (Al) + Mg2Si 594 (Al) 97.1 2.7 0.2 Mg2Si 0 66.7 33.3 L 81.5 5.4 13.1 (Al) 98.0 0.70 1.3 2. L (Al) + (Si) + Mg Si 557 2 (Si) 0 0 100.0 Mg2Si 0 66.7 33.3 L 46.1 53.8 0.1 3. L γ + Mg2Si 462.5 γ 46.1 53.9 0 Mg2Si 0 66.7 33.3 L 61.0 38.9 0.1 4. L β + Mg2Si 451.2 β 61.1 38.9 0 Mg2Si 0 66.7 33.3 L 64.0 36.3 0.1 (Al) 83.4 16.5 4.010-6 5. L (Al) + β + Mg Si 450 2 β 61.1 38.9 0 Mg2Si 0 66.7 33.3 L 57.4 42.5 0.1 β 61.1 38.9 0 6. L β + γ + Mg Si 449 2 γ 51.9 48.1 0 Mg2Si 0 66.7 33.3 L 30.9 69.0 0.1 (Mg) 11.6 88.4 5.510-5 7. L (Mg) + γ + Mg Si 435.6 2 γ 39.9 60.1 0 Mg2Si 0 66.7 33.3
Fe is the most common impurity and present in all Al alloys. It has a very low solubility in solid Al (~0.04 wt. %). Over the point of solubility Fe forms a series of insoluble intermetallic phases in combination with Al and some other elements (Si, Mn, Cr) that can initiate cracks, acting as stress concentrators and, therefore, must be kept at an absolute minimum in high ductility alloys.
28
As can be seen from Figure 2.3 the Fe-containing phases already appear even when Fe content converges to zero.
a) [123] b) [128] Figure 2.2. The effects of a) excess Si on ternary Al-15Mg2Si-Si system; b) excess Mg on the quasi-binary Al-Mg2Si system.
a) [122] b) [129] Figure 2.3. The equilibrium phase diagrams of calculated by Pandat software a) cross section of Al–5Mg–2Si–xFe. b) cross-section of Al–8Mg2Si–6Mg-xMn
For the Al–5Mg–2Si–xFe system, the calculated diagram shown in Figure 2.3 a can be divided according to different Fe contents [122]:
1) L α-Al + α-AlFeSi + Mg2Si, α-Al is a prior phase, at Fe < 0.22 wt. %;
1*) L α-Al + α-AlFeSi + Mg2Si, α-Al is a prior phase, β-AlFe may exist at 0.22 < Fe < 0.98 wt. %;
2) L β-AlFe + α-AlFeSi + α-Al + Mg2Si, β-AlFe is a prior phase at Fe > 0.98 wt. %. The fourth mandatory element for all commercial Al-Mg-Si casting alloys on the market is Mn. Mn forms intermetallic phases in Al-Mg-Si alloys due to its the low solubility. It is commonly used to inhibit formation of needle-shaped Fe-containing intermetallics and to promote the
29 formation of compact Mn,Fe-containing phases. Several formulas have been reported for the α-
AlFeMnSi intermetallics, such as Al16(Fe,Mn)Si3, Al15(Fe,Mn)3Si2, Al12(Fe,Mn)2Si [129]. It should be mentioned that the β-AlFe phase forms even with a very low Fe content. However, Mn in amount of 0.6 wt. % prevents the β-AlFe phase formation up to 1.4 wt. % Fe [122]. This agrees well with [129] where was reported that the hexagonal- or cubic-shaped α- AlFeMnSi is the most common when the weight ratio is Mn/Fe ≥ 0.5. The increase of Mn in the alloy results in a significant increase in the liquidus temperature
[122]. Similar results were obtained for Al–8Mg2Si–6Mg-xMn system [129]. Thus, the prior phase is changed from α-Al to α-AlFeMnSi when Mn concentration reaches 0.37 wt. % with increasing liquidus temperature. During solidification, the formed α-AlFeMnSi phase has a low volume fraction and is pushed apparently to the solidification front. Thus, the Fe- and Mn-containing intermetallics are located in the interdendritic space [129].
2.2. Effect of alloying elements and their proportions
All alloying elements in the design of Al-alloys can be divided into three groups: Main additions (define alloy groups and determine the main properties), Auxiliary additions (can improve special properties), Impurities (with neutral or negative influence). Depending upon the nature of an alloy, the same elements could play different roles [130]. The limits for the elements, which can form in Al-alloys coarse phases (during solidification), dispersoids (during high-temperature heat treatment) and precipitates (during artificial aging) are set by considerations such as cost, recyclability, and influence on mechanical behavior. The morphology, size and distribution of these particles are controlled by casting method and further thermomechanical processing schedule [62]. Conventional design of Al alloys is subject to the restriction of the solubility of the alloying elements. Only a few elements have the appreciable solubility (Table 1.2) (which are situated near A1 in the periodic table). This restricts the effective volume fraction of any precipitate phase to a value that converges to 1 vol. %. Thus, when creating Al alloys, it is necessary to strive to obtain the maximum possible volume fraction of a dispersed phase (which must be stable and difficult to shear by dislocations) with a crystal structure similar to the Al solid solution, and a low mismatch of the lattice parameters [50].
2.2.1. Basic alloying elements Mg, Si and their ratio
Mg is the main alloying element in the 5XXX series of wrought and 5XX.X of casting Al alloys. The maximum solubility of Mg in Al is 17.4 wt. %, but the Mg content in wrought alloys 30 does not exceed 5.5 wt. %. Casting alloys range from 4 to 10 wt. % Mg. Precipitation of Mg will occur even at room temperature in Al-10Mg casting alloys. Alloys containing less than about 7 wt. % Mg are substantially stable at room temperature, but not at higher temperatures. Mg precipitates mostly at grain boundaries in a form of a highly anodic phase that makes alloys susceptible to intergranular cracking and stress corrosion. The Mg addition significantly increases the strength of Al by solid solution strengthening (each 1 wt.% Mg increases UTS of Al alloy by 50MPa) without excessive decreasing the ductility [22,64]. Si, after Fe, has the highest impurity level in electrolytic commercial Al (0.01 to 0.15 wt. %). It is also the most common addition in casting alloys. Al-Si alloys have perfect castability due to better fluidity and good feeding characteristics. The addition of Si to Al reduces melting temperature and increases the strength of the alloys. Si does not form any chemical compounds with Al, thus, the eutectic composition can be distinguished as α-Al+Si. The Si content depends on the chosen casting process. For the processes with the high solidification rate (e.g. HPDC) and for the thin-wall casting, a usual Si content varies in a range of 8 wt. % to 12 wt. %. Si alone in Al produces a non-heat treatable alloy. Small amounts of Mg added to any Al-Si alloy make it heat treatable [64].
Al and Mg2Si form a quasi-binary-eutectic system. Mg2Si has an FCC (face-centered cubic) crystal structure with 12 atoms in a unit cell, 4 Si atoms at the corners and face-centered position, and 8 Mg atoms, which form a cubic sub-lattice inside the unit cell. The highest amount of Mg2Si that can be dissolved in the α-Al matrix is 1.85%, and it decreases with decreasing the temperature and with increasing excess of Mg or Si content. In the wrought alloys, Si is used with
Mg at levels up to 1.5 wt. % to produce Mg2Si strengthening phase in the 6XXX series of heat treatable alloys [22]. All alloys of the Al-Mg-Si system can be divided into several groups according to the Mg and Si concentrations: 1. Wrought alloys with an excess of Si: The total amount of Mg and Si ≤ 1.5 wt. %, with the Mg/Si ratio close to 1 or with a slight excess of Si. The total amount of Mg and Si ≥ 1.5 wt. % (up to 3 wt. %) and other additions such as 0.3 wt. % Cu. Such compositions increase strength in the T6 temper. Elements such as Mn, Cr, Zr are used for controlling grain structure. These alloys have values of strengths about 70 MPa higher than the first one. The significant excess of Si. An increase of the Si concentration by 0.2% increases the strength of alloys with 0.8 wt. % Mg+Si by about 70 MPa. Higher Si content is not so influential.
31 2. Wrought alloys with an excess of Mg. Mg can increase strength. However, it is of benefit
only at low Mg and Si contents because Mg lowers the solubility of Mg2Si. 3. Casting alloys have a substantial excess of Mg (the usual weight ratio of Mg:Si = 5:2). This leads to an increase in strength without significant reduction in ductility in the as-cast state. In alloys with excess Si, segregation of Si to the boundary leads to the grain boundary fracture in the recrystallized structures. Additions of Mn, Cr, Zr to the wrought alloys counteract the effect of Si by preventing recrystallization during heat treatment. In the works [82,84,131–133] was discussed that the Al-Mg-Si casting alloys with excess of Si (low Mg:Si ratio) have low tensile properties in the as-cast state due to the formation of brittle silicides and/or brittle secondary needle-shaped phases (containing Mn, Fe). However, such alloys show a significant increasing in the mechanical properties after heat treatment [82,84,131]. The excess of Si promotes a high amount of Si-rich clusters, which reduces diffusion. The microstructure becomes finer, the size of precipitates decreases with lowering their volume fraction. With the substantial excess of Si, a small number of the GP-zones can reach a critical size that allows them to grow and to affect hardness [134]. The aging response of the alloys with excess Si is utterly different from the Mg- rich alloys. Hardness of the alloy is low for both Mg:Si ratios (high and low). When the ratio converges to one (but still in excess of Si) a large number of GP-zones are formed, while an excess of Si stabilizes them against dissolution/transformation. This leads to a high level of hardness, but the size of the hardness peak is relatively small. When the Mg:Si ratio converges to 6:5 the alloy composition is closer to the composition of the β" precipitates (Mg5Si6) or U2 (MgAlSi) and a larger fraction of the particles of this type is formed [98]. When the Mg:Si ratio is reached to a factor 2, the GP-zones peaks are absent. More Mg and lattice vacancies are present and GP zones more easily dissolve or transform to β"-precipitates. The excess of Mg decreases the precipitates number density in the alloys. Alloys with an excess of Mg exhibit low strength in T4 temper and artificial aged tempers, but have excellent corrosion resistance. With further increase of Mg:Si ratio during the first stage of aging, no precipitation at all can be observed in the alloys. At peak hardness, the precipitate microstructures are much coarser and their density is 4 to 7 times lower than of the alloys with excess Si. In outline, the higher Mg content, the fewer β" precipitates are formed and more β’ and β can be found. Excess Mg alloys are also considered more recyclable than Si-rich alloys [98,134].
32
2.2.2. Common alloying elements on Al-Mg-Si alloys
Copper. Both cast and wrought Al-Cu alloys can be heat treated. Solution treatment with further artificial aging increases hardness and strength of the alloys and decreases their elongation. Alloys with 4 and 6 wt. % Cu have highest values of the strength, depending upon the influence of other alloying elements. The aging characteristics of binary Al-Cu alloys have been studied better than any other alloy system. However, there are only few commercial binary Al-Cu alloys. Most commercial Al- Cu-based alloys contain other alloying elements. Even a small amount of Mg changes aging characteristics of both cast and wrought Al-Cu alloys. Thus, the artificial aging of the Al-Cu-Mg alloys leads to further increase in strength and hardness, with significant decreasing in elongation. The most applicable high-strength alloy system is Al-Cu-Mg-Mn. The hardness and strength increase with a separate or simultaneous increase in Mg and Mn contents [22,64]. The most common alloying element for 6XXX wrought alloys is Cu. Cu has the highest impact on the strength and hardness of Al-Mg-Si alloys among all common alloying elements. Cu in presence of Mg usually improves precipitation strengthening, refines the alloy microstructure and reduces the negative effects of natural aging. The formed (during the decomposition ssss) Q' precipitates coexist with the β'' precipitates (Table 1.4.) and contribute to the significant strengthening effect [134]. Although the significantly increase in the hardness and strength values, Cu also decreases ductility and corrosion resistance [95,96].
Iron. Manganese. Chromium. Fe is the most common impurity found in Al and alloys. It has a high solubility in molten Al, but its solubility in the solid state is very low (~0.04 wt. %). Over this amount Fe forms in Al a series of insoluble intermetallics in combination with Al and other elements (such as Si, Mn, Cr) that can initiate cracks, acting as stress concentrators and, therefore, must be kept at an absolute minimum in high ductility alloys. In most commercial sand and gravity die casting alloys Fe is limited by 0.8% because higher Fe content reduces feeding characteristics and mechanical properties. Also, small amounts of Fe in Al alloys that are poor in Si, may reduce soldering during die casting [64,83]. The higher the Fe content in the Al-Mg-Si-(Mn) HPDC alloys, the more significantly the ductility reduces. This is followed by a slight increase of the YS. The UTS keeps a similar value when concentration of Fe is less than 0.6 wt. % but decreases significantly with the further increase of Fe contents [44,122]. For most HPDC Al-Si-Mg and Al-Mg-Si alloys, Fe is limited by 0.2 wt. % to achieve high levels of the ductility and toughness. Cr and Mn are sometimes introduced individually or together, in quantities generally of less than 1%, to increases slightly strength by means of solution strengthening (both casting and
33 wrought alloys), to improve the elevated temperature properties of 2XX.X and 3XX.X alloys. Also, Cr and/or Mn are added to the Al-Si-Mg or Al-Mg-Si alloys to change the morphology of platelet/acicular Fe-containing phases to a more compact cubic-shaped, thereby improving mechanical properties, especially ductility, at room temperature [64,83]. Both of them can form dispersoids in Al solid solution [135,136]. Mn increases the temperature of the recrystallization and promotes the fibrous structure formation during hot working. The Mn-containing dispersoids are effective in slowing recovery and prevent grain growth. The Mn precipitates increase the quench sensitivity of heat treatable alloys. Up to the 1.25 wt. % level, Mn is the main alloying addition of the Al-Mn-Mg wrought series. Alloys of this system after the work-hardening have good welding characteristics, high strength and high resistance to corrosion. Additions of Mn to Al-Mg-(Si) casting alloys increase hardness, slightly decrease ductility [22,64,83]. Mn (up to 0.6 wt. %) can be added to Al alloys with high Mg content to replace the Fe and prevent the die-sticking [137]. Mn stimulates the formation of the compact cubic α-AlFeMnSi phase and inhibits the formation of β-AlFe phase in the HPDC Al-Mg-Si alloys. The β-AlFe intermetallics are formed when the Mn/Fe ratio is less than 0.5. The strength slightly increase in the HPDC Al-Mg-Si alloy with Mn addition [122]. Small Mn addition enhances the formation of Mg-Si-Mn-vacancy and the diffusion of the solute atoms in the matrix. They are an effective cluster for the homogeneous nucleation of GP zones and thus, increase the density of the precipitates in the Al-Mg-Si alloys [42]. Wrought Al-Mg, Al-Mg-Si and Al-Mg-Zn alloys usually can be alloyed by Cr on purpose to the grain-size control and to prevent recrystallization during heat treatmen. It can also reduce stress corrosion sensitivity and improve material toughness [135,136]. Cr is usually added to Al alloys in amounts up to 0.35 wt. %. With a higher concentrations Cr forms a coarse intermetallic phase with other elements (such as Mn, Fe, Si) by peritectic reaction [22,64]. Cr has a slow diffusion rate in Al and can form fine dispersoids in the matrix. These dispersoids inhibit nucleation and grain growth. Cr slightly increases the strength by solid solution strengthening and by a finely dispersed phase [42,64,135].
In casting Al-Mg-Si alloys addition of up to 2 wt. % Cr refined the primary Mg2Si crystals as well as reduced ILS of Mg2Si eutectic. Cr addition changes Al-Mg2Si eutectic morphology from lamellae to fine rods. In hypereutectic Al-Mg-Si alloys, even a small amount of Cr provides formation of new coarse Cr-containing intermetallic phase. Cr-induced changes lead to the increased hardness, strength and elongation values [138]. Recently [92], a new Al-Mg-Si-Mn hypoeutectic alloy containing 0.10-0.30 wt. % Cr (Maxxalloy-Ultra SAG GmbH&Co) was developed, which possesses one of the highest hardness and strength levels among all established Cr-containing casting alloys.
34
Alloying elements with low solubility. Ni has very low solid solubility in Al that does not exceed 0.04%. With a higher amount, Ni is present as an insoluble intermetallic phase, usually in combination with Al and Fe. Ni (up to 2 wt. %) increases the strength of high-purity Al but reduces ductility. Binary Al-Ni alloys are no longer in use. However, Ni addition to the cast and wrought Al-Cu-(Mg), Al-Si-(Mg) and Al-Mg-Si alloys can affect strengthening properties at elevated temperatures and reduce the coefficient of expansion [20,64]. In hypereutectic Al-Mg-Si casting alloys Ni forms Al3Ni phase that together with Al-Mg2Si forms Al-Mg2Si-Al3Ni double rod (both
Mg2Si and Al3Ni rod-like rather than flake-like eutectics) ternary eutectic. Such alloys show an increase of the high-temperature tensile strength by 23% [139]. Ti has been frequently used in wrought Al–Zn–Mg and Al-Mg alloys to reduce cracking tendency, and has also been used in the manufacturing of sand and permanent mould casting alloys to refine the primary Al grain structure [64]. But for HPDC as well as for permanent mould Al- Mg-Si casting alloys Ti addition doesn’t lead to significant improvement of yield strength, tensile strength and hardness [10,140,141]. Moreover, an excess of Ti content over the peritectic concentration leads to the formation of coarse Al3Ti crystals that can act as stress concentrators and thus decrease tensile properties [10,83]. B is used as a grain refiner and for improving conductivity of Al alloys by precipitating V, Ti, Cr and Mo (all of which at their usual impurity level are harmful to electrical conductivity in commercial Al alloys). B can be also used alone (in amount 0.004 - 0.1 wt. %) as a grain refiner but is more effective together with Ti. Commercial grain refiners usually contain Ti and B with a ratio 5:l [22,64]. The most of commercially used Al-Mg-Si-Mn alloys [81,89,142,143] contain 0.15-0.2 wt. % Ti and 0.004 wt. % B. The main aim of Zr addition to Al-Mg and Al-Zn-Mg wrought alloys is to replace Ti as a refining additive and to increase the recrystallization temperature. The addition of Zr in casting alloys is useless due to the very slight grain refinement effect (compared to Ti or/and B). Rheinfelden Alloys [89] recently developed alloy Magsimal-59 Plus alloy with the addition of Zr. They reported that the addition of Zr into the AlMg5Si2Mn alloy promotes the formation of Zr- containing nanosized precipitates. Nevertheless, in the works [10,144] the Zr-containing precipitates after conventional heat treatment modes in the structure were not found.
On the other hand, the Zr is also used in conjunction with Sc (in order to form of Al3(Sc,Zr) strengthening precipitates) [145,146]. Alloying with Sc can bring a significant increasing of the mechanical properties of Al–alloys. Even a minor amount (no more than 0.3 wt. % of Sc) can improve the hardness and strength of wrought alloys by 100-150 MPa. Moreover, Sc addition improve the resistance to corrosion and weldability of alloys [147]. It was reported [148] that in
Al-Mg-Si casting alloys Sc can form Al3Sc particles that can act as nucleation particles for α-Al
35 dendrites as well as Al-Mg2Si eutectic, that causes a strong grain refining effect. Moreover, in Al-
Mg-Si alloys Sc modified Al-Mg2Si eutectic from lamellae to fine fibrous [148]. Since the Al-Mg-Si casting alloys are promising for high-temperature applications [139], the transition metals are interesting as additives for further design of high-temperature Al alloys.
A large number of transition metals can form a thermodynamically stable dispersed phase (Al3M trialuminides) with a similar crystal structure to Al and with a low lattice parameter mismatch with the Al. Trialuminides (Al3M-type) have many beneficial characteristics including low density, high elastic modulus, high melting points. The high-symmetry-cubic L12 and related tetragonal D022 and D023 structures are prevalent among the early transition elements (Groups 3 to 5) and elements of later groups have less-symmetry structures (Fe, Co, Ni, Re, Ir) [50]. It should be also mentioned that most of the lanthanides, some of the early actinides and some other elements such as Li (Al3Li, which is a potent strengthening phase in aerospace Al-based alloys) can also form trialuminides.
Sc is the only element from the transition metals that can form thermodynamically stable L12 trialuminide (Al3Sc) with Al.
Alloying elements with high solubility. Ag has an extremely high solid solubility in Al (up to 55%). Ag substantially increases the stress corrosion resistance and the strength of heat-treated and aged Al-Cu-Mg and Al-Zn-Mg alloys [64]. It was reported that there are also opportunities to increase the strength characteristics of Al-Mg-Si alloys with additional alloying by Ag [149]. Binary Al-Li alloys can be strengthened by artificial aging. However, binary alloys of this system are not commercially used. The most widely used Al-Li-base alloys are Al-Mg-Li alloys (especially in the aircraft industry). These alloys can be heat treated to strengths comparable with high-strength Al-Cu-Mg alloys but have lower density and higher modulus. Al-Mg-Li alloys have a high volume fraction of cubic-shaped coherent A13Li precipitates. Li addition also increases the elastic modulus and the fatigue cracks growth resistance maintaining intermediate levels of stress intensity [22,64]. Addition of Li to Al-Mg-Si alloys leads to modification of Al-Mg2Si eutectic from lamellae to fine fibrous. This together with the formation of strengthening precipitates strongly improved tensile properties of alloys [150–152]. Ge has a small industrial application due to its exotic nature. However, it is a prospective alloying element for Al alloys that affects precipitation process in the Al-Mg and Al-Mg-Si alloys. The similarity of the Ge and Si networks and its ability to form Ge containing hardening phases which are isostructural with the β’-phases in the Al-Mg–Si alloy system. When both Si and Ge are present in the system their networks are mixed [149].
36
2.2.3. Zn addition to Al-Mg-(Si) alloys
Zn belongs to the elements with good solubility in Al but deserves more detailed consideration. The addition of Zn alone to Al does not improve the mechanical properties of the Al alloys. Moreover, Zn addition increases susceptibility to cracking. Therefore, Al-Zn alloys are no longer in use. On the other hand, the addition of some other elements to the Al-Zn alloys can improve strength and hardness characteristic of both casting and wrought Al alloys. The most commonly used Zn-containing Al alloys are Al-Zn-Mg and Al-Zn-Mg-Cu that belong to the systems (both casting and wrought) with the highest hardness and strength and are widely used in aerospace and civil industry.
Mg and Zn in Al-alloys form the stable MgZn2 phase. Increasing the MgZn2 concentration from 0.5 to 12 wt. % continuously increases the tensile and yield strength. The addition of Mg in excess (100 and 200%) of that required to form MgZn2 further increases tensile strength [64,83]. Investigation of the complexity of the processes of structure evolution, precipitates formation of Al-Zn-Mg alloys, and their sensitivity to changes in the chemical composition and heat treatment parameters are very popular similarly to Al-Mg-Si alloys. Cu, the main alloying element for the Al-Zn-Mg (7XXX) system as well as for Al-Mg-Si (6XXX) wrought system makes the processes, which occur in the alloys, even more complicated [153–158]. At the Al corner at temperatures between 400ºC and the solidus, five intermetallic phases can occur: binary phase θ-Al2Cu, ternary S-Al2CuMg and three solid solutions (휂 , T, Z) with extended composition ranges containing all four elements. Of these five phases, in commercial 7XXX alloys, only 휂 and T (in Cu-less alloys) and additionally S-phase(in Cu- containing alloys) appear [159]. The strengthening of the Al-Zn-Mg alloys related to the processes of the decomposition of ssss with the formation of a number of strengthening precipitates during artificial aging. The simplified precipitation sequence in the Al-Zn-Mg system is: 푠푠푠푠 → 푐푙푢푠푡푒푟푠 → 퐺푃퐼, 퐺푃퐼퐼 → 휂′, 훵′ → 휂, 훵 The decomposition of ssss at low temperatures (< 70ºC) leads to the formation of GP- zones. The formation of the metastable η' precipitates takes place at higher aging temperatures (120-180ºC). While at still higher temperatures the formation of the equilibrium η phase precipitates can be observed. The highest impact on the hardness and strength of the Al-Zn-Mg alloys have the very fine fully coherent with the α-Al matrix GP-zones (with diameters around 2- 3 nm) and 휂 '-precipitates [156]. In the Table 2.2, the general characteristics of the precipitates that can be observed in Al-Zn-Mg-(Cu) system were summarized. GP-zones. According to the Berg et al. [153] there are two types of GP zones can be found in Al-Zn-Mg system (GPI and GPII). GP-I zones are the major type of GP zones formed during
37 aging at RT (natural aging). This type of GP-zones has a full coherency with the α-Al solid solution, with the internal order of Zn and Al/Mg on {001}Al planes. GP-I zones can be present in the matrix over a wide range of temperatures (from RT to 150ºC). Sha et al. [157] showed that the small GP-I zones are Mg-rich while large GP-I zones have higher Zn content with a composition of Zn/Mg~1. There also was found that Cu atoms can also participate in the structure of GP-I zones.
Table 2.2. Precipitation sequence and precipitate structure for different Al-Zn-Mg-(Cu) alloy compositions [153–158] System Peak age Overaging Equilibrium
Al-Zn2Mg bal. GP + η′ η' + η η
Al-Zn2Mg-Mg GP + η′, T′ η + T′, T η, T
Al- Zn2Mg-Cu GP + η′, T′, S′′ η + T + S' η, T, S η' (discs) η GPI-zone (spherical) 푀푔4푍푛13퐴푙2 P63/mmc 푍푛2푀푔 P63/mmc 푍푛/푀푔 ≈ 1 a=b=0.496; c=1.402 a=b=0.521, c=0.860
GPII-zone (thin layers) T' T 푍푛/푀푔 ≫ 1 푀푔32(퐴푙, 푍푛)49, bcc 푀푔32(퐴푙, 푍푛)49, bcc a=1.435 a=1.461 S" (Lath) S' (Lath) S (Lath) 퐴푙 퐶푢푀푔 Сmсm 퐴푙10퐶푢3푀푔3 Сmсm 퐴푙2퐶푢푀푔 Сmсm 2 a=0.405, b=0.405, c=0.81 a=0.405, b=0.906, c=0.724 а=0.401, b=0.923, с=0.714 Space group of the monoclinic crystal system: Сmсm Space group of the hexagonal crystal system: P63/mmc
GP-II zones are formed during aging at temperatures above 70 ºC. GP-II zones are more stable than GP-I in various conditions. The GP-II zones can be represented as Zn-rich layers on
{111}Al, with internal order in the form of three sets of elongated <110> domains. Generally, both types of GP-zones can be predecessors for the η' precipitates [153,157,160]. η' and η phases. η' precipitates are the main strengthening phases of Al-Zn-Mg as well as Al-Zn-Mg-Cu alloys. η' precipitates are a metastable phase with the morphology of small discs with a thicknesses of few nanometers, and have an intensely faulted structure. The details of the lattice structure of the η' phase remain discussed. However, it is generally established that the η' precipitates have a hexagonal lattice with a = 0.496 nm and c = 1.402 nm and a structure related to the stable η-MgZn2 phase [160,161]. η phase is stable equilibrium phase with the hexagonal lattice with a = 0.521 nm and c =
0.860 nm and the chemical composition of MgZn2. This phase is incoherent with the matrix, so can have numerous orientations. In the peak-age condition, the main precipitate phases are a mixture of η' and η, whilst in overaged conditions (T7) the main precipitate phase is η [160,161].
38
Other sequences involving S (Al2CuMg) and T (Mg3Zn3Al2) can occur in alloys with different compositions but generally do not contribute to strengthening as much as η'. T’ and T phases are formed during low-temperature ageing of alloys with high Mg/Zn ratio. The structure of the T phase is based on a body-centered cubic lattice with a = 1.461 nm. T phase has the nearly constant atomic percentage of Mg and the Al/Zn ratio can vary, which justifies the general formula Mg32(Al, Zn)49. The T’ phase has also a bcc crystal structure with a lattice parameter close to that of the T phase: a = 1.435 nm. The orientation relationships between matrix and precipitates are the same for both phases [162]. In the alloys with Mg/Zn ratios between about 1:2 to 1:3, the η phase is the only stable precipitate at lower temperatures and T and S at the higher (higher than 200°C) ones. Also, a small amount of Cu doesn’t lead to the formation of S-phase series but can replace Zn atoms from the
η-phase and it becomes a formula of Mg(Zn,Cu,Al)2, which means that the atomic percentage of Mg is 1/3 and Mg is replaced by Cu and Al. The actual composition of the quaternary phases η' and η is strongly dependent on the alloy composition and heat treatment modes [163]. Based on the above it can be concluded that Zn is a prospective element for alloying the Al–Mg–Si system (excess of Mg type), since it provides similarly precipitation strengthening. However, the effect of Zn on mechanical properties on the Al-Mg-Si system is ambiguous. Thus, in studies [164–166] no significant change in the mechanical properties of wrought Al-Mg-Si alloys after addition Zn has been achieved. In the works [13,167] small increase in hardness values after Zn addition was achieved. On the other hand in a few recent studies improvements in hardness and strength of the casting Al-Mg-Si-Zn after heat treatment were obtained [168,169].
2.2.4. Summary
As a base alloy, the commercially available Magsimal 59 (Rheinfelden Alloys) with a nominal composition AlMg5Si2Mn (ratio Mg/Si>2) as the most common casting alloy on the market was chosen. Based on the behavior of alloying elements in the α-Al solid solution they can be divided into 4 groups: 1) Elements that do not form strengthening precipitates. They strengthen alloys by solid solution strengthening or by forming primary phases: Mg, Mn (ssss), Si (eutectic); 2) Elements that can form dispersoids in Al during solution treatment: Fe, Mn, Cr; 3) Elements that can form nanosized precipitates in combination with Al. The general
formula of such precipitates Al3M: Ti, Zr, Ni, Sc, Li; 4) Elements that can form nanosized precipitates in combination with Mg in Al-Mg alloys: Cu, Si, Ge, Zn.
39 For further research one element from each of 2-4 groups (the first group is not interesting for current research) were chosen. In case, that Mn and Fe from the second group are the most studied elements for Al-Mg-Si alloys, as the most interesting element Cr was selected. Despite the fact that cast alloys are not subject to recrystallization, the effect of Cr on the current system of alloys has not been studied enough. Sc was selected from the third group as the only element from the transition metals that has thermodynamically stable L12 trialuminide - Al3Sc. Cu is already a major additive to Al-Mg-Si alloys. From the other hand, the possibility of the effect of additional alloying of Zn in casting Al-Mg-Si-Mn alloys on their mechanical properties is underestimated. Due to the low solubility, Cr and Sc were used as microalloying additives with concentrations 0.1 wt.% and 0.2 wt.% of each (similar to Al-Mg system). The Zn concentrations of interest were determined taking into account 2 facts: level of excess Mg and Mg/Zn ratio of the strengthening precipitates. There is around 1 at.% of excess Mg in the commercial alloys with the nominal composition AlMg5Si2Mn. The GP-I zones have the ratio Mg/Zn≈1 and GP-II zones have the ratio Mg/Zn>1. For these reasons several concentrations of Zn in a range 0-2 at.% (0-5 wt.%) were chose (to get the ratio Mg/Zn around 1 and several different composition to prove this assumption).
2.3. Thermodynamic calculations of Al-Mg-Si-Mn system
The temperature range for calculations was set from 800°C to 0°C and pressure was set as 250 bar (according to the pressure used during casting). The X-axis was set with an ascending of alloying element. The Al was set as the remaining element. For the base composition (Al-5.7Mg- 2.6Si-0.6Mn) the effects of Mn, Si and Mg were studied. For the studied alloys the effects of Zn, Sc and Cr in the Al-5.7Mg-2.6Si-0.6Mn system were studied. The volume fraction of each phase and the concentration variation of different elements were calculated using same parameters. The Scheil model was used for calculation solidification curves. The maximum temperature step was set at 1°C. The calculation was stopped when the solid fraction was higher than 99 %.
2.3.1. Al-Mg-Si-Mn base alloy
In order to understand the phase formation in the Al-5.7Mg-2.6Si-0.6Mn alloy the multicomponent equilibrium phase diagrams in the section Al-5.7Mg-2.6Si-(0…2.0)Mn (without and with 0.1 wt.%Fe) and Al-5.7Mg-(0…10.0)Si-0.6Mn were calculated and are shown in Figure 2.4.
40
a) b)
c) d) Figure 2.4. Sectional equilibrium phase diagrams calculated by Thermo-Calc software (TCAl2 database) on the cross-sections: a) Al-5.7Mg-2.6Si-(0…2.0)Mn including metastable Al6Mn and Al12Mn phases; b) Al-5.7Mg-2.6Si-(0…2.0)Mn (only stable phases); c) Al-5.7Mg-2.6Si-(0…2.0)Mn-0.1Fe; d) Al-5.7Mg-(0…10.0)Si-0.6Mn- 0.1Fe
It can be seen from the Figure 2.4 that there are 6 phases present in the system: α-Al, β-
Mg2Si, β-AlMg and three Mn-containing phases: α-AlMnFeSi (usually reported as
Al15(MnFe)3Si), Al6Mn, Al12Mn. However, Fan et al. [122,129] have confirmed by the results of microstructure analysis the absence of the phases Al6Mn and Al12Mn in the structure of the Al- Mg-Si-Mn alloys and suggested that α-AlMnFeSi is equilibrium and stable phase in the current system. Moreover, from the diagram, it can be seen that these phases occur in the low-temperature zone. Li et al. reported that Al6Mn and Al12Mn form a stable phase only in Al-Mn alloys. And if a small amount of Si and Fe is present in the alloy, then α-AlMnFeSi forms as a stable phase [170].
Figure 2.4 b presents a recalculated phase diagram without phases Al6Mn and Al12Mn. The solubility of Mn in the solid α-Al in the current system varies from 0.14 wt. % at 590°C (eutectic temperature) to 0.03 wt. % at 300°C. Before the concentration of Mn reaches 0.42 wt. % the first crystallized phase is α-Al and the crystallization range is stable. When the Mn content is above 0.42 wt. %, the primary α-AlMnFeSi intermetallic phase is formed as a prior phase,
41 followed by the formation of α-Al phase. The liquidus temperature is increased with further increasing Mn content that leads to the increase of the solidification range. The solidification range of the Al-5.7Mg-2.6Si-Mn alloy is increased from 30°C (below 0.42 wt. % Mn) to 40°C at 0.7 wt. % Mn and 120°C at 2.0 wt. % Mn. The presence of even a small amount of Fe (e.g. 0.1 wt.%) in the system significantly changes the appearance of the diagram Figure 2.4 c. Thus, the solubility of Mn in the solid state drops to almost 0. And the liquidus temperature increases even more. The solidification range of the Al-5.7Mg-2.6Si-0.1Fe-Mn alloy is increased from 30°C (below 0.1 wt.% Mn) to 145°C at 0.7 wt. % Mn and 265°C at 2.0 wt. % Mn. As a long solidification range leads to hot tearing formation, the concentrations of Mn and Fe should not be higher than necessary. In general, all alloys in the described range of Mn concentration have the same set of phases: α-AlMnFeSi+α-
Al+β-Mg2Si+β-AlMg. Where, α-AlMgFeSi is more frequently referenced as α-Al15(Mn,Fe)3Si2;
α-Al is a matrix with solute elements; β-Mg2Si appears as Al-Mg2Si eutectic; β-AlMg has a stoichiometry Al3Mg2. Figure 2.4 d shows the multicomponent phase diagram of Al-5.7Mg-(0…10.0)Si-0.6Mn. This section of the diagram can be divided into 3 areas: 1 - with the excess of Mg, 2 – near equilibrium composition; 3 - with the excess of Si. When the Si content is below 2.0 wt. %, the primary α-Al phase is formed as a prior phase, followed by the formation of α-AlMnFeSi phase and then β-Mg2Si. In excess of Mg area, the β-AlMg phase is formed. Alloys in this range have the set of phases: α-Al + α-AlMnFeSi + β-Mg2Si + β-AlMg. The second area lies between 2 wt. % and 3.3 wt. % of Si. The increasing of the Si content inhibits the formation of β-AlMg phase.
Alloys in this range have equilibrium quasi-binary hypoeutectic Al-Mg2Si structure. The first crystallized phase in this area is the α-AlMnFeSi phase, followed by the formation of α-Al. The third area starts at 3.3 wt. % Si and is characterized by the increase of the solidification range. The first crystallized phase is still α-AlMnFeSi. The next two phases are α-Al β-Mg2Si. The last is Si- reach δ-phase. This δ-phase is referenced as δ-Al4(Mn,Fe)Si2. From the literature it is known that this phase is brittle, acicular-shaped and leads to decreasing of mechanical properties of Al-Mg-Si alloys in the as-cast state [82,84]. The change in the volume fraction of the phases depending on the change in the content of the elements is given in Table 2.3. As can be seen, with increasing Mn content in the alloy volume fraction of α-AlMnFeSi increases proportionally (2:3). Obviously, since the α-AlMnFeSi is a prior phase and solidifies first in the system, it captures a part of Si content. Simultaneously, this leads to the decreasing of the volume fraction of Mg2Si and the increasing of the fraction of β-AlMg phase.
42
Table 2.3. Calculated volume fractions of existing phases in the systems, with an increasing amount of one component and solidification range of the system The volume fraction of phases [%] Phases Al-5.7Mg-2.6Si-XMn [wt.%] Al-5.7Mg-XSi-0.6Mn [wt.%] 0.2 0.4 0.6 2.0 0.5 2.0 2.6 3.4 4.0 5.7 α-Al 89.4 88.7 88.0 83.0 83.3 86.1 87.7 88.6 88.4 88.4 β-Mg2Si 7.3 7.1 6.9 5.8 1.1 5.2 6.9 9.5 9.5 9.5 α-AlMnFeSi 0.6 1.1 1.7 5.7 1.9 1.9 1.9 1.9 0.1 - δ-AlSiMnFe ------2.0 2.1 β-AlMg 2.7 3.0 3.3 5.3 13.3 6.7 3.4 - - -
Solidification 33 33 38 110 70 29 40 70 100 90 range [ºC]
With increasing Si in the system, the volume fraction of β-AlMg abruptly decreases and the volume fraction of Mg2Si increases. At the concentration 3.4 wt. % Si the volume fraction of
Mg2Si reaches a maximum, while β-AlMg becomes 0. In the area with the excess of Si the volume fraction of α-AlMnFeSi drops to 0. α-AlMnFeSi transforms to the metastable Si-rich δ-phase.
2.3.2. Al-Mg-Si-Mn + Sc
Figure 2.5 shows the phase diagram of the Al-5.7Mg-2.6Si-0.6Mn-(0…1.0)Sc for the temperature ranges of 450-750°C (with further decrease in temperature no new phases are formed). Sc has very low solubility (see Table 1.2) in Al and forms a new phase in the cast state. The peritectic concentration in the Al-5.7Mg-2.6Si-0.6Mn-Sc system amounts to 0.12 wt. % and takes place at 591°C. During cooling, the solubility of Sc sharply decreases and drops to 0.03 at 450°C and to 0 at RT. Thus, Sc addition to Al-5.7Mg-2.6Si-0.6Mn base composition leads to the formation of Al3Sc intermetallic phases. Before the concentration of Sc reaches 0.13 wt. % the first crystallized phase is α-Al and the crystallization range is stable. When the Sc content increases above 0.13 wt. %, primary Al3Sc intermetallic phase is formed as a prior phase, followed by the formation of α-AlMnFeSi and then
α-Al phase and Mg2Si phases. Sc doesn’t make any other phases in the observed concentration range in excess of Mg condition. Nevertheless, the liquidus temperature is increased with the further increase of the Sc concentration that leads to the increase of the solidification range. The solidification range of the Al-5.7Mg-2.6Si-0.6Mn alloy is increased from 33°C (up to 0.13 wt.% Sc) to 50°C at 0.2 wt. % Sc and 185°C at 1.0 wt. % Sc.
43
Figure 2.5. Sectional equilibrium phase diagram calculated by Thermo-Calc software (TCAl2 database) on the cross-section Al-5.7Mg-2.6Si-0.6Mn-(0…1.0)Sc
The change in the volume fraction of the phases depending on the change in the content of the elements is given in Table 2.4.
Table 2.4. Calculated the volume fraction of existing phases in the systems, with an increasing amount of one component and the solidification range of the system The volume fraction of phases, α-Al-bal. [%] Phases Al-5.7Mg-2.6Si-0.6Mn-XSc Al-5.7Mg-2.6Si-0.6Mn-XCr 0.1 wt. % Sc 0.2 wt. % Sc 0.1 wt. % Cr 0.2 wt. % Cr β-Mg2Si 7.3 7.3 7.3 7.3 α-AlMnFeSi 1.97 1.97 1.97 1.97 Al3Sc 0.24 0.48 - - Al7Cr - - 0.38 0.77
Solidification 33 51 59 84 range [°C]
2.3.3. Al-Mg-Si-Mn + Cr
Figure 2.6 shows the phase diagram of the Al-5.7Mg-2.6Si-0.6Mn-(0…1.0)Cr for the temperature ranges of 450-750°C (with a further decrease in temperature no new phases are formed). Cr has relatively middling solubility in Al and leads to the formation of the new Al45Cr7 intermetallic phase (more frequently referenced as Al7Cr [171]) in the Al-5.7Mg-2.6Si-0.6Mn system. The peritectic reaction in the Al-5.7Mg-2.6Si-0.6Mn-Cr system takes place at 0.35 wt. % Cr. During cooling, the solubility of Cr decreases rapidly and reaches the value 0.09 wt. % at 450°C and to near zero at RT.
44
α-Al crystalizes as first phase before the concentration of Cr reaches 0.06 wt. % (area with a stable solidification range). When the Cr content increases above 0.06 wt. %, primary Al7Cr intermetallic phase is formed as a prior phase, followed by the formation of α-AlMnFeSi and then
α-Al phase and Mg2Si phases. The liquidus temperature is increased with the further increase of the Cr concentration. The solidification range of the Al-5.7Mg-2.6Si-0.6Mn alloy is increased from 33°C (up to 0.06 wt. % Cr) to 84°C at 0.2 wt. % Cr, and 200°C at 1.0 wt. % Cr. The change in the volume fraction of the phases depending on the change in the content of the elements is given in Table 2.4. As it can be seen neither Sc nor Cr addition does not influence any of the intermetallic phases present in the base composition.
Figure 2.6. Sectional equilibrium phase diagram calculated by Thermo-Calc software (TCAl2 database) on the cross-section Al-5.7Mg-2.6Si-0.6Mn-(0…1.0)Cr
2.3.4. Al-Mg-Si-Mn + Zn
Figure 2.7 shows changes in the equilibrium phase diagram on the cross section of Al- 5.7Mg-(0…6)Si-0.6Mn with increasing Zn content (diagram on the section Al-5.7Mg-2.6Si- 0.6Mn-(0…6)Zn see in the Figure 2.11 i). Unlike Sc and Cr, Zn has very good solubility level in the Al at high temperatures (see Table 1.2). From the phase diagram Figure 2.7, it can be seen that Zn in the Al-5.7Mg-2.6Si-0.6Mn alloy is not involved in any high-temperature reactions, but precipitates as T-AlMgZn phase in the solid state. The T-phase (denoted T according to [172], or τ according to [173]) has a large region with the homogeneous composition (Al,Zn)49Mg32 and belongs to the cubic lattice with the space group 퐼푚3̅ [172–174]. In the alloys with a lack of Mg (Si-rich area) another Al,Zn-reach phase is precipitated (Figure 2.7).
45
Figure 2.7. Equilibrium phase diagram of Al5.7Mg-(0…6)Si-0.6Mn-XZn system calculated by Thermo-Calc software (TCAl2 database) on the cross-sections with different Zn concentrations: 0.6, 1.8, 5 wt. %
The Zn addition to the base composition does not significantly affect the formation of the main phases (-Al, -AlMgFeSi, Mg2Si). However, the addition of Zn inhibits the formation of the β-AlMg phase. Also, Figure 2.8 a shows that with increasing Zn content in the alloy concentration of Si in α-AlMnFeSi increases. The non-linearly increase of the volume fraction (Table 2.5) of the T-AlMgZn phase is associated with the changes in its composition simultaneously with the change of the Zn content in the alloy (Figure 2.8 b).
Table 2.5. Calculated the equilibrium volume fraction of existing phases in the Al-5.7Mg- 2.6Si-0.6Mn-(0…6)Zn system and the solidification range of the system Phases The volume fraction of phases [%] 0.6 wt. % Zn 1.2 wt. % Zn 1.8 wt. % Zn 3 wt. % Zn 4 wt. % Zn 5 wt. % Zn α-Al 87.5 87.6 87.3 87.0 86.7 86.6 β-Mg2Si 7.6 7.6 7.6 7.5 7.5 7.5 α-AlMnFeSi 1.9 1.9 1.9 1.9 1.9 1.9 β-AlMg 0.5 - - - - - T-AlMgZn 2.5 2.9 3.2 3.6 3.9 4.0
Solidification 35 39 43 47 52 56 range [°C]
The main reactions during solidification according to equilibrium diagram (Figure 2.7) are: 1) L α-AlMnFeSi, 2) L α-AlMnFeSi + α-Al
3) L α-AlMnFeSi + β-Mg2Si + α-Al,
4) L α-AlMnFeSi + β-Mg2Si + α-Al + T-AlMgZn,
46
5.1) L α-AlMnFeSi + β-Mg2Si + α-Al + T-AlMgZn + β-AlMg;
5.2) L α-AlMnFeSi + β-Mg2Si + α-Al + T-AlMgZn;
5.3) L α-AlMnFeSi + β-Mg2Si + α-Al + δ-AlMnSi +V-(Al,Si)Zn. The solidification curves of the alloys with different Zn contents were calculated using the Scheil model, and are shown in Figure 2.9.
Al Si Fe Mn Al Zn Cu Mg 80 60 Z1 Z2 Z3 Z4 Z5 Z6 Z1 Z2 Z3 Z4 Z5 Z6 55 70 50 60 45 40 50 35 40 30 25 30 20 20 15.1% 15.1% 15
10.6% 11.1% 10 Concentration, Concentration, [at.%] Concentration, Concentration, [at.%] 10 2.3% 2.3% 5 0 0 0 0,2 0,4 0,6 0,8 1 1,2 1,4 1,6 1,8 2 2,2 0 0,2 0,4 0,6 0,8 1 1,2 1,4 1,6 1,8 2 2,2 Zn in alloy, [at.%] Zn in alloy, [at.%] a) b) Figure 2.8. The composition of the α-AlMnFeSi (a) and T-AlMgZn (b) phases calculated by Thermo-Calc software (TCAl2 database)
630 Liq.+α-AlMnFeSi (ΔT=8.5°C) 610 Liq.+α-AlMnFeSi+α-Al (ΔT=28°C) 590 Base
] Liq.+α-AlMnFeSi+α-Al+Mg2Si (ΔT=1.5°C)
570 ºC 550 625 Liq.+α-AlMnFeSi (ΔT=10°C) 620 1.8%Zn 530 Liq.+α-AlMnFeSi+α-Al (ΔT=28°C) Base 615 510 1.8%Zn Liq.+α-AlMnFeSi+α-Al+Mg2Si (ΔT=48°C) Temperature Temperature [ 610 490 5%Zn Liq.+α-AlMnFeSi (ΔT=12.5°C) 605 470 0 0,001 0,002 0,003 0,004 0,005 Liq.+α-AlMnFeSi+α-Al (ΔT=30°C) 450 5%Zn Liq.+α-AlMnFeSi+α-Al+Mg2Si (ΔT=130°C) 0 0,2 0,4 0,6 0,8 1 Solid fraction [mol] Figure 2.9. Effect of Zn on the solidification behavior of the base composition calculated by Thermo-Calc software (TCAl2 database) (in brackets the temperature ranges of each reaction are given)
2.3.5. Discussion
Despite the fact that the modelling of thermodynamic processes is a very common method of predicting the structure of multicomponent alloys, nuances such as non-equilibrium solidification (high-pressure processes); the complexity of the crystallization of intermetallic
47 phases; the software and the databases used in the calculation should always be taken into account during simulation. In order to get a more complete picture, a comparative analysis with the results obtained by Fan et al. [122,128,129,168,175] for Al-Mg-Si system was done (related diagram provided by Fan et al. [122,128,129] are represented in Appendix H). As can be seen from Figure 2.10 and Figure 2.11 (in comparison to Appendix H) the diagrams calculated by different software have a similar appearance with insignificant differences in the values of critical points.
Al-Mg-Si system Al-Mg-Si-Mn system
a) b)
c) d) Figure 2.10. Comparison of phase diagrams for Al-Mg-Si system and Al-Mg-Si-Mn system calculated by Thermo-Calc software (TCAl2 database) in cross-sections: a) Al-5Mg- 2Si-XFe; b) Al-5.7Mg-2.6Si-0.6Mn-XFe; c) Al-XMg-2.4Si; d) Al-XMg-2.6Si-0.6Mn;
The addition of 0.6 wt. % Mn to the Al-Mg-Si system (Figure 2.10 a,b) leads to the formation new α-AlFeMnSi phase that forms as a prior phase above 4 wt. % Mg followed by the formation of α-Al phase. The addition Mn to the alloys with excess Si leads to the replacement of a region α-Al + Si + Mg2Si by α-Al + α-AlFeMnSi + Mg2Si + δ-AlFeMnSi.
48
For the Al-5Mg-2Si-0.6Mn-xFe system, the calculated diagrams represented in Figure 2.10 c,d can be divided into several parts according to the Fe concentration. The phase formation follows [122]:
1) L α-Al + α-AlFeMnSi + Mg2Si, α-Al phase is prior;
2) L α-AlFeMnSi + α-Al + Mg2Si, α-AlFeMnSi phase is prior;
3) L β-AlFe + α-AlFeMnSi + α-Al + Mg2Si, β-AlFe phase is prior at Fe>1.4 wt. %. The increase in the Mn content in the system (Figure 2.10 c) leads to the disappearance of the area of crystallization of α-Al as a prior phase and a significant increase in the solidification range. The positive effect is that all reactions of the formation of β-AlFeSi phase are shifted towards higher Fe concentrations, that expands the zone free from the undesirable β-AlFeSi phase.
a) b) Figure 2.11. Comparison of phase diagrams for Al-Mg-Si-Zn systems calculated by Thermo- Calc software (TCAl2 database) in cross-sections: a) Al-8.0Mg-2.6Si-XZn; b) Al- 5.7Mg-2.6Si-0.6Mn-XZn
Zn slightly decreases both the liquidus and solidus temperatures, expanding the solidification range (Figure 2.11). Zn addition leads to the formation in the structure of a new T- AlMgZn phase. At the same time, phase β-AlMg with reaching of 0.8 wt. % Zn concentration completely disappears. Excess of Mg (Figure 2.11 a) increases the volume fraction of T-AlMgZn phase and promotes solubility of Zn in the alloy. Thus, when the excess of Mg is 3.5 wt. % the solubility of Zn at 100ºC is 0.2 wt. %, while when the excess of Mg is 1.2% (Figure 2.11 b) the solubility of Zn at 100ºC is 0.
2.3.6. Summary
The use of the Thermo-Calc software in the alloy design to predict the phase composition is a convenient tool. Despite the imprecisions in the simulation of multicomponent systems, the results obtained by a Thermo-Calc environment are in good agreement with other similar programs. 49 The selected alloying elements affect the system differently. Thus, even small amounts of Sc and Cr sharply increase the liquidus temperature (without affecting the solidus temperature) that significantly expands the solidification range. In turn, Zn slightly reduces both liquidus and solidus and slightly expands the solidification range.
In all 3 systems, new phases Al3Sc, Al7Cr, Al2Mg3Zn3 are formed. At the same time, Sc and Cr do not affect the formation of the remaining phases while the increasing Zn concentration inhibits the formation of β-AlMg phase, which fully disappears when the Zn concentration achieves 0.79 wt. % level. The maximum amount of Sc that can dissolve in the system at the temperature of the eutectic reaction is 0.12 wt. %. The value for Cr is slightly higher and reaches 0.35 wt, %. Zn has one of the highest solubility in Al among all elements and in the studied system, this value reaches 87 wt. %. The concentration of Fe (in alloys without Mn addition) in the studied system should be tightly controlled to compromise sticking tendency and ductility level. Although the negative effect of the beta phase can be eliminated by increasing the Mn content, however, this increases the solidification range of the alloys, and, thereafter, the porosity level.
2.4. Selection of heat treatment parameters for studied alloys
Alloys with specific mechanical properties can be obtained using various heat treatment modes (such as hardness, strength, fracture toughness, ductility, residual stresses, thermal and dimensional stability, corrosion resistance and stress corrosion cracking). The most common types of heat treatment of cast Al alloys are homogenization, annealing and precipitation strengthening, involving solution heat treatment, quenching, and aging. Aluminum Association developed a designation system of heat treatment modes for both wrought and cast Al alloys (see Appendix F). The temperature of the solution treatment (Figure 2.12) is carried out between the solvus and the solidus temperatures (~450°C - 650°C). The temperatures depend on the alloy composition and the maximum solid solubility of the alloying elements. To avoid local melting (e.g. grain boundaries) heating rate has to be sufficiently slow to dissolve non-equilibrium phases and temperature has to be below the eutectic temperature. The main aim of the solution treatment is to improve workability by dissolving the non- equilibrium phases and brittle particles, providing a more homogeneous structure and formation of supersaturated solid solution (ssss) for further aging. The cooling rate should be high enough to retain the saturated solid solution and a high number of vacancies at room temperature. A low cooling rate could result in the precipitation of coarse particles that are decreasing mechanical and corrosion properties.
50
Figure 2.12. Scheme of heat treatment
Table 2.6 shows the common industrial-used temperatures for heat-treated Al alloys. As can be seen, the highest treatment temperatures are used for the Al-Mg-Si system, while the lowest for Al-Mg-Zn system. This is due to the fact that the dissolution temperature of the coarse β-Mg2Si phase (in Al-Mg-Si alloys) during the solution treatment is significantly higher (up to 593ºC) than the dissolution temperature of the T-phase (in the Al-Mg-Zn alloys; up to 382ºC). A similar situation is with the dissolution of the GP-zones in these alloys [40,51,52]. Heat-treatable Al alloys are usually subjected to aging in some cases immediately after casting (T1, T2, T10 tempers), but mostly after quenching (T3–T9 tempers). The wrought Al alloys can retain ssss and high concentration of the lattice vacancy without quenching by high-rate cooling after cold work at elevated temperatures.
Table 2.6. Industrial-used temperatures of the heat treatment for heat treated Al alloys [55] Temperature [°C] Alloy systems Solution treatment Artificial aging Al-Cu 515-535 170-190 Al-Cu-Mg 495-530 160-205 Al-Mg-Si 520-570 175-205 Al-Mg-Zn 465-510 120-150 Al-Cu-Li 530-540 160-190
Studied alloys belong to Al-Mg-Si system that can promote a good response to the precipitation strengthening (by Mg2Si precipitates) during full T6 treatment if it could be solution treated. However, as it was already noted, the studied alloys contain a significant amount of the
Mg2Si phase, which cannot dissolve at temperatures below solidus. Moreover, these alloys contain an excess of Mg that inhibits of further strengthening during aging. In addition, alloying elements make changes in the phase composition, form new phases, which must be analyzed before applying the heat treatment.
51 2.4.2. Solution treatment
It is well known that Al-Mg-Si alloys belongs to the alloys with one of the highest solidus temperature and can be heat treated using relatively high heat treatment temperatures (Table 2.6) [10,84,176,177]. Despite this, the temperature may be limited due to the features of the chosen casting process. Due entrapping of the gases in the alloy melt during forced filling, it is expected that the HPDC castings contain increased gas contents in comparison with the gravity castings. The normally accepted gas content inside HPDC castings is usually in the range 1 and 12cc/100g [30,31,54]. The gas volume is compressed up to 1000 times during casting under the pressure. The presence of the compressed gaseous phases limits the heat treatment possibilities at high temperatures. The compressed gases expand during heat treatment that leads to unacceptable surface blistering, gas porosities inside the castings and dimensional instability [31,169]. In the studies [31,169] the effect of temperature and heat treatment time on the surface blistering and pores formation in the HPDC Al-Si-Mg-(Cu) and Al-Mg-Si-(Zn) alloys were investigated. It was found that the higher the treatment temperature is, the faster the pores emerge in the specimens. Therefore, the time and temperature of the solution treatment should be controlled taking into account the following [169]: Dissolution of some equilibrium phases, dispersoids and precipitates; Uniform distribution of solute alloying elements; Avoiding the promotion of porosity growth. It has been confirmed [169] that when the specimen's surface has no visible blisters after treatment, the specimen also has an acceptable porosity level. In contrast, if the blisters were visible on the surface the specimen has an unacceptable porosity level. In the studies [31,169] it was also confirmed that the maximum applicable temperature for HPDC Al-Si-Mg-(Cu) and Al- Mg-Si-(Zn) is in the range 480-510ºС. Figure 2.13 a shows a comparison of the surfaces for the base Al-5.7Mg-2.6Si-Mn alloy with different solution treatment temperatures. As can be seen from Figure 2.13 a the only specimens with blistering are after treatment with 560ºС. It is well known that microporosity is more often formed in the middle part of the casting. Thus a similar experiment was conducted for specimens from which the surface was mechanically removed. As can be seen from Figure 2.13b the temperature 540ºС is also inapplicable for heat treatment of the studied alloys. It should be noted that the formation of porosity in the studied alloys is not a subject of the current study and this part refers only to preliminary work to determine the allowable treatment temperatures. From the thermodynamic calculations the required temperatures of the solution treatment can also be determined. Table 2.7 contains the results of solubility of the chosen alloying elements
52 in the current system. Thus, during heat treatment with 520ºС (maximum treatment temperature) the new phases in the alloyed alloys can be dissolved at 0.07 wt. % Sc, 0.21 wt. % Cr and up to 81 wt. % Zn.
a)
b) Figure 2.13. Surface appearances of the alloy Al-5.7Mg-2.6Si-Mn after different solution- treatment temperatures (solution treatment time was 2h): a,b) samples with the surface layer; c,d) samples after removal of the surface
Table 2.7. The solubility of chosen elements in the Al-5.7Mg-2.5Si-0.6Mn-X systems based on calculated diagrams Alloying Solubility at different temperatures [wt.%] element 591ºC 540ºC 520ºC 500ºC 480ºC 460ºC 420ºC 380ºC Sc 0.12 0.09 0.07 0.05 0.03 0.02 - - Cr 0.35 0.25 0.21 0.17 0.12 0.1 0.05 - Zn 87 84 81 78 75 58 10 7
It should be also mentioned, that for the studied system such low solution temperature doesn’t lead to a maximum of the hardness properties but in a combination of further overaging treatment leads to significant improvement of ductility level [178] with decreasing of the strength (in comparison to the as-cast properties).
2.4.3. Artificial aging
Since different alloying elements in Al alloys form different intermetallic compounds that have different properties (such as melting point, distribution across the grains, interaction with
53 dislocations etc.) and provide different mechanisms of strengthening, it is necessary to determine effective temperatures of age treatments for different systems. Table 2.8 summarizes available literature data concerning possible aging treatments in the Al alloys additionally doped by elements that we are interested in. Thus, Cr can form only dispersoids during solutionizing [135,136] and doesn’t form any strengthening precipitates during aging. So, the age hardening of these alloys is not required or aging similar to the Al-Mg-Si alloys can be applied. Sc-containing precipitates form at relatively high aging temperatures (compared to standard precipitates that are formed in the Al-Mg-Si system) [147]. This is a disadvantage for the alloying of Al-Mg-Si alloys. Since with temperatures higher than 250ºС β" and even β' precipitates dissolved with further formation equilibrium β phase [105,116,117]. However, in alloys with the high Mg content a very small amount of β" or β' is formed or even not at all, and the dominant phase is equilibrium β phase [98,134][134]. This means that precipitates are not the main hardening phase in these alloys.
Table 2.8. Temperatures for artificial aging of the alloys doped by chosen elements [ºC] Alloying element Peak hardness Overaged Sc 250-350 - Cr - - Zn 120-150 175-200 Al-Mg-Si 170-200 220-250
Artificial aging of the “excess of Mg” alloys with Zn addition can be carried out at temperatures similar to Al-Zn-Mg alloys. The aging temperatures for the Al-Zn-Mg system of alloys are slightly lower than for the system Al-Mg-Si. This can be a significant advantage since it makes possible to obtain combined effects (for example, use temperatures in the range of 180- 200ºC, which are the peak for the Al-Mg-Si system and overaging for the Al-Zn-Mg).
2.4.4. Summary
Based on the above, Cr-containing series of alloys can be heat treated at 520°C, followed by quenching, to form Cr-containing dispersoids in the a-Al matrix. Further, the alloys can be aged similarly to the Al-Mg-Si alloys. Solutionizing Sc-containing series of alloys does not make sense. However, Sc-containing alloys can be aged (with the formation of nanoscale Al3Sc precipitates in the Al-matrix) in the range of temperatures 250-350°C without previous solution treatment. Zn-containing alloys are most promising for subsequent heat treatment. Due to the low solution treatment and aging temperatures, an excellent combination of mechanical properties can be achieved.
54
CHAPTER 3 Experimental Details
Chapter 3 presents the experimental methods and procedures used in the present study. The present study has been focused on the detailed characterization of the hypoeutectic Al-Mg-Si-Mn alloys with Sc, Cr and Zn additions paying attention to changes in structure and in α-Al solid solution (formation of nanosized precipitates) and related changes in mechanical properties. For structure and element distribution investigation the following methods have been used: light microscopy, scanning electron microscopy (SEM), electron microprobe analysis (EMPA) and transmission electron microscopy (TEM), electron dispersive X-ray analysis (EDX). For investigation of mechanical properties, the following methods have been used: hardness measurements (Brinell), microhardness of α-Al dendrites, eutectic colonies and intermetallic phases (Vickers) and tensile tests.
3.1. High pressure die casting and alloys composition
As a starting material and as a base alloy (M alloy) the commercial Magsimal59® (“Rheinfelden Alloys GmbH”, Rheinfelden, Germany) was used. For achieving the aimed alloy compositions, pure Zn and commercial AlCr20, AlSc2 master alloys were used. To avoid moisture of the starting materials they were preheated at 250C for 24 hours. A vacuum density tester (MK 3VT) in conjunction with a density index balance (MK3000) was used to control the density index after melt degassing. A spark emission spectroscope (Ametek Spectromaxx) was used to measure and control the chemical composition of the casted alloys. The measured chemical compositions are presented in Table 3.1.
250 6 -
5 200 Pressure [bar]
Vacuum Vacuum Vacuum pressure [mbar] 4
150 Switchover point distance-time control [m/s]
Piston speed [m/s] 3 [bar] /
x x [mbar] 10 TK / 100
2 Velocity 50
Position 0,1 x [mm] 1 Pressure pressure 0 0 0 200 400 600 800 1000 Distance [mm] or time [ms]
Figure 3.1. Typical casting curve of studied series
A HPDC unit with a cold chamber (Frech DAK 450-54 with mould clamping force of 4500 kN) was used for castings. In order to achieve same cooling rate for all castings, all
55 parameters were controlled: melt temperature Tm = 720°C, mold temperature Td = 200°C, piston speed v = 2m/s, dwell pressure p = 250 bar (Figure 3.1 presents a typical casting curve of the studied series). Plates with the dimensions of 300x150x4 mm were cast as specimens (Figure 3.1) [144,179].
Table 3.1. Chemical composition of investigated alloys (Al-balance) [wt. %]
Chemical composition Alloying Impurity Mg/Si Zn Alloys Fe Ti Be Cu Si Mg ratio Mn Cr Sc [wt. %] [at. %] <0.2 <0.2 <0.004 <0.2 >1.73 C-series C1 2.62 5.90 2.2 0.68 0.10 - - 0.11 0.09 0.004 0.10 C2 2.68 5.91 2.2 0.68 0.19 - - 0.11 0.09 0.004 0.10 S-series S1 2.12 5.78 2.5 0.66 - 0.1 - 0.14 0.09 0.003 0.14 S2 2.14 5.69 2.5 0.66 - 0.2 - 0.14 0.09 0.003 0.14 Z-series Z1 2.77 5.84 2.1 0.66 - 0.58 0.24 0.14 0.09 0.003 0.15 Z2 2.73 5.78 2.1 0.66 - 1.20 0.50 0.14 0.09 0.003 0.15 Z3 2.70 5.72 2.1 0.66 - 1.83 0.76 0.14 0.09 0.003 0.15 Z4 2.65 5.57 2.1 0.65 - 2.89 1.21 0.14 0.09 0.003 0.15 Z5 2.62 5.51 2.1 0.65 - 3.99 1.69 0.14 0.09 0.003 0.15 Z6 2.60 5.47 2.1 0.65 - 5.01 2.11 0.14 0.08 0.003 0.15 Base alloy (base composition) M 2.7 5.67 2.1 0.66 - 0.14 0.08 0.003 0.15 (Magsimal59)
Figure 3.2. As-casted plate directly after withdrawing from the HPDC machine,
3.2. Differential scanning calorimetry
Differential scanning calorimetry (DSC) is a method to determine the heat flow between the specimen and a reference sample as a function of time and temperature. The phase transformations that occur during heating or cooling are followed by endo- or exothermic
56 reactions. The temperatures of these reactions and their heat exchanges can be determined from the DSC results. NETZSCH STA409 (heat flux type DSC) was used to perform the DSC studies. Specimens with a weight around 20 mg were used. Argon with 99.99% purity and with 10-20 cm3/min flow rate was used to protect specimens during studies. Measurements were made in the range from RT 25 to 710C with heating and cooling rates of 5 K/min. The results were processed with a NETZSCH Protheus software.
3.3. Microstructural investigation
3.3.1. Light microscopy and specimens preparation
Specimens for microstructural and local mechanical properties investigations for all alloys were prepared using standard techniques. Metallographic specimens were cut from the plates using a Discotom-2 (Struers) saw with water cooling. The specimens have dimensions 1054 mm. For handling, specimens were hot embedded. The grinding of the samples was done with a Labopol-21 (Struers) using following sequence of papers: 320, 500, 800, 1200, 2500, 4000 SiC-grit. The final step was polishing using Tegrapol-31 polishing pads (MD MOL, MD NAP, MD CHEM) with different diamond suspensions. An optical microscope Leica DM6000M with a Leica DFC450 digital microscope camera with a c-mount interface containing a high quality 5 Megapixel CCD sensor, was used for metallographic analysis.
3.3.2. Scanning electron microscopy and electron microprobe analysis
A SEM FEI QUANTA 450 (operated at 15keV) was used in this work to perform microstructural analysis. The SEM was equipped with EDAX Energy Dispersive Spectrometry system (EDS). The backscattered electron (to obtain material contrast) and secondary electrons (to obtain topology) detectors are used. A JEOL JXA-8100 (operated at 20 kV) was used to provide electron probe microanalysis (EPMA). PET (pentaerythritol), TAP (thallium acid phthalate) and LIF (lithium fluoride) detectors, spot size of 5 μm and a dwell time of 5 ms were used.
3.3.3. Transmission electron microscopy
A STRUERS Tenupol-5 electropolishing machine with an electrolyte A3 (Struers) was used for preparation thin-foil specimens for transmission electron microscopy (TEM) study. The
57 specimens have diameter 3 mm and a thickness <100 m. Electropolishing was performed at - 25 °C, a flow rate 15, a voltage of 50 V and a current of 120-140 mA. The Light-Stop-Value was 220. TEM study was performed using Philips CM20 (with accelerating voltage 200 kV). Bright field (BF) images were used to analyze the morphology of α-Al dendrites, intermetallic phases and nanoscale precipitates and dispersoids.
3.4. Partly automated image analysis for determination structure parameters
Based on the metallographic examination connected with quantitative image processing using Fiji/ImageJ software, a quantitative phase analysis was performed. Based on the image processing, precipitate size, their distribution and volume fraction average grain/dendrite size were determined. The quantitative image analysis was compared with the data from ThermoCalc software.
3.4.1. Particle size and volume fraction Volume fraction and the size of intermetallic phases were measured with the partly automated image analysis in the “Fiji”/“ImageJ” software. In the present work specialized algorithms were used for threshold determination (in order to minimize deviation from the manual thresholding). To isolate bright phases from the other microstructure, the “Otsu” algorithm [180] was used while for the dark Mg2Si phase was used the “Isodata” algorithm [181]. Otsu is one of the most referenced thresholding algorithms according to Sezgin et. al [182] and uses statistical calculations to identify the grey values that correspond to foreground and background of an image. Isodata is based on an iterative approach of repeating background-foreground separation until an optimum is found. Some images had to be subjected to pre-processing for the Otsu thresholding algorithm to work: they were darkened in steps of 10 grey values at a time until the algorithm correctly identified the bright intermetallic phase. This procedure was verified by testing the thresholding algorithm on a continuous sequence of images with decreasing darkness to make sure that once the algorithm “jumps” to correct identification of the phases its result is not sensitive to small brightness changes anymore. Other than that, no contrast adjustment or other image manipulation was done. As the masking of the T-AlMgZn intermetallic phase for the quantification of the α- AlMnFeSi intermetallic phase already included identifying both phases, the same thresholded
58 images were used to quantify both intermetallic phases. Every particle was identified as belonging to one of the two bright-appearing intermetallic phases and thus it was ensured that no particle was counted twice. After thresholding, the two different intermetallic phases were separated manually based on their location in the context of other microstructural features, their shape and “glow” into α-Al close to it. T-AlMgZn intermetallic phases have a brightening effect on α-Al in the neighbourhood while α-AlMnFeSi intermetallics darken the neighbourhood. To correctly count small, compact intermetallic particles, the Watershed algorithm of the image analysis software was used to identify and separate overlapping particles. Finally, the software is able to count all particles in one image and measure their surface area. This information is gathered for all particles and can be aggregated into the particle average size and overall volume fraction. The quantification was done ignoring particles with a size of under 0.05 µm² in order to ignore imaging artefacts of single bright pixels on the image. For the α-AlMnFeSi intermetallic coarse phase measurements a specimen surface area of 310000 µm² was analysed for each alloy. The measurement area spans one half of the specimen thickness from the surface to the center, therefore, in the results there is no specification for the distance from the specimen surface as it was done for the measurements on the other phases.
3.4.2. Dendrite arm spacing (DAS) In order to determine influence of the solidification rate on DAS the measurements were performed in two locations per specimen: at 600 µm from surface and 1400 µm from surface. The measurements were conducted on an area of 5476 µm² for each location. 44 line measurements were taken from each location. If a clear similarity in growth direction was identifiable for neighbouring dendrites, more than one dendrite was covered by a line measurement.
3.4.3. Interlamella spacing (ILS) High magnification SEM images were used for measuring interlamellae (or interfibre) spacing (ILS). 162 eutectic colonies for each specimens were analysed. As the fibre thickness increases when it approaches the border of the colony, it had to be made sure that an equal number of measuring points inside eutectic colonies and at their borders was taken into the final calculation of the resulting mean values to avoid statistical bias.
59 3.4.4. Strengthening precipitate density
The measurement of the precipitate density TEM images of α-Al dendrites with <100> direction were used. The precipitate density for 5 different α-Al dendrites for each alloy (and each HT state) to avoid the effect of non-homogeneous distribution were measured. A suitable magnification for TEM images allows 200-600 precipitates to be counted. The type of the precipitates was determined by the size and morphology according to literature data (see Table 1.4, Table 2.2). The foil thickness was determine using contamination spot technique [183,184].
3.5. Mechanical tests
3.5.1. Hardness measurements
A Brinell hardness tester with a ball diameter of 2.5 mm and a load of 306.56 N (31.25kp) was used to perform hardness measurements. 10 individual indents were measured for each specimen and the average value is presented. The size of ball-indents was in a range 0.6-0.8 mm (Figure 3.3) that may include about 4300 grains/dendrites with an average size of 12 μm (Table 4.5) Microhardness measurements were used as an indicator of mechanical response to the solid solution of α-Al strengthening after additional alloying in as-cast state and after heat treatment. The microhardness (Figure 3.4) was measured using a LECO M-400-G hardness tester. The load used was 0.01 kp and a dwell time of 15 seconds. The resulted hardness value for each specimen was taken as an Figure 3.3. SEM image of Brinell indent average for at least 10 measurements.
a) b) c) Figure 3.4. SEM images of indents in: a) α-Al dendrites; b) Al-Mg2Si eutectic; c) intermetallic particle
60
Due to the inhomogeneous distribution of hardness across the plate (Figure 3.5 a) for the microstructural analysis and for the local mechanical properties measurements the specimens were cut from the certain location from the different plates (Figure 3.5 b).
a) b) Figure 3.5. Hardness distribution across the plate (a) and locations for different type of the specimens
3.5.2. Tensile Tests
Tensile tests at RT were performed on an Instron 5967 testing machine at a cross head speed of 1 mm/min (initial strain rate of 8.3x10-4 1/s). The machine was equipped with an extensiometer MFA 25 (Mess- und Feinwerktechik GmbH). For the tensile testing specimens with a dog-bone shape and a rectangular cross section (gage length of 25 mm) according to ASTM E8_E8M_13a were used (Figure 3.6). The specimens for the tensile test were cut out from the middle part of the plates. Mechanical properties were measured on at least 5 specimens on samples in as-cast state and after heat treatment.
Figure 3.6. Specimens for the tensile tests
61 3.1. Heat treatment
For the heat treatment the preheated furnace Heraeus RL 200 E was used. The specimens with the size 1054 mm were positioned in the center of the furnace to achieve equal heating rate and temperature distribution. For the further studies of the microstructure and mechanical properties, two types of heat treatment were applied to the studied alloys: Two-step heat treatment (full T6) that includes solution treatment (ST) at different temperatures (380-580°C), quenching into the water and artificial aging (AA) (at 125°C-225°C) for various times (Table 3.2). One-step heat treatment (T5) - AA from the as-cast state at the temperature range 125°C - 325°C for various times (Table 3.2).
Table 3.2. Applied heat treatment modes for studied alloys Treatment mode Temperature [ºC] Base alloy C-series S-series Z-series One step AA 125, 175, 225, 325 175, 225, 325 175, 225, 325 125, 175, 225, ST 480, 500, 520, 570 520 - 350, 380, 480, 520 Two step AA 125, 175, 225 175, 225, 325 - 175, 225
62
CHAPTER 4 As-cast state of Al-Mg-Si-Mn-X alloys
In this chapter, the detailed characterisation of the microstructure analysis and mechanical properties of the studied series of alloys were presented. The effect of Sc, Cr and Zn additions on the solidification behavior and microstructure were investigated using deferential scanning calorimetry (DSC), scanning electron microscopy (SEM) with EDX analysis and electron-probe microanalysis (EPMA). Quantification of the structure components were done using partially automated image analysis. The structure of the α-Al solid solution was studied using transmission electron microscopy (TEM). To determine the effect of changes in the microstructure and in the composition of α-Al solid solution on the mechanical properties the common set of properties (HB, UTS, YS, A) and local properties of the α-Al dendrites and the eutectic cells were mesuared.
4.1. Differential scanning calorimetry
4.1.1. Cr and Sc containing alloys
Figure 4.1 shows the DSC curves for the C-series (Figure 4.1 a,b) and S-series (Figure 4.1 c,d) of the studied alloys in comparison to the M alloy (AlMg5.7Si2.6Mn). The upper trace of the DSC curves belongs to the endothermic reaction which occurs during heating. The lower part of the curves belongs to the exothermic reaction during cooling. The shapes of the DSC curves for Sc- and Cr- containing alloys are similar to the M alloy. This indicates that alloying of the alloys of Al-Mg-Si-Mn system with Sc and Cr does not influence the phase equilibrium. For each curve only one sharp endothermic peak was detected that can be characterized by three specific temperatures: Tonset (1), Toffset (2), Tpeak (3). Table 4.1 gives an overview of the specific temperatures of the investigated S- and C-series of alloys.
Table 4.1. Temperature details of the C- and S-series of alloys (rate of 5 K/min), [°C]
Alloy Tonset Tpeak Toffset Solidification range M 591.9 609.3 630.0 38.1 C1 591.3 608.7 628.9 37.6 C2 591.7 608.5 627.9 36.2 S1 588.2 610.6 629.9 41.7 S2 588.4 610.7 633.2 44.8
63 2,2 1,3 M C1 C2 1,1 M S1 S2 1,7 0,9 1,2 0,7 0,5 0,7 0,3 0,1 0,2 -0,1200 300 400 500 600 -0,3200 300 400 500 600 -0,3
-0,5 Heat Flow [mW/mg] FlowHeat -0,8 [mW/mg] FlowHeat -0,7 -0,9 -1,3 Temperature [ºC] -1,1 Temperature [ºC] -1,8 -1,3 2,2 1,3 1,8 3 1,1 4 4 0,9 3 1,4 0,7 1,0 0,5 1 2 1 2 0,6 0,3 0,1 0,2 -0,1540 560 580 600 620 640 -0,2540 560 580 600 620 640 -0,3
-0,6 -0,5 Heat Flow [mW/mg] FlowHeat Heat Flow [mW/mg] FlowHeat -0,7 -1,0 -0,9 -1,4 -1,1 Temperature [ºC] Temperature [ºC] -1,8 -1,3 a) b) Figure 4.1. DSC curves of investigated alloys in comparison with the M alloy: a) C-series of alloys; b) S-series of alloys
The onset temperature (denoted 1 on the Figure 4.1), which is equal for the M alloy to 591.9°C close to those that can be found in the literature and corresponds to the melting of the Al-
Mg2Si eutectic (see Table 2.1). The second endothermic peak (denoted 4 on the Figure 4.1) related to the melting point of the -Al dendrites. The offtset temperature (denoted 2 on the Figure 4.1) has specified the end of melting. For the M alloy, the offtset temperature was found to be 630C. Any other thermal effects on the DSC curve of the M alloy were not found. Addition of Cr with the studied concentrations has no significant effect on the solidification range of the alloys. On the other hand, Sc increases solidification range from 38ºC to 44ºC. Any additional thermal peaks on the DSC curves of the C-series as well as S-series alloys were not found.
4.1.2. Zn containing alloys
Figure 4.2 a,b show the DSC curves for the Z-series of alloys in comparison to the M alloy.
The shapes of the DSC curves for Z-series alloys are similar to the M alloy and consist of Tonset
(1), Toffset (2), Tpeak (3). The process of the melting in all alloys from Z-series starts by melting Al- 64
Mg2Si eutectic (Tonset (1), Tpeak (3)) than turning into the melting of dendrites (Tpeak (4) and ends with complete meltdown Toffset (2). However, the position of the curves varies while the zinc concentration changes. Thus, both temperatures (Tonset, Toffset) go down with increasing of the Zn content in the alloy (Figure 4.2 c).
The rate of decrease in the values of Tonset is higher than Toffset. This leads to an enlargement in the solidification range. Any additional thermal peaks on the DSC curves of the Z-series of alloys were not found.
1,5 1,5 M Zn1 Zn2 3 Zn3 Zn4 Zn5 4 1,0 1,0 Zn6 2 0,5 0,5 1
0,0 0,0 100 200 300 400 500 600 540 560 580 600 620 640
-0,5 -0,5
Heat Flow [mW/mg] FlowHeat [mW/mg] FlowHeat
-1,0 -1,0 Temperature [ºC] Temperature [ºC] -1,5 -1,5 a) b) 650 47 640 630,0 46,7 46 628,1 627,3 44,8 46,1 46,5
630 45 C] ° 620 43,2 624,8 44 622,2 620,0 610 617,8 43 600 591,9 41,0 42 584,1 590 580 41 576,1 573,5 Temperature [ Temperature 580 587,1 571,1 40
570 38,1 39 [ºC] range Solidification 560 38 Tonset Toffset Solidification range 550 37 0,00 0,24 0,50 0,76 0,98 1,20 1,45 1,70 1,90 2,10 Zn, at.%
c) Figure 4.2. DSC curves of Z-series of alloys (a,b) and changes in solidification range (c)
4.2. Microstructure and element distribution
4.2.1. Al-Mg-Si-Mn base alloy
Figure 4.3 shows the microstructure of the cross-section of the M alloy. The structure divides into 3 zones: outer skin, band zones, central region. The band zones are parallel to the surface of the casting plate. The zone is located at a distance of 750-900 μm from the surface and has a width 220- 280 μm. The band separates the outer skin and the central region of the casting. In the band zone, the volume fraction of α-Al dendrites is lower and volume fraction of eutectic is higher than the
65 surroundings. The similar effect of the band-zone formation in the die-cast Al-Mg-Si alloys has been observed by Ji et al. [20] and Otarawanna et al. [90]. However, the band zones in the investigated castings are not so clear than in the cases [20,90]. Figure 4.4 represents a general microstructure of the base AlMg5.7Si2.6Mn alloy (chemical composition is reported in the Table 3.1). The structures of the M alloy is composed of the dendrites of -Al, Al+Mg2Si eutectic and -AlMnFeSi phase.
Figure 4.3. Optical image with low magnification of the M alloy
Figure 4.4. SEM images of the M alloy
Al+Mg2Si eutectic has a lamellar and/or fibers morphology, where long Mg2Si plates and fibers alternate with -Al (Figure 4.4, Figure 4.5). The Mg2Si fibers have an arbitrary location and orientation. In different images thay can take the form of separate spheres or long curved lines.
The lamellar Mg2Si plates (Figure 4.5 d) are coarser with higher ILS values. Such structures are formed in the shot-sleeve of the HPDC machine (with lower solidification rate [175,185]). The Mn-containing phase can be present in the alloy in different forms such as: hexagonal, cubic and star-like (Figure 4.5 a-c). Despite the different morphologies, the phases have the same
66 composition and belong to the a-AlMnFeSi phase. The a-AlMnFeSi phases are mostly located in the interdendritic region and inside eutectic colonies. They can be found apart (Figure 4.5 c) or as a small colonies (Figure 4.5 b).
a) b)
c) d) Figure 4.5. Microstructures of M alloy a) hexagonal -AlMnFeSi phases; b) cubic -AlMnFeSi * phases; c) star-like -AlMnFeSi phases; d) lamellar Al-Mg2Si eutectic
According to [20,185] there are two possible types of the α-Al dendrites in the as-cast state: the α1-Al phase (solidified in the shot sleeve) and the α2-Al phase (solidified in the die cavity). Due to the differences in the cooling rates in the shot sleeve (102 °C/s) and in the die cavity 3 (10 °C/s) [20,185], the size and morphology of the primary α1-Al and α2-Al phases are significantly different. The sizes of the coarse α1-Al vary from 20 to 40 μm (Figure 4.6 a); the sizes of the fine α2-Al with globular-rosette morphology vary from 5 to 15 μm (Figure 4.6 b). The EDX-analysis of -Al dendrites and -AlMnFeSi intermetallics is presented in the Table 4.2. Based on the literature data and the average composition, this intermetallic phase is usually identified as α-Al15(Mn,Fe)3Si2. The segregation of elements for the M alloy was studied using electron probe microanalysis (EPMA). The EPMA maps are shown in Figure 4.7.
* The SEM images (Fig. 4.5. a, b) have been made at the Technische Universität Darmstadt, Darmstadt, Germany 67 a) b) Figure 4.6. Morphology of the -Al in M alloy in as-cast state: a) α1-Al; b) α2-Al
The segregation of alloying elements in the α-matrix is determined by partition coefficients (the coefficient is defined as the slope of the liquidus over the slope of the solidus in a binary phase diagram [24]).
Table 4.2. Average composition of the -Al solid solution and the Mn-containing phase in the M alloy Phase Elements Al Mg Si Mn Fe Cu Zn Ti -Al dendrites [at.%] 96.4 2.8 0.3 0.3 0.1 0.1 0.1 0.1 [wt.%] 95.8 2.5 0.3 0.6 0.2 0.2 0.2 0.2 -AlMnFeSi [at.%] 75.9 3.5 7.3 10.1 3.0 0.1 0.1 0.1 [wt.%] 66.6 2.7 6.6 18.1 5.4 0.2 0.2 0.2
a) b)
c) d) Figure 4.7. SEM-image with EPMA maps of the M alloy a) SEM image of the selected area, and segregation of the b) Mg c) Si d) Mn across selected area
68
Mg as well as Si (k~0.5 and ~0.1 respectively), enrich the dendritic edge. They are also form the stable Mg2Si compound that solidifies as eutectic in the interdendritic region (Figure 4.7. b and c). Mn (k ~ 1) is accumulated mostly in the dendritic edge and forms the α-AlMnFeSi phase in the interdendritic space (Figure 4.7. d).
4.2.2. Al-Mg-Si-Mn with Sc and Cr additions
Figure 4.8 and Figure 4.9 show the microstructure of the studied alloys after Cr and Sc additions. These alloys have mainly the same structure as the M alloy: -Al grains, Al-Mg2Si eutectic and α-AlMnFeSi intermetallic phase.
a) b)
c) d) Figure 4.8. SEM images of the C-series of alloys a) C1 (0.1 wt. %Cr); b) C2 (0.2 wt. %Cr); b,c) Cr-containing intermetallics in the C2 alloy *
Additionally to basic set of the phases a new Cr-containing coarse intermetallic phase in C2 alloy (with 0.2 wt. % Cr) (Figure 4.8 b,c,d) was found (as it was predicted by the thermodynamic calculations on Figure 2.6). This phase belongs to the Al45Cr7 or more frequently referenced as Al7Cr. Al7Cr phase has an irregular polygonal blocky-like shape (Figure 4.8 b,c,d), with the inhomogeneous sizes (in the range from 5 μm to 30 μm). This phase also has a significant content of Mn. In the C1 alloy (with 0.1 wt. % Cr) Cr doesn’t form new phase but only enriched
* The SEM images (Fig. 4.8. c, d) have been made at the Technische Universität Darmstadt, Darmstadt, Germany 69 the solid solution. Cr is also present in the Mn-phase in the C1 and C2 alloys, thus forming stable
α-Alx(Fe,Mn,Cr)ySiz (and/or Al12(Cr,Mn)) intermetallics [171]. In the S1 and S2 alloys (Figure 4.9), a small bright phase with an irregular shape is found. According to the calculated phase diagram (Figure 2.5) and literature data this phase can be identified as Al3Sc. Table 4.3 shows the average chemical compositions of the -Al grains and intermetallic phases in the studied alloys. The main alloying elements (Mg, Si, Mn) presented in the solid solution remain at the same level as in the M alloy. The average content of the addition elements (Sc and Cr) are close to the amount of Sc and Cr added to the M alloy.
a) b) Figure 4.9. SEM images of the S-series of alloys in as-cast state a) S1 (0.1 wt. % Sc); b) S2 (0.2 wt. % Sc)*
EMPA maps of the studied series of alloys are given in Figure 4.10 and Figure 4.11. As it can be seen from the Figure 4.10, the interdendritic region in the C-series of alloys is slightly poor with Cr (k~2.0) while α-Al dendrites are enriched with Cr. Over and above a high amount of Cr in the C1 and C2 alloys is present in the Mn-containing phases. The EMPA maps of the S-series of alloys are presented in the Figure 4.11. Sc (k~1.0) is mainly located in the interdendritic region. Sc with Al forms trialuminid Al3Sc. The Al3Sc phase has an irregular morphology and is located in the interdendritic region similarly to the α-AlMnFeSi phase, morover it has a similar contrast to the α-AlMnFeSi phase on the SEM images, thus, its identification becomes more complicated. The mean size of the intermetallics and their volume fraction are given in Table 4.4. The volume fraction of the α-AlMnFeSi phase as well as the Al7Cr and Al3Sc intermetallics (for C2 and S2 alloy) are rather close to those obtained from the thermodynamic calculations for each of the proposed compositions. Nevertheless, the volume fractions of the Al7Cr and Al3Sc in the alloys
* The SEM images have been made at the Technische Universität Darmstadt, Darmstadt, Germany 70 with 0.1 wt.% of the alloying element (C1 and S1) are not measurable due to the low amount and/or very inhomogeneous distribution across the sample.
Table 4.3. Average composition of the -Al solid solution and the Mn-containing phase in the Sc- and Cr- containing alloys [at.%][144] Alloy S1 S2 C1 C2 S1 S2 C1 C2 C2
Element α-Al dendrites Mn-phase Al7(Cr,Mn) Mg 2.8 2.7 2.8 2.8 3.5 2.4 1.3 1.8 1.5 Al 95.9 96.3 96.1 96.0 75.9 78.6 77.0 75.8 85.2 Si 0.4 0.4 0.4 0.4 7.3 6.0 7.7 7.0 1.2 Mn 0.4 0.3 0.2 0.2 10.1 10.3 10.1 10.2 5.5 Fe 0.1 0.1 0.1 0.1 2.9 2.4 2.5 2.4 0.4 Cr - - 0.1 0.2 - - 1.4 2.6 5.0 Sc 0.1 0.2 - - 0.1 0.4 - - - Ti 0.1 0.1 0.1 0.1 0.1 0.1 0.1 0.1 1.2
Table 4.4. Volume fraction and mean size of the intermetallics in the investigated alloys. Alloy Mn-containing phase Intermetallic phase Size, [μm] V, [%] V, [%] (calc) Size, [μm] V, [%] V, [%] (calc) М 1.0±0.2 1.5±0.2 1.87 - - - Cr-containing phase C1 0.8±0.3 1.7±0.3 1.87 - - 0.38 C2 0.7±0.3 1.8±0.3 1.87 13.5±7.8 0.7±0.2 0.77 Sc-containing phase S1 1.3±0.3 1.4±0.2 1.87 - - 0.24 S2 1.3±0.4 1.4±0.3 1.87 - 0.5±0.1 0.48
Table 4.5 shows the grain size and the interlamellar space (ILS) of the present series of alloys. Cr and Sc additions in such small amount do not effect the grain size. Meanwhile, the ILS decreases with the addition of Cr and Sc.
Table 4.5. DAS and ILS of the eutectic measured for the investigated alloys. Alloy DAS, [μm] ILS, [μm] M 6.1±2.5 0.21±0.06 C1 5.9±1.5 0.15±0.05 C2 5.7±2.2 0.13±0.06 S1 5.8±1.5 0.11±0.04 S2 5.5±2.5 0.10±0.06
71 a) b)
c) d) Figure 4.10. SEM-images with EPMA maps of the distribution of Cr in a-b) C1, c-d) C2 alloys
a) b)
c) d) Figure 4.11. SEM-images with EPMA maps of the distribution of Sc in a-b) S1, c-d) S2 alloys
72
4.2.3. Al-Mg-Si-Mn-Zn alloys
Figure 4.12 shows morphology of the studied alloys with different Zn content. The structure of alloys consists of mainly the same phases as in the M alloy, namely, -Al grains, Al-
Mg2Si eutectic and α-AlMnFeSi intermetallics. Even a high Zn content in the alloys (up to 5 wt. %) has insignificant impact on the size and morphology of the main phases. Quantification of the structure components is presented in Table 4.6. Bright irregular-shaped phases are also found after Zn addition. According to Figure 2.7 and Figure 2.11 this phase can be identified as T-AlMgZn phase. The composition of the phase was measured for alloys Z2-Z6 (measurements for the Z1 alloy are not possible due to the small size of the phase). The results are presented in Figure 4.15.
a) b)
c) d)
e) f) Figure 4.12. SEM image of the Z-series f alloys in as-cast state: a) Z1 (0.2 at. % Zn); b) Z2 (0.5 at. % Zn); c) Z3 (0.7 at. % Zn); d) Z4 (1.2 at. % Zn); e) Z5 (1.7 at. % Zn); f) Z6 (2.1 at. % Zn)
73 The EMPA maps of the Z-series of alloys are presented in the Figure 4.13. In contrast to Cr and Sc, Zn, with k ~0.45 [24], is mainly distributed in the interdendritic area, similarly to Si and Mg. The dendritic core is poor in Zn. With the increasing in the alloy amount of Zn its concentration increases in both the interdendritic space and in the α-Al dendrites (Figure 4.15). Starting with the Z2 alloy, a bright white phase with a high Zn content appears in the interdendritic space. T-AlMgZn intermetallics is mainly precipitated on the border between eutectic and α-Al dendrites. With increasing Zn content morphology of the T-AlMgZn phase changes. Thus, in alloy with 1.8 wt.% Zn (Z3) small eutectic T-AlMgZn phases were found (Figure 4.14). The next increasing Zn content in the alloy leads to the further growth of the eutectic T-phase. In the Z4-Z6 alloys (1.2-2.1 at. % Zn) new irregular coarse Mn-containing particles were observed (Figure 4.14 c). The composition (Figure 4.15 d) and the hardness of these phases (1.3 ± 0.1 GPa) is similar to fine Mn-containing phases, but size is higher (5.6-24.7μm [186]). The coarse Mn-containing phase has negligible volume fraction for accurate measurement, and its values for Z4, Z5 and Z6 alloys are in the range of 0.012-0.035% [186].
Table 4.6. Parameters of structure components of the studied alloys [179]
α-Al dendrites Al-Mg2Si-eutectic α-AlMnFeSi T-AlMgZn
DAS, V , V , [%] ILS, V , V 2 , [%] Size, V, [%] V, V, [%] d Al eut. Mg Si V, [%] [μm] [%] (calc) [μm] [%] (calc) [μm] (calc) [%] (calc) M 6.1±2.5 53 90.61 0.21±0.06 45 7.52 1.0±1.0 1.5±0.1 1.87 - - Z1 5.2±1.5 52 87.92 0.26±0.06 46 7.75 0.8±0.5 1.3±0.1 1.87 - 2.45 Z2 5.3±1.5 51 87.57 0.33±0.06 47 7.66 0.9±0.5 1.2±0.3 1.87 0.2 2.89 Z3 5.2±1.7 54 87.33 0.29±0.06 44 7.60 0.9±0.5 1.3±0.1 1.88 0.1 3.19 Z4 5.9±3.0 58 87.00 0.33±0.06 40 7.51 0.9±2.0 1.3±0.7 1.88 0.9 3.60 Z5 5.6±2.5 57 86.74 0.27±0.06 40 7.48 1.3±0.6 1.2±0.3 1.88 1.0 3.90 Z6 5.2±3.0 53 86.70 0.28±0.06 44 7.46 0.9±1.0 1.2±0.5 1.88 1.3 3.96
The change of the chemical compositions of -Al solid solution and intermetallic phases of studied alloys with increasing Zn content is presented in Figure 4.15a. -Al solid solution of M alloy (without Zn addition) has high level of Mg (up to 3 at. %) and small Si, Mn and Fe content. As it can be seen from Figure 4.15 a with increasing Zn content in the alloys (from 0.2 to 2.1 at. % Zn) Zn concentration in a-Al dendrites increases proportionally (from 0.2 to 1 at. %). The composition of the α-AlMnFeSi phase in all alloys is mostly stable and consists of Al, Mn, Fe and Si. The concentration of Al, Mn and Fe in all alloys in the α-AlMnFeSi phase is mostly the same, only a little increase of Si can be distinguished in the alloys with high Zn content. The size and amount of the T-AlMgZn phase in the Z1 alloy is too small for determining its chemical composition. Figure 4.15 b represents the changes in composition of the T-phase. It mostly
74 consists of Al, Mg and Zn. Furthermore, the presence of Cu in the EDX-spectrum of T-phase is caused by its presence in the alloy as an impurity.
a) b) c)
d) e) f) Figure 4.13. EPMA maps of the dendritic area: a) Z1, b) Z2, c) Z3, d) Z4 e) Z5, f) Z6 alloys [179]
a) b) c) Figure 4.14. New phases in Z-series of alloys in as-cast state: a) eutectic T-AlMgZn in Z3 alloy; b) eutectic T-AlMgZn in Z5 alloy; c) сoarse Mn-containing sludge particles [179]
75
a) b)
c) d) Figure 4.15. EDX analysis of phases in studied alloys in as-cast state (as a function of Zn content): a) α-Al dendrites, b) T-AlMgZn phase, c) fine α-AlMnFeSi, d) coarse α- AlMnFeSi
4.3. Structure of α-Al dendrites
As reported in the Chapter 1, Al-Mg-Si alloys belong to the heat-treatable series of alloys, as they are capable of forming highly dispersed nanoscale β”, β’-Mg2Si precipitates in the α-Al solid solution. Therefore, the structure of the Al-dendrites of the studied alloys was studied by transmission microscopy. Figure 4.16 a represents TEM bright-field image of the general structure of the M alloy, where all the main structural components are clearly distinguishable, namely α-Al dendrites, Al-
Mg2Si eutectic and α-AlMnFeSi intermetallic. In the structure of α-Al dendrites of M alloy two types of precipitates were distinguished:
β' and β-Mg2Si phases (Figure 4.16 b-f). The β-Mg2Si is the equilibrium phase in the alloys of Al- Mg-Si system and has square platelet and cubic shapes. In the M alloy the side length of the β-
Mg2Si phase ranges from tens of nanometers to 1-2 µm (Figure 4.16 b). The β' precipitates have rod-like shape with length 50-200 µm elongated in the <100>Al. Additionally to β' and β precipitates the α-AlMnFeSi dispersoides with a cubic shape (mean size in a range of 10-20 nm) 76 in the α-Al dendrites were found (Figure 4.17). Any other types of strengthening precipitates (e.g. β'' or GP-I zones) were not detected.
a) b)
c) d)
e) f) Figure 4.16. TEM bright-field images of the α-Al dendrites in M alloy in as-cast state
a) b) Figure 4.17. Mn-containing dispersoids in the M alloy in as-cast state
77 The as-cast structure of α-Al dendrites in the S- and C-series of alloys is similar to the M alloy. Any significant changes in all C1, C2, S1 and S2 alloys in comparison to the M alloy were not found. Only small amounts of Sc-containing precipitates (Al3Sc) can be found in S-series of studied alloys. In addition to the two already described types of precipitates the studied alloys from Z-series have one more (Figure 4.18 – Figure 4.20 ). Highly dispersed precipitates (Figure 4.18) have a “coffee-bean” morphology with a mean size of 1-5 nm and are slightly elongated in the
<100>Al. These precipitates belong to the Zn-containing GP-I zones. For comparison of the effects of the existing dispersed phases their density was calculated (Table 4.7). As can be seen Z1 alloy (0.2 at. % Zn) has the lowest density of the GP-I zones among all alloys from Z-series, while Z3 (0.7 at. % Zn) has the highest density (Table 4.7). With further increasing Zn content (Z4-Z6 alloy) the density of the precipitates changes insignificantly. All alloys from Z-series have similar levels of β' precipitates with M alloy. Any other types of the strengthening precipitates (휂 ’ or GP-II zones) in the Z-series of alloys were not found.
a) b) Figure 4.18. TEM bright-field images of the GP-I zones in Z2 alloy
Table 4.7. Density of the precipitates in the Zn-containing series of alloys Density, [μm-3] Type Size, [nm] Z1 Z2 Z3 Z4 Z5 Z6 GP-I (Zn,Mg) 2-5 7 430 20 580 33 200 28 890 28 940 28 860 β'-Mg2Si 50-200 270 280 280 230 260 270
78
a) b)
c) d)
e) f) Figure 4.19. TEM bright-field images of the α-Al dendrites in Z-series of alloys in as-cast state: a) Z1 (0.2 at. % Zn); b) Z2 (0.5 at. % Zn); c) Z3 (0.7 at. % Zn); d) Z4 (1.2 at. % Zn); e) Z5 (1.7 at. % Zn); f) Z6 (2.1 at. % Zn)
79
a) b)
c) d)
e) f) Figure 4.20. TEM bright-field images of the α-Al dendrites in Z-series of alloys in as-cast state [179]: a) Z1 (0.2 at. % Zn); b) Z2 (0.5 at. % Zn); c) Z3 (0.7 at. % Zn); d) Z4 (1.2 at. % Zn); e) Z5 (1.7 at. % Zn); f) Z6 (2.1 at. % Zn)
4.4. Mechanical properties
The mechanical properties of the M alloy as well as after Cr, Sc and Zn addition were measured by Brinell hardness (HB) and tensile test. The average values of the mechanical
80 properties of the M alloy, which were used for comparison with the properties of the studied alloys, are presented in the Table 4.8.
Table 4.8. Mechanical properties of the M alloys in as-cast state HB YS, [MPa] UTS, [MPa] A, [%] 80 163 299 8
4.4.1. Al-Mg-Si-Mn with Sc and Cr additions
Figure 4.21 a represents results of the hardness (HB) values as a function of the concentration of alloying element. Sc has stronger effect on the hardness of alloys than Cr. The hardness of the S2 alloy (0.2 wt. % Sc) is 14 % higher than the M alloy, but only 1% higher than the S1 alloy. The hardness of the C1 alloy (0.1 wt. % of Cr) is around 6% higher than the M alloy. The hardness of the C2 alloy is on the same level as S1 and S2 alloys.
Cr Sc Cr Sc 95 200
190 90
180 85
170 YS [MPa] YS
Hardness[HB] 80 160
75 150 0.0 0.1 0.2 0.0 0.1 0.2 Sc, Cr [wt.%] Sc, Cr [wt.%] a) b) Cr Sc Cr Sc 350 11
330 10
310 9
290 [%] A 8 UTS [MPa] UTS
270 7
250 6 0.0 0.1 0.2 0.0 0.1 0.2 Sc, Cr [wt.%] Sc, Cr [wt.%] c) d) Figure 4.21. Effect of the alloying elements on the mechanical properties of the studied alloys [144]. a) results of the Brinell hardness tests, b) yield strength c) ultimate tensile stress and d) elongation to failure.
81 Figure 4.21 b-d represents the tensile properties of the studied alloys (YS, UTS and A) as a function of the concentration of alloying elements. It can be seen that the alloying by Cr does not produce noticeable effect on the tensile strength (the maximum increase for the Cr containing series of alloys is around 5%). The greatest improvement in the whole set of the tensile properties is obtained in the S-series of alloys (Figure 4.21 b-d). Alloy S1 (0.1 wt. % Sc) shows an increasing in UTS and elongation by 11% and YS up to 16%. Alloy S2 (0.2 wt. % Sc) shows an increasing in the elongation by 9%, in the UTS and YS up to 3% in comparison with the S1 alloy and in the UTS up to 15%, in the YS and elongation by 21% in comparison with the M alloy.
Furthermore, the indentation behaviour of the α-AlMnFeSi and Al7Cr intermetallic phases was investigated [144]. Thus, Al7Cr has a hardness value of 9.57±0.31GPa and α-AlMnFeSi has the highest hardness value of 14.55±0.9 GPa among all studied intermetallic phases.
4.4.2. Al-Mg-Si-Mn-Zn alloys
The hardness and tensile properties (yield stress, ultimate tensile stress and elongation to failure) are plotted as a function of Zn content and are presented in Figure 4.22. As it can be seen the Zn addition has a strong effect on all sets of the studied properties. At the first stage of increasing the Zn concentration (up to 0.75 at. %), a sharp increase in both UTS and YS (up to 30%) can be observed. Upon reaching a concentration of 1.2 at.% Zn, a slight decrease in UTS is observed, a further increase in the Zn concentration does not lead to significant changes in UTS and YS values. The behaviour of the change in in the values of elongation is inversely proportional to the change in the UTS values. Z1 and Z2 alloys have same to the M alloy levels of elongation and in general, elongation gradually decreases in the range of concentration 0-1.2 at. % Zn (Z1- Z4 alloys). The elongation for Z4-Z6 alloys, similarly to UTS and YS, remains at the same level (around 3%). Figure 4.22 b represents the results of the hardness (HB) as a function of Zn content. As can be seen, the hardness values strictly depend on the Zn content in the alloy. As can be seen from the analysis of the results of whole set of studied properties, there is a rather large concentration range without significant changes in the properties (alloys Z4-Z6 in the range 1.2- 2.1 at. % Zn). It was found that Z1 alloy has the maximum elongation, Z3 alloy - the maximum of UTS, Z6 alloy - the maximum of YS. Figure 4.22 c represents the local mechanical properties (HV0.01) of the studied alloys, such as microhardness of α-Al dendrites and Al-Mg2Si eutectic cells. The increasing of the Zn concentration in the alloys leads to the rising of the microhardnes of the dendrites as well as the eutectic cells. However, these changes have different behaviour. The increasing Zn content up to 0.7 at. % (Z3 alloy) leads to the growth of the microhardness of the α-Al dendrites up to 18%.
82
With the further increasing in Zn content (Z4-Z6 alloys) the rate of this growth decreases. Thus, the microhardness of the α-Al dendrites in an Z6 alloy (2.1 at. % Zn) is higher than in Z3 alloy (0.7 at. % Zn) by only 5%. The microhardness of the eutectic cells has a completely different trend. The microhardness of the eutectic cells increases proportionally with increasing Zn content. The microhardness of the eutectic cells in Z6 alloy is almost twice as high as the microhardness in the M alloy.
UTS YS A% 400 10 9 350 8 7 300 6
5 % A,
250 4 UTS, YS, [MPa] YS, UTS, 3 200 2 1 150 0 0 0,5 1 1,5 2 Zn, [at.%] a) 130 HV0,01(a-Al) HV0,01(eutectic) 210 120 190 110 170 100 150 90
130
Hardness, HV0.01 Hardness, Hardness, [HB] Hardness, 80 110
70 90 0 0,5 1 1,5 2 0 0,5 1 1,5 2
Zn, [at.%] b) Zn, [at.%] c)
Figure 4.22. Mechanical properties of the Z-series of alloys [179]: a) tensile properties; b)
hardness (HB); c) microhardness of α-Al dendrites and Al-Mg2Si eutectic
4.5. Discussion
The results obtained in the Thermo-Calc software show that the structure of the base composition (M alloy) consists of several phases: -Al, Mg2Si, -AlMnFeSi, β-AlMg. The presented set of phases is directly related to the specific Mg/Si ratio ( >2 for all investigated alloys, see Table 3.1). The results of microstructural investigation have shown that the structure of the M alloy consists of -Al dendrites, Mg2Si (that is presented as the part of Al-Mg2Si eutectic colonies) and -AlMnFeSi intermetallic shaped phase (cubic, hexagonal or star-like). The β-AlMg phase was
83 not found. This may depend on a high crystallization rate (due to the use of HPDC), the small volume fraction of the phase as well as the close contrast to -Al. In [129], the β-AlMg phase was detected in the alloy Al–10.75Mg–2.86Si–0.59 Mn, which contains 5.8 wt. % of excess Mg, whereas the studied M alloy contains only 1.1 wt. % of excess Mg. The specific ratio of Mg/Si leads to the fact that the bulk of the Si is needed for the formation of the Mg2Si phase (eutectic lamellae) and consequently only a small amount of Si can be present in the -Al solid solution. From the literature, it is known that the -Al contains around 0.2 wt. % Si and around 2.7 wt. % of Mg at the eutectic temperature (~594°C) [126]. These values are close to those obtained in Table 4.2 and Figure 4.15. Because Fe is always present in Al die-cast alloys, the main reason to add Mn is to prevent formation in the structure brittle, acicular β-AlFeSi phase and to promote formation more compact and less harmful to ductility -AlMnFeSi phase. As can be proved by Thermo-Calc calculations and by microstructural analysis β-AlFeSi phase was not found in the studied alloys. Moreover, in the structure of all studied alloys the most common types (with the ratio Mn/Fe>0.5) of intermetallic phases (cubic, hexagonal, star-like) were found [122].
4.5.1. Influence of the Sc- and Cr- additions to the base alloy in the as-cast state
a) Microstructure. -Al contains around 0.2 wt. % Si and around 2.7 wt. % of Mg at the eutectic temperature (~594°C) [126]. These values are close to the EDX results of the current research obtained from the analysis at RT (Table 4.2). Cr and Sc additions do not have a noticeable effect on the segregation of the main elements in the dendrites (Table 4.3). The alloys from C- and
S-series (Cr and Sc additions) show the formation of Al7Cr and Al3Sc intermetallics, which are not present in the M alloy. The Al7Cr intermetallic phases with an polygonal blocky-like shape were found in the C2 alloy, but no Al7Cr phase was found in the C1 alloy. This can be due to the inhomogeneous distribution, low volume fraction and fast solidification during HPDC. Cr (with segregation coefficient in Al higher than 1) enriches the α-Al dendritic core (Figure 4.10). In the alloys from C-series, some amount of Cr is present in the α-AlMnFeSi intermetallics (Table 4.3 and Figure 4.10). Cr joins with trace elements and forms stable α-AlFeMnCrSi (and/or
Al12(Cr,Mn) intermetallics [171,187]. This may lead to the absence of the coarse Cr-containing phase in the C1 alloy. On the other hand, Sc does not affect the chemical composition of the α- AlMnFeSi. Sc has a little tendency for a segregation (k≈1) in the Al, similarly to Mn, and is mainly homogeneously distributed across the α-Al dendrites, however, the interdendritic area has a higher Sc content than the dendritic core (Figure 4.11) [129,188]. Apparently, during solidification Mn and Sc are pushed to the front of solid-liquid interface. As a result, the α-AlMnFeSi and Al3Sc
84 phases can be found mainly in the interdendritic area. According to literature [40,189] in binary
Al-Sc system Al3Sc intermetallics are formed at 0.52 wt. % Sc concentration. For the studied system (Figure 2.5) the peritectic reaction take place with the concentration of ~ 0.12 wt. % Sc and Al3Sc phase is formed. Microstructural and EPMA analysis confirm the presence of Al3Sc in the alloys from S-series. b) Mechanical properties. Similar concentrations of Cr and Sc additions positively affect the tensile properties (UTS, YS, A) in the wrought Al-Mg-Si alloys and Al-Mg2Si composites, mainly due to grain refinement effect [138,145]. However, according to Table 4.5, it was found that the addition of Cr and Sc up to 0.2 wt.% to the Al-Mg-Si casting alloy does not affect the DAS. Consequently, the increasing of the mechanical properties in the studied alloys cannot be ascribed to the grain refinement effect. On the other hand, the alloying by Cr and Sc decrease ILS
(Table 4.5). Thus, ILS of the Al-Mg2Si eutectic structure in the M alloy is equal to 0.21 μm; Cr and Sc additions reduce the ILS up to twice.
The local mechanical properties 160 Cr (α-Al) Sc (α-Al) (microhardness of the Al-Mg2Si eutectic and Cr (eutectic) Sc (eutectic) 150 α-Al dendrites) were measured by 140 microhardness tester. The microhardness of the dendrites as well as the eutectic (Figure 130 4.23) of the alloys from C- and S-series is 120 Al-Mg2Si eutectic significantly higher than the M alloy. Alloys HV0.01 Hardness, 110 with Cr and Sc show an increase of the 100 microhardness of α-Al dendrites up to 16% 90 α-Al dendrites 0 0.1 0.2 and 28% respectively. The highest hardness Sc, Cr wt.% values of α-Al dendrites as well as of Al- Figure 4.23. Microhardness HV0.01 of α-Al dendrites and Al-Mg2Si eutectic Mg2Si eutectic were recorded (118 and 153 cells as a function of the respectively) for the alloys with the 0.2 wt. % concentration of alloying element Sc addition (Figure 4.23). However, the increase in the hardness values is not proportional to the increasing in concentration. The addition of 0.1 wt. % Sc leads to increasing hardness of the α-Al up to 25% and eutectic up to 18%. The hardness of α-Al in alloy S2 (with 0.2 wt. % Sc) is higher than in alloy S1 (0.1 wt. %) by only 3% and eutectic 5% respectively. Similar situation occurs in the Cr-containing series. This means, that the lower amount (0.1 wt.% both Cr and Sc) of alloying elements enriches enough the α-Al solid solution. The excess amount of element that was not solved in the α-Al, forms the Al7Cr and Al3Sc intermetallic phases. Similar trends were obtained with nanoindentation technique [144].
85 It is known that intermetallic phases present in the structure can act as obstacles to dislocation motion that can improve hardness and strength of alloys. However, needle-shaped, large-sized, brittle, fusible intermetallic particles (silicides, AlxFe binary or AlxFeySiz ternary phases, etc.) can act as stress concentrators, crack initiators or impede filling of the mold [82,122,187,190]. Addition of Mn and/or Cr can eliminate this harmful effect by a substitution of
Fe and formation of a fine and compact α- Alx(FeMnCr)ySiz stable phases.
Al7Cr in the C2 alloy forms rather compact but coarse intermetallics that are uniformly distributed. Despite the high strength of the coarse Al7Cr phase [144] it is brittle and therefore deteriorates the yield stress and elongation of the C2 alloy (Figure 4.24). On the other hand, Al3Sc particles interact with the Al-Mg2Si eutectic seemingly without acting as stress concentrators and do not have a negative effect on tensile properties (Figure 4.21). As shown in Figure 4.21a, the maximum hardness values have S1, S2 and C2 alloys. Nevertheless, the alloys from C-series have the lowest values of elongation.
a) b)
Figure 4.24. Coarse Cr-containing phase with cracks in the fracture of the C2 alloy (b-c).
4.5.2. Influence of the Zn- addition to the base alloy in as-cast state
a) As it can be seen from the EDX-analysis the α-AlFeMnSi phase does not contain Zn. However, with an increase Zn content in the alloy, the chemical composition of the α-AlFeMnSi phase changes. The α-AlFeMnSi phase in Z4-Z6 alloys (1.2 - 2.1 wt.% Zn) contains a higher Si content. Such changes can be caused by the T-phase formation and the redistribution of the elements. The changes in the composition of α-AlFeMnSi intermetallic phase are also in good agreement with the results calculated in the Thermo-Calc software (Figure 2.8). The formed coarse Mn-containing phase (in the Z4-Z6 alloys) has a higher Fe and Si content, compared to the fine α- AlFeMnSi phase. This type of the intermetallic phase forms during holding in the furnace (with insufficient temperature) and is commonly called sludge [187,191]. However, during casting the temperature was controlled at 720ºC, so it was high enough to prevent formation of the sludge
86 particles. On the other hand, sludge particles can also nucleate in the process of the transport of molten metal to the shot sleeve at the walls of the crucible or in the cold chamber [187]. The effects of the Cr, Mn, Fe [187,191] or Si, Cu [192,193] on the sludge formation are well studied, on the contrary there is no literature about the effect of Zn. As already mentioned, Zn has an effect on the redistribution of elements in the alloys (especially Si). Zn addition also increases the solidification range (Figure 4.2 c), incl. increase of the range of the solidification of primary α-AlFeMnSi phase (Figure 2.9). This may be the indirect cause of the sludge particles formation in the Z-series of alloys. The calculations from the Thermo-Calc software (Figure 2.7) show that the low solubility of the Zn in the base composition (M alloy) at low temperatures leads to the formation of a new T-AlMgZn intermetallic phase even at a small Zn concentration. Analysis of the microstructure and EPMA results confirm the presence of the T-phase in the alloys from Z-series. The T-AlMgZn intermetallics have a bright contrast, an irregular morphology and are mainly precipitated in the interdendritic area that is in good agreement with the results [129,168]. The identification and quantification of T-AlMgZn phase are not an easy tasks, due to the T-AlMgZn and α-AlMnFeSi intermetallis have a similar contrast on the SEM images (T-AlMgZn is slightly brighter) and can usually be found in the similar sites. The equilibrium diagrams calculated using FactSage software [13,167] show the possible appearance of several more Zn-containing intermetallic compounds
(Mg2Zn3, MgZn2) in the system even at small Zn concentrations. However, such phases were not found in the studied alloys. EPMA analysis (Figure 4.13) shows that the Zn concentration in the interdendritic area is significantly higher than its concentration in the α-Al dendrites. The increasing of the Zn content in the alloy its concentration in α-Al dendrites increases proportionally. The results of EDX analysis (Figure 4.15) show that Zn concentration in α-Al in Z6 alloy (2.1 at. % Zn) amount to 1 at.%. The volume fraction of the T-AlMgZn intermetallic phase grows with an increase in the Zn content in the alloys (Table 4.6). The Z3 alloy has the critical Zn concentration (0.75 at.% Zn) after which the volume fraction of the T-AlMgZn phase sharply increases. The differences between the calculated (in the Thermo-Calc) and measured (by the partially automated image analysis) values of the volume fraction of the T-AlMgZn phase may be caused by the low levels of the diffusion rate at low temperatures that make further growth of the phase difficult. Figure 4.25 represents a comparison of the structure, distribution of the T-AlMgZn phase and Zn segregation in the middle part of the casted plate and in the region near surface of the plate of Z5 alloy. As can be seen the T-AlMgZn phase has non-homogeneous distribution throughout the plate. The region near surfaces is almost free from the T-AlMgZn phase (Figure 4.25 a-c), since crystallization rate is higher at the walls of the mold (Td = 200ºC), in contrast to the middle part of the plate (Figure 4.25
87 d-f). The results of the EPMA analysis (Figure 4.25 c, f) also show that Zn is distributed near the surfaces more homogeneously, enriches the dendritic core without formation of the T-AlMgZn phase.
Under the surface the Under
a) b) c)
Middle part
d) e) f) Figure 4.25. Differencies in the structure near the surface of the casting (a-c) and in the middle part (d-f) of Z5 alloy [179]: a,c – SEM images; b,e – results of image analysis (T- AlMgZn is black, α-AlMnFeSi – is outlined); c,f – EPMA maps.
Table 4.9 represents changes in the volume fraction of T-AlMgZn phase during cooling (Thermo-Calc calculations). The results from the image analysis can be compared with the values of the volume fraction at the initial stages of the formation of the T-AlMgZn phase calculated by Thermo-Calc. The growth of the T-AlMgZn phase stops in the temperature range 250-300ºC depending on Zn content in the alloy.
Table 4.9. Volume fraction of the T-phase at different temperatures at solid state [%] [179]: Measured Calculated with Thermo-Calc software RT 70ºC 100ºC 140ºC 180ºC 220ºC 250ºC 280ºC 300ºC 310ºC 320ºC 330ºC Z1 - 2.45 1.76 1.41 1.02 0.62 0.11 ------Z2 0.2 2.89 2.45 2.16 1.76 1.31 0.76 0.24 - - - - - Z3 0.1 3.19 2.89 2.65 2.28 1.84 1.28 0.75 0.10 - - - - Z4 0.9 3.60 3.40 3.24 2.94 2.55 2.02 1.49 0.83 0.30 - - - Z5 1.0 3.90 3.81 3.70 3.48 3.15 2.65 2.14 1.49 0.98 0.69 0.38 0.05 Z6 1.3 3.96 3.94 3.90 3.78 3.52 3.08 2.62 2.00 1.50 1.22 0.92 0.60
88
Significant inhibition of the diffusion processes at the temperatures below 300°C also explains the absence of the β-AlMg phase in the structure of M and Z1 alloys (the β-AlMg precipitates even at lower temperatures than T-AlMgZn, see Table 4.10).
Table 4.10. Volume fraction of the β-AlMg phase at different temperatures at solid state [%] Alloys Calculated with Thermo-Calc software RT 30ºC 40ºC 50ºC 60ºC 70ºC 80ºC 90ºC 100ºC 110ºC M 2.55 2.37 2.17 1.95 1.70 1.42 1.12 0.79 0.43 0.04 Z1 0.42 0.25 0.06 ------Z2-Z6 ------
b) Precipitation strengthening TEM-study of the M-alloy, S- and C-series shows that the structure of the α-Al dendrites is similar to the wrought 6XXX alloys with excess of Mg [98,134]. As already mentioned, the excess of Mg decreases the precipitates number density and leads to coarsen of the precipitates (less amount of fine GP-I and β" and more coarse β’ and β precipitates are formed) [98,134]. Thus, in the M-alloy as well as in S- and C-series of alloys in the structure of the α-Al dendrites only β’- and β-type precipitates are observed. The precipitates observed in the Z-series of alloys have size, morphology and distribution similar to GP-I zones in the alloys of Al-Zn-Mg system [153]. The observed precipitates have size 1-5 nm with rounded “coffee-bean” morphology. Increasing of the Zn content changes neither size nor morphology of the GP-I zones. On the other hand, an increase in the Zn concentration in the alloy affects the density of GP-I zones in the α-Al. The density strongly increases until concentration 0.75 at. % Zn. However, further growth of Zn concentration doesn’t affect the density of GP-I zones. It can be concluded that concentration of 0.75 at. % Zn content for studied alloys is the critical point for the formation of the strengthening precipitates.
c) Mechanical properties. It is known that the mechanical properties of Al-alloys can be affected by grain size, ILS, intermetallics and precipitates. The reduction of ILS and DAS (or grain size) have a positive effect on the mechanical properties (UTS, A) [194]. Zn addition does not significantly affect either ILS or DAS in the Z-series of alloys (Table 4.6). In this regard, only formation of new T-AlMgZn phase and nanoscale precipitates can explain all changes in the mechanical properties observed in the Z-series of alloys. A significant increase in the HB, YS and UTS in comparison with M alloy (and even with C-series of alloys) can be explained by the influence of Zn on the eutectic cells as well as on the α-Al dendrites (Figure 4.15). The microhardness of the α-Al dendrites normally increases with increase of the density of the GP-I zones (in Z4-Z6 alloys with a constant density of the GP-I zones
89 microhardness increases insignificantly) while the increasing of the microhardness of the eutectic cells is extreme (Figure 4.22 c). Such a strong strengthening effect can be explained by a presence of a large amount of Zn in the eutectic (Figure 4.13), thus producing a strengthening effect. The values of the UTS and elongation show good correlation with the density of the GP-I zones in the α-Al solid solution. The Z3 alloy (0.75 at. % Zn) has highest UTS value (Figure 4.21) that coincides with the maximum density of the GP-I zones. On the other hand, the level of elongation is significantly lower than in the M alloy. Such behavior of these mechanical characteristics can be best explained by precipitation strengthening mechanism that, as known [60], increases strength while deteriorating ductility. Further degradation of elongation is associated with a sharp increasing in the volume fraction of the fusible T-AlMgSi phase, and the formation of the coarse Mn-containing sludge particles lead to a further degradation of the levels of elongation. Sludges have relatively high hardness values (up to 14.5 GPa [144]), however, unlike the compact α-AlMnFeSi intermetallics the coarse sludges are brittle (similarly to coarse Cr-containing particles [144]) and thus can initiate cracks under tension (Figure 4.26). From the literature [192,193] it is well known that the coarse -AlMnFeSi phase (sludge) has a negative effect on the whole set of tensile properties while effect on hardness is insignificant.
a) b) c) Figure 4.26. Fine compact Mn-containing phases in the fracture of the M alloy (a) and the coarse Mn-containing phase with cracks in the Z4 alloy (b-c) [179]
4.6. Summary
In this study, the microstructural evolution and the mechanical properties of Al-Mg-Si-Mn alloys with additional alloying by Cr, Sc and Zn were studied. The main conclusions are summarized as follows: 1) The as-cast structure of M alloy consists of three phases, namely the α-Al dendrites,
Al+Mg2Si eutectic and α-AlMnFeSi phase, which is distributed in the interdendritic space. The composition of α-Al dendrites consists of 2.8 at.% Mg, 0.3 at.% Si, 0.3 at. %Mn and the
amount of Fe, Cu, Ti, Zn is less than 0.1 at. % of each. Al+Mg2Si eutectic has a lamellar and/or fibers morphology. The eutectic melting point determined by DSC is 591°C that is in
90
good agreement with calculated phase diagrams. The α-AlMnFeSi phase can take hexagonal, cubic and star-like shape. The composition of α-AlMnFeSi intermetallics consists of 75 at. % Al, 10 at.% Mn, 7 at. %Si, 3 at. % Fe. 2) The M alloy with the composition Al-5.7Mg-2.6Si-0.6Mn has good combination of strength-
ductility level. Al+Mg2Si eutectic is the main strengthening phase in the structure. The
hardness of Al+Mg2Si eutectic is 35% higher than hardness of α-Al dendrites. Unlike Al- Mg-Si wrought alloys, cast alloys are not predisposed to a significant strengthening of α-Al phase by β’’-precipitates (due to excess of Mg). 3) The additional alloying the M alloy by Cr did not lead to the expected increase in mechanical properties. 0.1 wt. % Cr very slightly increased the strength of the alloy and slightly decreased the elongation. The higher amount of Cr (0.2 wt. %) leads to the formation of the brittle sludge particles that decrease elongation even more. The Cr addition leads to the formation of semicoherent Mn,Cr-containing (non-strengthening) dispersoids in the α-Al dendrites. 4) Sc addition significantly increase strength and hardness of the alloy in the as-cast state (up to 20%) with simultaneous increase in the elongation. The concentration 0.1 wt.% Sc in the as-cast state is enough for enriching α-Al solid solution and further increasing Sc content doesn’t lead to further significant strengthening. Sc addition leads to the formation of new
Al3Sc intermetallics in the as-cast state. 5) Zn addition promotes precipitation strengthening by formation of η-type precipitates in α- Al dendrites. The maximum of strength (380 MPa) was achieved with 1.8 wt. % Zn, which coincides with the maximum density of GP-II zones. Further increasing of Zn concentration doesn’t lead to the further increasing of the strength, while hardness still growth (up to 120 HB in alloy with 5 wt. %). In alloys with Zn concentration 1.2 wt. % or higher, new T- AlMgZn phase is formed. Precipitation and new intermetallics lead to the significant decreasing of elongation. The summary of the mechanical properties of the studied alloys in as-cast state is shown in Figure 4.27.
91 400 260 1.8-5 380 1.2-1.8 240 wt.% Zn wt.% Zn 1.2-1.8 360 1.8-5 0.1-0.2 220 0.6-1.2 wt.% Zn
wt.% Zn 0.1-0.2 340 wt.% Zn wt.% Sc 200 0.6-1.2 wt.% Zn wt.% Sc
320 0.1-0.2 180 , , [MPa]
wt.% Cr 0.1-0.2 Base YS UTS, [MPa] UTS, 300 160 Base wt.% Cr alloy 280 alloy 140
260 120 1 2 3 4 5 6 7 8 9 10 11 12 1 2 3 4 5 6 7 8 9 10 11 12 Elongation, [%] a) Elongation, [%] b)
Figure 4.27. Summary of the mechanical properties of the all studied alloys in the as-cast state: a) tensile strenght vs. elongation; b) yield strength vs. elongation
92
CHAPTER 5 Heat treatment of Al-Mg-Si-Mn-X alloys
In this chapter, the detailed characterisation of the microstructure and mechanical properties of the studied series of alloys after heat treatment were presented. Two types of heat treatment were applied to the studied alloys: artificial aging from as-cast state (one-step heat treatment) and solution treatment with subsequent artificial aging (two-step heat treatment). The changes after heat treatment in the microstructure were investigated using scanning electron microscopy (SEM) with EDX analysis and electron-probe microanalysis (EPMA). The structure of the α-Al solid solution was studied using transmission electron microscopy (TEM). To determine the effect of changes in the microstructure and in the composition of α-Al solid solution on the mechanical properties the common set of properties (HB, UTS, YS, A) and local properties of the α-Al dendrites and the eutectic cells were measured.
5.1. One-step heat treatment
In this section, the changes that occur during artificial aging from as-cast state were considered. Artificial aging as a one-step heat treatment can be applied for parts that were casted with a relatively fast solidification. The processes occur in the solid solution, and consist of forming nanoscale precipitates via decomposition of a supersaturated solid solution (ssss) [10,84].
5.1.1. Artificial aging of the base alloy
As already mentioned in previous paragraphs, despite the fact that the M alloy includes elements capable of forming precipitates (Mg, Si), however, the features associated with the ratio of these alloying elements (>1.73) do not allow the formation of a sufficient amount of precipitates to obtain significant strengthening effect as a result of artificial aging. Figure 5.1 a shows changes in hardness of the M alloy with different temperatures and aging times. Thus, from the diagram it is clear that the heat treatment of the M alloy does not have a significant effect on the hardness. The most noticeable changes were recorded at the aging temperature of 175ºC and time in the range of 2-3 hours. However, the hardness in this mode is only 5-7% higher than in the cast state. Figure 5.1 b shows results of tensile properties of the M alloy with different temperatures of aging for time with peak hardness (3 h). As it can be seen, the increase in tensile strength is insignificant, but yield strength has changes that are more prominent. Nevertheless, the elongation is decreased.
93 125ºC 175ºC 225ºC 325ºC YS, [MPa] UTS, [MPa] A, [%] 88 330 Peak hardness 86 280 84 230
82 Hardness, [HB] Hardness,
180 [%] Elongation,
80 Strenght, [MPa] Strenght,
78 130 0 3 6 9 12 15 0 50 100 150 200 250 300 Aging time, [h] a) Aging temperature, [ºC] b) Figure 5.1. Mechanical properties of the M alloy as a function of aging temperature: a) Brinell hardness (different aging time); b) Tensile properties of the M alloy (time – 3 h)
a) b) c) Figure 5.2. SEM-images of the microstructure of the M alloy after aging at: a) as cast state; b) 175ºC; c) 325 ºC
a) b) c)
d) e) f) Figure 5.3. TEM bright-field images of the α-Al dendrites in M alloy after artificial aging: a-c) aging at 175ºC; d-f) aging at 325ºC; (aging time 3 h)
94
Results of SEM and TEM investigations of the M alloy don’t show any significant changes neither in microstructure nor in the structure of the α-Al dendrites during treatment in temperature range of 125-325ºC. The microstructure consists of α-Al dendrites, Al-Mg2Si eutectic colonies, and α-AlMnFeSi intermetallics. The aging in such temperature range and time does not lead to changes in the morphology of all phases (Figure 5.2). The structure of α-Al dendrites consists of β’and stable β phases similarly to the as-cast state (Figure 5.3). However, in some α-Al dendrites small amounts of β”-precipitates were found after aging at 175ºC (Figure 5.3 c).
5.1.2. Artificial aging of the Al-Mg-Si-Mn alloys with Sc and Cr additions
Similarly to the M alloy, alloys additionally doped by Cr and Sc don’t show any changes in the microstructure after heat treatment in temperature range of 125-325ºC. However, alloys with Sc addition show increasing density of nanoscale precipitates after high-temperature aging treatment (325ºC) (Figure 5.4). The alloys from the C-series as well as S-series show approximately similar values of hardness after aging with temperatures 125-225ºC (Figure 5.5). However, the aging with temperature 325ºC shows different behavior. Thus, the hardness of Sc-containing series of alloys, especially S2 alloy (with 0.2 wt.% of Sc) increases significantly and reaches values about 120 HB.
a) b)
c) d) Figure 5.4. TEM bright-field images of the α-Al dendrites in S-series of alloys after artificial aging: a) S1 after 225ºC; b) S1 325ºC; c) S2 225ºC; d) S2 325ºC;
95 125 125 125ºC 125ºC 120 120 175ºC 175ºC 115 225ºC 115 225ºC 110 325ºC 110 325ºC
105 Peak hardness 105 Peak hardness 100 100
Hardness, [HB] Hardness, 95 95 Hardness, [HB] Hardness, 90 90 85 85 80 80 0 3 6 9 12 15 0 3 6 9 12 15 Aging time, [h] a) Aging time, [h] b) 125 125 125ºC 125ºC 120 175ºC 120 175ºC 115 225ºC 115 225ºC 325ºC 325ºC 110 110 Peak hardness 105 105
100 100 Hardness, [HB] Hardness,
95 [HB] Hardness, 95 90 90 Peak hardness 85 85 80 80 0 3 6 9 12 15 0 3 6 9 12 15 c) Aging time, [h] Aging time, [h] d) 125 125ºC 120 175ºC 115 225ºC 325ºC 110 105 100
Hardness, [HB] Hardness, 95 Peak hardness 90 85 80 0 3 6 9 12 15 Aging time, [h] e) Figure 5.5. Brinell hardness of the studied alloys as a function of aging temperature and time: a) C1 (0.1 wt.% Cr); b) C2 (0.2 wt.% Cr); c) S1 (0.1 wt.% Sc); d) S2 (0.2 wt.% Sc) e) M alloy (for comparison)
Figure 5.6 represents results of tensile tests of S- and C-series of alloys for different aging temperatures for the time of maximum hardness values (3 h). Both alloys from C-series show similar changes and similar levels of all tensile properties. Thus, maximum values of UTS and YS belong to the aging temperature range 175-225ºC. That coincides with the lowest values of elongation. S1 alloy (0.1 wt. % Sc) shows approximately similar results of all tensile tests for all aging temperatures. Only minor increasing of YS can be observed for aging temperature higher than 125ºC. S2 alloy shows the highest values of UTS and YS, especially for aging at 325ºC.
96
YS, [MPa] UTS, [MPa] A, [%] YS, [MPa] UTS, [MPa] A, [%] 330 9 330 9 310 310 290 290 7 7 270 270 250 250 230 5 230 5 210 210
190 190
Strenght, [MPa] Strenght, [MPa] Strenght, Elongation, [%] Elongation, 3 [%] Elongation, 3 170 170 150 150 130 1 130 1 0 100 200 300 0 100 200 300 Aging temperature, [ºC] a) Aging temperature, [ºC] b)
YS, [MPa] UTS, [MPa] A, [%] YS, [MPa] UTS, [MPa] A, [%] 11 380 10 380 9
330 8 330 7
280 6 280
5
Strenght, [MPa] Strenght, [MPa] Strenght, Elongation, [%] Elongation, 230 4 [%] Elongation, 230 3
180 2 180 1 0 100 200 300 0 100 200 300 Aging temperature, [ºC] c) Aging temperature, [ºC] d)
Figure 5.6. Tensile properties of the studied alloys as a function of aging temperature for peak- hardness aging time: a) C1; b) C2; c) S1; d) S2
Figure 5.7 shows results of tensile properties for aging with 325ºC as a function of the time. The increasing of the aging time reduces both tensile and yield strength simultaneously with increasing elongation for all studied alloys.
300 400 13 M M S1 280 S1 380 S2 C1 S2 12 C2 260 C1 360 C2 240 340 11
220 320 10 200 300 9
180 280 YS. [MPa] YS.
160 [MPa] UTS, 260 Elongation, [%] Elongation, 8 140 240 M S1 7 120 220 S2 C1 C2 100 200 6 1 2 5 15 1 2 5 15 1 2 5 15 Aging time, [h] a) Aging time, [h] b) Aging time, [h] c) Figure 5.7. Tensile properties of the studied alloys as a function of the aging time for 325 ºC: a) yield strength; b) tensile strength; c) elongation
97 5.1.3. Structural changes in Al-Mg-Si-Mn-Zn alloys during artificial aging
According to the equilibrium phase diagram (see paragraph 2.3) T-AlMgZn is a fusible phase and can be dissolved at relatively low heat-treatment temperatures. In Z1 (0.2 at.% Zn) T- AlMgZn phase dissolves at 230ºC; in Z2 (0.5 at.% Zn) - at 260ºC; in Z3 (0.7 at.% Zn) - at 290ºC; in Z4 (1.2 at.% Zn) - at 320ºC; in Z5 (1.7 at.% Zn) - at 340ºC; in Z6 (2.1 at.% Zn) - at 350ºC. Thus, heat treatment of the Zn-containing group of the studied alloys, in contrast to the Sc- and Cr- containing groups, should be carefully selected. However, heat treatment with certain temperatures (according to the results of Thermo- Calc simulation) and even a bit higher is not enough for full dissolution of T-AlMgZn phase. Thus, even heat treatment at 325ºC doesn’t dissolve T-phase even in Z2 alloy (Figure 5.8 a). Nevertheless, after 3 h aging at 350ºC T-AlMgZn phase starts disintegrating even in Z6 alloy with higher Zn content (Figure 5.8 b). Heat treatment with 375ºC leads to dissolution of the T-AlMgZn phase in the alloys with Zn content up to 5 wt.% (2.1 at. %). Figure 5.9 shows selected EPMA maps for the studied alloys.
a) b) Figure 5.8. SEM-image of the microstructure of the alloys after aging: a) Z2 alloy at 325ºC during 3 h; b) Z6 alloy at 350ºC; (aging time 3 h)
Figure 5.10 shows changes in the concentrations of the main elements in the α-Al dendrites during heat treatment with temperatures in the range 225-520ºC. Thus, Mg (Figure 5.10 a) and Si (Figure 5.10 b) have similar behavior and their concentration change stepwise from 2.8 to 2.6 at.% for Mg and from 0.4 to 0.2 at.% for Si. Zn concentration, in contrast to Mg and Si increases with increasing treatment temperature and in Z6 alloys for 400 ºC reaches 1.7 at.%. Further increasing treatment temperature does not lead to further enriching Al dendrites by Zn. It should be mentioned that heat treatment does not lead to any significant changes in the Mn concentration (as well as Fe, Ti and other impurities).
98
Z3 Z6
350ºC
a) b)
375ºC
c) d) Figure 5.9. SEM images and EPMA maps of Z-series of alloys after 3 h aging: a) Z3 after 350ºC; b) Z6 after 350ºC; c) Z3 after 375ºC; d) Z6 after 375ºC
After aging at 175ºC several types of nanoscale precipitates (in addition to the already described for as-cast state) were found in the α-Al dendrites of the alloys from Z-series (Figure 5.11, Figure 5.12): Highly dispersed needle-shaped precipitates (Figure 5.11 a,b) with a mean size of 1-
4 nm and elongated in the <100>Al (belongs to the Zn-containing GP-I zones),
99 Semi-coherent η' with a mean size of 10-25 nm (Figure 5.11 c,d). For comparison of the effects of the existing dispersed phases, their density was measured (Table 5.1). The highest density of the coherent precipitates (GP-I, II) have Z3 and Z4 alloys. In alloys with higher Zn content, the density of GP-zones decreases. However, density of the η' steadily increases with increasing zinc concentration in the alloys.
Figure 5.10. Changes in the concentrations of elements in the α-Al dendrites during heat treatment as a function of temperature and Zn content (values are rounded to tenths) [at.%]: a) Mg; b) Si; c) Zn
a) b)
c) d) Figure 5.11. TEM bright-field images of different type of the precipitates in Z3 alloy (artificial aging at 175ºC, aging time - 3 h)
100
a) b)
c) d)
e) f) Figure 5.12. TEM bright-field images of studied alloys (artificial aging at 175ºC during 3 h): a) Z1 (0.2 at. % Zn); b) Z2 (0.5 at. % Zn); c) Z3 (0.7 at. % Zn); d) Z4 (1.2 at.% Zn); e) Z5 (1.7 at. % Zn); f) Z6 (2.1 at. % Zn)
Table 5.1. Density of the precipitates in the Zn-containing series of alloys Size, Density, [μm-3] Type [nm] Z1 Z2 Z3 Z4 Z5 Z6 GP-I (Zn,Mg) 2-5 4740 10500 10380 10860 9020 10020 Coherent GP-II (Zn,Mg) 1-3 21960 27600 45820 36180 24340 15440 Semi- η'-Zn2Mg 10-25 - - 500 1500 2370 3050 coherent β'-Mg2Si 50-200 1700 1770 1640 1150 1100 920
5.1.4. Mechanical properties of Al-Mg-Si-Mn-Zn alloys after artificial aging
Results from the hardness measurements as a function of aging time are presented in Figure 5.13 (only temperatures that do not lead to phase transformations were chosen). As can be seen
101 from the diagrams, the hardness of the alloys strictly depends on the Zn concentration in the alloy. Aging at 125 ºC and 175ºC (Figure 5.13 a, b) lead to the presence of peak hardness values, after which the hardness of the alloys decreases. The maximum hardness values are achieved in the range of 2-5 hours of aging. Effect of aging at 125ºC is lower than at 175ºC. Aging at higher temperatures (225ºC and higher) decreases hardness for all studied alloys from Z-series (Figure 5.13 c,d).
125 125
120 Z6 120 Z6 Z5 Z5 115 115 Z4 110 Z4 110 Z3
105 105 Z3 Z2
100 Z2 100 Hardness. [HB] Hardness. Z1
95 [HB] Hardness. 95 Z1 M 90 90 Peak hardness Peak hardness M 85 85
80 80 0 3 6 9 12 15 0 3 6 9 12 15 Aging time, [h] Aging time, [h] 125 125
120 120
115 115
110 110 Z6 Z6 105 105 Z5 Z5 100 100
Z4 Z4 Hardness. [HB] Hardness. 95 Z3 [HB] Hardness. 95 Z3 Z2 Z2 90 90 M M 85 85 Z1 Z1 80 80 0 3 6 9 12 15 0 3 6 9 12 15 Aging time, [h] Aging time, [h] Figure 5.13. Brinell hardness of the Z-series of alloys as a function of aging time after aging at a) 125ºC; b) 175ºC; c) 225ºC; d) 325ºC
Figure 5.14 represents results of tensile tests of Z-series of alloys for different aging temperatures for the time of maximum hardness values (3 h). All alloys show similar changes of all tensile properties. Thus, maximum values of strength, similarly to hardness, belong to the aging
102 temperature range 125-175ºC. The lowest values of the strength (tensile and yield) belongs to the heat treatment at 225 ºC.
Z1 Z2 Z3 Z4 Z5 Z6 380 370 360 350 340 330 320 310 T=125ºC 300 T=175ºC Tensile strength, MPa 290 T=225ºC 280 0 0,5 1 1,5 2 Zn content in alloy, at.% a) Z1 Z2 Z3 Z4 Z5 Z6 290 270 250 230 210 190 T=125ºC
Yield strength, Yield MPa T=175ºC 170 T=225ºC 150 0 0,5 1 1,5 2 Zn content in alloy, at.% b) Z1 Z2 Z3 Z4 Z5 Z6 10 T=125ºC 9 T=175ºC 8 T=225ºC 7 6 5 4 Elongation, Elongation, % 3 2 1 0 0,5 1 1,5 2 Zn content in alloy, at.% c) Figure 5.14. Tensile properties of the studied alloys as a function of Zn concentration a) tensile strength; b) yield strength; c) elongation; (aging time – 3h)
Hardness and tensile tests give integral values of properties, while the structure of the alloy is non-homogeneous and consists of two main phases – α-Al dendrites and Al-Mg2Si eutectic. Thus, the microhardness (Figure 5.15) of these phases was measured to understand the contribution of the separate phases to the mechanical properties of the alloys.
103 As can be seen from the Figure 5.15 the behavior of the microhardness of α-Al dendrites; and Al-Mg2Si eutectic is different. Thus, the microhardness of the Al-Mg2Si eutectic (Figure 5.15 a) strictly depends on the Zn concentration in the alloy as well as aging temperature. The highest values of the microhardness for all alloys belongs to the aging at temperatures in the range 125-175ºC (similarly to the hardness and strength). With increasing of the aging temperature microhardness of the eutectic decreases. On the other hand, microhardness of the α-Al dendrites aging at 175ºC leads to the presence of the obvious peak in the hardness values.
160-180 140-160 120-140 100-120 80-100 200-220 180-200 160-180 140-160
Peak hardness
220 Peak hadness 220 Overaging Overaging 200 200
180 180
160 160
140 140 Hardness HV0.01 Hardness
Hardness HV0.01 Hardness Z6 Z6 120 Z5 120 Z5 Z4 Z4
100 Z3 100 As-cast Z3 Alloys
Z2 Z2 Alloys As-cast 80 Z1 80 Z1 0 125 0 125 175 225 175 225 325 325 350 375 a) 350 375 b)
Figure 5.15. Microhardness HV0.01 of a) α-Al dendrites; b) Al-Mg2Si eutectic cells as a function of Zn-content and aging temperature. Aging time – 3h
5.2. Two-step heat treatment
Full T6 heat treatment consists of two steps: high temperature solution treatment, quenching and artificial aging. The main aim of the solution treatment is to dissolve non- equilibrium phases and saturate the solid solution with dissolved elements. Immediately after the solution treatment, artificial aging of the samples follows. The aim of the artificial aging is the formation of nanodispersed strengthening precipitates in the solid solution.
5.2.1. Two-step heat treatment of the base alloy
The changes of the structure and properties of Al-Mg-Si-Mn alloys during heat treatment are the result of several processes that occur simultaneously during heating. The first process is the spheroidization of eutectic fibrous (Figure 5.16). A higher solution-treatment temperature leads to a faster eutectic cells decomposition into separate Mg2Si particles and faster coarsen of these
104 particles. Table 5.2 shows the average size of these particles depending on the treatment temperature.
a) b)
c) d) Figure 5.16. Microstructure* of the M alloy in a) as-cast state and b) after solution treatment at 480°C during 1.5 h; c) after solution treatment at 520°C during 1.5 h; d) after solution treatment at 570°C during 1.5 h
Table 5.2. Average size of separate Mg2Si particles as a function of treatment temperature (treatment time – 1.5 h) Temperature, [°C] 480 520 570 Diameter, [μm] 0.59 0.66 0.82
During the solution treatment, several processes occur in the solid solution. First of all, high temperature treatment leads to dissolving all strengthening precipitates (including GP-zones,
β’, η’) that form during natural and artificial aging (see 4.3). However, equilibrium β-Mg2Si phase after solution treatment is still present in the microstructure (Figure 5.17 a-d). The second process is formation of Mn-containing dispersoids (Figure 5.17 c-f). Figure 5.17 c-f show that dispersoids after high temperature heat treatment form and attach to spherical Mg2Si particles and to the equilibrium β-Mg2Si phase.
* The SEM images have been made at the Technische Universität Darmstadt, Darmstadt, Germany 105 a) b)
c) d)
e) f) Figure 5.17. TEM images of the α-Al dendrites in M alloy after solution treatment at: a-d) 520ºC; e-f) 570ºC (without further aging)
It can be seen from Figure 5.18 a-b that during aging after solution treatment at the temperature of 520ºC, strengthening phases are practically not formed. However, solution treatment at 570ºC with further aging (Figure 5.18 c-d) leads to formation of GP-II zones with the high density in the solid solution. Figure 5.19 and Table 5.3 show mechanical properties of the M alloy after treatment with different temperatures. Results of the hardness measurements (Figure 5.19) show that aging after low temperature solution treatment (e.g. temperature range 480-520ºC) doesn’t lead to the significant increasing in the hardness of the M alloy. On the other hand, aging after solution treatment at 570ºC increases the hardness by 25% in comparison to solutionized state and by 12% in comparison to as-cast state.
106
a) b)
c) d) Figure 5.18. TEM images of the α-Al dendrites in M alloy after solution treatment at: a-b) 520ºC; c-d) 570ºC; (with further aging at 175ºC; ST time – 1.5 h; aging time – 3 h)
Results of tensile test (Table 5.3) confirm the concerns predicted in subchapter 2.4. Thus, despite the high values of hardness after solution treatment at 570ºC (and further aging) the values of tensile strength and elongation are extremely low. The values of the yield strength remain high.
Table 5.3. Tensile properties of the M alloy after full T6 treatment Temperature [ºC] UTS [MPa] YS [MPa] A [%] Solution treatment Artificial aging 175 291.3 ± 2.8 235.8 ± 1.1 3.8 ± 0.4 570 225 293.7 ± 0.5 201.3 ± 9.3 6.0 ± 2.6 175 280.6 ± 8.7 184.8 ± 2.6 9.5 ± 2.3 520 225 242.0 ± 11.5 137.0 ± 16.6 11.5 ± 1.3 175 241.6 ± 2.1 98.3 ± 0.8 16.8 ± 3.8 480 225 228.1 ± 1.0 93.7 ± 0.9 18.0 ± 2.9
The situation is different with treatment at lower temperatures (480-520ºC). Thus, despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18%. The following paragraphs consider the effect of heat treatment on alloys alloyed with Cr and Zn. Sc-containing series of alloys did not show any weighty results (that is in good agreement with the results and conclusions made in the CHAPTER 2), however, some results are reported in the work. 107 100 ST570ºC/AA225ºC 95 ST570ºC/AA175ºC 90 ST570ºC/AA125ºC 85 ST520ºC/AA225ºC 80 ST520ºC/AA175ºC 75 ST520ºC/AA125ºC Hardness, [HB] Hardness, 70 ST480ºC/AA225ºC 65 Solution treatment at 480ºC 60 ST480ºC/AA175ºC
55 ST480ºC/AA125ºC 0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 Aging time, [h]
Figure 5.19. Brinell hardness of the M alloy as a function of aging time and temperatures after full T6 treatment with solution treatment temperatures a) 480ºC; b) 520ºC; c) 570ºC
5.2.2. Two-step heat treatment of the Al-Mg-Si-Mn-Cr alloys
According to the equilibrium phase diagram (see paragraph 2.3) Al7Cr is a refractory phase and can be dissolved at relatively high heat-treatment temperatures. In C1 (0.1 wt. % Cr) Al7Cr phase dissolves at 460ºC; in C2 (0.2 wt. % Cr) at 520ºC. Thus, heat treatment of the Cr-containing group of the studied alloys should be carefully selected (due to the problems described in 2.4). As can be seen form Figure 5.20 b heat treatment with temperature of 520 ºC even for 3h
(according to the results of Thermo-Calc simulation) is not enough for the dissolution of Al7Cr phase. Applied heat treatment doesn’t increase Cr concentration in the α-Al solid solution. Al-
Mg2Si eutectic in the both alloys from C-series after heat treatment changes in the same way as the M alloy (eutectic cells disintegrated into separate spherical particles). Heat treatment also doesn’t affect morphology and distribution of Mn-containing intermetallics. Cr addition doesn’t affect aging behavior of the β-type precipitates in the studied alloys. Similarly to the M alloy, solution treatment at 520ºC with further aging doesn’t lead to the formation high density β-type precipitates (GP-zones and β’’). However, Cr addition affects formation of (Mn,Cr)-containing dispersoids (Figure 5.21). Thus, in the C-series of alloys dispersoids after solution treatment can be found inside α-Al dendrites and with increasing Cr content in the alloy the density of the dispersoids increases (Table 5.4).
108
a) b) Figure 5.20. SEM images of C1 (a) and C2 (b) alloys after solution treatment at 520ºC (treatment time – 3 hour)
a) b)
c) d) Figure 5.21. TEM images of the α-Al dendrites in C-series of alloys after solution treatment at 520ºC (time – 1.5 h) and further aging at 175ºC (aging time – 3 h). a,b) C1; c,d) C2
Table 5.5 represents results of the hardness and tensile tests of the alloys from C-series alloys for different aging temperatures for the time of maximum hardness values (3 h). The higher aging temperature the lower strength (and hardness) and the higher elongation. In general, the behavior of changes is similar to the M alloy (Table 5.3).
Table 5.4. Density of the (Mn,Cr)-dispersoids Alloys M C1 C2 Density, [μm-3] * 1700 5400 * Cannot be measured due to the specific particles distribution (particles can be found only close to boundaries between α-Al dendrites and Al-Mg2Si eutectic but not inside the dendrites Figure 5.17 c-f)
109 Table 5.5. Mechanical properties of the C-series of alloys after full T6 heat treatment (solution treatment at 520ºC for 1.5 hour, artificial aging for 3 hour) Alloy Aging temperature [ºC] HB UTS [MPa] YS [MPa] A [%] 175 78.3 301.4 ± 5.2 174.5 ± 4.5 12.8 ± 1.0 C1 225 75.1 265.1 ± 3.3 131.1 ± 7.3 13.7 ± 1.3 325 69.1 255.2 ± 3.8 101.8 ± 1.2 14.5 ± 3.3 175 81.2 287.5 ± 5.1 171.1 ± 9.7 8.4 ± 0.8 C2 225 77.3 256.2 ± 3.6 133.9 ± 11.3 8.7 ± 1.0 325 67.3 242.9 ± 6.5 102.5 ± 0.4 9.5 ± 1.5
5.2.3. Structural changes in Al-Mg-Si-Mn-Zn alloys
Similarly to the M alloy and alloys from C-series, Al-Mg2Si eutectic in all alloys from Z- series during solution treatment disintegrate to separate spherical particles (Figure 5.22). Table 5.6 represents the size of the Mg2Si particles in alloys with different Zn content after solution treatment at 480ºC for 1.5 h. As it can be seen, increasing Zn content in the alloy leads to the faster coarsening of Mg2Si particles during solution treatment (Table 5.6). EPMA maps (Figure 5.22) show that heat treatment at 480ºC leads to dissolution of the T- AlMgZn phase in the alloys with Zn content up to 5 wt.% (2.1 at. %) and redistribution of Zn with enriching (Figure 5.10 c, Figure 5.23 a) the Al dendrite core and impoverishment of the interdendritic space. Heat treatment also doesn’t affect morphology, size and composition (Figure 5.23 b) of Mn-containing phase. After solution treatment at 480ºC and aging at 175ºC in the α-Al dendrites of the alloys from Z-series several types of nanoscale precipitates were found (Figure 5.24, Table 5.7). For comparison of the effects of the existing dispersed phases, their density was measured (Table 5.7). Unlike aging from as-cast state of the studied alloys, when several types of the precipitates influence on the hardness and tensile properties, when aging after solution treatment one precipitate type dominates. Thus, in the alloys Z1 and Z2 the main strengthening phase (with highest density) is GP-I zones, in alloys Z3-Z5 GP-II, in alloy Z6 semi-coherent η'-Zn2Mg. In all alloys equilibrium β-Mg2Si phase is still present. Nevertheless, β’-Mg2Si was not found.
Table 5.6. Measured average circular Mg2Si particle diameter after ST at 480 ºC, [µm] M Z1 Z2 Z3 Z4 Z5 Z6 0.59 0.58 0.68 0.68 0.76 1.05 0.98
Table 5.7. Density of the precipitates in the Zn-containing series of alloys, [μm-3] Type Size, nm Z1 Z2 Z3 Z4 Z5 Z6 GP-I (Zn,Mg) 2-5 12 100 23 600 * - - - GP-II (Zn,Mg) 1-3 - - 30 200 30 300 43 400 * η'-Zn2Mg 5-25 - - - * 1 870 34 200 * Density is too low for measuring, but such kinds of the precipitates may exist in the alloy
110
a) b) c)
d) e) f) Figure 5.22. SEM images with EPMA maps of the Zn distribution after solution treatment at 480ºC (treatment time – 1.5 h)
3,5 20 80 Mg Zn Si Mn 3,0 Si Mn Zn Fe Al 2,5 15 60 2,0 10 40 1,5 1,0
5 20
Elements, at.% Elements, at.%
0,5 Aluminium,at.% 0,0 0 0 0,0 0,5 1,0 1,5 2,0 a) 0,0 0,5 1,0 1,5 2,0 b) Zn in alloy, at.% Zn in alloy, at.% Figure 5.23. EDX analysis of phases in studied alloys after solution treatment at 480ºC, 1.5 h: a) α-Al, b) α-AlMnFeSi
111 a) b)
c) d)
e) f) Figure 5.24. TEM bright-field images of the α-Al dendrites in Z-series of alloys (solution treatment at 480ºC, 1.5 h and aging at 175ºC, 3 h): a) Z1 (0.2 at. % Zn); b) Z2 (0.5 at. % Zn); c) Z3 (0.7 at. % Zn); d) Z4 (1.2 at. % Zn); e) Z5 (1.7 at.% Zn); f) Z6 (2.1 at. % Zn)
5.2.4. Mechanical properties of Al-Mg-Si-Mn-Zn alloys
Results from the hardness measurements as a function of aging time are presented in Figure 5.25. As can be seen from the diagrams, higher solution temperature leads to higher hardness values. The hardness of the alloys strictly depends on the Zn concentration in the alloy. Aging at 175ºC temperatures (Figure 5.25 a, b) leads to the presence of peak hardness, after which the hardness of the alloys decreases. The maximum hardness values are achieved in the range of 2-3 hours of aging. Aging at 225ºC temperatures (Figure 5.25 c, d) leads to lower hardness values
112 in comparison with aging at 175ºC. The maximum hardness was achieved in the range of 3-5 hours of aging.
Artificial aging 175ºC Artificial aging 225ºC
110 110
105 Z5 105 Z6 Z4
100 100 Z6 95 95 Z5 Z3 90 90 Z4 Z2 Z3 85 Z1 85
M Z2 Hardness, [HB] Hardness, Hardness, [HB] Hardness, 80 80 Z1 Peak hardness M 75 75 Peak hardness
70 70
Solution treatment 520ºC Solution treatment 65 65 ST520ºC/AA175ºC ST520ºC/AA225ºC 60 60 0 3 6 9 12 15 0 3 6 9 12 15 Aging time, [h] a) Aging time, [h] b)
110 110 ST480ºC/AA175ºC ST480ºC/AA225ºC 105 105 Peak hardness
100 100
95 95 Z6 Peak hardness 90 Z5 90 85 85 Z6
Z4 Hardness, [HB] Hardness, Hardness, [HB] Hardness, Z5 80 80 Z4 75 75 Z3 Z3 Z2 70 70 Z2
Solution treatment 480ºC Solution treatment Z1 65 65 Z1 M M 60 60 0 3 6 9 12 15 0 3 6 9 12 15 Aging time, [h] c) Aging time, [h] d)
Figure 5.25. Brinell hardness of the Z-series of alloys after full T6 treatment as a function of aging time at a) ST520ºC+AA175°C, b) ST520ºC+AA225°C, c) ST480ºC+AA175°C, d) ST480ºC+AA225°C
Figure 5.26 represents results of tensile tests of Z-series of alloys for different heat treatment tempers for the time of maximum hardness values (3 h of aging). As it can be seen from Figure 5.26 that temperature of ST has the lower influence on the mechanical properties than aging temperature. Combination of different ST and aging temperatures influences on alloys with different Zn concentration differently. Thus, alloys can be divided into two groups (similar to as- cast state). First group of alloys with Zn concentration is in the range 0-0.75 at. %, and second with Zn concentration in the range 0.75-2.1 at. %. Each alloy from the first group (Z1, Z2, and Z3)
113 doesn’t show significant changes in the UTS no matter what aging and solution treatment temperatures are used. UTS values for each alloy vary in the range of 10%. On the other hand, the yield strength increases with increasing both solution treatment and aging temperatures. The values of the elongation change insignificantly no matter what aging and solution treatment temperatures are used.
Artificial aging 175ºC Artificial aging 225ºC 380 380 Z1 Z2 Z3 Z4 Z5 Z6 Z1 Z2 Z3 Z4 Z5 Z6 360 360 340 340 320 320 300 300 347 346 321 322 280 280 312 313 311 312 311 303 305 302 300 304 295 291 295 296 292 283 286 281
260 260 272 274
Tensile strenght, Tensile [MPa] strenght, Tensile [MPa] strenght, 240 240 ST480+AA175 ST520+AA175 a) ST480+AA225 ST520+AA225 b) 300 300 Z1 Z2 Z3 Z4 Z5 Z6 Z1 Z2 Z3 Z4 Z5 Z6
250 250
200 200 174 242 241 161 148 152 217 143 217 150 191 197 150 137 180 185 185 185
122 167 173 129 Yield strenght, Yield [MPa] Yield strenght, Yield [MPa] 114 117 114 108 100 100 ST480+AA175 ST520+AA175 c) ST480+AA225 ST520+AA225 d) 22 22 Z1 Z1 20 20 18 Z2 18 Z2 16 Z3 16 Z3 14 Z4 14 Z4 12 Z5 12 Z5 20,8 10 19,519,0 Z6 10 19,5 Z6 16,716,717,0 16,7 16,5 15,7 8 8 14,5
Elongation, Elongation, [%] 16,5
Elongation, Elongation, [%] 6 6
4 7,0 4 7,3 7,7 7,5 6,0 6,3 6,5 6,7 2 5,7 5,5 5,5 2 5,6 0 0 ST480+AA175 ST520+AA175 e) ST480+AA225 ST520+AA225 f) Figure 5.26. Tensile properties of the studied alloys a,b) tensile strength; c,d) yield strength; e,f) elongation. Aging time – 3h
Tensile strength of the alloys from the second group (Z4, Z5, and Z6) are more dependent on aging temperature compared to the first group. Thus, higher aging temperature leads to lower UTS and YS values. Aging at 175ºC after ST at 480ºC and 520ºC has the stronger effect on the strength of the second group of alloys than aging at 225ºC. The higher Zn content is, the stronger
114 is the effect of aging at 175ºC. The elongations for all three alloys remain at the same level for all treatment tempers. In general, higher strength values correspond to lower elongation values.
5.3. Discussion
5.3.1. Heat treatment of the Al-Mg-Si-Mn base alloy and after Cr addition
a) Artificial aging from as-cast state. The heat treatment of the Al alloys usually applies to obtain specific mechanical properties. There are various heat treatment modes for this reason. The most common types of heat treatment of cast Al alloys are homogenization, annealing and precipitation strengthening, involving solution heat treatment, quenching, and aging (see 1.2). The main prerequisite of aging is the presence of ssss in the alloy that is why Al-Mg-Si alloys are usually subject to aging after high temperature solution treatment and quenching. Nevertheless, some results were published on possibilities of artificial aging Al-Mg-Si alloys from as-cast state [166,195], moreover, rapid cooling after HPDC makes prospective aging from the as-cast state [196,197]. Figure 5.1 shows changes in the mechanical properties of the M alloy with different aging temperatures. Such heat treatment of the M alloy does not have a significant effect on the hardness. The most noticeable changes were recorded at the aging temperature of 175ºC and time in the range of 2-3 hours. However, the increase in hardness in this mode is only 5-7% higher than the as-cast state. Such an insignificant effect is concerned with the fact that the M alloy belongs to alloys with an excess of Mg content. As mentioned above, excess Mg drastically reduces the solubility of Mg2Si in the α-Al, which in turn reduces the hardening potential (by the formation of nanoscale precipitates) and also increases the stability of the β’ and equilibrium β phases (which do not have a significant hardening effect) [98,134]. In general, it can be argued that excess Mg makes the properties of the M alloy very stable at the elevated temperatures (which makes the material promising for further research at elevated temperatures). The results of the mechanical properties are also confirmed by the microstructural investigation, which shows that the aging treatment in the studied temperature range (125-325ºC) and time (up to 15 hours) does at least not change the morphology of the phases (Figure 5.2). The structure of α-Al dendrites mostly consists of β’and stable β phases similarly to the as-cast state (Figure 5.3). However, in some α-Al dendrites only small amounts of β”-precipitates were found after aging at 175ºC (Figure 5.3 c). Similarly to the M alloy, alloys additionally doped by Cr have changes neither in microstructure nor in precipitation sequence after heat treatment in temperature range of 125- 325ºC. Thus, any new Cr-containing precipitates were not found. Both C1 and C2 alloys show similar to the M alloy behavior of all studied mechanical properties after aging (Figure 5.5 - Figure
115 5.7). Despite this, aging from the as-cast state remains promising for Al-Mg-Si-Mn alloys additionally alloyed by other elements [198]. b) Full T6 heat treatment. The changes of the properties of Al-Mg-Si-Mn alloys during high temperature solution treatment (as first stage of the full T6 treatment) are the result of the changes of structural components that occur during heating. The first process is the spheroidization of eutectic Mg2Si phase (Figure 5.16). A higher solution-treatment temperature leads to a faster eutectic cells decomposition into separate Mg2Si particles and faster coarsening of these particles
(Table 5.2). Due to the fact, that Mg2Si eutectic is the main strengthening phase in this type of alloys, the process of its disintegration reduces the hardness and strength that coincides with the increase of elongation of the alloys [178]. Similar tendency was reported for the casting Al-Mg- Si-Mn alloys casted into the permanent mold [10,82]. Together with eutectic disintegration, the solution treatment usually leads to dissolution the strengthening β’’-Mg2Si precipitates that might be formed in the as-cast state of Al-Mg-Si alloys after solidification [25,26]. However, as it was already discussed in CHAPTER 4, due to the specific Mg/Si ratio of the M alloy, the amount of the strengthening β’’-Mg2Si precipitates in as-cast is negligible. Solution treatment of the Al5Mg2SiMn permanent mold casting leads to the formation of β-AlMnFeSi dispersoids [10,82]. Their lack of coherence with the matrix leads to a reduction in the hardness of the alloys still more (along with decomposition of eutectic lamellas). On the other hand, in the studied HPDC Al-Mg- Si-Mn alloys α-AlMnFeSi dispersoids were found. This type of dispersoids is semi-coherent with the Al-matrix [170] that can slightly affect the hardness of the alloys. This explains the higher hardness values of the HPDC compared to the regular permanent casting alloys [10,82]. Artificial aging after solution treatment causes formation of the strengthening precipitates again [10] that improve strengthening properties of the alloys. Nevertheless, different solution treatment temperatures have different effects. Thus, aging after low temperature solution treatment (480-520ºC) doesn’t lead to the significant increasing in the hardness of the M alloy (Figure 5.19). On the other hand, aging after solution treatment at 570ºC has a more noticeable effect on the hardness. This may be due to the formation of a larger number of lattice vacancies by quenching (“quench-in” vacancies) that promote diffusion of the solute atoms, which in turn, promote formation of atoms clusters and GP-zones [100,101,199,200]. Although the higher hardness values after solution treatment at 570ºC, this kind of treatment is not recommended. Thus, the values of tensile strength and elongation are extremely low. Analysis of the stress-strain curves of the specimens after 570ºC solution treatment shows that tensile specimens break prematurely. This indicates the appearance of defects (porosity) in alloys after heat treatment. Also, such a heat treatment is not recommended by the manufacturer [89].
116
The situation is different with treatment at lower temperatures (480-520ºC). Thus, despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18% as confirmed by the results presented in [178]. Similarly to the situation with aging from as-cast state, alloys from C-series after full T6 treatment have same changes as in the M alloy. Moreover, solution treatment at 520ºC during 1.5 hour (and even 3 hour) doesn’t lead to dissolution of the Cr-containing phase in C2 (as it was predicted by Thermo-Calc software). A primary cause of this is the low diffusion rate of Cr. The dissolution of this phase could improve the tensile properties of the alloy, but this did not happen. On the other hand, heat treatment with higher temperatures (for dissolution of Cr-containing particles) is not recommended. Nevertheless, solute Cr in α-Al dendrites during solution treatment promote formation of (Mn,Cr)-containing dispersoids, which, however, do not significantly affect the mechanical properties of the alloy.
5.3.2. Heat treatment of the Al-Mg-Si-Mn-Sc alloys
In addition to the high cost, Sc has some other disadvantages as alloying element for Al- Mg-Si alloys. Even addition of only 0.1 wt. % Sc makes it meaningless to use a full T6 treatment for Al5Mg2SiMn (a full research is presented in [201]). Also, aging temperatures for the formation of nano-dispersed Al3Sc precipitates are significantly higher than the temperatures used to obtain maximum hardness values in Al-Mg-Si alloys by the formation of β’’-Mg2Si precipitates [147]. Such temperatures (250-350ºС) lead to overaging and make more stable β’- and equilibrium β-
Mg2Si phases (which do not significantly affect the strength). On the other hand, as already mentioned, the M alloy has rather stable properties and structure during aging up to 325ºС. Knipling et al. [50] reported the criteria for designing alloys for application at elevated temperatures. Sc was chosen as one of the most promising alloying additives for this reason. This supposition is also confirmed by the work Lenczowski et al. [198] where AlMg3Si1 alloy containing Sc+Zr in T5 state shows ultimate tensile strength (UTS) of 270 MPa at RT and 265 MPa at 250ºC. Current research has shown that aging of the alloys from S-series at a temperature 325ºС from the as-cast state is a promising heat treatment. Thus, strength of S2 alloy (with 0,2 wt. % Sc) increased up to 380 MPa (11% higher than as-cast state of S2 and 30% higher than M alloy). Such temperature does not lead to any structural changes in the alloy and even does not change the distribution of Sc in the alloy. Accordingly, such a significant increase in hardness and strength can only be explained by the formation of nano-dispersed precipitates Al3Sc. Such an increase in properties is in good agreement with the existing studies concerning the alloying of Al-alloys by Sc [198,202].
117
5.3.3. Heat treatment of the Al-Mg-Si-Mn-Zn alloys
a) Artificial aging from as-cast state. First of all it should be mentioned that the hardness of the alloys even after aging treatment (Figure 5.13) strictly depends on the Zn concentration in the alloy. Different aging temperatures and different duration of the treatment lead to different effects on the hardness, strength and elongation properties. Aging at 125ºC and 175ºC temperatures lead to the presence of peak hardness values (in the range of 2-5 hours of aging), after which the process of overaging occurs and hardness of the alloys decreases. The treatment at the higher temperatures, 225ºC and 325ºC (Figure 5.13 c,d), leads to a very rapid processes of overaging and softening of the alloys (already after the first hour of aging). Aging at temperatures of 125-175ºC increases the strength of the Z1 and Z2 alloys in comparison with the as-cast state (Figure 5.14). At the same time, the strength of the Z3-Z6 alloys practically did not change. Aging at a temperature of 175ºC increased the yield strength of the alloys, while aging at 125ºC did not change it, and aging at 225ºC lowered it. Aging at a temperature of 225ºC lowered the strength of the entire series of alloys. It should be noted that the general behavior of the properties (as a function of the Zn concentration) and values of the tensile strength of the studied alloys after heat treatment is similar to the as-cast state. Analysis of local mechanical properties (Figure 5.15) confirms the results of the hardness tests and the fact of the rapid overaging at temperatures above 175°C. The nature of the changes in the hardness values of the α-Al and Al-Mg2Si eutectic depending on the temperature of aging is different. Thus, unlike Al-Mg2Si eutectic (Figure 5.15 b), α-Al has a pronounced peak of the hardness at the temperatures of 125-175°C (Figure 5.15 a). The microhardness of the Al-Mg2Si eutectic practically does not change in the temperature range 125-175°C and remains at the as-cast state level. This, as well as the absence of other structural changes in the alloys during aging at the studied temperatures, indicates that the main influence on the properties is exerted by changes in the solid solution. TEM investigations of the studied alloys have shown a significant change in the structural composition of the solid solution (Figure 5.11, Figure 5.12). Thus, aging at 175°С resulted in a reduction in the number of GP-I zones and the formation of GP-II zones (Table 5.1), which were not found in the as-cast state (Figure 4.20, Table 4.7). This corresponds to the theory that GP-II zones are formed in Al-Mg-Zn alloys after aging at the temperatures above 70°C [153,157,160]. GP-II zones after aging were the most high-density precipitates in the alloys (Table 5.1). In addition, in alloys with a high Zn content (0.7 at. % and above) also the η'-precipitates were detected. Their density increases with an increase in the Zn content in the alloy and a simultaneous decrease in the density of the GP-II zones. This is caused by the increase in the Zn concentration
118 in the α-Al, as a result the precipitates with a higher Zn content are formed (GP-zones have ratio Zn/Mg~1, n'-precipitates have ratio Zn/Mg~3, see Table 2.2). The best combination of the mechanical properties (strength-elongation) show alloys with higher density of the GP-zones - Z2 and Z3 alloys (Table 5.1). At the same time, Z4 alloy, which has a similar to Z2 alloy level of GP- zones density, shows significantly lower values of the elongation. Similar to the as-cast state, this is due to the presence of the brittle phases (T-phase and coarse Mn-containig phase) that are not dissolved during the aging process. b) Full T6 heat treatment. Solution treatment at 480 and 520°C for 90 minutes lead to decomposition of Al-Mg2Si eutectic (similarly to the M alloy and in good agreement with a large number of studies on structural changes in Al-Mg-Si alloys in the solution treatment process), dissolving of the T-AlMgZn phase and the GP-zones in the solid solution (that also agrees well with minor studies of Al-Mg-Si-Zn casting alloys [129,168]). Such treatment decreases hardness and increases elongation [178]. The mechanical properties after the full T6 treatment (that includes following artificial aging) is significantly affected by the Zn concentration. Hardness (Figure 5.25), tensile strength (Figure 5.26 a) and the yield strength (Figure 5.26 b) increase with increasing Zn concentration. The formation of certain precipitates is influenced by the concentration of Zn in a-Al (the higher the treatment temperature, the higher the Zn content in α-Al, see Figure 5.10), the aging temperature (the higher temperature, the faster the overaging occurs) and the aging time. Despite the fact that the overaging temperature for Al-Mg-Si alloys is usually 225°C and higher [178], the results of the current study have shown that this temperature is not suitable for overaging for alloys additionally doped with Zn. Aging at such temperature after solution treatment at temperatures of 480°C and 520°C leads to ultra-fast overaging (similarly to aging from as-cast state). On the other hand, overaging at a temperature of 175°C occurs after 5 hours of treatment.
The microhardness (Figure 5.27) of the α-Al dendrites and Al-Mg2Si eutectic was measured to understand the contribution of the separate phases to the mechanical properties of the alloys. As can be seen from the Figure 5.27 the behavior of the microhardness of α-Al dendrites and Al-Mg2Si eutectic after full T6 heat treatment is different in comparison to T5 treatment
(Figure 5.15). Thus, the differences in the microhardness between α-Al and Al-Mg2Si eutectic are not as significant as after T5. This is primarily due to the disintegration of the eutectic, as well as the uniform Zn distribution in the sample after solution treatment (Figure 5.9, Figure 5.22).
119 130-140 120-130 130-140 120-130 110-120 100-110 110-120 100-110 90-100 80-90 90-100 80-90
140 140 130 130 120 Z6 120 Z6 Z5 Z5 110 110 Z4 Z4 100 Z3 100 Z3 90 Z2 90 Z2
Z1 Z1 Microharness, HV0,01 Microharness, 80 HV0,01 Microharness, 80 225 225 175 175 0 a) 0 b)
Figure 5.27. Microhardness HV0.01 of a) α-Al dendrites; b) Al-Mg2Si eutectic cells as a function of Zn-content and aging temperature. (ST: 520ºC, 1.5 h; AA time – 3h)
ST 350ºC ST 375ºC 380 380 Z1 Z2 Z3 Z4 Z5 Z6 Z1 Z2 Z3 Z4 Z5 Z6 360 360 340 340 320 320 300 300
325 323 280 316 280 312 315 306 306 306 307 307 311 303 305 303 299 300 300 305 302 293 291
260 260 289 291 289
Tensile strenght, [MPa] Tensile strenght, [MPa] 240 240 ST350+AA175 ST350+AA225 a) ST375+AA175 ST375+AA225 b) 300 300 Z1 Z2 Z3 Z4 Z5 Z6 Z1 Z2 Z3 Z4 Z5 Z6
250 250
200 200
150 155 148 215 221 150 192 150 133 174 175 183 127
117 116 Yield Yield strenght, [MPa] Yield Yield strenght, [MPa] 112 115 142 112 113 113 122 122 131 132 132 100 100 ST350+AA175 ST350+AA225 c) ST375+AA175 ST375+AA225 d) 22 22 20 Z1 Z2 Z3 Z4 Z5 Z6 20 Z1 Z2 Z3 Z4 Z5 Z6 18 18 16 16 14 14 12 12 10 10 17,5 8 8 15,3 14,014,3 12,3 Elongation, [%] 12,212,0 Elongation, Elongation, [%] 6 6 15,5 11,5 10,0 10,7 9,0 9,7 4 7,3 4 7,0 7,3 7,2 5,5 4,5 2 4,0 4,0 3,5 3,2 3,1 2 0 0 ST350+AA175 ST350+AA225 e) ST375+AA175 ST375+AA225 f) Figure 5.28. Tensile properties of the studied alloys a,b) tensile strength; c,d) yield strength; e,f) elongation. Aging time – 3h
120
The obtained tensile test results are in good agreement with the results of the works [129,168]. Li et. al [181] obtained after solution treatment at 490°C and further aging at 180°C an increase in yield strength and tensile strength with increasing Zn content, similar to the results from this study. The differences in the results are associated with the used casting method (Li et al. used conventional permanent mold casting), higher Mg content and the absence of Mn in the studied alloys. The latter led to a more gradual decrease in the elongation values (with increasing Zn concentration) since a coarse brittle Mn-containing phase did not form. In the studied alloys (Z4-Z6), the Mn-containing phase does not dissolve even during solution treatment at a temperature of 520°С. Despite the fact that these alloys after heat treatment have the elongation values increased by more than two times, however, it remained significantly lower than in alloys Z1-Z3. From the Figure 5.9 it can be seen that the T-AlMgZn phase in the studied alloys dissolves during heat treatment at temperatures in the range of 350-375°C. This transformation has a significant effect on the tensile properties of the alloys (Figure 5.28). A significant increase in the elongation of the all alloys confirms the assumption that the T-AlMgZn phase has a harmful effect on the ductility of the alloys. The values of the mechanical properties of the alloys after solution treatment at a temperature of 375°С are comparable to the mechanical properties obtained after treatment at a temperature of 480°С. Regardless of the solution treatment temperature for the studied alloys, the preferred temperature of aging is 175°C, at which the most interesting combinations of properties can be achieved.
5.4. Summary
The effect of aging from as-cast state and after solution treatment on the microstructure and mechanical properties were investigated. The main conclusions are summarized as follows: 1) Artificial aging from as-cast state (T5 temper) of the M alloy doesn’t promote the precipitation strengthening and in the studied temperature range doesn’t affect microstructure. 2) Artificial aging at 325°C from as-cast state of the Sc-containing alloys has the most significant improvement in the mechanical properties. Alloy with 0.2 wt.% Sc exhibit the
highest strength with value of 380 MPa. Nano-dispersed precipitates Al3Sc have been found in Sc-series after artificial aging. 3) Artificial aging at 125-225ºC from as-cast state of the Zn-containing alloys doesn’t lead to any changes in the microstructure of alloys, however, leads to formation of several types of the Zn-containing precipitates in the α-Al solid solution: GP-I, GP-II, η’. The density of the
β’-Mg2Si precipitates increases in comparison to the as-cast state. The optimal temperature
121 for artificial aging from as-cast state is in the range 125-175ºC for 2-5 h. Higher temperature leads to overly rapid overaging and lowering all set of the mechanical properties. 4) The solution treatment leads to the eutectic spheroidization and to strong degradation of UTS and YS and an to increase in ductility. The further aging at the temperatures 125-325ºC doesn’t lead to intense precipitates formation, and in particular the full T6 treatment of the M alloy doesn’t improve its properties. 5) The solution treatment changes the microstructure of the studied alloys. The T-AlMgZn phase dissolves into the α-Al phase during solution treatment at temperatures higher than 375ºC, and subsequently precipitated during aging to form nanosized strengthening phase. The preferred temperature of aging is 175°C, at which the most interesting combinations of properties can be achieved. Higher temperature leads to overly rapid overaging and lowering strength, especially for alloys with higher Zn content (2-5 wt. % Zn).
T5 (AA from as-cast state) Full T6 (ST+AA) a) 400 c) 340 3-5 wt.% Zn 380 0.2 wt.% 1.2-1.8 320 Sc wt.% Zn
360 300 1.2-1.8
3-5 wt.% 0.6-1.2 wt.% Zn 340 Zn wt.% Zn 0.1 wt.% 280 0.6-1.2 0.1-0.2 Sc wt.% Zn 0.1-0.2 wt.% Sc 320 260 wt.% Cr 0.1-0.2 M alloy
wt.% Cr UTS, [MPa] UTS, UTS, [MPa] 300 240
280 220 M alloy 260 200 1 2 3 4 5 6 7 8 9 10 11 12 9 10 11 12 13 14 15 16 17 18 19 20 21 Elongation, [%] Elongation, [%] b) 300 d) 220 3-5 wt.% 3-5 wt.% 280 Zn 200 Zn 0.2 wt.%
260 Sc 180
240 1.2-1.8 160 0.1-0.2 wt.% Zn wt.% Sc 1.2-1.8 220 0.6-1.2 0.1 wt.% 140 0.1-0.2 wt.% Zn
wt.% Zn Sc wt.% Cr S, S, [MPa] S, S, [MPa] 0.1-0.2
0.6-1.2 Y Y 200 120 wt.% Cr M alloy wt.% Zn 180 100
M alloy 160 80
1 2 3 4 5 6 7 8 9 10 11 12 9 10 11 12 13 14 15 16 17 18 19 20 21 Elongation, [%] Elongation, [%] Figure 5.29. Summary of the mechanical properties of the all studied alloys in the different tempers: a, b) T5 (aging from as-cast state): M alloy and C-series - aging at 175-225°C, 3 h; S-series – aging at 325°C, 2 h; Z-series – aging at 125-175°C, 3 h c, d) T7 (aging after solution treatment): ST at 480-520°C aging at 175-225°C, 3 h
122
SUMMARY
In this study, the microstructural evolution and the mechanical properties of Al-Mg-Si-Mn alloys with additional alloying by Cr, Sc and Zn were studied. As can be seen from the analysis of all mechanical properties of the studied alloys, despite the implementation of a commercial analogue of the C-series of alloy in the production (Maxxalloy-Ultra), Cr has the smallest effect among the studied elements in all tempers. Sc has the most significant influence on the base composition (M alloy) in as-cast state among all elements (considering a very small alloying amount). Subsequent artificial aging from the as-cast state further increases the strength of the alloys, however, it decreases elongation. The biggest drawback of alloying the studied system with Sc is its very high price (addition of 0.2 wt.% Sc increases the price per kg Al-alloy by 4 €). Zn on the contrary has an affordable price. Alloys with the Zn concentration in the range of 1.2-1.8 wt.% (Z2 and Z3 alloys) are the most promising, have good properties both in the as-cast state and after heat treatment in the T5 and T6 modes. However, the Zn concentration must be controlled, since the higher the concentration, the more amount of the harmful T-phase is formed and less-hardening precipitates are formed. The reasons for the change in properties in the studied alloys were the formation of new primary phases in the as-cast structure and the formation of nanosized precipitates during the heat treatment. It was found that coarse phases (such as Al7Cr in C-series and Mn-containing sludge particles in Z-series of alloys) are brittle and significantly reduce tensile properties of alloys, especially ductility. The Al3Sc phase (S-series of alloys) did not negatively affect the ductility of the alloys. As expected, M alloy showed a good response to additional alloying with elements that enhance the precipitation strengthening. Thus, the greatest impact on the mechanical properties was affected by nanoscale strengthening precipitates formed in α-Al solid solution. Both elements
Sc (forms nanosized precipitates Al3M in combination with Al) and Zn (forms nanosized precipitates in combination with Mg) improve hardness and strength of the M-alloy. Both elements increase response of the M alloy on the heat treatment, however, they affect it differently (due to different solvability of the primary Al3Sc phase and T-AlMgZn phase). The T-AlMgZn phase dissolves during relatively low solution treatment temperature and with subsequent aging promotes formation of nanoscale Zn-containing GP-zones. The primary Al3Sc phase doesn’t dissolve even during high temperature solution treatment. However, Sc that remains in α-Al during solidification is sufficient for the formation of a nanoscale strengthening precipitates during aging.
123 Despite the fact that the Cr addition to the M-alloy leads to the formation of a dispersed Cr-containing phase in a-Al, it does not have a significant strengthening effect (similar to Cr- containing wrought alloys).
Perspective for further work The current thesis expands the understanding of the mechanisms of the influence of the additional alloying of the Al-Mg-Si-Mn system by Sc, Cr, and Zn and can be used as a basis for further alloy development. Potential topics for further development are.
- Additional alloying of the base composition with elements that can produce Al3M precipitates in α-Al matrix (Hf, Ni or other). - Study of the Al-Mg-Si-Mn alloys as a potential candidate for high temperature application.
124
REFERENCES
[1] International Aluminum Institute, Global Aluminium Recycling : A Cornerstone of Sustainable Development, International Aluminum Institute. (2009) 1–36. http://www.world-aluminium.org/media/filer_public/2013/01/15/fl0000181.pdf. [2] European Aluminium Association, Market Overview, (2018). https://www.european- aluminium.eu/activity-report-2017/market-overview/. [3] A. Kelkar, R. Roth, J. Clark, Automobile Bodies: Can Aluminum Be an Economical Alternative to Steel?, JOM. (2001) 28–32. https://doi.org/10.1002/cite.201650064. [4] I.J. Polmear, Light Alloys. From Traditional Alloys to Nanocrystals, 4th ed., Elsevier/Butterworth-Heinemann, 2006. https://www.elsevier.com/books/light- alloys/polmear/978-0-7506-6371-7. [5] European Aluminium Association, Aluminum Content in Cars. Summary Report. Public version., 2016. https://european-aluminium.eu/media/1721/european-aluminium-ducker- study-summary-report_sept.pdf. [6] The International Aluminium Institute, World Aluminium — Primary Aluminium Production, (2018). http://www.world-aluminium.org/statistics/#linegraph (accessed November 23, 2018). [7] P. Zapp, G. Rombach, W. Kuckshinrichs, The future of automotive aluminium, Light Metals: Proceedings of Sessions, TMS Annual Meeting (Warrendale, Pennsylvania). (2002) 1003-1010. [8] J.G. Kaufman, E.L. Rooy, Aluminum Alloy Castings: Properties , Processes , and Applications., ASM International. (2004) 340. https://doi.org/10.1017/CBO9781107415324.004. [9] J. Campbell, Castings, 2nd ed., Elsevier/Butterworth-Heinemann, Oxford, 2003. [10] O. Trudonoshyn, M. Puchnin, K. Mykhalenkov, Features of structure formation and changes in the mechanical properties of cast Al-Mg-Si-Mn alloy with the addition of (Ti + Zr), Acta Polytechnica. 55 (2015) 282–290. https://doi.org/10.14311/AP.2015.55.0282. [11] J.G. Kaufman, Introduction to Aluminum Alloys and Tempers, ASM International, 2000. https://doi.org/10.1361/iaat2000p001. [12] The Aluminium Association, International Alloy Designations and Chemical Composition Limits for Wrought Aluminum and Wrought Aluminum Alloys With Support for On-line Access From: Aluminum Extruders Council Use of the Information, (2015). https://www.aluminum.org/sites/default/files/Teal Sheets.pdf. [13] X. Wang, M. Guo, J. Luo, J. Zhu, J. Zhang, L. Zhuang, Effect of Zn on microstructure, texture and mechanical properties of Al-Mg-Si-Cu alloys with a medium number of Fe-rich phase particles, Materials Characterization. 134 (2017) 123–133. https://doi.org/10.1016/j.matchar.2017.10.012. [14] M. Nakai, T. Eto, New aspect of development of high strength aluminum alloys for aerospace applications, Materials Science and Engineering A. 285 (2000) 62–68. https://doi.org/10.1016/S0921-5093(00)00667-5. [15] C. Giummarra, B. Thomas, R.J. Rioja, New Aluminium Lithium alloys for aerospace applications, Proceedings of the Light Metals Technology Conference. (2007) 2–5. [16] T. Warner, Recently-Developed Aluminium Solutions for Aerospace Applications, Materials Science Forum. 519–521 (2006) 1271–1278. https://doi.org/10.4028/www.scientific.net/MSF.519-521.1271. [17] E.A. Starke, J.T. Stalcy, Application of Modern Aluminum Alloys to Aircraft, 32 (1996) 131–172. https://doi.org/https://doi.org/10.1016/0376-0421(95)00004-6. [18] N.E. Prasad, R.J.H. Wanhill, Aerospace Materials and Material Technologies. Volume 1: Aerospace Materials, Springer, Singapore, 2017. https://doi.org/10.1007/978-981-10-2134- 3. [19] J.T. Staley, D.J. Lege, Advances in aluminium alloy products for structural applications in
125 transportation, Le Journal de Physique IV. 03 (1993) C7-179-C7-190. https://doi.org/10.1051/jp4:1993728. [20] S. Ji, D. Watson, Z. Fan, M. White, Development of a super ductile diecast Al-Mg-Si alloy, Materials Science and Engineering A. 556 (2012) 824–833. https://doi.org/10.1016/j.msea.2012.07.074. [21] Alcan Marin, Aluminium Alloys in marine application, in: Aluminium Users’ Guide, Alcan, Carquefou Cedex, 2017: pp. 33–52. http://www.almet- marine.com/images/clients/EN/Aluminium-users-guide/Ch03-aluminium-alloys-in- marine-applications.pdf. [22] J.R. Davis, Alloying: understanding the basics, ASM International, Ohio, 2001. https://doi.org/10.5860/choice.39-5209. [23] M.H. Avazkonandeh-Gharavol, M. Haddad-Sabzevar, H. Fredriksson, Effect of partition coefficient on microsegregation during solidification of aluminium alloys, International Journal of Minerals, Metallurgy and Materials. 21 (2014) 980–989. https://doi.org/10.1007/s12613-014-0999-1. [24] R. Nadella, D.G. Eskin, Q. Du, L. Katgerman, Macrosegregation in direct-chill casting of aluminium alloys, Progress in Materials Science. 53 (2008) 421–480. https://doi.org/10.1016/j.pmatsci.2007.10.001. [25] F. Czerwinski, Heat Treatment Conventional and, n.d. https://doi.org/10.1039/c2dt31079a. [26] H. Rolling, Heat Treatment – A Brief Introduction Heat, AluReport. 1 (2014) 10–15. https://www.amag-al4u.com/fileadmin/user_upload/amag/Downloads/AluReport/EN/AR- 2014-1-EN-S10-15-Waermebehandlung.pdf. [27] D.M. Stefanscu, J.R. Davis, J.D. Destefani, Metals Handbook, 9th ed., ASM International, 1998. [28] European Aluminium Association, Aluminium in Cars – Unlocking the Light-Weighting Potential, Brussels, 2012. https://www.european-aluminium.eu/media/1326/aluminium-in- cars-unlocking-the-lightweighting-potential.pdf. [29] M.S. Salleh, M.Z. Omar, J. Syarif, M.N. Mohammed, An Overview of Semisolid Processing of Aluminium Alloys, ISRN Materials Science. 2013 (2013) 1–9. https://doi.org/10.1155/2013/679820. [30] E.J. Vinarcik, High Integrity Die Casting Processes, John Wiley & Sons, Inc, New York, 2002. https://doi.org/10.1186/1472-6963-14-S2-O26. [31] R.N. Lumley, R.G. Odonnell, D.R. Gunasegaram, M. Givord, Heat treatment of high- pressure die castings, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science. 38 (2007) 2564–2574. https://doi.org/10.1007/s11661-007-9285-4. [32] M. Schwankl, Al-Al-Verbundguss – Technologische Grundlagen und werkstoffkundliche Charakterisierung der stofflichen Anbindung von Aluminiumknetlegierungen im Aluminium-Druckguss, FAU University Press, Erlangen, 2017. [33] L. Chauke, H. Möller, U.A. Curle, G. Govender, Industrial Heat Treatment of R-HPDC A356 Automotive Brake Callipers, Solid State Phenomena. 192–193 (2013) 533–538. https://doi.org/10.4028/www.scientific.net/SSP.192-193.533. [34] E. Macherauch, H.-W. Zoch, Praktikum in Werkstoffkunde, Springer Vieweg, 2014. https://doi.org/10.1007/978-3-658-05038-2. [35] M.A. Meyers, K.K. Chawla, Mechanical Behavior of Materials, 2nd ed., Cambridge University Press, Cambridge, 2008. [36] B. Holmedal, E. Nes, H. Ekstro, Strengthening Mechanisms in Solid Solution Aluminum Alloys, 37 (2006) 1999–2006. [37] O. Prach, C. Minnert, K.E. Johanns, K. Durst, A new nanoindentation creep technique using constant contact pressure, Journal of Materials Research. 34 (2019) 2492–2500. https://doi.org/10.1557/jmr.2019.188. [38] R.W. Cahn, P. HAASEN, Physical Metallurgy. Volume 3, North-Holland, Amsterdam, 1996. https://doi.org/0 444 89875 1.
126
[39] L.F. Mondolfo, Metallography of aluminum alloys, ВUTTER WORTHS, London, 1976. [40] L.F.F. Mondolfo, Aluminum alloys : structure and properties, Butterworths, London, 1979. [41] H. Baker, ASM Handbook: Alloy Phase Diagrams, ASM International, 1992. [42] S. Wang, S. Chen, K. Matsuda, T. Kawabata, J. Nakamura, S. Ikeno, K. Kawakita, H. Takagi, T. Yamashita, Effect of Mn or Fe Addition on Age-Hardening Behaviour of Al - Mg 2 Si Alloys, 53 (2012) 1521–1528. https://doi.org/10.2320/matertrans.MAW201219. [43] Y. Meng, J. Cui, Z. Zhao, Y. Zuo, Effect of vanadium on the microstructures and mechanical properties of an Al-Mg-Si-Cu-Cr-Ti alloy of 6XXX series, Journal of Alloys and Compounds. 573 (2013) 102–111. https://doi.org/10.1016/j.jallcom.2013.03.239. [44] S. Ji, W. Yang, F. Gao, D. Watson, Z. Fan, Effect of iron in Al-Mg-Si-Mn ductile diecast alloy, Minerals, Metals and Materials Series. (2016) 317–322. https://doi.org/10.1007/978- 3-319-65136-1_55. [45] I. Charit, R.S. Mishra, Low temperature superplasticity in a friction-stir-processed ultrafine grained Al-Zn-Mg-Sc alloy, Acta Materialia. 53 (2005) 4211–4223. https://doi.org/10.1016/j.actamat.2005.05.021. [46] M.S. Kabir, T.I. Minhaj, E.A. Ashrafi, M. Islam, The Influence of Ageing Time and Temperature on the Structure and Properties of Heat Treated A201 . 0 Aluminum Alloy, International Journal of Recent Technolofy and Engineering. 3 (2014) 78–83. [47] F. Wei, J. Li, R. Hu, H. Kou, Influence of 1.0 wt% Li on precipitates in Al-Zn-Mg-Cu alloy, Chinese Journal of Aeronautics. 21 (2008) 565–570. https://doi.org/10.1016/S1000- 9361(08)60175-2. [48] T.B. Bhat, V.S. Arunachalam, Strengthening mechanisms in alloys, Prec. Indian Acad. Sci. (Engg. Sci.). 3 (1980) 275–296. [49] J. Roesler, H. Harders, M. Baeker, Mechanical Behaviour of Engineering Materials, Springer, Wiesbaden, 2007. https://doi.org/10.1007/978-3-540-73448-2. [50] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep- resistant aluminum-based alloys – A review, Z. Metallkd. 97 (2006) 246–265. [51] G.A. Edwards, K. Stiller, G.L. Dunlop, M.J. Couper, The precipitation sequence in Al-Mg- Si alloys, Acta Materialia. 46 (1998) 3893–3904. https://doi.org/https://doi.org/10.1016/S1359-6454(98)00059-7Get. [52] N. Afify, A.-F. Gaber, G. Abbady, Fine Scale Precipitates in Al-Mg-Zn Alloys after Various Aging Temperatures, Materials Sciences and Applications. 02 (2011) 427–434. https://doi.org/10.4236/msa.2011.25056. [53] V. Boyko, O. Prach, O. Trudonoshyn, K. Mykhalenkov, Natural Aging of AlMg5Si2Mn Casting Alloy, Research Bulletin of the National Technical University of Ukraine “Kyiv Polytechnic Institute.” (2014) 47-53 (In Ukrainian). https://doi.org/10.20535/1810- 0546.2014.1.26582. [54] I.J. Polmear, Aluminium alloys - A century of age hardening, Materials Forum. 28 (2004) 1–14. https://doi.org/7F6104775CD4BCE9E9D087602166B700. [55] T. Arai, S. Harper, ASM Handbook. Heat-treating., ASM International, 1998. https://doi.org/10.1016/S0026-0576(03)90166-8. [56] K. Matsuda, S. Ikeno, H. Gamada, K. Fujii, Y. Uetani, T. Sato, A. Kamio, High-resolution electron microscopy on the structure of Guinier-Preston zones in an Al-1.6 mass Pct Mg2Si alloy, Metallurgical and Materials Transactions A. 29 (1998) 1161–1167. https://doi.org/10.1007/s11661-998-0242-7. [57] R. Bjørge, C.D. Marioara, S.J. Andersen, R. Holmestad, Precipitation in Two Al-Mg-Ge Alloys, Metallurgical and Materials Transactions A. 41 (2010) 1907–1916. https://doi.org/10.1007/s11661-010-0301-8. [58] M.X. Guo, X.K. Zhang, J.S. Zhang, L.Z. Zhuang, Effect of Zn addition on the precipitation behaviors of Al – Mg – Si – Cu alloys for automotive applications, Journal of Materials Science. 52 (2016) 1390–1404. https://doi.org/10.1007/s10853-016-0433-3. [59] M.F. Ashby, Oxide Dispersion Strengthening, Gordon and Breach, New York, 1968.
127 [60] P. Mukhopadhyay, Alloy Designation, Processing, and Use of AA6XXX Series Aluminium Alloys, ISRN Metallurgy. 2012 (2012) 1–15. https://doi.org/10.5402/2012/165082. [61] S.N. Samaras, G.N. Haidemenopoulos, Modelling of microsegregation and homogenization of 6061 extrudable Al-alloy, Journal of Materials Processing Technology. 194 (2007) 63– 73. https://doi.org/10.1016/j.jmatprotec.2007.03.126. [62] G.B. Burger, A.K. Gupta, P.W. Jeffrey, D.J. Lloyd, Microstructural control of aluminum sheet used in automotive applications, Materials Characterization. 35 (1995) 23–39. https://doi.org/10.1016/1044-5803(95)00065-8. [63] S.K. Kairy, P.A. Rometsch, K. Diao, J.F. Nie, C.H.J. Davies, N. Birbilis, Exploring the electrochemistry of 6xxx series aluminium alloys as a function of Si to Mg ratio, Cu content, ageing conditions and microstructure, Electrochimica Acta. 190 (2016) 92–103. https://doi.org/10.1016/j.electacta.2015.12.098. [64] John .E. Hatch, Aluminum Properties and Physical Metallurgy, ASM International, Ohio, US, 1984. https://doi.org/10.1361/appm1984p001. [65] A. Simar, Y. Bréchet, B. De Meester, A. Denquin, C. Gallais, T. Pardoen, Integrated modeling of friction stir welding of 6xxx series Al alloys: Process, microstructure and properties, Progress in Materials Science. 57 (2012) 95–183. https://doi.org/10.1016/j.pmatsci.2011.05.003. [66] J.R. Davis, Corrosion Of Aluminum and Aluminum Alloys, ASM International, Ohio, 1999. https://doi.org/10.1361/caaa1999p001. [67] W.J. Liang, P.A. Rometsch, L.F. Cao, N. Birbilis, General aspects related to the corrosion of 6xxx series aluminium alloys: Exploring the influence of Mg/Si ratio and Cu, Corrosion Science. 76 (2013) 119–128. https://doi.org/10.1016/j.corsci.2013.06.035. [68] M.H. Larsen, J.C. Walmsley, O. Lunder, R.H. Mathiesen, K. Nisancioglu, Intergranular Corrosion of Copper-Containing AA6xxx AlMgSi Aluminum Alloys, Journal of The Electrochemical Society. 155 (2008) C550. https://doi.org/10.1149/1.2976774. [69] V. Guillaumin, G. Mankowski, Localized corrosion of 6056 T6 aluminium alloy in chloride media, Corrosion Science. 42 (2000) 105–125. https://doi.org/10.1016/S0010- 938X(99)00053-0. [70] G. Mrowka-Nowotnik, J. Sieniawski, Influence of heat treatment on the microstructure and mechanical properties of 6005 and 6082 aluminium alloys, Journal of Materials Processing Technology. 162–163 (2005) 367–372. https://doi.org/10.1016/j.jmatprotec.2005.02.115. [71] G. Mrowka-Nowotnik, J. Sieniawski, A. Nowotnik, Effect of heat treatment on tensile and fracture toughness properties of 6082 alloy, Journal of Acheivements in Materials and Manufacturing Engineering. 32 (2009) 162–170. [72] L. Liu,Y, B. Kang, S, The solidification process of Al – Mg – Si alloys, Journal of Materials Science. 2 (1997) 1443–1447. https://doi.org/10.1023/A:1018545732009. [73] Y.L. Liu, S.B. Kang, H.W. Kim, The complex microstructures in an as-cast Al–Mg–Si alloy, Materials Letters. 41 (1999) 267–272. https://doi.org/10.1016/S0167-577X(99)00141-X. [74] G. Mrówka-Nowotnik, J. Sieniawski, M. Wierzbińska, Analysis of intermetallic particles in AlSi1MgMn aluminium alloy, Journal of Achievements in Materials and Manufacturing Engineering. 20 (2007) 7–8. [75] J. Banhart, C.S.T. Chang, Z. Liang, N. Wanderka, M.D.H. Lay, A.J. Hill, Natural aging in Al-Mg-Si alloys - A process of unexpected complexity, Advanced Engineering Materials. 12 (2010) 559–571. https://doi.org/10.1002/adem.201000041. [76] C.S.T. Chang, I. Wieler, N. Wanderka, J. Banhart, Positive effect of natural pre-ageing on precipitation hardening in Al-0.44 at% Mg-0.38 at% Si alloy, Ultramicroscopy. 109 (2009) 585–592. https://doi.org/10.1016/j.ultramic.2008.12.002. [77] W.F. Miao, D.E. Laughlin, Precipitation hardening in aluminum alloy 6022, Scripta Materialia. 40 (1999) 873–878. https://doi.org/10.1016/S1359-6462(99)00046-9. [78] W. Chrominski, S. Wenner, C.D. Marioara, R. Holmestad, M. Lewandowska, Strengthening mechanisms in ultrafine grained Al-Mg-Si alloy processed by hydrostatic extrusion -
128
Influence of ageing temperature, Materials Science and Engineering A. 669 (2016) 447– 458. https://doi.org/10.1016/j.msea.2016.05.109. [79] T. Khelfa, M.A. Rekik, J.A. Muñoz-Bolaños, J.M. Cabrera-Marrero, M. Khitouni, Microstructure and strengthening mechanisms in an Al-Mg-Si alloy processed by equal channel angular pressing (ECAP), International Journal of Advanced Manufacturing Technology. 95 (2018) 1165–1177. https://doi.org/10.1007/s00170-017-1310-1. [80] S. Ji, Y. Wang, D. Watson, Z. Fan, Microstructural evolution and solidification behavior of Al-Mg-Si alloy in high-pressure die casting, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science. 44 (2013) 3185–3197. https://doi.org/10.1007/s11661-013-1663-5. [81] P. Krug, H. Koch, R. Klos, Magsimal-25 - A new High-Ductility Die Casting Alloy for Structural Parts in Automotive Industry, 59 (n.d.) 1–8. [82] O. Prach, O. Trudonoshyn, M. Puchnin, Effects of chemical composition on mechanical properties of Al-Mg-Si-Mn based alloys, Materials Engineering - Materiálové Inžinierstvo (MEMI). 24 (2017) 11–20. http://ojs.mateng.sk/index.php/Mateng/article/download/215/399. [83] A. Slyudova, O. Trudonoshyn, O. Prach, V. Lisovskii, Morphology and nucleation of intermetallic phases in casting Al-Mg-Si alloys, Metallography, Microstructure, and Analysis. 9 (2020). https://doi.org/10.1007/s13632-020-00702-w. [84] O. Trudonoshyn, M. Puchnin, O. Prach, Use of the ABI technique to measure the mechanical properties of aluminium alloys: Effect of heat-treatment conditions on the mechanical properties of alloys, Materiali in Tehnologije. 50 (2016). https://doi.org/10.17222/mit.2014.295. [85] E. Georgatis, A. Lekatou, A.E. Karantzalis, H. Petropoulos, S. Katsamakis, A. Poulia, Development of a cast Al-Mg2Si-Si in situ composite: Microstructure, heat treatment, and mechanical properties, Journal of Materials Engineering and Performance. 22 (2013) 729– 741. https://doi.org/10.1007/s11665-012-0337-6. [86] O. Trudonoshyn, O. Prach, A. Slyudova, V. Lisovskii, Structure formation and multistep nucleation in casting Al-Mg-Si alloys, International Journal of Cast Metals Research. 33 (2020) 184–193. https://doi.org/10.1080/13640461.2020.1822632. [87] M. Di Sabatino, L. Arnberg, S. Brusethaug, D. Apelian, Fluidity evaluation methods for Al– Mg–Si alloys, International Journal of Cast Metals Research. 19 (2006) 94–97. https://doi.org/10.1179/136404606225023345. [88] Z. Hu, L. Wan, S. Lü, P. Zhu, S. Wu, Research on the microstructure, fatigue and corrosion behavior of permanent mold and die cast aluminum alloy, Materials and Design. 55 (2014) 353–360. https://doi.org/10.1016/j.matdes.2013.10.012. [89] Rheinfelden Alloys GmbH & Co. KG, Ductile Aluminum Alloys for Automotive Structural Applications, Rheinfelden, 2017. http://rheinfelden-alloys.eu/wp- content/uploads/2015/04/Handbook-Alloys-for-Structural-Applications_RHEINFELDEN- ALLOYS_2017_EN.pdf. [90] S. Otarawanna, C.M. Gourlay, H.I. Laukli, A.K. Dahle, Microstructure formation in AlSi4MgMn and AlMg5Si2Mn high-pressure die castings, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science. 40 (2009) 1645–1659. https://doi.org/10.1007/s11661-009-9841-1. [91] Rheinfelden Alloys GmbH & Co., Primary Aluminium Alloys for Pressure Die Castings, Rheinfelden, 2007. [92] W. Leis, Druckguss (43. Folge) Teil 1: Werkstoffe, Giesserei. 94 (2007) 48–59. https://www.giesserei.eu/e-paper/. [93] H. Koch, Al-Mg-Si-Aluminium-Gusslegierung mit Scandium Al-Mg-Si, EP 1 682 688 B1, 2010. https://data.epo.org/publication-server/rest/v1.0/publication- dates/20100106/patents/EP1682688NWB1/document.pdf. [94] M.H. Jacobs, The structure of the metastable precipitates formed during ageing of an Al- Mg-Si alloy, Philosophical Magazine. 26 (1972) 1–13.
129 https://doi.org/10.1080/14786437208221015. [95] D.J. Chakrabarti, D.E. Laughlin, Phase relations and precipitation in Al-Mg-Si alloys with Cu additions, Progress in Materials Science. 49 (2004) 389–410. https://doi.org/10.1016/S0079-6425(03)00031-8. [96] D.J. Chakrabarti, Y. Peng, D.E. Laughlin, Precipitation in Al-Mg-Si alloys with Cu additions and the role of the Q′ and related phases, Aluminium Alloys 2002 Their Physical and Mechnaical Properties: Proceedings of the 8th International Conference ICAA8. 396– 402 (2002) 857–862. https://doi.org/10.4028/www.scientific.net/MSF.396-402.857. [97] S.J. Andersen, C.D. Marioara, R. Vissers, A. Frøseth, H.W. Zandbergen, The structural relation between precipitates in Al-Mg-Si alloys, the Al-matrix and diamond silicon, with emphasis on the trigonal phase U1-MgAl2Si2, Materials Science and Engineering A. 444 (2007) 157–169. https://doi.org/10.1016/j.msea.2006.08.084. [98] C.D. Marioara, S.J. Andersen, H.W. Zandbergen, R. Holmestad, The influence of alloy composition on precipitates of the Al-Mg-Si system, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science. 36 (2005) 691–702. https://doi.org/10.1007/s11661-005-1001-7. [99] M. Torster, H.S. Hasting, W. Lefebvre, C.D. Marioara, J.C. Walmsley, S.J. Andersen, R. Holmestad, The influence of composition and natural aging on clustering during preaging in Al-Mg-Si alloys, Journal of Applied Physics. 108 (2010). https://doi.org/10.1063/1.3481090. [100] C.D. Marioara, S.J. Andersen, J. Jansen, H.W. Zandbergen, Atomic model for GP-zones in a 6082 Al-Mg-Si system, Acta Materialia. 49 (2001) 321–328. https://doi.org/10.1016/S1359-6454(00)00302-5. [101] S.J. Andersen, H.W. Zandbergen, J. Jansen, C. Træholt, U. Tundal, O. Reiso, The crystal structure of the β″ phase in Al-Mg-Si Alloys, Acta Materialia. 46 (1998) 3283–3298. https://doi.org/10.1016/S1359-6454(97)00493-X. [102] R. Vissers, M.A. van Huis, J. Jansen, H.W. Zandbergen, C.D. Marioara, S.J. Andersen, The crystal structure of the β′ phase in Al-Mg-Si alloys, Acta Materialia. 55 (2007) 3815–3823. https://doi.org/10.1016/j.actamat.2007.02.032. [103] K. Matsuda, Y. Ishida, I. Mullerova, L. Frank, S. Ikeno, Cube-phase in excess Mg-type Al- Mg-Si alloy studied by EFTEM, Journal of Materials Science. 41 (2006) 2605–2610. https://doi.org/10.1007/s10853-006-7819-6. [104] K. Matsuda, T. Yoshida, T. Wada, M. Yoshida, Y. Uetani, T. Sato, A. Kamio, S. Ikeno, Precipitation Sequence in Al-Mg-Si Alloys with Excess Magnesium, J．JapanInst． Metals. 62 (1998) 718–726. [105] L.C. Doan, K. Nakai, Y. Matsuura, S. Kobayashi, Y. Ohmori, Effects of Excess Mg and Si on the Isothermal Ageing Behaviours in the Al – Mg2Si Alloys, Materials Transactions. 43 (2002) 1371–1380. [106] K. Matsuda, S. Ikeno, T. Sato, Y. Uetani, Metastable phases in an Al-Mg-Si alloy containing copper, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science. 32 (2001) 1293–1299. https://doi.org/10.1007/s11661-001-0219-2. [107] D.W. Pashley, M.H. Jacobs, J.T. Vietz, The basic processes affecting two-step ageing in an Al-Mg-Si alloy, Philosophical Magazine. 16 (1967) 51–76. https://doi.org/10.1080/14786436708229257. [108] M. Murayama, K. Hono, Pre-Precipitate Clusters and Precipitatation Processes in Al-Mg- Si Alloys, Acta Mater. 47 (1999) 1537–1548. [109] M. Murayama, K. Hono, M. Saga, M. Kikuchi, Atom probe studies on the early stages of precipitation in Al–Mg–Si alloys, Materials Science and Engineering A. 250 (1998) 127– 132. https://pdfs.semanticscholar.org/0d14/674887b7db9d238c565b45e1ae586c0b2087.pdf. [110] G. Thomas, The aging characteristics of aluminum alloys electron transmission studies of Al-Mg-Si alloys, California, 1961. https://doi.org/10.1007/978-3-319-46448-0.
130
[111] H.S. Hasting, A.G. Frøseth, S.J. Andersen, R. Vissers, J.C. Walmsley, C.D. Marioara, F. Danoix, W. Lefebvre, R. Holmestad, Composition of Β″ precipitates in Al-Mg-Si alloys by atom probe tomography and first principles calculations, Journal of Applied Physics. 106 (2009). https://doi.org/10.1063/1.3269714. [112] P.H. Ninive, A. Strandlie, S. Gulbrandsen-Dahl, W. Lefebvre, C.D. Marioara, S.J. Andersen, J. Friis, R. Holmestad, O.M. Løvvik, Detailed atomistic insight into the β″ phase in Al-Mg- Si alloys, Acta Materialia. 69 (2014) 126–134. https://doi.org/10.1016/j.actamat.2014.01.052. [113] S.I. K. Matsuda, Y. Sakaguchi, Y. Miyata, Y. Uetani, T. Sato, A.K. D, Precipitation sequence of various kinds of metastable phases in Al-1.0mass% Mg2Si-0.4mass% Si alloy, Journal of Materials Science. 35 (2000) 179– 189. [114] C.D. Marioara, H. Nordmark, S.J. Andersen, R. Holmestad, Post-β″ phases and their influence on microstructure and hardness in 6xxx Al-Mg-Si alloys, Journal of Materials Science. 41 (2006) 471–478. https://doi.org/10.1007/s10853-005-2470-1. [115] S.J. Andersen, C.D. Marioara, A. Frøseth, R. Vissers, H.W. Zandbergen, Crystal structure of the orthorhombic U2-Al4Mg4Si4precipitate in the Al-Mg-Si alloy system and its relation to the β′ and β″ phases, Materials Science and Engineering A. 390 (2005) 127–138. https://doi.org/10.1016/j.msea.2004.09.019. [116] Y. Ohmori, L.C. Doan, Y. Matsuura, S. Kobayashi, K. Nakai, Morphology and Crystallography of b-Mg2Si Precipitation in Al-Mg-Si Alloys, Materials Transactions. 42 (2001) 2576–2583. [117] Y. Ohmori, L. Doan, K. Nakai, Ageing Processes in Al-Mg-Si Alloys during Continuous Heating, Materials Transactions. 43 (2002) 246–255. https://doi.org/10.2320/matertrans.43.246. [118] O.R. Myhr, O. Grong, H.G. Fjær, C.D. Marioara, Modelling of the microstructure and strength evolution in al-mg-si alloys during multistage thermal processing, Acta Materialia. 52 (2004) 4997–5008. https://doi.org/10.1016/j.actamat.2004.07.002. [119] Y.S. Lee, J.H. Cha, S.H. Kim, C.Y. Lim, H.W. Kim, H.S. Kim, Modification of eutectic Mg2Si in AlMg5Si2Mn alloy by pre-homogenization deformation treatment with different reduction conditions, Materials Characterization. 141 (2018) 388–397. https://doi.org/10.1016/j.matchar.2018.04.045. [120] H. Feufel, T. Godecke, H.L. Lukas, F. Sommer, Investigation of the Al-Mg-Si system by experiments and thermodynamic calcuiations, Journal of Alloys and Compounds. 247 (1997) 31–42. https://doi.org/Pll S0925-8388(96)02655-2. [121] V. Raghavan, Al-Mg-Si (Aluminum-Magnesium-Silicon), Journal of Phase Equilibria and Diffusion. 28 (2007) 189–191. https://doi.org/10.1007/s11669-007-9027-8. [122] S. Ji, W. Yang, F. Gao, D. Watson, Z. Fan, Effect of iron on the microstructure and mechanical property of Al-Mg-Si-Mn and Al-Mg-Si diecast alloys, Materials Science and Engineering A. 564 (2013) 130–139. https://doi.org/10.1016/j.msea.2012.11.095. [123] F. Yan, S.X. Ji, Z.Y. Fan, Effect of Excess Mg on the Microstructure and Mechanical Properties of Al-Mg2Si High Pressure Die Casting Alloys, Materials Science Forum. 765 (2013) 64–68. https://doi.org/10.4028/www.scientific.net/MSF.765.64. [124] J. Zhang, Z. Fan, Y.Q. Wang, B.L. Zhou, Microstructural evolution of the in situ Al- 15wt.%Mg2Si composite with extra Si contents, Scripta Materialia. 42 (2000) 1101–1106. https://doi.org/10.1016/S1359-6462(00)00338-9. [125] Y. Tang, Y. Du, L. Zhang, X. Yuan, G. Kaptay, Thermodynamic description of the Al-Mg- Si system using a new formulation for the temperature dependence of the excess Gibbs energy, Thermochimica Acta. 527 (2012) 131–142. https://doi.org/10.1016/j.tca.2011.10.017. [126] K.C. Hari Kumar, N. Chakraborti, H.-L. Lukas, O. Bodak, L. Rokhlin, Al-Mg-Si (Aluminium - Magnesium - Silicon), Light Metal Systems. Part 3. 11A3 (2005) 165–177. https://doi.org/10.1007/10915998\textunderscore 18.
131 [127] S.P. Li, S.X. Zhao, M.X. Pan, D.Q. Zhao, X.C. Chen, O.M. Barabash, R.I. Barabash, Solidification and structural characteristics of alpha(Al)-Mg2Si eutectic, Materials Transactions Jim. 38 (1997) 553–559. https://doi.org/10.2320/matertrans1989.38.553. [128] X. Zhu, H. Yang, X. Dong, S. Ji, The effects of varying Mg and Si levels on the microstructural inhomogeneity and eutectic Mg2Si morphology in die-cast Al–Mg–Si alloys, Journal of Materials Science. 54 (2019) 5773–5787. https://doi.org/10.1007/s10853- 018-03198-6. [129] S. Ji, F. Yan, Z. Fan, Development of a high strength Al-Mg2Si-Mg-Zn based alloy for high pressure die casting, Materials Science & Engineering A. 626 (2015) 165–174. https://doi.org/10.1016/j.msea.2014.12.019. [130] V.S. Zolotorevsky, N.A. Belov, M.V. Glazoff, Casting Aluminum Alloys, Elsevier, Oxford, 2007. [131] J. Zhang, Y. Wang, B. Yang, B. Zhou, I. Introduction, Effects of Si content on the microstructure and tensile strength of an in situ Al/Mg2Si Si composite, (1999). [132] M. Puchnin, O. Trudonoshyn, O. Prach, Use of the ABI technique to measure the mechanical properties of aluminium alloys: Effect of chemical composition on the mechanical properties of the alloys, Materiali in Tehnologije. 50 (2016) 247–252. https://doi.org/10.17222/mit.2014.294. [133] M. Puchnin, O. Trudonoshyn, O. Prach, F. Pešlová, Comparison of ABI technique and standard methods in measuring mechanical properties of aluminium Al-aloys, Manufacturing Technology. 16 (2016). [134] L. Ding, Z. Jia, Z. Zhang, R.E. Sanders, Q. Liu, G. Yang, The natural aging and precipitation hardening behaviour of Al-Mg-Si-Cu alloys with different Mg/Si ratios and Cu additions, Materials Science & Engineering A. 627 (2015) 119–126. https://doi.org/http://dx.doi.org/10.1016/j.msea.2014.12.086. [135] L. Lodgaard, N. Ryum, Precipitation of chromium containing dispersoids in Al–Mg–Si alloys, Materials Science and Technology. 16 (2000) 599–604. https://doi.org/10.1179/026708300101508315. [136] L. Lodgaard, N. Ryum, Precipitation of dispersoids containing Mn and / or Cr in Al – Mg – Si alloys, Materials Science & Engineering A. 283 (2000) 144–152. https://doi.org/https://doi.org/10.1016/S0921-5093(00)00734-6. [137] P. Zhang, Z. Li, B. Liu, W. Ding, L. Peng, Improved tensile properties of a new aluminum alloy for high pressure die casting, Materials Science and Engineering A. 651 (2016) 376– 390. https://doi.org/10.1016/j.msea.2015.10.127. [138] M.R. Ghorbani, M. Emamy, N. Nemati, Microstructural and mechanical characterization of Al-15%Mg2Si composite containing chromium, Materials and Design. 32 (2011) 4262– 4269. https://doi.org/10.1016/j.matdes.2011.04.020. [139] C. Li, Y. Wu, H. Li, Y. Wu, X. Liu, Effect of Ni on eutectic structural evolution in hypereutectic Al – Mg 2 Si cast alloys, Materials Science & Engineering A. 528 (2010) 573–577. https://doi.org/10.1016/j.msea.2010.09.056. [140] S.X. Ji, D. Watson, Y. Wang, M. White, Z.Y. Fan, Effect of Ti Addition on Mechanical Properties of High Pressure Die Cast Al-Mg-Si Alloys, Materials Science Forum. 765 (2013) 23–27. https://doi.org/10.4028/www.scientific.net/MSF.765.23. [141] O. Trudonoshyn, O. Prach, V. Boyko, M. Puchnin, K. Mykhalenkov, Design of a new casting alloys containing Li or Ti+Zr and optimization of its heat treatment, in: METAL 2014 - 23rd International Conference on Metallurgy and Materials, Conference Proceedings, 2014: pp. 1399–1404. http://metal2014.tanger.cz/files/proceedings/17/reports/2750.pdf. [142] SAG Aluminium Lend GmbH & Co KG, Progress in Aluminium, Land, n.d. [143] J.L. Jorstad, Excellent Die Casting Alloy System Made Better by Rheocasting, Die Casting Engineer. (2008) 18–20. [144] O. Prach, O. Trudonoshyn, P. Randelzhofer, С. Körner, K. Durst, Effect of Zr, Cr and Sc on
132
the Al-Mg-Si-Mn high-pressure die casting alloys, Materials Science and Engineering: A. 759 (2019) 603–612. https://doi.org/https://doi.org/10.1016/j.msea.2019.05.038. [145] Y. Meng, Z.H. Zhao, J.Z. Cui, Effect of minor Zr and Sc on microstructures and mechanical properties of Al-Mg-Si-Cu-Cr-V alloys, Transactions of Nonferrous Metals Society of China (English Edition). 23 (2013) 1882–1889. https://doi.org/10.1016/S1003- 6326(13)62673-4. [146] O. Prach, O. Trudonoshyn, P. Randelzhofer, С. Körner, K. Durst, Multi-alloying effect of Sc, Zr, Cr on the Al-Mg-Si-Mn high-pressure die casting alloys, Materials Characterization. (2020) 110537. https://doi.org/10.1016/j.matchar.2020.110537. [147] J. Royset, J. Røyset, Scandium in Aluminium alloys overview: Physical Metallurgy, Properties and applications, Metallurgical Science AndTechnology. 25 (2007) 11–21. https://www.gruppofrattura.it/ors/index.php/MST/article/view/1126/1078. [148] O. Trudonoshyn, O. Prach, Multistep Nucleation And Multi-Modification Effect Of Sc In Hypoeutectic Al-Mg-Si Alloys, Heliyon. 5 (2019) 1–12. https://doi.org/10.1016/j.heliyon.2019.e01202. [149] E.V.A.A. Mørtsell, C.D. Marioara, S.J. Andersen, J. Røyset, O. Reiso, R. Holmestad, Effects of Germanium, Copper and Silver Substitutions on Hardness and Microstructure in Lean Al-Mg-Si Alloys, Metallurgical and Materials Transactions A. 46 (2015) 4369–4379. https://doi.org/10.1007/s11661-015-3039-5. [150] O. Trudonoshyn, Studying the structure of Al-Mg-Si casting alloys doped by lithium, Physics of Metals and Metallography. 121 (2020) 701–707. https://doi.org/10.1134/S0031918X2007011X. [151] A.I. Trudonoshin, Studying the Structure of Al–Mg–Si Casting Alloys Doped by Lithium, Fizika Metallov i Metallovedenie. 121 (2020) 771–778. https://doi.org/10.31857/S0015323020070116. [152] O. Prach, J. Hornik, K. Mykhalenkov, Effect of the addition of Li on the structure and mechanical properties of hypoeutectic Al-Mg2Si alloys, Acta Polytechnica. 55 (2015) 253– 259. https://doi.org/10.14311/AP.2015.55.0253. [153] L.K. Berg, G. Waterloo, D. Schryvers, L.R. Wallenberg, GP-zones in Al-Zn-Mg alloys and their role in artificial aging, Acta Materialia. 49 (2001) 3443–3451. [154] C. Wolverton, Crystal structure and stability of complex precipitate phases in Al-Cu-Mg- (Si) and Al-Zn-Mg alloys, Acta Materialia. 49 (2001) 3129–3142. [155] T. Jian-guo, C. Hui, Z. Xin-ming, L.I.U. Sheng-dan, L.I.U. Wen-jun, O. Hui, L.I. Hong- ping, Influence of quench-induced precipitation on aging behavior of AlíZníMgíCu alloy, Transactions of Nonferrous Metals Society of China. 22 (2012) 1255–1263. https://doi.org/10.1016/S1003-6326(11)61313-7. [156] N.Q. Chinh, J. Lendvai, D.H. Ping, K. Hono, The effect of Cu on mechanical and precipitation properties of Al -Zn-Mg alloys, Journal of Alloys and Compounds. 378 (2004) 52–60. https://doi.org/10.1016/j.jallcom.2003.11.175. [157] G. Sha, A. Cerezo, Early-stage precipitation in Al-Zn -Mg-Cu alloy (7050), Acta Materialia. 52 (2004) 4503–4516. https://doi.org/10.1016/j.actamat.2004.06.025. [158] L.I. Chunmei, C. Zhiqian, Z. Sumin, C. Nanpu, C. Tianxiao, Intermetallic phase formation and evolution during homogenization and solution in Al-Zn-Mg-Cu alloys, SCIENCE CHINA Technological Sciences. 56 (2013) 2827–2838. https://doi.org/10.1007/s11431- 013-5356-5. [159] X. Li, M.J. Starink, Effect of compositional variations on characteristics of coarse intermetallic particles in overaged 7000 aluminium alloys, Materials Science and Technology. 17 (2001) 1324–1328. [160] Y. Li, L. Kovarik, W. Wang, M.J. Mills, High-resolution characterization of the precipitation behavior of an Al – Zn – Mg – Cu alloy, Philosophical Magazine Letters. 92 (2012) 166– 178 Downloaded. https://doi.org/10.1080/09500839.2011.652682. [161] H. Lofler, I. Kovacs, J. Lendavi, Review Decomposition processes in Al-Zn-Mg alloys,
133 Journal of Materials Science. 18 (1983) 2215–2240. [162] A. Bigot, P. Auger, S. Chambreland, D. Blavette, A. Reeves, Atomic Scale Imaging and in Al-Mg-Zn Alloys Analysis of T ’ Precipitates, Microsc. Microanal. Microstruct. 8 (1997) 103–113. [163] X. Fang, M. Song, K. Li, Y. Du, Effects of Cu and Al on the crystal structure and composition of g (MgZn2) phase in over-aged Al-Zn-Mg-Cu alloys, Journal of Materials Science. 47 (2012) 5419–5427. https://doi.org/10.1007/s10853-012-6428-9. [164] T. Saito, F.J.H. Ehlers, W. Lefebvre, D. Hernandez-Maldonado, R. Bjørge, C.D. Marioara, S.J. Andersen, R. Holmestad, HAADF-STEM and DFT investigations of the Zn-containing β’’ phase in Al-Mg-Si alloys, Acta Materialia. 78 (2014) 245–253. https://doi.org/10.1016/j.actamat.2014.06.055. [165] T. Saito, S. Wenner, E. Osmundsen, D. Calin, S.J. Andersen, J. Røyset, The effect of Zn on precipitation in Al-Mg-Si alloys, Philosophical Magazine. 94 (2014) 2410–2425. https://doi.org/10.1080/14786435.2014.913819. [166] Y. Cai, C. Wang, J. Zhang, Microstructural characteristics and aging response of Zn- containing Al-Mg-Si-Cu alloy, International Journal of Minerals, Metallurgy and Materials. 20 (2013) 659–664. https://doi.org/10.1007/s12613-013-0780-x. [167] X.P. Ding, H. Cui, J.X. Zhang, H.X. Li, M.X. Guo, Z. Lin, L.Z. Zhuang, J.S. Zhang, The effect of Zn on the age hardening response in an Al-Mg-Si alloy, Materials and Design. 65 (2015) 1229–1235. https://doi.org/10.1016/j.matdes.2014.09.086. [168] L. Li, S. Ji, Q. Zhu, Y.U.N. Wang, X. Dong, W. Yang, S. Midson, Y. Kang, Effect of Zn Concentration on the Microstructure and Mechanical Properties of Al-Mg-Si-Zn Alloys Processed by Gravity Die Casting, Metallurgical and Materials Transactions A. 49 (2018) 3247–3256. https://doi.org/10.1007/s11661-018-4684-2. [169] F. Yan, W. Yang, S. Ji, Z. Fan, Effect of solutionising and ageing on the microstructure and mechanical properties of a high strength die-cast Al-Mg-Zn-Si alloy, Materials Chemistry and Physics. (2015) 1–9. https://doi.org/10.1016/j.matchemphys.2015.10.014. [170] Y.J. Li, A.M.F. Muggerud, A. Olsen, T. Furu, Precipitation of partially coherent α- Al(Mn,Fe)Si dispersoids and their strengthening effect in AA 3003 alloy, Acta Materialia. 60 (2012) 1004–1014. https://doi.org/10.1016/j.actamat.2011.11.003. [171] J.L. Murray, The Al-Cr (aluminum-chromium) system, Journal of Phase Equilibria. 19 (1998) 368–375. https://doi.org/10.1361/105497198770342102. [172] V. Raghavan, Al-Mg-Zn (Aluminum-Magnesium-Zinc), Journal of Phase Equilibria and Diffusion. 31 (2010) 293–294. https://doi.org/10.1007/s11669-010-9684-x. [173] P. Liang, T. Tarfa, J.. Robinson, S. Wagner, P. Ochin, M.. Harmelin, H.. Seifert, H.. Lukas, F. Aldinger, Experimental investigation and thermodynamic calculation of the Al–Mg–Zn system, Thermochimica Acta. 314 (2002) 87–110. https://doi.org/10.1016/s0040- 6031(97)00458-9. [174] V. Raghavan, Al-Mg-Si-Zn (aluminum-magnesium-silicon-zinc), Journal of Phase Equilibria and Diffusion. 32 (2011) 72–74. https://doi.org/10.1007/s11669-010-9803-8. [175] F. Yan, Development of High Strength Al-Mg2Si-Mg Based Alloy for High Pressure Diecasting Process, Brunel University, Uxbridge, United Kingdom, 2013. [176] O. Trudonoshyn, O. Prach, V. Boyko, K. Mykhalenkov, Selection and Optimization of Heat Treatment Process to Improve the Mechanical Properties of Casting Alloys of the Al-Mg- Si System, Casting Processes. (2013) 12-21 (In Russian). http://dspace.nbuv.gov.ua/bitstream/handle/123456789/159852/03-Trudonoshin.pdf. [177] A.I. Trudonoshin, K.V. Mikhalenkov, Eutectic spheroidization in cast Al-Mg-Si-Mn alloys during homogenization, Metallurgy of Machinery Building (Metallurgiya Mashinostroyeniya). 4 (2014) 17-20 (In Russian). http://elibrary.ru/item.asp?id=22827954. [178] D. Dragulin, M. Belte, O. Hoffman, AlMg5Si2Mn – Aluminiumdruckgusslegierung unter dem Aspekt einer Wärmebehandlungsperspektive, DRUCKGUSS. 3 (2010) 875–876. [179] O. Trudonoshyn, S. Rehm, P. Randelzhofer, С. Körner, Improvement of the high-pressure
134
die casting alloy Al-5.7Mg-2.6Si-0.7Mn with Zn addition, Materials Characterization. 158 (2019) 109959. https://doi.org/https://doi.org/10.1016/j.matchar.2019.109959. [180] N. Otsu, A Threshold Selection Method from Gray-Level Histograms, IEEE Transacations on Systems, Man, and Cybernetics. SMC-9 (1979) 62–66. https://doi.org/10.1109/TSMC.1979.4310076. [181] T.W. Ridler, S. Calvard, Picture Thresholding Using an Iterative Selection Method, IEEE Transacations on Systems, Man, and Cybernetics. SMC-8 (1978) 630–632. [182] M. Sezgin, B. Sankur, Survey over image thresholding techniques and quantitative performance evaluation, Journal of Electronic Imaging. 13 (2004) 146–165. https://doi.org/10.1117/1.1631316. [183] Z. Pan, C.K.L. Davies, R.N. Stevens, Measurement of foil thickness in transmission electron microscopy, Journal of Materials Science. 29 (1994) 1920–1924. https://doi.org/10.1007/BF00351314. [184] V.D. Scott, G. Love, Foil thickness measurements in transmission electron microscope, Materials Science and Technology (United Kingdom). 3 (1987) 600–608. https://doi.org/10.1179/mst.1987.3.8.600. [185] D. Watson, Microstructure and Mechanical Properties of Ductile Die-cast Al-Mg-Si-Mn Alloys, Brunel University London, Uxbridge, United Kingdom, 2015. [186] S. Rehm, Effect of Zn content on the structural changes in the base alloy Al-5,7Mg-2,6Si- Mn in the as-cast state and after solution treatment, Friedrich-Alexander-University of Erlangen-Nürnberg, 2019. https://www.researchgate.net/publication/332371554_Effect_of_Zn_content_on_the_struc tural_changes_in_the_base_alloy_Al-57Mg-26Si-Mn_in_the_as- cast_state_and_after_solution_treatment. [187] G. Timelli, S. Ferraro, A. Fabrizi, Effects of chromium and bismuth on secondary aluminium foundry alloys, International Journal of Cast Metals Research. 26 (2013) 239– 246. https://doi.org/10.1179/1743133613Y.0000000061. [188] K.E. Knipling, D.N. Seidman, D.C. Dunand, Ambient- and high-temperature mechanical properties of isochronally aged Al-0.06Sc, Al-0.06Zr and Al-0.06Sc-0.06Zr (at.%) alloys, Acta Materialia. 59 (2011) 943–954. https://doi.org/10.1016/j.actamat.2010.10.017. [189] L.S. Toropova, D.G. Eskinm, M.L. Kharakterovam, T. V. Dobatkina, Advanced Aluminum Alloys Containing Scandium: Structure and Properties, Routledge; 1 edition (June 1, 1998), 1998. [190] O.V. Berdova-Bashura, E.L. Prach, A.I. Trudonoshin, K.V. Mikhalenkov, The effect of chemical composition on the structure and mechanical properties of alloys of the Al-Mg-Si system, Casting Processes. (2015) 59-69 (In Russian). http://dspace.nbuv.gov.ua/bitstream/handle/123456789/160427/08-Berdova.pdf. [191] A.M. Samuel, F.H. Samuel, Effect of alloying elements and dendrite arm spacing on the microstructure and hardness of an Al-Si-Cu-Mg-Fe-Mn (380) aluminium die-casting alloy, Journal of Materials Science. 30 (1995) 1698–1708. https://doi.org/10.1007/BF00351598. [192] S. Ferraro, G. Timelli, Influence of Sludge Particles on the Tensile Properties of Die-Cast Secondary Aluminum Alloys, Metallurgical and Materials Transactions B: Process Metallurgy and Materials Processing Science. 46 (2015) 1022–1034. https://doi.org/10.1007/s11663-014-0260-3. [193] I. Outmani, L. Fouilland-Paille, J. Isselin, M. El Mansori, Effect of Si, Cu and processing parameters on Al-Si-Cu HPDC castings, Journal of Materials Processing Technology. 249 (2017) 559–569. https://doi.org/10.1016/j.jmatprotec.2017.06.043. [194] D. Shimosaka, S. Kumai, F. Casarotto, S. Watanabe, Effect of Cooling Rates during Solidification of Al-5.5%Mg-2.3%Si-0.6%Mn and Al-13%Mg2Si Pseudo-Binary Alloys on Their Secondary-Particle Morphology and Tear Toughness, Materials Transactions. 52 (2011) 920–927. https://doi.org/10.2320/matertrans.L-MZ201125. [195] K. Li, M. Song, Y. Du, X. Fang, Effect of minor cu addition on the precipitation sequence
135 of an as-cast Al-Mg-Si 6005 alloy, Archives of Metallurgy and Materials. 57 (2012) 457– 467. https://doi.org/10.2478/v10172-012-0047-y. [196] H. Möller, G. Govender, W. Stumpf, The T5 Heat Treatment of Semi-Solid Metal Processed Aluminium Alloy F357, Materials Science Forum. 618–619 (2009) 365–368. https://doi.org/10.4028/www.scientific.net/msf.618-619.365. [197] J.B. Patel, Y.Q. Liu, G. Shao, Z. Fan, Rheo-processing of an alloy specifically designed for semi-solid metal processing based on the Al-Mg-Si system, Materials Science and Engineering A. 476 (2008) 341–349. https://doi.org/10.1016/j.msea.2007.05.046. [198] K. Eigenfeld, A. Franke, S. Klan, H. Koch, B. Lenzcowski, B. Pflege, New developments in heat resistant aluminum casting materials, Casting Plant and Technology International. 4 (2004) 4–9. [199] H. Seyedrezai, D. Grebennikov, P. Mascher, H.S. Zurob, Study of the early stages of clustering in Al-Mg-Si alloys using the electrical resistivity measurements, Materials Science and Engineering A. 525 (2009) 186–191. https://doi.org/10.1016/j.msea.2009.06.054. [200] E. Ozawa, H. KIMURA, Excess vacancies and the nucliation of precipitates in Aluminium- Silicon alloys, Acta Metallurgica. 18 (1970) 995–1004. [201] O. Trudonoshyn, O. Prach, P. Randelzhofer, K. Durst, С. Körner, Heat Treatment of the New High-Strength High-Ductility Al-Mg-Si-Mn Alloys with Sc, Zr and Cr Additions, Materialia. (2020). https://doi.org/10.1016/j.mtla.2020.100981. [202] O.I. Trudonoshyn, A.A. Slyudova, V.A. Lisovskiy, O.L. Prach, The effect of scandium and chromium on the mechanical properties of alloys of the Al–Mg–Si–Mn system, in: А.А. Popov (Ed.), Actual Problems of the Physical Metallurgy of Steels and Alloys, 2020: pp. 14–16 (In Russian). http://elar.urfu.ru/bitstream/10995/83109/1/978-5-7996-2926- 7_2020_005.pdf. [203] F.S. M. Leonardo, EN 1706:2010 “Aluminium and aluminium alloys - Castings - Chemical composition and mechanical properties,” Brussels, 2010. [204] BS 1490:1988, Aluminium and aluminium alloy ingots and castings for general engineering purposes, 1998. [205] GOST 1583-93, Splavy alyuminiyevyye liteynyye [Aluminum casting alloys], Minsk, 1993. [206] GB/T 15115-2009 (Chinese Standard), Die casting aluminum alloys, 2009. [207] JIS H2212 (Japanese Standart), Aluminium-base alloy in ingot for die castings, 1965. [208] D.M. Stefanescu, Casting Handbook, 9th ed., ASM International, 2001. https://doi.org/10.1016/S0026-0576(03)90166-8. [209] M.C. Flemings, Behavior of Metal Alloys in the Semisolid State, Metallurgical Transaction B. 22B (1991) 269–293. https://doi.org/10.1007/BF02651227. [210] E. Dağdelen, A. Ulus, Aluminum Sheet Production : Heat Treatment of Aluminium and Temper Designations of Aluminium Alloys, Uluslararası Metalurji ve Malzeme Kongresi. (2016) 681–684.
136
APPENDICES Appendix A. List of acronyms Symbol Definition AA Artificial aging; A Elongation, %; b Burgers vector; c Foreign atom content in at. %;
CL Concentration of solute atoms in liquid;
CS Concentration of solute atoms in solid; DAS Dendrite arm spacing; DSC Differential scanning calorimetry; EDS Energy Dispersive Spectrometry system; EDX Energy Dispersive X-Ray analysis; EPMA Electron-probe microanalysis; FCC Face-centered cubic crystal structure; G Shear modulus; GP-zones Guinier-Preston zones (GP-I, GP-2); HB Brinell hardness HPDC, High-pressure die casting; HT Heat treatment; HV Vickers hardness; IADS International Alloy Designation System; ILS Interlamella spacing; ISO International Organization for Standardization; k Partition coefficient; k' Material-dependent constant; l Free distance (average free spacing) (distance between pracipitates); LIF Liquid impact (metal) forging; N Number of precipitates; p Dwell pressure, bar; PM Permanent mold casting; R Deformation resistances; r Radius of spherical precipitate; RT Room temperature;
137 SEM Scanning electron microscopy; S-L Solid-liquid interface; SSM Semi-solid metal casting; ssss Supersaturated solid solution; ST Solution treatment; TC Thermo-Calc software; TEM Transmission electron microscopy;
Td Mold temperature, ºC;
Toffset, Temperature of the end of a reaction, ºC;
Tonset, Temperature of the begin of a reaction, ºC;
Tm Melt temperature, ºC;
Tpeak Temperatures of the peak of a reaction, ºC; UTS Ultimate tensile strength, MPa; v Piston speed, 2m/s; V Volume fraction, %;
VAl Volume fraction of α-Al (incl. Al in the Al-Mg2Si eutectic), %;
VB Volume fraction of the second phase (B), %;
Vd Volume fraction of α-Al dendrites, %;
Veut. Volume fraction of Al-Mg2Si eutectic, %;
Vf, Volume fraction of the phase, %;
VMg2Si Volume fraction of Al-Mg2Si phase, %; YS, Yield strength MPa; λ Mean free particle spacing (ILS); τ, Shear stress critical;
τm Shear stress of the matrix material; υ Poisson's ratio of the matrix;
γeff Interface energy decisive for cutting;
138
Appendix B. Classification of Al-alloys Table B.1. Characterization of the wrought series of Al alloys [11,22] Age Properties System Series Elements harde- A Aplication HB ning [МPа] [%] Al 1XXX - - 70-175 20 30 Heat insulators 1-3%Cu 170-310 15 100 Shock absorbers, handles, spokes, conrods, pistons, rotors, Al-Cu-Mg 2XXX 3-6%Cu + 380-520 10 150 brake components, compressor Mg, Cu, Li 280-560 - 150 wheels Piping, cowl grilles, heat Al-Mn-Mg 3XXX Mn - 140-280 10 40 insulators Pistons, valve lifters, power Al-Si 4XXX Si ± 105-350 10 120 steering, rotors Body panels, airbag inflators, 1-3%Mg 140-280 15 60 covers, bonnets, roofs, doors, pillars, floors, seat frames, Al-Mg 5XXX - wheels, suspension, drivetrain, 3-6%Mg, engine cylinders, brake pistons, 280-380 10 120 Mn fuel pipes, tanks, underbody components Suspension arms, bumpers, ABS, side sills, shock absorbers, space frames, wheels, propeller shafts, arms, space frames, engine Al-Mg-Si 6XXX Mg, Si, Cu + 150-380 20 80 brackets, seat frames, roof railings, piping, power plant frames, airbags, joists, receiver tanks, bonnets, fenders, pillars, wheel housings, roofs, doors Zn, Mg 380-520 10 150 Keys, jacks, seatbelt hinges, bobbins, bumper reinforcement, Al-Zn-Mg 7XXX + seat sliders, impact beams, Zn, Mg, Cu 520-620 10 180 steering components, crossmembers, retractors Li, Mg, * 8ХХХ series is not limited to Other* 8XXX ± 280-560 - 150 Cu the elements Li, Mg, Cu Unused 9XXX - - - - - Xxxx - indicates the major alloy group; xxXX – indicate different Al alloys in the group (for 1XXX unalloyed wrought Al, the last 2 digits indicate the purity of the Al); xXxx – indicates alloy modifications: the original alloy belongs to 0 and digits from 1 to 9 indicate consecutive alloy modifications.
139 Table B.2. Characterization of the cast series of Al alloys [11,22] Age Mechanical properties Application in automobiles System Series harde- , A, В HB ning [МPа] [%] Al 1XX.Х - 75 30 30 Rotors Structural members; cylinder heads and pistons; pump, valve bodies, bushings; Al-Cu 2XX.Х + 400-600 3-5 65-140 meter parts; bearings; generator housings; compressor connecting rods; Crankcases; tanks; oil pans; engine parts; transmissions; pump bodies. chassis parts; Al-Si-Cu outboard motor parts; instrument cases; 3XX.Х + 240-320 2-5 65-85 Al-Si-Mg cover plates; irrigation system parts; hinges; air brakes; gear cases; pistons, blocks, manifolds, cylinder heads Al-Si 4XX.Х - 160-280 2-5 60-85 cylinder blocks, pistons, Al-Mg - 150 5 55-75 handling equipment; tire molds; fittings and ornamental hardware; frame sections 5XX.Х of cars, trucks and buses; instrument parts Al-Mg-Si + 240 3-10 90 and other applications with dimensional stability Unused 6XX.Х - - - - - 210 1-5 90 machine tool parts, furniture, trailer parts, Al-Zn-Mg 7XX.Х + mining equipment parts Al-Sn 8XX.Х - 180-190 3 60-80 bushings and bearings Other 9XX.Х Xxx.x - indicates the major alloy group; xXX.x - is given to identify a specific alloy in the series (or indicate the purity for 1xx.x group of pure Al); xxx.X - is used to identify alloys in the form of castings or foundry ingot: .0 indicates castings, .1 indicates ingot with limits for alloying elements very similar to those for the castings, .2 indicates that the ingot has composition limits that differ from but fall within the .1 limits. A serial letter before the numerical designation indicates modifications to original casting alloys (omitting I, O, Q, and X).
140
Appendix C. Common Al-alloys for automotive application
Table C.1. Mechanical properties of wrought Al alloys for automotive body structure [19–21]
Alloy temper Yield strength [MPa] Ultimate strength [MPa] Elongation [%] 5030-T4 138 276 - 5086-O 100 240 17 5182-O 131 276 24 5754-O 190 240 18 6009-T4 124 221 21 6010-T4 172 290 21 6013-T4 186 317 23 6016-T4 139 248 21 6061-T6 240 260 10 6082-T6 250 290 10 6111-T4 159 290 22
Table C.2. Typical mechanical properties of Al die cast alloys commonly used in the automotive industry [20,27] Alloy Alloy system UTS [MPa] YS [MPa] Elongation [%] 360 Al-Si 324 172 3 356 Al-Si 280 145 4 380 Al-Si-Cu 331 165 3 384 Al-Si-Cu 324 172 1 390 Al-Si-Cu-Mg 279 241 1 392 Al-Si-Mg-Cu 290 262 1 413 Al-Si 296 145 3 443 Al-Si-Cu 228 110 9 518 Al-Mg 310 186 8
141 Appendix D. Commercial Al-Mg-Si alloys Table D.1. The chemical composition of wrought Al alloys of 6XXX series, [wt.%] [21,22,64] Designation Si Mg Cu, max Mn, max Cr, max Others* 6003 0.3-1.0 0.8-1.5 0.10 0.8 0.35 - 6005 0.6-1.0 0.4-0.6 0.10 0.10 0.10 - 6060 0.3-0.6 0.3-0.6 0.10 0.10 0.05 6061 0.4-0.8 0.8-1.2 0.40 0.15 0.35 Fe 0.7% 6063 0.2-0.6 0.4-1.0 0.10 0.10 0.10 - 6066 0.9-1.8 0.8-1.4 1.20 1.10 0.40 - 6070 1.0-1.6 0.5-1.2 0.40 1.00 0.10 - 6082 0.7-1.3 0.6-1.2 0.40 1.00 0.25 Fe 0.5% 6101 0.3-0.7 0.3-0.8 0.10 0.03 0.03 В up to 0.06% 6105 0.6-1.0 0.4-0.8 0.10 0.10 0.10 - 6151 0.6-1.2 0.4-0.8 0.35 0.20 0.35 - 6162 0.4-0.8 0.7-1.1 0.20 0.10 0.10 - 6201 0.5-1.0 0.6-1.0 0.10 0.03 0.03 В up to 0.06% 6253 0.4-1.0 1.0-1.5 0.10 - 0.35 Zn 1.6-2.4% 6262 0.4-0.8 0.8-1.2 0.40 0.15 0.15 Pb, Bi 0.4-0.7% each 6351 0.7-1.3 0.4-0.8 0.10 0.80 - - 6463 0.2-0.6 0.4-1.0 0.20 0.05 - - * All alloys can contain Zn and Ti 0.1-0.2 wt. % each and Fe up to 0.3 wt. %
Table D.2. The chemical composition of the commercial casting Al-Mg-Si alloys Manufacturer / Chemical composition, Al-bal. [wt. %] Name Country1 Ref. Standard Si 142 Table D.3. Mechanical properties of the commercial casting Al-Mg-Si alloys [81,89,142,143] Mechanical properties Name Country Manufacturer Temper YS, [MPa] UTS, [MPa] HB δ, [%] For application in as-cast state or/and after T5 Magsimal22 F/T5 120 - 145 206 - 18 - 21 Magsimal25 F/T5 105 - 130 200-210 - 17 - 21 F 140 - 170 250 - 320 80-90 9 - 14 Magsimal59 4mm F3mm 160 – 220 310 – 340 - 12 - 18 DEU Rheinfelden F3mm 200 - 220 340 - 360 - 9 - 12 Magsimal59 Plus T5 230 - 250 350 - 380 - 8 - 12 F 190 - 220 350 - 380 80 - 100 7 - 10 Thermodur72 T5150ºC 220 - 245 260 - 290 - ~15 T5225ºC 150 - 175 180 - 205 - ~20 Maxxalloy59 F 130 - 175 250 - 280 75 - 95 9 - 16 AUT SAG Maxxalloy59 Ultra F 200 330 85 - 95 8 - 10 Aural 11 CAN Rio Tinto (Alcan) F 220 310 80 12 For application after full T6 heat treatment F 70 - 100 160 - 210 50 - 65 6 - 14 Peraluman36 T6 160 - 220 250 - 300 75 - 90 5 - 16 DEU Rheinfelden F 110 - 150 180 - 240 65 - 85 3 - 5 Peraluman56 T6 110 - 160 210 - 260 75 - 85 3 - 18 F 110 - 150 180 - 240 85 3 - 5 Perfoundal56 AUT SAG T6 110 - 160 210 - 260 75 - 85 3 - 18 143 Appendix E. Short characteristic of the casting methods Expendable Mold Casting Permanent Mold Casting Special casting processes Permanent pattern Die casting Vacuum Continuous casting Expendable pattern Liquid impact forming Centrifugal Chilled casting Low pressure die casting Liquid Metal 3D Printing Hot chamber Rheocasting Semi-solid metal High pressure die casting Thixocasting processes Cold chamber Thixomolding Figure E.1 Classification of the casting techniques [8] Table E.1. Characteristics of the gravity casting methods [9,208] Advantages Disadvantages Properties - Low cost of castings production; - Poor sanitary conditions; Surface roughness Rа, 25... - Ability to produce castings of - Large surface roughness; [µm] 12,5 large mass; - Limited minimum wall thickness; Admitted weight, [kg] 300 - Ability to produce castings of - High probability of defects Wall thickness min 5 alloys with a wide crystallization formation. [mm] max - Sand castings interval (Al-Cu, Al-Zn-Mg) Cooling rate [ºC/s] 0.1-2 d l - High dimensional accuracy; - The complexity of moulds Surface roughness Rа, 3,2... - Low surface roughness; production, limited service life (for [µm] 1,6 - Saving material; refractory alloys) Admitted weight, [kg] 50 - Better sanitary conditions; - The intractability of the metal mould Wall thickness min 3 - Fine-grained structure and metal rods [mm] max 100 Permanent mo - Difficult gas removal from the mould Cooling rate [ºC/s] 5-30 Table E.2. Characteristics of the high-pressure casting methods [30,31,209] Advantages Disadvantages Properties - High productivity and automatization of the - Size and weight Surface roughness Rа, 1,6... process, and high sanitary conditions; restriction; [µm] 0,8 - Minimum allowances for machining, minimum - The complexity of Admitted weight, [kg] 50 surface roughness, and high dimensional accuracy; molds production, low Wall thickness min 0,5 HPDC - The high accuracy of the resulting relief allows to mold stability; obtain castings with a minimum wall thickness; - Gas porosity in castings [mm] max 10 - Fine-grained casting structure can reduce tightness Cooling rate [ºC/s] 20-100 - Preventing the formation of shrinkage defects and - The low service life of Surface roughness Rа, 1,6... removing gases from the melt; casting moulds; [µm] 0,8 - There is no need for a gating system; - The difficulty of dosing - High solidification and cooling rates promote more the liquid metal into the Admitted weight, [kg] 300 uniform and fine-grained structure and improve the matrix; mechanical properties of products; - Restriction on: LIF min 2 - Reducing the consumption of used materials; · the types of alloys; Wall thickness - Possibility to produce: · the complexity, size, [mm] · both thin- and thick-walled parts; and configurations of the max 100 · castings from both wrought and casting alloys; parts; · parts with minimum allowances for machining; Cooling rate [ºC/s] 40-100 - Ability to obtain castings of complex shape with Compositions of alloys Surface roughness Rа, 1,6... high dimensional accuracy; need to be improved and [µm] 0,8 - High utilization of material, reducing energy costs, adapted for SSM Admitted weight, [kg] 50 reducing further mechanical processing. conditions. SSMC Wall thickness min 2 [mm] max 10 144 Appendix F. Designation system of the heat treatment tempers Table F.1. Designation system of the heat treatment of Al alloys (W-wrought, C- cast) [55,210] Tempers Description Alloys F As-fabricated (as-cast state for the casting alloys). There are no specified W, C mechanical properties and strength levels may vary greatly. O Annealing. Such products are annealed fully to reduce hardness and W, C strength, highest ductility temper. H Strain hardening by cold work. Applies for products that can be strengthened by strain hardening, with or without supplementary thermal treatments. H1 Strain hardening only; W H2 Strain hardening and partial annealing; W H3 Strain hardening and stabilization; W H4 Strain-hardening and painting or lacquering W T Heat treatment to stable condition, excluding annealing (O) T1 Natural aging after rapid cooling from elevated temperatures shaping W process (such as extruding) T2 Cold work after rapid cooling from elevated temperatures shaping process W and subsequent natural aging T3 Solution treatment, quenching, cold work and natural aging W T4 Solution treatment, quenching and natural aging C T5 Artificial aging after rapid cooling from elevated temperatures C T6 Solution treatment, quenching and subsequent artificial aging C T7 Solution treatment, quenching and overaging C T8 Solution treatment, cold work and artificial aging; W T9 Solution treatment, artificial aging and cold work; W T10 Cold work after rapid cooling from elevated temperatures shaping process W and subsequent artificial aging W Solution treatment. An unstable temper applicable only to alloys that W, C naturally aged at RT, after solution heat treatment. This designation is specific only when the period of natural aging is indicated. 145 Appendix G. Binary equilibrium diagrams a) b) c) d) e) f) Figure G.1 Binary diagrams calculated by Thermo-Calc software (TCAl2 database) a) Al-Mg, b) Al-Si, c) Al-Mn, d) Al-Zn, e) Al-Sc, f) Al-Cr 146 Appendix H. Equilibrium phase diagrams of the studied system a) [128] b) [122] c) [168] Figure H.1 The equilibrium phase diagrams calculated by Pandat software: a) Al-XMg-2.4Si; b) Al-5Mg-2Si-XFe; c) Al-8.0Mg-2.6Si-XZn 147