Effect of Cu, Mg and Fe on solidification processing and
microstructure evolution of Al-7Si based foundry alloys
Thèse
Mousa Javidani
Doctorat en génie des matériaux et de la métallurgie
Philosophiae doctor (Ph.D.)
Québec, Canada
© Mousa Javidani, 2015
Résumé
Au cours de la dernière décennie, les alliages de fonderie Al-Si ont été utilisés de plus en
plus comme une alternative appropriée à la fonte dans la fabrication de composants de
moteurs (par exemple les culasses). Les objectifs du projet étaient d'étudier l'effet des
éléments tels que le cuivre, le magnésium et le fer sur les défauts de solidification, et sur
l'évolution des phases poste-eutectiques les alliages de fonderie Al-Si.
Tout d’abord, les travaux antérieurs sont soigneusement examinés afin de mieux
comprendre les charges de fatigue thermomécanique, les caractéristiques, les exigences et
les matériaux applicables dans les composantes du moteur. Par la suite, les défauts de
solidification (tendance de fissuration à chaud (HTS) et microporosité) des alliages à base
d’Al-Si ont été évalués. En augmentant la teneur en Cu et en Fe des alliages, la valeur de
HTS et de microporosité ont été augmentées. Les indices théoriques de fissuration à chaud
ont été simulés avec un modèle de microségrégation multiphasique avec rétrodiffusion dans
la phase primaire «multiphase back diffusion model». La corrélation obtenue entre les
résultats expérimentaux (HTS) et les résultats simulés est excellente.
L’effet de la composition chimique (Cu, Mg et Fe contenu) dans les alliages Al-Si sur
l'évolution de la microstructure ont donc été étudiées. Les microstructures à l'état de coulée
et à l'état de traitement thermique de mise en solution (SHT) ont été évaluées par les
microscopies optique/électronique. Deux intermétalliques contenant du Mg (QAl5Cu2Mg8Si6, π-Al8FeMg3Si6) qui apparaissent avec une couleur grise sous le microscope
optique ont été discriminés par des attaques chimiques que nous avons développées.
L’analyse calorimétrique différentielle à balayage (DSC) a été utilisée pour examiner les
transformations de phase survenant au cours du processus de chauffage et de
refroidissement. Les calculs thermodynamiques ont été effectués pour évaluer la formation
de la phase à l'état d'équilibre et hors-équilibre.
Les résultats ont démontré que la séquence de solidification et la stabilité des
intermétalliques contenant du Cu/Mg ont été fortement influencée par la composition
chimique des alliages. La phase Q-Al5Cu2Mg8Si6 a été solidifiée soit à la même température
ou plus tôt que la phase θ-Al2Cu en fonction de la teneur en Cu de l'alliage. Par ailleurs, les
iii
phases Q-Al5Cu2Mg8Si6 et π-Al8FeMg3Si6 qui étaient solubles à 505
dans l'alliage Al-
7Si-1.5Cu-0.4mg, sont restées presque intactes dans l'alliage Al-7Si-1.5Cu-0.8mg wt.-%.
Bien que l’intermétallique-AlCuFe a été à peine observé dans la microstructure de coulée,
la réaction entre la phase primiare α-Al avec la phase β-Al5FeSi a causé la formation de la
phase N-Al7Cu2Fe au cours de la mise en solution. La transformation de phase à l'état
solide de la phase β-Al5FeSi à la phase N-Al7Cu2Fe a également été étudiée.
iv
Abstract
Over the past decade, Al-Si based foundry alloys have increasingly been used as a suitable
alternative for cast iron in the fabrication of engine components. This project was aimed to
study the effect of Cu, Mg and Fe elements on solidification defects (hot rearing tendency
and microporosity), and on evolution of post eutectic phases in the Al-7Si (wt.-%) based
alloys.
Initially, the previous works and the most pertinent literatures were thoroughly reviewed to
elaborate the thermo-mechanical fatigue loads, characteristics, requirements and materials
applicable in engine components (mainly cylinder-head). Subsequently, the solidification
defects of the Al-Si based alloys were evaluated. By increasing Cu and Fe content of the
alloys, the hot tearing sensitivity and the microporosity content of the alloys were both
enhanced. Multiphase back diffusion model was utilized to simulate the theoretical hot
tearing indices. A very good correlation was obtained between the experimental and the
theoretical hot tearing indices.
Effect of the chemistry (Cu, Mg and Fe content) on microstructure evolution of the Al-Si
foundry alloys was consequently studied. As-cast and solution heat treated (SHT)
microstructures of the alloys were evaluated by optical- and electron-microscopy. Two
etchants were developed to discriminate the Mg-bearing intermetallics (Q-Al5Cu2Mg8Si6,
π- Al8FeMg3Si6) under optical microscope. Differential scanning calorimetry (DSC) was
utilized to examine the phase transformations occurring during heating/cooling process.
Thermodynamic computations were carried out to assess the phase formation in the
equilibrium/non-equilibrium conditions.
According to the predicted/experimental results, the solidification sequence and the
stability of Cu/Mg bearing intermetallics are strongly influenced by the chemistry of the
alloys. Q-Al5Cu2Mg8Si6 phase was solidified either at the same temperature or earlier than
θ-Al2Cu phase depending the Cu content of the alloy. Moreover, Q-Al5Cu2Mg8Si6 and πAl8FeMg3Si6 which were soluble at 505
in the alloy Al-7Si-1.5Cu-0.4Mg, remained
almost intact in the alloy Al-7Si-1.5Cu-0.8Mg wt.-%.
v
Tough the AlCuFe- intermetallic was barely observed in the as-cast microstructure, the
reaction of α-Al with the β-Al5FeSi phase caused the formation of the N-Al7Cu2Fe phase
during SHT. The solid state phase transformation (precipitation temperature and
mechanism) of β-Al5FeSi to the N-Al7Cu2Fe phase was also investigated.
vi
Table of Content
RÉSUMÉ ................................................................................................................................................ III ABSTRACT ............................................................................................................................................ V TABLE OF CONTENT ............................................................................................................................. VII LIST OF TABLES .................................................................................................................................... X LIST OF FIGURES ................................................................................................................................... XI ACKNOWLEDGMENTS ......................................................................................................................... XVII PREFACE .......................................................................................................................................... XVIII CHAPTER 1 INTRODUCTION ...................................................................................... 1 Background ...................................................................................................................................... 2 Objectives ........................................................................................................................................ 5 Structure of thesis ............................................................................................................................ 6 CHAPTER 2 LITERATURE REVIEW ............................................................................ 9 “APPLICATION OF CAST AL-SI ALLOYS IN INTERNAL COMBUSTION ENGINE COMPONENTS” ..................... 9 Thermomechanical fatigue ............................................................................................................ 10 Engine characteristics and requirements ...................................................................................... 15 2.2.1. Engine components and requirements ....................................................................................................... 16 2.2.2. Magnesium alloys...................................................................................................................................... 18 2.2.3. Aluminium alloys ...................................................................................................................................... 19 Description of Al–Si based alloys .................................................................................................. 22 2.3.1. The binary Al–Si system ........................................................................................................................... 22 2.3.2. Influence of iron as impurity ..................................................................................................................... 23 Solidification sequence in 356 and 319 Al alloys .......................................................................... 25 2.4.1. 356-type Al alloys ..................................................................................................................................... 25 2.4.2. 319-type Al alloys ..................................................................................................................................... 26 Effect of microstructural features on TMF strength ...................................................................... 27 2.5.1. Porosity...................................................................................................................................................... 28 2.5.2. Secondary dendrite arm spacing ................................................................................................................ 29 2.5.3. Segregation ................................................................................................................................................ 30 2.5.4. Cracking/debonding of Si particles............................................................................................................ 30 2.5.5. Slip bands .................................................................................................................................................. 31 Strengthening of cast aluminium alloys ......................................................................................... 32 2.6.1. Heat treatment of AlSiCuMg alloys .......................................................................................................... 32 vii
Dispersion hardening..................................................................................................................... 40 Recent developments in Al–Si alloys and applications in engine components ............................... 44 Summary ........................................................................................................................................ 48 CHAPTER 3 MATERIALS AND METHODS. .............................................................. 51 Alloy making and melting: ............................................................................................................. 52 3.1.1. Alloy making and melting procedures to evaluate hot tearing susceptibility ............................................ 52 3.1.2. Alloy making and melting procedures for microstructure evolution ......................................................... 52 Thermodynamic Prediction: .......................................................................................................... 54 Atomic absorption spectroscopy .................................................................................................... 55 Microstructural Analysis: .............................................................................................................. 55 Differential Scanning Calorimetry (DSC): .................................................................................... 56 Heat Treatment: ............................................................................................................................. 57 CHAPTER 4 . ............................................................................................................... 59 “HOT TEARING SUSCEPTIBILITY OF AL-SI BASED FOUNDRY ALLOYS CONTAINING VARIOUS CU, MG AND
FE CONTENT”...................................................................................................................................... 59 Résumé: ....................................................................................................................................................... 59 Abstract: ...................................................................................................................................................... 60 Introduction: .................................................................................................................................. 60 Materials and Method:................................................................................................................... 63 4.2.1. Hot tearing indexation: ............................................................................................................................. 65 4.2.2. Samples preparation and characterization ................................................................................................. 67 4.2.3. Thermodynamic Prediction: ...................................................................................................................... 67 Experimental results and discussion .............................................................................................. 67 4.3.1. Microstructural constituents ...................................................................................................................... 67 4.3.2. Characterization of microporosity............................................................................................................. 70 4.3.3. Hot tearing sensitivity ............................................................................................................................... 71 4.3.4. Hot tear surface analyses........................................................................................................................... 72 4.3.5. Prediction Hot Tearing Susceptibility: ...................................................................................................... 74 Conclusion: .................................................................................................................................... 78 CHAPTER 5 . ............................................................................................................... 81 “EVOLUTION OF INTERMETALLIC PHASES IN MULTICOMPONENT AL-SI FOUNDRY ALLOYS CONTAINING
DIFFERENT CU, MG AND FE CONTENT” ................................................................................................ 81 Résumé: ....................................................................................................................................................... 81 Abstract: ...................................................................................................................................................... 82 Introduction ................................................................................................................................... 82 Experimental procedure................................................................................................................. 84 Results and discussion ................................................................................................................... 85 5.3.1. As-cast microstructure .............................................................................................................................. 85 5.3.2. Microstructure of the solution treated specimens ...................................................................................... 86 5.3.3. Time period of solution treatment ............................................................................................................. 88 5.3.4. High temperature solution heat treatment ................................................................................................. 91 5.3.5. Stability of Q-phase .................................................................................................................................. 92 viii
Conclusion ..................................................................................................................................... 95 CHAPTER 6 . ............................................................................................................... 97 “ASSESSMENT OF POST-EUTECTIC REACTIONS IN MULTICOMPONENT AL-SI FOUNDRY ALLOYS
CONTAINING CU, MG AND FE” ............................................................................................................. 97 Résumé: ....................................................................................................................................................... 97 Abstract: ...................................................................................................................................................... 98 Introduction ................................................................................................................................... 98 Experimental Procedure .............................................................................................................. 101 Results and discussion ................................................................................................................. 101 6.3.1. Microstructure of the alloys ..................................................................................................................... 102 6.3.2. Thermal analysis of as-cast specimens .................................................................................................... 105 6.3.3. The N-phase ............................................................................................................................................ 108 6.3.4. Sequence of the θ- and Q-phases transformation in heating/cooling processes ....................................... 116 6.3.5. Effect of Cu content on the post-eutectic phases ..................................................................................... 119 Conclusion ................................................................................................................................... 120 CHAPTER 7 . ............................................................................................................. 123 “SOLUBILITY/ STABILITY OF CU/MG BEARING INTERMETALLICS IN AL-SI FOUNDRY ALLOYS CONTAINING
DIFFERENT CU AND MG CONTENT” .................................................................................................... 123 Résumé: ..................................................................................................................................................... 123 Abstract: .................................................................................................................................................... 124 Introduction: ................................................................................................................................ 124 Materials and methods................................................................................................................. 127 Results and Discussion ................................................................................................................ 129 7.3.1. Characterizing the microconstituents under OM: .................................................................................... 129 7.3.2. Stoichiometry of the phases after etching: ............................................................................................... 134 7.3.3. Effect of Cu/Mg content of the alloys on evolution of as-cast microstructure ......................................... 134 7.3.4. Effect of Cu/Mg content on maximum applicable SHT temperature....................................................... 137 7.3.5. Microstructure evolution and age hardening after SHT at 505 : ........................................................... 139 7.3.6. Effect of high temperature SHT on dissolution of intermetallics ............................................................ 142 General discussion....................................................................................................................... 144 Conclusion: .................................................................................................................................. 149 7.4.1. Stability of the Cu/Mg bearing intermetallics:......................................................................................... 145 CHAPTER 8 PERSPECTIVE AND GENERAL CONCLUSIONS.............................. 151 General conclusions .................................................................................................................... 152 Recommendations for future works: ............................................................................................ 157 CHAPTER 9 APPENDIX ........................................................................................... 159 Appendix (1): calculation of R (ratio of solidification shrinkage)........................................................... 159 Appendix (2): Back diffusion model (BDM) .............................................................................................. 161 BIBLIOGRAPHY.............................................................................................................. 165 ix
List of Tables
Table 2-1: Weight reduction results for CGI vs. grey cast iron cylinder blocks 65 ........................................... 18 Table 2-2: Chemical composition (wt-%) of 356-type, 319-type and 390-type Al alloys ................................ 21 Table 2-3: Some major properties of the Al 319-, 356-, and 390- type alloys91 ............................................... 21 Table 2-4: Reactions occurred during solidification of A356.2 7 ...................................................................... 26 Table 2-5: Summary of reactions occurring during solidification of 319.1Al alloys ........................................ 27 Table 2-6: Chemical composition, mechanical strength and creep properties of Al-Si alloys233 ...................... 48 Table 3-1: chemical composition of the alloys used to evaluate hot tearing susceptibility (wt.%) ................... 52 Table 3-2: chemical composition of the alloys (wt.%) used for microstructure evolution................................ 54 Table 4-1: chemical composition (wt.%) and SDAS of the alloys .................................................................... 63 Table 4-2: Mould temperature of the alloys ...................................................................................................... 65 Table 4-3: Crack size parameters for hot tearing index .................................................................................... 66 Table 5-1: Chemical composition of the Al alloys (wt.%) ................................................................................ 84 Table 6-1: chemical composition of the alloys (wt.%).................................................................................... 101 Table 6-2: the solidification temperatures K ( ) of the post eutectic phases predicted by the MBD model 1.
....................................................................................................................................................... 104 Table 7-1: chemical composition of the alloys (wt.%).................................................................................... 128 Table 7-2: stoichiometry of Q- and π-phases measured before and after etching. .......................................... 134 Table 7-3: concentrations of Mg element in α‐Al after different SHT conditions in the studied alloys. ............................................................................................................................... 145 Table 8-1: mass density of the secondary phases in Al-Si based foundry alloys. ........................................... 160 x
List of Figures
Figure 2-1 a) out-of-phase and b) in-phase thermo mechanical loading16 ....................................................... 11 Figure 2-2 Active damaging mechanisms during an OP- TMF cycle16, 23, 24 .................................................... 11 Figure 2-3 Photographs of a typical crack initiation area in the cylinder head, a) reprinted with permission
from American Foundry Society 27, b) copyright © 2006 SAE International; reprinted with
permission 28 .................................................................................................................................... 12 Figure 2-4 Hoop-stress vs. hoop-strain at valve bridge 36 ................................................................................. 13 Figure 2-5 Typical thermal loading in valve bridge (e.g. point A) of cylinder head (maximum operating temperature 573K), (reprinted with permission from Elsevier)20, 21 ........................................... 14 Figure 2-6 Increasing the peak firing pressure in truck engines to fulfil emissions standards requirements58 16 Figure 2-7 a) Cross-section of a cylinder head b) engine block, (reprinted with permission from Taylor &
Francis)55 .......................................................................................................................................... 17 Figure 2-8 The relation between vehicle mass and fuel consumption68, 69 ........................................................ 17 Figure 2-9 Material properties of compacted graphite iron (CGI-400), grey cast iron (GJL-250), Al-A390,
AlSi9Cu and Mg-MRI 230D 65, 74, 81. ............................................................................................... 20 Figure 2-10 Peak firing pressure limits for various materials in diesel engine cylinder block58, 81 ................... 20 Figure 2-11 The equilibrium phases diagram of the Al–Si alloy system92, 93.................................................... 22 Figure 2-12 The role of Al5FeSi in the formation of shrinkage porosity, (reprinted with permission from
Taylor and France)137 ....................................................................................................................... 25 Figure 2-13 SEM micrograph of A356 as-cast Al alloy showing the close association between Al5FeSi and πAl8FeMg3Si6 phase, (reprinted with permission from Springer)140 .................................................. 26 Figure 2-14 Effect of cooling rate on the formation of β-Al5FeSi brittle phase38 ............................................ 29 Figure 2-15 SEM images of: a) debonded (reprinted with permission from Elsevier)150, and b) fractured Si
particle (reprinted with permission from Springer)94 ....................................................................... 31 Figure 2-16 Temperature ranges for heat treatment and relevant solvus line for binary aluminum alloys 193 .. 33 Figure 2-17 DSC curves of AlSiCuMg alloy solution treated at 773K for different times (reprinted with
permission from Elsevier)6............................................................................................................... 35 Figure 2-18 Cu concentration measured across dendrite arms in different solutionising times at 768 K (495 )
for various samples: (a) SDAS 50 μm, (b) SDAS 25 μm and (c) SDAS 10 μm, (reprinted with
permission from Elsevier)192 ............................................................................................................ 36 Figure 2-19 DSC curve of Al7Si3Cu0.4Mg alloy, solution-treated 10hrs@773K, water-quenched and aged for
different times at 443K, (reprinted with permission from American Foundry Society)121 .............. 38 Figure 2-20 a) L12, (b) D022, and (c) D023crystal structures, (reprinted with permission from Elsevier)252 . 42 Figure 2-21 Calculated diffusivities for different solute elements at 573K (300 ), 673 K (400 ), and 933 K 660 (Tm of Al), (reprinted with permission from Carl Hanser Verlag) 240.............................. 43 Figure 3-1: Pyrex tubes and propipette used in sampling ................................................................................. 54 xi
Figure 3-2: SDAS mesurement of the specimens ............................................................................................. 56 Figure 3-3: cooling curve of the alloy RC3(M0) during quenching, recorded by Data-Logger (OM-DAQPRO5300). ............................................................................................................................................... 57 Figure 4-1: Schematic view of the ring mould used for the investigation of hot tearing tendency ................... 64 Figure 4-2: macrographs to illustrate the different severity levels of hot tearing in the alloys ......................... 66 Figure 4-3: As-cast microstructures of the four alloys studied. ........................................................................ 69 Figure 4-4: DSC cooling curves of the alloys RC3 and RC3F0.7, with a scanning rate of 10 K/min.. ............ 69 Figure 4-5: evolution of the main intermetallic phases (calculated by MBD model) in alloys: (a) RC3, (b)
RC3F0.7. .......................................................................................................................................... 70 Figure 4-6: microporosity content in the alloys. ............................................................................................... 71 Figure 4-7: Microstructure of 3 alloys showing the micropores formed near the hot spot. The amount of
porosity in these microstructure are in ascending order from left to right and can be representative
of the 3 categories of alloys.............................................................................................................. 71 Figure 4-8: Hot tearing index (HTS) of the studied alloys ................................................................................ 72 Figure 4-9: SEM micrographs of the hot tear section in the alloys. .................................................................. 73 Figure 4-10: Optical micrograph across the tear region of the alloys a) RC3(M0) and b) RC3F0.7(M0), solidarrow: β-Al5FeSi phase, dash-arrow: Si-particles. ........................................................................... 74 Figure 4-11: Physically blocking the metal feeding by β-Al5FeSi phase ......................................................... 74 Figure 4-12: Plot of HTS versus HCS assuming ocr= 0.02 and fraction liquid =0.78 at the dendrite coherency
point. ................................................................................................................................................ 75 Figure 4-13: Plot of porosity area versus HCS calculated with cr = 0.02 and fraction liquid =0.78 at the
dendrite coherency point. ................................................................................................................. 76 Figure 4-14: Plot of porosity area versus R calculated with cr = 0.02 and fraction liquid =0.78 at the dendrite
coherency point. ............................................................................................................................... 77 Figure 4-15: Plot of HTS versus R calculated with cr = 0.02 and fraction liquid =0.78 at the dendrite
coherency point. ............................................................................................................................... 78 Figure 5-1: As-cast microstructures of: a) alloy (RC0.5), b) alloy (RC3). ........................................................ 86 Figure 5-2: Heating DSC curves of the alloys in as-cast condition................................................................... 86 Figure 5-3: a) Comparison of DSC curves of alloy (RC0.5) in as-cast and after solution treatment
(2h@502C+8h@540C), b) remaining of π-phase after the first step of solution treatment
(2h@502C). ...................................................................................................................................... 87 Figure 5-4: DSC curves of a) alloy (RC1.5) and b) alloy (RC3) after different solution treatment times at 502C (935F). .............................................................................................................................................. 88 Figure 5-5: Remnant of θ‐ and Q- phase in alloy (RC3) after 8 hours of solution treatment............................ 89 Figure 5-6: Remnant of Q-phase in alloy (RC3) after 15 hours of solution treatment at 502C (935F)............. 90 Figure 5-7: Calculated mass fraction of phases vs. temperature (in equilibrium state) for Al-7Si-3.5Cu0.35Mg containing a) 0.15 and b) 0.75 wt % Fe. ............................................................................. 90 Figure 5-8: Area fraction of intermetallic phases containing Cu (EPMA mappings), and the area under DSC
curves (mW/mg) corresponding to peaks II and III in alloy (RC3F0.7). ......................................... 91 Figure 5-9: Incipient melting after five hours of solution treatment of alloys (RC3F0.7) at 535C (995F). ...... 92 xii
Figure 5-10: DSC curves of alloys (RC3) and (RC3F0.7) in as-cast condition and after solution treatment
(5h@535C). ...................................................................................................................................... 92 Figure 5-11: Variation of the dissolution temperature of Q-phase with chemical composition (calculated with
Thermo-Calc). .................................................................................................................................. 94 Figure 5-12: Phase field distribution of Al-7Si-xCu-xMg-0.15Fe at 500C (932F), calculated by ThermoCalc).
......................................................................................................................................................... 94 Figure 6-1: as-cast microstructures of the experimental alloys: a) alloy#1 (RC0.5), b) alloy#4 (RC3), c)
alloy#5 (RC3F0.7), d) alloy#6 (RC3(M0)). ................................................................................... 103 Figure 6-2: mass fraction of phases vs. temperature as calculated with the MBD model 1. ........................... 104 Figure 6-3: DSC heating curves of the as-cast specimens with scanning rate of 10K/min. The solidus
temperatures (Ts) given above were calculated with the MBD model 1......................................... 107 Figure 6-4: DSC cooling curves of alloys #3 to #5 (RC1.5, RC3 and RC3F0.7) with a scanning rate of 5
K/min. The starting temperature of the DSC cooling tests was 933 K (660 . ............................ 108 Figure 6-5: morphology of the AlCu-eutectic and block-like AlCuFe- intermetallic phases in alloy RC3
(prepared with the permanent mould). ........................................................................................... 109 Figure 6-6: a) alloy #7 (RC3F0.7(M0)) in as-cast condition; b) alloy #7 heated up to 800 k (527 , just beyond
peak II) and rapidly cooled (a and b were taken at the same location). ......................................... 110 Figure 6-7: a) alloy #7 (RC3F0.7(M0)) in as-cast condition; b) alloy #7 heated up to 810 K (537 , just
beyond peak III) and rapidly cooled (a and b were taken at the same location). ........................... 110 Figure 6-8: elemental mapping of alloy #5 (RC3F0.7) in as-cast condition. .................................................. 111 Figure 6-9: elemental mapping of alloy #5 (RC3F0.7) solution heat treated 15 min. at 775 K (502
and
quenched. ....................................................................................................................................... 112 Figure 6-10: elemental mapping of alloy #5 (RC3F0.7) solution heat treated 20 hours at 502 ................... 112 Figure 6-11: DSC curves of alloy #5 (RC3F0.7) in as-cast and solution heat treated conditions; scanning rate
10 K/min. ....................................................................................................................................... 112 Figure 6-12: phase morphology, EBSD pattern and simulation results for the N-Al7Cu2Fe phase in alloy #5RC3F0.7 (MAD=0.2). .................................................................................................................... 114 Figure 6-13: N-Al7Cu2Fe phase under a) optical microscope and b) SEM. .................................................... 114 Figure 6-14: microstructure of alloy #7 (RC3F0.7(M0)) a) 10 hours solution treated at 775 K (502 ), b) 10
hours solution treated at 775 K (502 ), quenched and 10 min. solution treated at 810 K (537 );
(a and b were taken at the same location); (solid black arrows are AlCuFe-, dotted red arrows are
AlFeSi- and dashed arrows are AlCu- intermetallics).................................................................... 115 Figure 6-15: microstructure of alloy #5 (RC3F0.7) a) 10h solution treated at 775 K (502 ), b) 10h solution
treated at 775 K (502 ), quenched and 10 min. solution treated at 803 K ((530 )); (a and b were
taken at the same location). ............................................................................................................ 116 Figure 6-16: effects of Cu and Mg contents on the precipitation temperature of the Q- and θ-phases; predicted
by the MBD model. ....................................................................................................................... 116 Figure 6-17: a) alloy #4 (RC3) in the as-cast condition. b) alloy #4 heated up to 787 K (514 , just beyond
peak I) and quenched (a and b were taken at the same location); (prepared with the permanent
mould). ........................................................................................................................................... 118 Figure 6-18:a) alloy #3 (RC1.5) in the as-cast condition. b) alloy #3 heated up to 787 k (514 , just beyond
peak I) and quenched (a and b were taken at the same location); (prepared with the permanent
mould). ........................................................................................................................................... 118 xiii
Figure 6-19: a) alloy #4 (RC3) in the as-cast condition. b) alloy #4 heated up to 803 K (530 , just beyond
peak III) and quenched (a and b were taken at the same location); (prepared with the permanent
mould). ........................................................................................................................................... 118 Figure 6-20: DSC cooling curves of alloys #1 (RC0.5), #3 (RC1.5) and #4 (RC3), with a scanning rate of 5
K/min. ............................................................................................................................................ 120 Figure 6-21: evolution of mass fraction and temperature formation of the π-phase with Cu content (in Al-7SixCu-0.35Mg-0.15Fe), predicted by the MBD1. .............................................................................. 120 Figure 7-1: microstruture of alloy #2-RC1M0.8 (SHTed 14h @ 505 : a) microstruture before tretment with
the etchant, b) after treatment with (HNO3), c) after treatment by (HNO3+HCl); the micrographs
were taken at the same coordinate.................................................................................................. 130 Figure 7-2: as-cast microstructure of alloy #4-RC1.6M0.8: a) before tretment with the etchant, b) etched with
HNO3 c-e) EPMA elemental mapping of the Cu, Mg and Fe elements; the dashed lines in EPMA
micrographs correspond to the same coordinate of the OM micrographs. ..................................... 132 Figure 7-3: as-cast microstructure of alloy #4-RC1.6M0.8: a) before tretment with the etchant, b) etched with
HNO3+HCl c-e) EPMA elemental mapping of the Cu, Mg and Fe elements; the dashed lines in
EPMA micrographs correspond to the same coordinate of the OM micrographs. ......................... 133 Figure 7-4: SHTed (6h@505 ) microstructure of alloy #2-RC1M.08 a) before treatment with the etchant,
and b) after being etched with HNO3; the micrographs were taken at the same coordinate; gray
colour of Q-phase was changed to dark colour (like Mg2Si) after being etched. ........................... 133 Figure 7-5: a EPMA elemental mapping of the Mg element correspond to the dashed area in Figure 7-2-d, b)
the area correspond to Q-phase (white area) was manually masked, c) the hue in threshold color of ImageJ 134, 168 and the counted area fraction is 2.9% d the hue 134, 169 and the counted area fraction is 10.7%. ................................................................................................... 134 Figure 7-6: The effect of Mg content on phase fractions in Al-7Si-1Cu-0.07Fe-xMg (wt.%); the dashedvertical-red lines correspond to the chemistry of the #3-RC1.6M0.4 and #4-RC1.6M0.8 (the results
were predicted by MBD1). ............................................................................................................. 135 Figure 7-7: as-cast microstructures of the experimental alloys: a) alloy#1 (RC1M0.4), b) alloy#2 (RC1M0.8),
4 c) alloy#3 (RC1.6M0.4), d) alloy#4 (RC1.6M0.8)...................................................................... 136 Figure 7-8: the quantified area fractions and predicted volume fraction by MBD1 of the phases Q Mg2Si and π β in as‐cast condition vs. ratio of Cu/Mg............................................... 137 Figure 7-9: microstructure of alloy #2 (RC1M0.8), a) in as-cast condition, b) after 5 hours SHT at 525 ; the
micrographs were taken at the same coordinate; Cu-bearing intermetallics were melted but Mg2Si
and π-phases were remained almost intact. .................................................................................... 138 Figure 7-10: evolution of as-cast microstructure (a, d & g) after applying 1st step SHT (6h@505 ) (b, e & h)
and after 2nd step SHT (8h@525 ) (c, f & i); a-c: alloy #2 (RC1M0.8), d-f: #3 (RC1.6M0.4) and
g-i: alloy #4 (RC1.6M0.8); the micrographs of each alloy were taken at the same coordinate...... 139 Figure 7-11: Phase fractions (in as-cast condition, after SHT (14h@505 ) and predicted@505 ) in the
alloys vs. ratio of Cu/Mg a) Q+Mg2Si b) π β . .......................................................................... 141 Figure 7-12: Concentration of the elements [after SHT (14h@505 ) and predicted@505 in the α-Al
matrix vs. ratio of Cu/Mg a) Cu b) Mg c Si. ................................................................................. 141 Figure 7-13: microstructure of alloy #4 (RC1.6M0.8); a) in as-cast condition, b) after 1st step SHT
(6h@505 ); the micrographs were taken at the same coordinate; the phase inside the solid black
line represents Mg2Si phase which shrunk in favour of Q-phase after applying the SHT, Q- and πphases were remained almost intact. .............................................................................................. 141 Figure 7-14: Hardness vs. aging time relation; (specimens were SHTed (14h@505 ), quenched and aged at
180 . ........................................................................................................................................... 142 xiv
Figure 7-15: concentration of the elements (prediction & experiment) in the α-Al matrix of alloy #3
(RC1.6M0.4) vs. the different SHT processes [(14h@490 , (14h@505 ,
(6h@505+8h@520 and 6h@505+8h@530
...................................................................... 143 Figure 7-16: Hardness of alloy #3-RC1.6M0.4 vs. the different SHT processes [(14h@490 , (14h@505 ,
(6h@505+8h@520 and 6h@505+8h@530
...................................................................... 144 Figure 7-17: maximum dissolved Mg (mass %) in α-Al vs. Mg content of Al-7Si-yCu-xMg-0.1Fe alloy
(y=0.5, 1 & 1.5; x= 0.00 to 0.8%) at two different temperatures (505 and 530 ). The
microconstituents in sections: a=(Al+ Si+ β-Al5FeSi), b=( Al+ Si+ β-Al5FeSi+ π-Al8FeMg3Si6),
c=( Al+ Si+ π-Al8FeMg3Si6); for Cu= 0.5 and 1%: d=( Al+ Si+ π-Al8FeMg3Si6+ Mg2Si) and for
Cu= 1.5: d=( Al+ Si+ π-Al8FeMg3Si6+ Q-Al5Cu2Mg8Si6). ............................................................ 145 Figure 7-18: the green plan corresponds to TTS of each phase (the predicted equilibrium precipitation
temperature for θ, Q and π-phases in Al-7Si-yCu-xMg-0.1Fe), the red-plan corresponds to the
(=505 , and the vertical (blue & red) lines correspond to the chemistry of the studied
alloys. ............................................................................................................................................. 147 Figure 7-19: maximum Mg content of the Al-Si alloys soluble at the corresponding temperature; three Al-7Si0.1Fe-yCu-xMg (y=0.5-1.5 & 4%Cu) are compared. .................................................................... 147 Figure 7-20: alloy #3 (RC1.6M0.4), a) in as-cast condition, b) after 1st step SHT (6h@505 ); the
micrographs were taken at the same coordinate; θ-phase and the majority of Q-phase were
dissolved. ....................................................................................................................................... 148 Figure 7-21: alloy #4(RC1.6M0.8), a) in as-cast condition, b) after 10 hours SHT at 505 ; the micrographs
were taken at the same coordinate; Q- and π- phases remained almost intact after SHT and Mg2Si
was transformed to Q-phase........................................................................................................... 148 Figure 7-22: alloy #3 (RC1.6M0.4), a) in as-cast condition, b) after two step SHT (2 hours@505 +
5hour@525 ); the micrographs were taken at the same coordinate; as shown in dashed-dotted
area, the Cu-bearing intermetallics (Q & θ) were melted after the SHT........................................ 149 Figure 9-1: Calculated composition profiles of a specimen obtained at 3 different solidification steps (solid
fractions: 0.25, 0.50 and 0.75), a) in equilibrium condition, b) in Scheil condition, c) in BDM
condition. ....................................................................................................................................... 163 xv
To all my loved ones
xvi
Acknowledgments
I would like to express my sincere gratitude to my supervisor, Professor Daniel Larouche,
for having confidence in me to conduct this project, for his great availability for meetings
and discussions and for his valuable comments and suggestions. His encouragement,
patience, knowledge and advices were very helpful and appreciated all through my studies.
I am also thankful to my co-supervisor, Professor X. Grant Chen, for his insight, support,
and his valuable comments and discussions throughout this project. This dissertation would
not have happened without you both. I am also thankful to my thesis evaluation committee.
Thanks to all of the staff of Mining, Metallurgical and Materials Engineering Department
of Laval University for their help and support. Special thanks go to Marc Choquette,
Maude Larouche and André Fernand for their help with microstructural analyses, Daniel
Marcotte and Vicky Dodier for their availability, collaboration and technical assistance in
the laboratory. I am grateful to Amir R. Farkoosh (from McGill) and Honoré Kamguo
Kamga† for fruitful discussions, Zhan Zhang, Mohammad Shakiba and Kun Liu (from
UQAC) for their assistance in Scanning Electron Microscopy studies. Many thanks to all
my colleagues and friends in the department for their kind support, help, suggestions, and
making a joyful environment.
Finally, and most importantly I would like to thank my family for their encouragement,
sacrifices and patience. I am grateful to my parents, brothers and sisters for their dedication,
support and love in all and every stage of my life. Above all, I would like to thank my
loving wife, Sheida, for her endless understanding, encouragement and patience with me.
This dissertation would not have been possible without the support and love of my family.
xvii
Preface
To reveal the performance requirements for the engine components (engine blocks and
cylinder heads), the operating service conditions need to be thoroughly reviewed. Three
different loads that are applied on the cylinder head have to be considered: the assembly
load, the load produced by combustion pressure and the thermal load. The effects of
thermal load on the fatigue lifetime of a cylinder head are overwhelmingly greater than
those of the other loads. In a start–stop cycle, an engine might be warmed up from 243 K (30 ) in a cold winter to over 523 K (250 ). During such a thermal cycle, large
thermal/mechanical loads are applied on the engine components because of non-uniform
thermal expansion/contraction of different engine parts.
Engine components have historically been manufactured in cast iron owing to its inherent
high-temperature strength; however cast iron is a very dense material (~7.5 g cm−3).
Demands to improve fuel economy and to reduce emissions necessitate replacement of cast
iron with lighter metals. Excellent thermal conductivity and lower density make Al–Si
foundry alloys a suitable alternative for cast iron in the fabrication of engine components.
The increase in the maximum operation temperature and pressure of engines necessitates
improving the thermomechanical fatigue (TMF) performance of Al-Si alloys. Casting
defects are of the major parameters to affect the TMF performance of Al-Si alloys. In
defect-limited specimens, crack initiation can be significantly delayed.
Copper and Mg play a vital role in the strengthening of Al-Si alloys. To maximize the
efficiency of strengthening, the large post-eutectic phases (e.g. θ-Al2Cu and QAl5Cu2Mg8Si6) must be dissolved and re-precipitated by applying appropriate heat
treatment. The temperature(s) and reaction(s) of the last solidified eutectic phases are
critical parameters in the optimization of the solution heat treatment. Moreover, the Fe
content of the alloys, by which the solidification process and the overall mechanical
properties of the alloys are significantly affected, must be taken into account.
This doctoral thesis is presented to the department of mining, materials and metallurgical
engineering of Laval University. Financial assistance received from the Natural Sciences
xviii
and Engineering Research Council of Canada (NSERC), Rio-Tinto-Alcan (RTA) and
Fonds de recherche du Québec- Nature et technologies (FRQ-NT) by the intermediary of
the Aluminium Research Centre (REGAL) is gratefully acknowledged. The project was
carried out under supervision of Professor Daniel Larouche and co-supervision of Professor
X. Grant Chen. This thesis has been prepared as an article insertion thesis and includes five
articles, which at the time of the thesis submission, were mostly published or submitted for
publication.
My contribution to these articles was: define the objective of each article, prepare the plan
of experiments, design/assembly of the experimental set-ups and perform the experiments
as follow: modify the design of the ring mould test, present a semi-quantitative indexation
method, study the post eutectic reactions and evaluate the stability/solubility of the post
eutectic phases in the Al-Si hypoeutectic alloys. A computational algorithm developed by
Larouche1, was used in the thermodynamic computations to calculate the mass fraction of
phases and to simulate the theoretical hot tearing index. I subsequently prepared the first
draft of the articles, which were all revised by the co-author(s) before submission.
The first article titled: “Application of Cast Al-Si Alloys in Internal Combustion Engine
Components” co-authored by Professor Daniel Larouche, is a literature review paper and
has been published in the journal of International Materials Reviews, 2014, Vol. 59, No. 3,
pp. 132-158.
The second article titled: “Hot Tearing Susceptibility of Al-Si foundry alloys containing
vairous Cu, Mg and Fe content”, co-authored by Professor Daniel Larouche, has been
written and is ready to submit.
The third article titled: “Evolution of Intermetallic Phases in Multicomponent Al-Si
Foundry Alloys Containing Different Cu, Mg and Fe Content”, co-authored by Professor
Daniel Larouche and Professor X. Grant Chen, has been published in American Foundry
Society (AFS) Transactions, 2014, Vo. 122, No. 14-056.
The forth article titled: “Assessment of post-eutectic reactions in multicomponent Al-Si
foundry alloys containing Cu, Mg and Fe”, co-authored by Professor Daniel Larouche and
Professor X. Grant Chen, has been published in Metallurgical and Materials Transactions
A, 2015, Vol. 46, No. 7, pp. 2933-2946.
xix
The fifth article titled: “Solubility/ Stability of Cu/Mg Bearing Intermetallics in Al-Si
Foundry Alloys Containing Different Cu and Mg Content”, co-authored by Professor
Daniel Larouche and Professor X. Grant Chen, has been written and is ready to submit.
I collaborated on the following published/accepted articles, as well:
1) Farkoosh A., Javidani M., Hoseini M., Larouche D., Pekguleryuz M., “Phase formation
in as-solidified and heat-treated Al–Si–Cu–Mg–Ni alloys: Thermodynamic
assessment and experimental investigation for alloy design”, Journal of Alloys and
Compounds, 2013, 551(0): p. 596-606.
2) Larouche D., Javidani M., “Mathematical analysis of the heat measured by a power
compensated differential scanning calorimeter during the solidification of a
multiphase alloy”, Journal of Thermal Analysis and Calorimetry, (accepted on 16
May 2015).
xx
Chapter 1
Introduction
Background
Engine blocks and cylinder heads are the fatigue critical automotive components which
experience two distinct types of fatigue failure in service: mechanical fatigue, as a high
cycle fatigue (HCF), is initiated by the variation of pressure within the combustion
chamber; and thermal fatigue, as a low cycle fatigue (LCF), is originated by the start-stop
cycles of the engine. The effect of thermal load on fatigue life is overwhelmingly greater
than those of mechanical loads (combustion pressure). Thermal fatigue strength is not an
inherent property of the alloy and many parameters are generally involved to improve the
thermal fatigue resistance in the Al-Si alloys: high thermal conductivity, low thermal
expansion coefficient, low porosity level, high/room temperature tensile strength, high
ductility, high creep resistance, high fatigue strength, microstructural stability, small
secondary dendrite arm spacing (SDAS), and low content of coarse intermetallic phases.
Over the past decade, Al-Si casting alloys have increasingly been used in the automotive
industry as a suitable alternative for cast iron in fabrication of engine components. The
major advantage of the Al-Si alloys, besides their high strength to weight ratio, is their
excellent thermal conductivity, which allows the combustion heat to be extracted more
rapidly compared to cast iron. On the other hand, the automotive industry has been ever
facing the challenge of improving efficiency and overall performance of engines. To
increase the efficiency, the maximum operation temperature and pressure of the engine
must be raised. The increase of operation temperature, which leads to softening of
hypoeutectic Al–Si alloys, necessitates high-temperature strengthening of the Al–Si alloys.
Two main categories of the commercial aluminum alloys are commonly used in fabrication
of engine components: 1) Al-7Si-3Cu alloys (e.g. as A319) containing Mg (<0.4 wt.%), and
2) Al-7Si-0.3Mg alloys containing Cu (e.g. A356+0.5Cu-wt.%). Presence of Mg and Cu in
the Al-Si based alloys is required to improve the mechanical strength; while Fe is usually
present as an impurity element.
Performance and fatigue lifetime of Al-Si based alloys (319- and 356-type Al alloys) can
be more influenced by the actual casting processes than by alloy chemistry. Defects (e.g.
porosity and inclusions), which are associated with casting processes, strongly impair the
mechanical strength (in particular fatigue strength). In defect-limited specimens, crack
2
initiation can be significantly delayed. Therefore finding parameter(s), which affect the
casting defects (e.g. porosity and hot-tearing), reserve particular importance to improve the
quality characteristics of Al-Si based alloys.
The soundness of cast Al-Si alloys can be strongly influenced by content of the impurities
(e.g. Fe) and alloying elements (e.g. Cu and Mg). Iron as the most common impurity in AlSi alloys is generally appeared as lamellar intermetallic phases; these iron-bearing
intermetallics reduce the fluidity and enhance shrinkage porosity by physically blocking the
metal feeding. The alloying elements (e.g. Cu) promote the porosity formation by
increasing both the solidification interval (∆T) and the solidification shrinkage. The
solidification interval (∆T) has been reported to increase from 59K (59 , in Al-7Si-0.3Mg
wt.%) to 117K (117 , in Al-7Si-1Cu-0.3Mg wt.%); by further increasing the Cu content
(Al-7Si-4Cu-0.3Mg wt.%), the ∆T was decreased to 109K (109 )2. The overall
solidification shrinkage in Al-Cu binary alloys is ~8%, and in of Al-Si is ~4%3-5.
Precipitation hardening is one of the major strengthening mechanisms of the Al–Si
hypoeutectic alloys. The large eutectic phases (e.g. θ-Al2Cu and Q-Al5Cu2Mg8Si6)
precipitated during solidification weaken the strengthening role of the alloying elements
(Cu and Mg). To maximize the strengthening, the as-solidified large eutectic phases must
be dissolved by applying an appropriate solution heat treatment (SHT), and are reprecipitated as fine evenly distributed metastable phases.
The solution heat treatment is a heating process at a temperature range between the solvus
and the solidus line of the specimen. The time period of the heating process must be long
enough to entirely/ partially dissolve some certain microconstituents. Spheroidisation of the
eutectic Si particles and homogenisation of the alloying elements are the other objectives of
the solution treatment.
The temperature of solution heat treatment (TSHT) must be limited to melting temperature
(Tmp) of the last solidified eutectic phases. Applying SHT at higher temperature causes
incipient melting of the eutectic phases through which the mechanical properties
deteriorates. The last solidified eutectic reaction in Al-Si-Cu-Mg alloys, which involves θAl2Cu and Q-Al5Cu2Mg8Si6 phases, generally reported to occur at ~507
(780 K) 6, 7.
Therefore, solution heat treatment of Al-Si-Cu-Mg alloys is generally restricted to ~500 3
(773 K). It has been reported that the single step SHT is neither able to maximize the
dissolution of Cu rich intermetallic phases, nor is able to homogenize the microstructure
and modify the Si particle. Thus, a two-step SHT has been proposed, by which the Cubearing eutectic phases solidified at the last stage of solidification could be dissolved at the
first step of SHT. The second step of SHT, which could be ~10-35
higher than the TSHT
of the first step, assists to dissolve the remaining Cu-bearing intermetallic phase and further
homogenize the microstructure. It is worth mentioning that the solubility/stability of some
phases is strongly influenced by the content of the alloying elements. Fairly sluggish
dissolution rate or even stability of Q- Al5Cu2Mg8Si6 phase has been reported in the alloys
having high Mg content. Therefore, the content of the alloying elements plays vital role in
the solution heat treatment temperature (TSHT) and in the possibility of applying the second
step of SHT.
The solidification temperature of θ-Al2Cu phase was disputed in literature. Mulazimoglu et
al. 8 reported the precipitation temperature of θ-Al2Cu phase is at ~549
alloy. Samuel
9, 10
in 319.2 foundry
reported the appearance of the θ-Al2Cu phase with two distinct
morphologies, viz. eutectic-like and block-like morphology, at ~520
and at ~533 ,
respectively. The temperature of the reaction reported by Mulazimoglu et al. was neither
confirmed by Samuel 9, 10 nor by other authors 6, 7.
The mechanical properties of Al-Si alloys are significantly influenced by the iron-bearing
intermetallics. Their detrimental effect is directly proportional to their size, density and
morphology. β-Al5FeSi and π-Al8FeMg3Si6 phases are of the major iron-bearing
intermetallics which are frequently observed in Al-Si based foundry alloys. The latter can
be entirely/ partially soluble during solution treatment. Therefore, precipitation/dissolution
temperature of this phase can also influence optimization of the SHT. N-Al7Cu2Fe phase is
another Fe-bearing intermetallic which has been observed in the solution heat treated
(SHTed) specimens by a few studies
11-14
; but the detail of the phase transformation, its
effect on thermal analysis and the influence of chemical composition has never been
studied.
4
Objectives
The major purpose of this work is to investigate the parameters by which the thermomechanical fatigue strength of Al-Si based alloy can be influenced. The TMF loads are
cyclically exerted on the components (e.g. cylinder head) within a certain temperature
range with varying status of stress (tensile and or compression). The TMF strength is not an
intrinsic property to be studied and various mechanical properties must be considered to
improve it. Therefore, it is required to thoroughly review the literature to better understand
the TMF stresses/strains and temperature ranges in engine components, and the parameters
which affect the TMF strength in Al-Si foundry alloys.
Casting defects (e.g. porosity and hot tearing) is one of the major parameters to deteriorate
the TMF strength of Al-Si foundry alloys. The defects are correlated with the solidification
interval (∆T) of the alloys, which in turn, is affected by Cu and Mg content. Moreover,
mechanical properties of the secondary
(i)
319-type Al alloys have been often compared in
literature with the primary A356 alloys containing Cu. However, the high Fe content can be
of the major factor to promote the defects which, in turn, influence the mechanical
properties. Therefore, the first part of this work was designated to study the effect of the
elements (Cu, Mg and Fe) on casting defects (porosity and hot tearing).
In order to enhance the efficiency of precipitation strengthening, the microstructure
evolution of the post eutectic phases must be profoundly investigated. Solution heat
treatment (SHT) is generally limited to ~500
to avoid localized melting of eutectic Q-
and θ- phases. However, there is a controversy between the melting/solidification
temperatures of Q-phase reported in literature with the results predicted by Thermo-Calc.
According to literature, Q-phase is started to melt at ~507 in Al‐Si foundry alloys containing Cu and Mg; but according to Thermo-Calc the melting/solidification
temperatures of Q-phase can be varied by the alloying elements (Cu and Mg). Moreover, in
some references Q-phase has been reported to be soluble, but there are some other
references which reported stability of Q- phase after applying hours of SHT.
i- Recycled aluminum alloys
5
Thermodynamic computation can be a valuable tool to find the correlation of the alloying
elements (mainly Cu and Mg content) with the stability/solubility of Q-phase.
θ-Al2Cu phase has been reported to appear with two distinct morphologies; eutectic-like
morphology with Cu concentration of ~28 wt.% and block-like morphology with Cu
concentration of ~40 wt.% 9, 15, but they have not talked about the rest of the concentration
in the block-like θ-phase. Moreover, they observed the signatures of the two θ-phase
morphologies in thermal analysis during heating process; but they have not reported
presence of the signature (of blocky θ-phase) during cooling process9, 15. Thermodynamic
calculation predicts only one type of θ-phase in Al-Si-Cu-(Mg) system.
π-Al8FeMg3Si6 is an iron-bearing intermetallics which can be entirely/ partially soluble
depending the chemistry of the alloys. Therefore, precipitation/ dissolution temperature of
this phase and its correlation with the chemistry can be one of the criteria in optimization of
SHT. N-Al7Cu2Fe phase is another iron-bearing intermetallics, which has been rarely
reported in solution heat treated microstructures; but the presence of this phase in as-cast
microstructure has never been observed. Evaluating the effect of chemistry of the alloys
and SHT parameters on appearance of N-phase, and the signature of N-phase in thermal
analysis are the other purposes of this work.
Structure of thesis
The PhD dissertation was written in the form of a collection of scientific publications,
which were either published or submitted at the time of the thesis submission. The thesis is
presented in eight chapters:
Chapter 1 is allocated to the general introduction, the problem identification, the
objectives, and the structure of the thesis. In Chapter 2, a literature review on the
application of Al-Si based foundry alloys in the engine components is presented. The TMF
and the structural stress–strain in engine components are initially elaborated. The physical
and mechanical properties of the suitable alternative alloys in manufacturing of engine
components are compared with cast-iron. A detailed review on solidification sequence and
strengthening mechanisms of cast Al-Si alloys are presented. The effect of microstructural
6
features on TMF strength is thoroughly reviewed. The advantages/disadvantages of
application of various Al-Si foundry alloys containing different elements (e.g. Ni, Sc, etc.)
in cylinder heads, which has been studied in last decades, is elaborated.
It is worth
mentioned that this chapter is the first part of the paper published in journal of International
Materials Review. The second part of this paper which was allocated to the characteristics
of the engine block (requirements, applicable materials, procedures to reinforce the cylinder
block wall, etc.) was out of the scope of the thesis.
Chapter 3 This chapter gives a detailed description on experimental methodologies and
procedures. In this chapter, the procedures of preparation of the Al-Si alloys melt and hot
tearing indexation are provided. Procedures of metallography (mounting, grinding and
polishing) for microstructural characterization and differential scanning calorimetry (DSC)
analysis are detailed. Heat treatment applied to evaluate the stability/solubility of post
eutectic phases is described. Thermodynamic computations to calculate the mass fraction of
the phases and to simulate the theoretical hot tearing indices are also explained.
Chapter 4 The results of the second article are presented in this chapter. Solidification
defects (microporosity and hot tearing susceptibility) of seven different Al-Si foundry
alloys (356- and 319-based alloy) were compared. The hot tearing susceptibility (HTS) of
the alloys was ranked by a new semi-quantitative indexation. The HTS and microporosity
were correlated with the combined amount of the Cu and Fe of the alloys. The theoretical
hot tearing indices of the alloys were simulated by multiphase back diffusion (MBD) model
developed by Larouche1. The correlation between the experimental and the theoretical hot
tearing indices was excellent.
Chapter 5 The purpose of this article was to elucidate the evolution of Cu/Mg bearing
intermetallics in Al-Si-Cu-Mg alloys. Four Al-Si alloys containing Cu, Mg and Fe were
investigated. The SHT parameters were optimized to maximize the dissolution of θ-Al2Cu,
π-Al8Mg3FeSi6 and Q-Al5Cu2Mg8Si6 phases while minimizing the loss of Cu into NAl7Cu2Fe phase.
Chapter 6 The effect of Cu, Mg and Fe content on post eutectic reactions occurring in Al-
Si based foundry alloys was studied in the third article, and presented in chapter 6. Seven
different Al-7Si based alloys containing various Cu, Fe and Mg content were investigated.
7
The solidification temperature of Mg bearing intermetallics (Q-Al5Cu2Mg8Si6 & πAl8FeMg3Si6) was correlated with the Cu content of the alloy. The AlCuFe-intermetallic
compound, which was barely found in the as-cast microstructure, significantly enhanced
after SHT; this intermetallic compound was mostly detected as N-Al7Cu2Fe phase after
applying SHT.
Chapter 7 This article was aimed to specify the chemistry of Al-Si alloys for which the
Cu/Mg bearing intermetallics (θ, Q, Mg2Si and π) are all soluble. Four Al-Si based alloys
containing various Cu (1 and 1.6 wt.%) and Mg (0.4 and 0.8wt.%) contents were
investigated to assess with further details the effect of chemistry on evolution of Cu/Mg
bearing intermetallics. Two etchants were developed to distinguish the Mg bearing
intermetallics (Q-Al5Cu2Mg8Si6 & π- Al8FeMg3Si6) under optical microscope. The
chemistries of Al-Si alloys (the range of Cu and Mg content of the alloys), for which the
whole Cu/Mg bearing intermetallics are soluble, were predicted by Thermo-Calc.
Chapter 8 summarizes the major achievements and concludes the obtained results in this
project. In addition, recommendations are provided for future work.
8
Chapter 2
Literature review
“Application of Cast Al-Si Alloys in Internal Combustion Engine Components”
This chapter, which is parts of the paper published in journal of International Materials
Review, summarizes the literature most pertinent to the subject of this thesis. It has been
composed of eight main sections: the first section describes the thermo-mechanical fatigue
(TMF) in engine components. The second section elaborates characteristics, requirements
and materials applicable in engine components. The sections three and four deal with the
specifications and solidification sequence of the Al–Si foundry alloys. The fifth section
introduces the microstructural features of Al-Si foundry alloys which affect TMF strength.
The sixth section presents the strengthening mechanisms of Al-Si alloys. The seventh section
lists the Al-Si alloys used in engine components and their developments in last decades.
Thermomechanical fatigue
The cyclic stresses required to cause fatigue failure at elevated temperature (0.3Tm  T 
0.7Tm) do not necessarily result from the application of external loads; they could also be
created by cyclic thermal stresses. Thermal stresses are produced when the change in
dimensions of a member, which is in turn the result of a temperature change, is restricted
by some kind of constraint. For instance, in a fixed end bar, the thermal stress produced by
a temperature change (T) can be expressed as (if no plastic strain):
 
d d

 T    E  T
dT d
(1)
where  is the linear coefficient of thermal expansion and E is the elastic modulus.
Under thermomechanical conditions, the total strain (tot) is the sum of thermal strain (th)
and mechanical strain (mech) components, the latter being composed of elastic (el) and
inelastic strain (in) components:
tot  th mech  (T T0 ) el in
(2)
where T0 is the reference temperature and T is the test temperature.16, 17
In thermomechanical fatigue (TMF), thermal and mechanical strains with different phasing
might be applied to specimens.18 Two major cycles are generally employed in a TMF test:
(a) in-phase cycle, where the mechanical strain and thermal strain are at the same phase
(e.g. maximum strain at maximum temperature); and (b) out-of-phase cycle, where
mechanical strain is maximum at minimum temperature. Variations of strain components
(thermal/mechanical and total strain) with time corresponding to OP TMF (out-of-phase
TMF) and IP TMF (in-phase TMF) cases are illustrated in Figure 2-1.16, 19
10
Strain
a)
Hot
Hot
Δεmech
εth
εtot
t (sec)
εmech
Cold
Cold
Hot
Strain
b)
εth
εtot
εmech
Hot
Hot
Δεmech
t (sec)
Cold
Cold
Figure 2-1 a) out-of-phase and b) in-phase thermo mechanical loading16
The governing damage mechanism in engine components (e.g. cylinder heads) has been
reported to be OP TMF cycles.20, 21 In each cycle of OP TMF, since a specimen crosses a
temperature range, it can be affected by a variety of thermally activated processes (as
illustrated in Figure 2-2). The damage mechanisms can affect the specimen either
individually or in mutual interactions. The major damage mechanisms in TMF processes
are activated by fatigue, environment (oxidation) and creep.19, 22
n
atio
orm
f
e
d
y
stic
ion
er Pla
v
iat ation
t
co ess
i
e
n g
R roc
k i pa
p
ac pro
r
C nd
a
Cold
Cr
effe eep
cts
O
xi
da
tio
n
σ
Hot
Hardening
process
Cyclic ageing
ε
Coarsening effects Plastic deformation
Figure 2-2 Active damaging mechanisms during an OP- TMF cycle16, 23, 24
Because of the complex geometry, thermal/mechanical strains in a cylinder head are known
to be larger than in an engine block; therefore the former is more susceptible to failure by
TMF. Detailed information about geometry, constituent parts and applied conditions on
cylinder heads can be found elsewhere.25, 26 Figure 2-3 shows two pictures of typical crack
initiation areas in cylinder heads.
11
Figure 2-3 Photographs of a typical crack initiation area in the cylinder head, a) reprinted with permission from
American Foundry Society 27, b) copyright © 2006 SAE International; reprinted with permission 28
Several studies28-30 have been done to simulate/measure the thermal/mechanical stress–
strain variations and temperature gradient in cylinder heads. As mentioned above, three
loads on the cylinder head must be taken into account: the assembly load, the load
produced by combustion pressure and the thermal load. The assembly load is generated by
the screws connecting the cylinder head to the engine block, press fitting of valve seats and
hot plug. The peak firing pressure, which is generated by combustion pressure, can reach
values up to 200 bar in diesel engines.20, 31-33 In a start–stop cycle, the engine is warmed up
to over 523 K (250
), with a strong temperature gradient being created during operation
between the water cooled flame deck (from 373 to 393 K, (100 to 120
combustion chamber face (from 523 to 573 K, (250 to 300
)) and the
)). The constraint imposed on
thermal expansion creates the most significant operating stresses at the critical flame-face
sections of the cylinder head (e.g. valve bridge). The thermal load affects the fatigue
lifetime to a far greater extent than the other two loads mentioned.33-35
Figure 2-4 illustrates the calculated hoop strain–hoop stress for the first through third hot–
cold cycle in the valve bridge area of a cylinder head. At the beginning, assembly loading
generates a tensile hoop stress. The stress is compressive during heating which becomes
tensile upon cooling of the assembly. As illustrated, the mean hoop strain is compressive
while the mean stress is rather tensile during the temperature cycle.34, 36
12
Hoop stress (MPa)
200
100
0
-100
-200
-0.8
-0.6
-0.4
-0.2
0
0.2
Hoop strain (%)
Figure 2-4 Hoop-stress vs. hoop-strain at valve bridge 36
Two distinct fatigue modes control the lifetime of engine cylinder heads: mechanical
fatigue and thermal fatigue. Mechanical fatigue, as a high cycle fatigue (HCF) in cylinder
heads, is driven by the fluctuation of pressure in the combustion chamber. The thin walls
(thickness ~10 mm), adjacent to the water ducts in the valve bridge of a cylinder head, are
the critical locations for mechanical fatigue crack initiation. The temperature range in these
areas has been reported to be 393 K (120
, at lower engine speed) up to 443 K (173
, at
higher speed).37 The design of cylinder heads, the intrinsic fatigue strength of the alloy and
residual stresses induced by heat treatment are the three major factors which significantly
affect the mechanical fatigue resistance.37, 38 Thermal fatigue, as a low cycle fatigue (LCF),
is driven by the start–stop cycles of the engine. The typical thermal stress and strain cycles
in the valve bridge (i.e. point A) of a cylinder head are illustrated in Figure 2-5. The
thermomechanical loading factor KTM = −(mech/th) is ~0.75 in the cylinder head. It seems
that the influence of HCF loadings on the lifetime is small; the typical ignition pressure is
less than 200 bar, and the time of the HCF loading occurring is superimposed with the
heating period and dwell time during which the stress is compressive.20, 21, 23
13
T
Transient
temperature
Engine start
t
σ
A
t
Plastic deformation
ɛ
superimposed HCF-loading
Stress relaxation
ɛth
ɛmech
t
superimposed HCF-loading
Figure 2-5 Typical thermal loading in valve bridge (e.g. point A) of cylinder head (maximum operating temperature 573K), (reprinted with permission from Elsevier)20, 21
The mechanism of fatigue failure can be explained as follows. After ignition of the engine,
the valve bridge is heated up and the temperature becomes quite high (exceeds 523 K
(250 )) relative to the circumference of the combustion chamber. The bridge section tends
to expand but cannot do so freely, since it is constrained by the water cooled flame deck
across which it is suspended. This creates a local compressive stress field within the bridge
section and induces compressive yielding. The most severe stress is created when the
temperature difference between the combustion chamber and the water cooled flame deck
is the largest (i.e. at the maximum speed). It is important to note that plastic deformation,
which occurs at high temperature, does not cause fatigue cracking (because of being in a
compressive state) as long as the engine is running.30, 31, 34 When the engine is turned off,
the bridge section tends to contract while cooling back to room temperature. The yielded
regions cannot return to the initial condition and tensile stresses are generated in these
regions.34, 39, 40 Therefore, the stress field for the yielding regions of the cylinder head is
compressive at high temperatures, but becomes tensile at low temperatures (as shown in
Figure 2-5). The repetition of these compressive–tensile stress cycles is considered to cause
the cracking in the radial direction. As a result, the number of engine start–stop cycles
could be a better indicator of TMF failure than the mileage of a vehicle.21, 37
Therefore, to prevent crack initiation, the alloy must have either high yield strength to
accommodate stress elastically, or high ductility to delay crack formation.38,
14
41, 42
The
former is required to prevent gas leakage, and the latter is required to prevent cracking in
the valve bridge area of a cylinder head. Another factor that must be taken into account is
the degradation of strength owing to overaging, which makes plastic deformation easier.37,
41, 43
Moreover, there are some other parameters that improve TMF resistance such as:
narrow thermal stress hysteresis loop,41,
44
high thermal conductivity, low thermal
expansion coefficient,45-47 microstructural stability,41,
spacing (SDAS),38,
48
43
small secondary dendrite arm
low porosity level49-51 and low content of coarse intermetallic
phases.41, 52
Engine characteristics and requirements
Diesel engines have become a suitable alternative to gasoline engines over the last decade.
Cars powered by diesel engines account for approximately 50% of the total market share in
Europe (60% in France). Less fuel consumption, lower CO2 emissions and larger power
output and torque of diesel engines are the main reasons for this progress.53, 54 The major
difference between diesel and gasoline engines is their fuel combustion method, which has
been elaborated by Denton.55 Diesel engines operate at a higher compression ratio (between
14:1 and 25:1 compared to gasoline engines at between 8:1 and 12:1) because of the higher
temperature and pressure of the mixture in a diesel cycle.
To increase engine efficiency and fulfil emission standard requirements (Euro legislation),
the maximum operation temperature and pressure of the engine must be raised, in particular
in diesel engines. For instance, the combustion pressure in truck engines was about 125 bar
in 1992 and met the Euro I regulations; but it had to rise above 200 bar to fulfil the Euro V
regulations (see Figure 2-6). This has increased the maximum operating temperature of
cylinder heads from below 443 K (170
K (250
) in earlier engines42, 56 to temperatures above 523
) in recent engines.26, 36 These operating service conditions enhance the specific
power of diesel engines from ~25 kW L−1 up to 75 kW L−1.37, 57
15
Peak firing pressure (bar)
220
200
180
160
140
120
ca.
?
>200
180
160
145
125 125
135
1989 1992 1995 2000
Euro
0
I
II
III
2005 2008
IV
V
Figure 2-6 Increasing the peak firing pressure in truck engines to fulfil emissions standards requirements58
Andersson59 stated that only ~12% of the total vehicle power is transferred to the wheels.
About 15% of the energy is consumed by mechanical losses (mainly frictional) in
powertrain system, the rest of the energy being dissipated in cooling and exhaust systems.59,
60
Funatani et al.61 stated that friction in the engine system can lead to a loss of over 40% of
total power. The major sources of these frictional losses are attributed to the contact
between the piston assembly and cylinder bore.60-62 Therefore, surface modifications of the
cylinder bore could contribute to significant friction reduction, with further benefits for
emissions and fuel economy.61,
63
A 10% decrease in frictional losses could reduce fuel
consumption by about 3%. A volume of 600 L of petroleum could therefore be saved for
each vehicle having an average fuel consumption of 10 L/100 km and running a distance of
200 000 km over its entire lifetime.59
2.2.1. Engine components and requirements
The engine block and cylinder head, which are shown in Figure 2-7, are the two major
components of an engine; both components have historically been manufactured in cast
iron owing to its inherent high-temperature strength. Nevertheless, cast iron is a dense
material (~7.5 g cm−3) and the engine is the single heaviest component within the
powertrain group (~14% of total vehicle mass64). About 3–4% of the total mass of an
average vehicle is generally assigned to the engine block. The improved specifications and
legislations for fuel economy and emissions oblige car manufacturers to make a significant
weight reduction in their products. It has been reported that each 100 kg in weight
reduction could contribute to ~0.5 L of petrol being saved per 100 km driven.64-67 As
illustrated in Figure 2-8, weight reduction of a vehicle by a certain amount could result in
16
significant improvement in fuel economy.68,
69
Social impetus, for instance the US
Partnership for a New Generation of Vehicles (PNGV) program, demands car
manufacturers produce vehicles having a fuel consumption lower than 1 L/30 km.70
Angle of Valve
Angle of Valve
spark
Plug
drilling
Exhaust Port
Inlet Port
a)
b)
Distance / Fuel consumption
(Km/Litre)
Figure 2-7 a) Cross-section of a cylinder head b) engine block, (reprinted with permission from Taylor &
Francis)55
26.50
21.25
17
12.75
8.50
4.25
0
54
908
1362
1816
2770
Vehicle mass (Kg)
Figure 2-8 The relation between vehicle mass and fuel consumption68, 69
Using materials with higher strength and stiffness, such as compacted graphite iron (CGI)
instead of grey cast iron, contributes to increase in power and decrease in size of an engine
by reducing the main bearing thickness (see Table 2-1).65 Another alternative is to replace
cast iron with lightweight materials (e.g. aluminium and magnesium alloys). Owing to the
considerable difference of the density between cast iron (~7.5 g cm−3) and aluminium (~2.7
g cm−3) and magnesium (~1.74 g cm−3) alloys, the substitution of cast iron by one of these
alloys could make a significant weight reduction.
17
Table 2-1: Weight reduction results for CGI vs. grey cast iron cylinder blocks 65
Engine size
(Litres)
1.6
1.8
2.5
4.6
9.2
Engine type
I-4 Petrol
I-4 Diesel
V-6 (Racing)
V-6 Petrol
V-6 Diesel
Grey iron
weight (kg)
35.4
38.0
56.5
72.7
158
CGI weight
(kg)
25.0
29.5
45
59.6
140
Weight reduction (%)
29.4
22.4
20.4
18.0
11.4
2.2.2. Magnesium alloys
Magnesium is ~75% and ~33%, respectively, lighter than iron and aluminium. It has
attracted great interest in the automotive industry. However, the specific stiffness of
aluminium and iron has been reported to be slightly (~0.69% and 3.752%, respectively)
higher than that of Mg; but the specific strength of Mg is significantly greater than that of
aluminium and iron (14.075% and 67.716% for aluminium and iron, respectively).68, 71
The regular commercial cast Mg alloys (e.g. AZ91 and AM50), which are widely used in
the automotive industry, suffer from poor creep resistance.72 The creep resistance of the
magnesium alloys (e.g. AM50: Mg–5Al–0.3Mn–0.2Zn (approximate wt-%ii)) has been
reported to be ~15% less than that of aluminium alloys (e.g. A380: Al–8.5Si–3.5Cu–3Zn
(approximate wt-%)) at 293 K, and ~65% less at 403 K.73 Therefore, new Mg alloys (e.g.
MRI 201, MRI 230) have been developed to improve the creep resistance and hightemperature strength. These alloys could compete with the commercial Al alloys (e.g. A380
and A319) in terms of creep resistance and high-temperature strength.74-76
Despite these advantages, application of magnesium alloys in the automotive industry has
been very limited: the average application of Al alloys has been reported to be over 100 kg
per car, while that of Mg alloys has been reported as ~6 kg.77 The higher total cost of Mg
alloys is one of the major reasons for impeding their widespread application in the
automotive industry.77-79 It is worth noting that the price of magnesium has been
considerably reduced in the last few years.71 Lower thermal conductivity and higher
thermal expansion are the other disadvantages of Mg alloys compared with Al alloys.72
ii
All chemical compositions are given in weight percent (wt-%) hereafter, unless otherwise stated.
18
2.2.3. Aluminium alloys
In the late 1970s, the generation of aluminium engine blocks was introduced to be used in
gasoline engines. However, because of technical requirements, application of aluminium
alloys was very limited in diesel engines until the mid-1990s. Nowadays, blocks for
gasoline engines are generally cast in aluminium alloys; and the use of aluminium in diesel
engines is continuing to increase. Also, most cylinder heads are cast in aluminium alloys.
Substitution of cast iron by aluminium in engine blocks could result in a weight reduction
of 15–35 kg.80 Inline cylinder blocks made in aluminium are noticeably lighter than
corresponding cylinder blocks produced with CGI. For an engine weighing 35 kg in CGI,
the weight of the inline cylinder block should be 28 kg using an aluminium alloy.65
However, if the design of the engine is adapted to CGI (V-8 instead of inline), a marginal
weight saving can be made with CGI.
Some important properties of Al alloys, Mg alloy, grey cast iron (GJL-250) and CGI
(CGV-400) are compared in Figure 2-9. As shown in this figure, another advantage of
aluminium alloys compared to cast iron is their excellent thermal conductivity, which
accelerates cooling of engine. In spite of all these advantages, softening of the commercial
foundry aluminium alloys at service temperature restricts their application in engine
components. For instance, as shown in Figure 2-10, some studies from AVL reported that
the application of aluminium engine blocks must be restricted for those passenger car
engines with 150 bar peak firing pressure.81, 82
19
Normalized Value (%)
450
CGI (GJV-400)
400
Cast Iron (GJL-250)
350
Aluminum-A390
300
Aluminum (AlSi9Cu)
Mg-MRI 230D
250
200
150
100
50
0
Density
Young Modulus
UTS
Thermal Expansion
Thermal Conductivity
peak firing pressure (bar)
Figure 2-9 Material properties of compacted graphite iron (CGI-400), grey cast iron (GJL-250), Al-A390,
AlSi9Cu and Mg-MRI 230D 65, 74, 81.
200
V-Engines
In Line Engines
100
0
Figure 2-10 Peak firing pressure limits for various materials in diesel engine cylinder block58, 81
Table 2-2 presents the chemical compositions of the most common aluminium alloys used
in engine applications. Alloys 356+Cu and 319 have been extensively studied for use in
engine components, in particular in cylinder heads. For instance, they were studied by
BMW,28,
83
VAW Aluminium AG,84 Ford Motor Company85 and General Motors.86,
87
Considering their importance, special emphasis will therefore be given to the 356- and 319type alloys in the following sections. Hypereutectic Al–Si alloys could be another
alternative for cast iron in production of engine blocks. Jorstad,88 who is often credited as
the pioneer of 390 hypereutectic Al–Si alloys, has thoroughly reviewed the application of
these alloys in the manufacture of engine block from inception until now. Mercedes, BMW,
Porsche, Audi and Volkswagen are some of the companies which have used hypereutectic
Al–Si alloys in the production of engine blocks.
20
Table 2-2: Chemical composition (wt-%) of 356-type, 319-type and 390-type Al alloys
Composition
356.0
A356.2
356+Cu
319.0
A319.1
B319.1
390
A390
B390
Si
Cu
Mg
Fe
Mn
Ti
Zn
Ni
Al
Ref.
6.5-7.5
6.8
7.1
5.5-6.5
5.5-6.5
5.5-6.5
16-18
16-18
16-18
< 0.25
0.04
0.5
3.0-4.0
3.0-4.0
3.0-4.0
4-5
4-5
4-5
0.25-0.45
0.35
0.36
< 0.1
0.1
0.1-0.5
0.45-0.65
0.45-0.65
0.45-0.65
<0.6
0.08
0.12
<1
0.8
0.9
< 1.3
< 0.5
< 1.3
<0.35
0
0.05
<0.5
0.5
0.8
< 0.1
< 0.1
< 0.5
<0.25
0.15
0
<0.25
0.25
0.25
<0.2
<0.2
<0.2
<0.1
0.01
0
<1
<1
1
<0.1
<0.1
<1.5
0
0
0
<0.35
0.35
0.50
-
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
7, 89
7, 89
47, 85
7, 89
7, 89
7, 89
88
88
88
Table 2-3 presents some major mechanical/physical properties of three Al–Si (319-, 356and 390-type) alloys. The symbols F (as cast, without heat treatment), T4 (quenched and
naturally aged), T (artificially aged after casting), T (quenched and artificially aged for
maximal strength) and T7 (quenched and overaged), which represent the most common heat
treatment condition of Al–Si alloys, have been designated by the Aluminium Association of
the USA.90
131
186
--186
165
--414
359
----90
76
--117
100
Shrinkage
tendency* (h)
74
---72
72
82
----
2
2
2
1
1
1
3
3
3
Fluidity*
(g)
85
95
--80
70
110
110
145
120
Resistance to
hot cracking*
(f)
2.5
3
5
2
5
6
<1
<1
<1
<1
Modulus of
elasticity
KPa*106 (e)
131
186
124
138
186
165
200
200
310
262
Endurance
limit (MPa) (d)
Hardness BHN
(c)
234
276
179
186
262
221
200
200
310
262
Compressive
Yield (MPa)
(b)
Elongation %
(in 50 mm)
390.0
Tensile Yield
(MPa) (b)
356.0
F
T6
F
T5
T6
T7
F
T5
T6
T7
Ultimate
tensile
strength- UTS
(MPa)
319.0
Temper
Alloy Number
Table 2-3: Some major properties of the Al 319-, 356-, and 390- type alloys91
(a) These nominal properties are useful for comparing alloys, but they should not be used for design purposes. (b) Offset: 0.2%. (c) 500-kg load
on l0-mm ball. (d) Endurance limits based on 500 million cycles of completely reversed stresses using rotating beam-type machine and specimen.
(e) Average of tension and compression moduli. (f) Ability of alloy to withstand stresses from contraction while cooling through hot-short or
brittle temperature range. (g) Ability of molten alloy to flow readily in mould and fill thin sections. (h) Decrease in volume accompanying
freezing of alloy and measure of amount of compensating feed metal required in form of risers. (*) For ratings of characteristics, 1 is the best and
3 is the poorest of the alloys listed.
The 356-type aluminium alloys present good combinations of strength and ductility, but
their strength reduces rapidly above 473 K (200°C). The 319-type aluminium alloys present
relatively higher yield and creep strength at elevated temperatures (~523 K), although
prolonged exposure at such temperatures could result in softening. Therefore, to achieve the
increasingly exacting requirements of engine components (higher pressure and
temperature) without new material inventions, the existing capabilities of Al–Si
21
hypoeutectic alloys have to be improved by optimisation of either production process (e.g.
casting and heat treatment) or chemical composition.
Description of Al–Si based alloys
2.3.1. The binary Al–Si system
The phase diagram of the Al–Si system is illustrated in Figure 2-11. There is a eutectic
reaction at 850.75 K (577.6
) and 12.6 wt-% silicon, where the liquid phase is in
equilibrium with the α-Al solid solution phase and nearly pure Si (L → α-Al + Si).92, 93 The
maximum solubility of silicon in aluminium is ~1.5 at.-% at the eutectic temperature and
decreases down to ~0.05 at.-% at 573 K (300
). Generally, the morphology of the eutectic
microconstituent tends to be fibrous if the volume fraction of the minor phase is less than
25%. However, in Al–Si binary alloys, the typical Al–Si eutectic morphology is usually
lamellar. This could be ascribed to the low interfacial energy between Al and Si and the
strong growth anisotropy of silicon.93
Silicon (wt.%)
1500
0
10 20 30 40
70
80
90 100
1414°C
1300
Temperature (°C)
50 60
Liq.
1100
Liq.+Si
900
Liq.+Al
700
577.6°C
12.6
500
(Al)
(Si)
(Al)+(Si)
300
0
Al
10
20 30 40
50 60 70
Silicon (at.%)
80 90 100
Si
Figure 2-11 The equilibrium phases diagram of the Al–Si alloy system92, 93
The morphology of the eutectic silicon particles (i.e. particle size and shape) can
appreciably affect the mechanical properties of Al–Si alloys. The coarse lamellar silicon
particles, which appear under normal solidification conditions, may act as stress
concentration sites and crack propagation paths.87,
94, 95
alleviated by imposing higher solidification rates,96,
97
This negative effect can be
carrying out solution heat
treatment41, 98 or by alloying with certain elements (e.g. Sr, Na, etc.), which can change the
morphology of Si particles from plate-like form to fine fibrous form.99,
22
100
During the
solution heat treatment, the unmodified Si particles undergo: (a) necking at several places
along the length of the Si particles resulting in their fragmentation, (b) gradual
spheroidisation of the fragments and (c) coarsening by the Ostwald ripening process.41, 98
There are numerous elements which can modify the Al–Si eutectic microstructure, such as
Sr,99, 101 Na,102 Ca,103, 104 Sb,105 Sc106, 107 and several rare earth metals.108 It was proposed
that the modifier agent is adsorbed at the silicon/liquid interface and results in the growth of
twins and branching of silicon particles.100,
109, 110
Such modifications could reduce the
solution treatment time and improve the overall mechanical properties.52, 111 Nevertheless,
some studies112,
113
have shown that the addition of the modifier elements is often
associated with increased porosity. Gruzleski114 and Lados115 stated that chemical
modification by Sb and Sr did not have a considerable impact on fatigue lifetime of
AlSiMg alloys; meanwhile Gundlach et al.41 reported the beneficial effect of eutectic Si
modification on thermal fatigue resistance. Therefore, an optimum content of the modifier
agent is required to yield an acceptable level of modification without affecting the porosity
level. The optimum content can be varied depending on the constituents of each alloy. For
instance, the modifying effect of Sr can be somewhat nullified by the presence of other
elements, namely P, Bi, Sb116 and Mg.87,
116
The reader wanting more details on the
modification of Al–Si casting alloys may consult various publications.97, 116, 117
Silicon significantly improves castability (fluidity, metal-feeding)118,
119
and wear
resistance120 and contributes to reduce the density and the coefficient of thermal expansion
of aluminium alloys.118 In addition, dissolution of Si in α-Al matrix (e.g. ~0.7 wt-% at 773
K (500
)) can significantly improve the age hardenability of AlSiCuMg alloy by
combining with Mg.121
2.3.2. Influence of iron as impurity
Al–Si binary alloys, even prepared from pure materials (~99.99%), can contain more than
50 ppm of iron. The presence of iron can considerably affect the solidification process of
Al–Si alloys.92 Iron, as the most common impurity in Al–Si alloys, strongly reduces the
fluidity and the overall mechanical properties through the formation of brittle intermetallic
phases. Primary Al–Si alloys typically contain between 0.05 and 0.2 wt-% Fe; but, in
secondary Al–Si alloys, it can reach up to 1 wt-%. Economically, there is no known way to
23
further reduce Fe from primary Al–Si alloys. Owing to a relatively high solubility of Fe in
liquid Al, it can readily enter into the melt from unprotected steel tools, furnace equipment
and addition of low-purity alloying materials.122 The amount of Fe exceeding the solid
solubility limit appears in the form of iron-bearing intermetallic phases such as β-AlSiFe,
α-AlSiFe and π-Al FeMg Si . The α-AlSiFe phase, which appears in the form of Chinese
script particles, has the composition of Al Fe Si (~31.6% Fe, ~7.8% Si). The stoichiometry
of the β-AlFeSi phase is Al5FeSi (~25.6% Fe, ~12.8% Si), with a probable range of 25–
30% Fe and 12–15% Si. The β-Al5FeSi phase has a platelet morphology (in three
dimensions) which appears as a needle in micrographs.123, 124
Many studies122, 125, 126 found that as Fe levels increase, the ductility and tensile strength of
Al–Si alloys strongly decrease; however, the yield strength remains in general almost
unaffected by iron. The iron-bearing compounds are much more easily fractured under
tensile load compared to the Al matrix or the modified silicon particles. Their detrimental
effect is directly proportional to the morphology, size and volume fraction. The platelet
morphology of β-Al FeSi phase explains why it is the most deleterious intermetallic phase
in cast Al–Si alloys.127, 128
The size and density of iron-bearing compounds (particularly β-phase) increase with iron
content. Moreover, intermetallic phases that can form prior to (or with) the solidification of
the aluminium dendrite network (pre-dendritic particles) are much larger than those that
form during or after the period of Al–Si eutectic solidification.129 More available time for
growth at a slower solidification rate also leads to enlarged intermetallic particles.122
Furthermore, it has been reported that the amount and size of porosity in the microstructure
are strongly enhanced by increasing Fe content. This behaviour is mainly related to the
increased amount of β-phase, since it promotes shrinkage porosity during solidification by
physically blocking the metal feeding, as shown in Figure 2-12.
The β-platelets are much more susceptible to crack linkage and fracture than the α-iron
Chinese script particles, so the formation of the α-iron phase instead of the β-phase can be
less detrimental to mechanical properties owing to its compact morphology. According to
Mondolfo,123 low Mn and Cr concentration and a low cooling rate (~0.8 K s−1) are the main
factors that favour the crystallisation of Al FeSi phase. Hence, chemical modification (by
24
Mn, Cr and Ni addition), high solidification rate128, 130-132 and superheating of the melt133
contribute to the formation of the α-iron phase. The amount of Mn needed to convert all of
the β phase is not yet well known. Several researchers9, 134 reported that a Mn/Fe ratio of
0.5 seems to be sufficient for complete substitution of Al5FeSi by α-Al15(Fe,Mn)3Si2 phase.
However, other researchers122, 135 stated that even at these levels of Mn addition some βphase could still form. It should be noted that an undesired amount of Mn in AlSiCu/Mg
alloy could lead to the precipitation of Al–Cu–Mn particles (T-Al20Cu2Mn3 phase136)
during solution treatment, which in turn decreases the Cu content in -Al matrix.121 Kim et
al.130, 132 reported that the combined addition of Mn and Cr to modify β-phase could be
more effective which considerably improved tensile properties (ultimate tensile strength
(UTS) and elongation); the improved mechanical properties were attributed to the
precipitation of α-Al(Mn,Cr,Fe)Si nanoparticles in the microstructure of A356 Al alloy.
β phase
Figure 2-12 The role of Al5FeSi in the formation of shrinkage porosity, (reprinted with permission from
Taylor and France)137
Solidification sequence in 356 and 319 Al alloys
2.4.1. 356-type Al alloys
Backerud et al.7 studied the solidification sequence in various Al alloys using a thermal
analysis technique, followed by a subsequent metallographic examination of specimens.
Their results on solidification of A356.2 alloy with a cooling rate of 0.7 K s−1 are
summarised in Table 2-4. Reactions (2b) and (3b) were not observed by Arnberg et al.138
25
and Mackay et al.139 in their investigation of almost the same chemical composition. They
stated that no pre-eutectic (Al5FeSi) phase could be crystallised with such low Fe contents,
although their specimens contained 0.08% Fe as did those of Backerud et al.7 Backerud et
al.7 stated that the Fe is strongly partitioned in the liquid phase which results in
precipitation of the pre- or co-eutectic Al5FeSi phase. Subsequently, the Al5FeSi phase is
partly transformed into the Al8FeMg3Si6 phase through a quasi-peritectic reaction (3b).
Wang et al.140 confirmed the Backerud et al.7 results on solidification sequence by scanning
electron microscopy (SEM) analysis. As illustrated in Figure 2-13, the -Al8FeMg3Si6
phase was directly grown from the Al5FeSi phase, which could imply the occurrence of
reactions (3a) and (3b).
No.
1
2
3
4
5
Table 2-4: Reactions occurred during solidification of A356.2 7
Reaction
Temp., with 0.7K/s
Development of dendritic network
888-883 K (615- 610
a) Liq → Al + Si
883-835 K (610- 562
b) Liq → Al + Al FeSi
837-831 K (564- 558
a) Liq → Al + Si + Al FeSi
b) Liq + Al FeSi → Al + Si + Al FeMg Si
831-822 K (558- 549
Liq → Al + Mg Si + Si
819-814 K (546- 541
Liq → Al + Si + Mg Si + Al FeMg Si
)
)
)
)
)
β
π
Si
10 μm
Figure 2-13 SEM micrograph of A356 as-cast Al alloy showing the close association between Al5FeSi and πAl8FeMg3Si6 phase, (reprinted with permission from Springer)140
2.4.2. 319-type Al alloys
The solidification sequences of two 319-type aluminium alloys with chemical compositions
of (Al–5.7Si–3.4Cu–0.62Fe–0.36Mn–0.10Mg (wt-%))7 and (Al–6.23Si–3.8Cu–0.46Fe–
0.14Mn–0.06Mg (wt-%))9 are listed in Table 2-5. The precipitation of Al15Mn3Si2 (possibly
together with Al5FeSi) which was observed by Backerud et al.7 was not detected by Samuel
et al.9 This is presumably because of the smaller Mn content of the alloy studied by the
26
latter authors. The presence of Mg (even in the small amount of ~0.06 wt-%) leads to the
transformation of the Al5FeSi phase to -Al8FeMg3Si6 phase as well as precipitation of
Mg2Si phase during solidification, attributed to reaction (C) in Table 2-5.9, 141 Furthermore,
precipitation of Q-Al5Cu2Mg8Si6 phase, corresponding to reaction (E), is caused by the
addition of Mg.9, 142 The Q-Al5Cu2Mg8Si6 phase grows out of -Al2Cu particles during the
complex eutectic reaction in the final stages of solidification.87, 143 The morphology of the
-Al2Cu phase, which can be blocky or eutectic form, strongly depends on solidification
rate and Sr modification. It has been reported that high solidification rate leads to fine
eutectic Al–Al2Cu phases,9,
144
while Sr modification increases the proportion of blocky
Al2Cu phase.145-147
Table 2-5: Summary of reactions occurring during solidification of 319.1Al alloys
Samuel et al. 9
Temp.
1
609
A Development of α-Al dendrite network
Formation of α-Al dendrite network
2 a) L → (Al) + Al Mn Si
590
Bäckerud et al. 7
Temp.
608
b) L → (Al)+ Al15Mn3Si2+(Al5Fesi)
3
4
5
L → (Al) + Si+ Al FeSi
L → (Al) + Al Cu + Si + Al FeSi
L → (Al) +Al Cu +Si+Al Cu Mg Si
575
525
507
B
C
D
E
Precipitation of eutectic Si
Precipitation of Al Mg FeSi + Mg Si
Precipitation of Al Cu
Precipitation of Al Cu Mg Si
557
544
505
496
Effect of microstructural features on TMF strength
It is largely accepted that fatigue lifetime of Al–Si based alloys (319- and 356-type Al
alloys) is more affected by the actual casting processes than by alloy chemistry. Crack
initiation can be greatly delayed in defect-limited specimens.27,
38, 148
Porosity and oxide
inclusions are the most deleterious metallurgical defects associated with casting processes
and both strongly impair the fatigue strength. There is a critical size of the pores and
inclusions below which the impact of these defects is not the root cause of fracture, and
cracks can be initiated by other microstructural features like large eutectic constituents
(fractured/detached Si particles) or persistent slip bands.83, 149-151
The transition from one mode of failure to another is of importance in predicting the service
lifetime of engineering components. For instance, transition from transgranular to
intergranular fracture is usually followed by a dramatic reduction in ductility and fatigue
lifetime. Creep damage, which is in the form of intergranular cracking, is generally
27
observed in in-phase TMF test specimens. No detectable intergranular damage in
isothermal and out-of-phase TMF tests was reported by a majority of researchers.85, 152, 153
Therefore, to develop a new alloy, the creep/fatigue failure mechanisms have to be clarified
in terms of intrinsic material properties and microstructure.
2.5.1. Porosity
The combined effect of volumetric shrinkage and dissolved gas leads to the formation of
porosity.125,
154
In alloys with low fluidity, the shrinkage of the melt between dendrites
cannot be fully filled by the liquid phase remaining, which leads to porosity being spread
out along these dendrites. The only gas which is sufficiently soluble in aluminium alloys
and leading to porosity is hydrogen.155,
156
The solubility of hydrogen decreases with
decreasing temperature and hydrogen atoms precipitate and form molecular hydrogen
during solidification.
Porosity formation in Al–Si hypoeutectic alloys can be affected by alloying elements via a
few mechanisms. Addition of Cu to Al–Si alloys assists porosity formation by increasing
both the solidification range and the solidification shrinkage.3,
4, 157
The overall
solidification shrinkage in Al–Cu binary alloys is ~8.4% while it is ~4.5% for Al–7% Si.3-5
Moreover, increasing the copper content enhances the activity coefficient of hydrogen
which, in turn, decreases the solubility of hydrogen. Therefore, the alloys containing copper
can be more prone to form porosity during solidification.158 Caceres4,
157
stated that ‘the
addition of only 1% Cu causes the development of a significant level of porosity in
comparison with the Cu-free A356.2 alloy, while increasing the levels of Cu beyond 1%
and up to about 4% results in a relatively small increase in porosity level’. The iron-bearing
platelets (e.g. β-AlSiFe phase) reduce permeability and restrict the flow of liquid metal at
the latter stage of the solidification process,159 which was elaborated in the “Influence of
iron as impurity” section. Grain refinement obtained by alloying elements such as titanium
and boron reduces the volume fraction and size of porosity.159, 160 It is worth pointing out
that Mg4, 159 and Si3, 4 can have a positive impact in reducing both pore size and density.
Tensile and fatigue properties are made significantly poorer by increasing porosity.51, 57, 161
Surappa et al.162 found that the decrease in the elongation to fracture could be correlated to
the pores on the fracture surface. Ma163 showed that increasing metal soundness, in terms of
28
porosity, resulted in a higher elongation to fracture in alloys A319 and A356. The effect of
porosity on fatigue strength is strongly dependent on a number of factors, such as
morphology, size and position of the pores within the cast part. Skallerud et al.164 reported
that a shrinkage pore could be more deleterious than a gas pore. Fatigue cracks are
generally initiated from shrinkage pores at or near the free surface of a specimen. The effect
of large pores far away from the free surface of specimens on the fatigue lifetime can be
very small, while even a small pore (or inclusion) located near the free surface can be very
deleterious to fatigue lifetime.51, 165, 166
2.5.2. Secondary dendrite arm spacing
In an alloy microstructure, the SDAS generally characterises the solidification rate.
Increasing the solidification rate substantially improves the fatigue and tensile properties
(except modulus).49 This improvement is generally attributed to the influence of
solidification rate on the number density and size of porosity, and to the refinement of
grains and secondary phase microconstituents.38, 167 Several authors27, 38, 148 reported that
reducing SDAS strongly decreased both the number density and the average pore size in
Al–Si alloy castings. Chen et al. stated that ‘in A356.2 Al alloy as the SDAS increases from
15 m to 50 m, the fatigue lifetime decreases about three times under LCF and over six
times under HCF’, since for the alloy with SDAS greater than ~30 μm, pores act as fatigue
crack initiation sites.168,
169
Furthermore, the content of β-Al5FeSi as the least desirable
secondary phase was significantly reduced by the increasing cooling rate (see Figure
2-14).38 So, one can say that the influence of SDAS cannot be separated from the influence
that solidification rate has on the size and distribution of all microconstituents.
β Al5FeSi content (%)
3.0
2.5
2.0
1.5
1.0
0.5
0.00
20
40
60
SDAS (μm)
80
100
Figure 2-14 Effect of cooling rate on the formation of β-Al FeSi brittle phase38
29
2.5.3. Segregation
Segregation is another important phenomenon which can considerably affect fatigue
lifetime. In casting, heat is transferred through the mould walls and this causes higher
volume fraction of the α-Al phase to be located in the outer surface of the casting and a
larger volume fraction of eutectic phases and shrinkage porosity to be located in the centre.
Consequently, local fatigue resistance could vary with the location within a casting. Seniw
et al.165,
170
reported interesting results on the effect of segregation of Si on fatigue
properties of A356 cast alloy. They revealed that specimens taken from the outer surface of
a cast bar, which was the first zone to be solidified, survived 106 cycles without failure,
while specimens taken from the part to solidify last failed after only 150 000 cycles. This
illustrates how fatigue lifetime can be reduced down the solidification path.
2.5.4. Cracking/debonding of Si particles
Crack propagation in Al–Si based alloys depends on the size, orientation and local
distribution of the Si particles.115,
171, 172
In modified Al–Si alloys (fine Si particles, size
~1.5–2.5 μm), fatigue cracking progresses by decohesion of the Si particles from the Al
matrix. But, with increasing Si particle size, the tendency to particle cracking increases,
such that in unmodified alloys (coarse Si particles, ~3–9 μm) particle cleavage is the
dominant feature.115, 173, 174 Figure 2-15 illustrates the debonding of a Si particle from the Al
matrix and a fractured Si particle caused during a fatigue test. Plastic deformation in TMF
loading can cause debonding of Si particles.83,
95, 175
This is a result of significant
thermal/mechanical misfit between the brittle Si particles and the surrounding ductile
matrix, which leads to separation during thermal/mechanical loading.83, 150
30
Fractured Si particle
a)
2µm
b)
Al-1% Si Matrix
2µm
5000X
Figure 2-15 SEM images of: a) debonded (reprinted with permission from Elsevier)150, and b) fractured Si
particle (reprinted with permission from Springer)94
2.5.5. Slip bands
Several researchers176-179 reported that in the absence of casting defects (e.g. porosity) or in
castings with small SDAS,180 cracks initiated from persistent slip bands on the surface.
Nyahumwa et al.178 observed a faceted transgranular appearance on the fatigue fracture
surface of hot isostatic pressed A356.2 aluminium alloy. They reported that the faceted
transgranular fracture mode of some specimens was by the slip mechanism. Jiang and coworkers181 observed slip band cracking only in naturally aged and underaged samples. Zhu
et al.177 also reported that twin boundary initiated failures in 319 Al alloy could occur only
at elevated temperature. Jang et al.176 hypothesised that, with increasing temperature, the
critical effective stress for fatigue crack initiation (CESFCI) value at slip band would be
comparable to the CESFCI value at porosity, while it is lower at porosity than at slip band at
room temperature. Moreover, Jang et al.176 reported interesting results on TMF crack
initiation in cast 319-T7 aluminium alloy. The crack initiation of 11 specimens (out of 29
specimens) occurred at near surface porosity, but, for those specimens with relatively small
porosity near the surface, coarse transgranular facets were observed at the crack initiation
site. They proposed that the slip band mechanism was responsible for crack initiation.
Owing to the presence of oxide films in these transgranular facet areas, the authors176, 182
concluded that these oxide films were formed as a result of fretting damage under fatigue
cyclic loading, rather than pre-existing oxide films. However, Campbell183,
184
criticised
31
their idea and proposed that the oxide film on the fatigue initiation site was created as an
inclusion during the solidification, and was a prerequisite for slip band crack initiation.
Gundlach et al.41 investigated TMF of 319 and 356 Al alloys and reported the occurrence
of stress relaxation in 356 Al alloy on heating above ~505 K (232
stated that stress relaxation started at ~493 K (220
). Takahashi et al.44
) in the TMF process of Al–6Si–
2.5Cu–0.3Mg (wt-%) alloy. At this temperature, which is ~0.56Tmiii, diffusion creep and
dislocation creep can occur;169 therefore, they concluded that these creep micromechanisms could be responsible for softening of the alloys.44 Angeloni185 also reported
that the aforementioned creep micro-mechanisms could be responsible for plastic
deformation of Al–9Si–3Cu–0.3Mg (wt-%) alloy in elevated temperature fatigue tests
(~553 K). Strengthening of cast aluminium alloys
The principal objective in the design of aluminium alloys is to improve their tensile
strength, hardness, creep resistance and fatigue resistance. The strengthening of cast
aluminium alloys relies on several different mechanisms based on restricting/hindering the
motion of dislocations. The two major methods used to strengthen cast Al alloys are
precipitation hardening and dispersoid hardening; the latter refers to precipitates formed
with transition elements and stable at higher temperatures. The works dedicated to applying
and optimising these methods will be described in this section.
2.6.1. Heat treatment of AlSiCuMg alloys
The common thermal treatments, which are generally applied for AlSiCuMg cast alloys,
involve either age hardening of the as cast alloy (T5 type) or solution treatment followed by
age hardening (T6, T7 type).12, 118 If peak mechanical properties are not required, castings
with sufficiently high cooling rates and artificially aged (T5 type) may meet the intended
strength requirements. This allows a reduction of production costs since the solution heat
treatment is not made. However, T6 (‘peak-aged’) and T7 (‘overaged’) are the most
common heat treatments made on AlSiCuMg alloys. The T6 heat treatment is generally
iii
Tm is the absolute melting point of the aluminium alloy (Tm = 888 K).
32
used for room temperature applications,186, 187 while for high temperature applications, and
especially in the case of 319-type Al alloys, the T7 treatment is recommended.37,
43, 188
These heat treatment processes, which involve the following three consecutive stages, have
to be optimised: (1) solution treatment, (2) quenching and (3) aging.189-191
2.6.1.1.
Solution treatment
The solution heat treatment (SHT) is achieved by heating the alloy at a temperature range
between the solvus and the solidus line (see Figure 2-16). The soaking period must be long
enough to cause one or more constituents to enter into solid solution. Homogenisation of
the alloying elements and spheroidisation of the eutectic Si particles are the other purposes
of the solution treatment.192, 193
Temperature
Liquidus
Solidus
Homogenizing
Solvus
GP (2)
Annealing
GP (1)
Aging
Concentration
Figure 2-16 Temperature ranges for heat treatment and relevant solvus line for binary aluminum alloys 193
The dissolution rate of intermetallic compounds is strongly dependent on the solutionising
temperature (TSHT). Samuel10 has reported that increasing the solutionising temperature
from 753 to 773 K in Al–6.17Si–3.65Cu–0.45Mg (wt-%) alloy improved the yield strength
from 330 to 410 MPa and the UTS from 340 MPa to 420 MPa. On the other hand, the
maximum applicable solution treatment temperature (
of the last solidified phases.10,
194, 195
) is limited by incipient melting
Incipient melting deteriorates the mechanical
properties as a result of void formation.144, 189 According to Samuel,10 the TSHT of a cast Al–
6Si–3Cu (wt-%) alloy containing 0.04% Mg can be ~792 K (519
content to 0.5% restricts the TSHT temperature to ~778 K (505
), but increasing the Mg
) to avoid incipient
melting.10 It has been reported that even a small amount of Mg (0.1 wt-%) can reduce the
solidus temperature of a 319.0-type Al alloy down to 780 K (507
) under non-equilibrium
33
solidification conditions.7,
9
Moreover, Fuoco et al.196 pointed out that the TSHT for
AlSiCuMg alloys must not exceed 773 K (500
) to avoid incipient melting. Therefore, the
melting point of the last solidified phase must be known accurately to optimise the
solutionising temperature. This can be achieved by using a microsegregation model or by
conducting a thermal analysis.
Sokolowski et al.197,
198
reported that single-step SHT of Al–7Si–3.7Cu–0.23Mg (wt-%)
alloy, which must be at less than 768 K (495
), is neither able to maximise the dissolution
of Cu rich phases nor able to homogenise the microstructure and modify the Si particles. As
a result, they proposed a two-step SHT (i.e. 8 h at 768 K (495
) + 2 h at 793 K (520
)).
By doing so, the Cu-containing phase (Q-Al5Cu2Mg8Si6) with the lowest melting point
(Tm~ 780 K (507
))6,
10
would be dissolved at the first step of SHT. The higher
solutionising temperature of the second step could dissolve the remaining Cu-bearing phase
and further homogenise the microstructure.6,
198
Nevertheless, some authors reported the
stability or very slow dissolution rate of Q-Al5Cu2Mg8Si6 phase at ~773 K (500°C)199, 200
when the magnesium content is sufficiently high. The holding time period of the first step
and the TSHT of the second step are very critical parameters to avoid incipient melting.189,
201, 202
Therefore, to achieve an effective dissolution while avoiding coarsening of the
constituents, the solutionising parameters (namely time and temperature) have to be
optimised.203, 204 In this regard, differential scanning calorimetry (DSC) and electron probe
microanalysis are powerful tools, which are discussed in more details below.
Wang et al.6 used DSC analysis to optimise the SHT of Al–11Si–4Cu–0.3Mg (wt-%) alloy.
Figure 2-17 displays the DSC curves of the alloy for different solution times at 773 K
(500 ). Peaks (1), (2) and (3) correspond to the following reactions:
Reaction of peak (1):
α (Al) + Al2Cu + Si + Al5Cu2Mg8Si6 → Liquid
Reaction of peak (2):
α (Al) + Al2Cu + Si → Liquid
Reaction of peak (3):
α (Al) + Si (+ Al5FeSi+ …) → Liquid
As illustrated in this figure, with increasing solution heat treatment time (tSHT), the height of
peaks (1) and (2) gradually decreased. After 10 hours SHT, peak (1) completely
disappeared, which indicates the complete dissolution of eutectic phases (α-Al+ Al2Cu+
Si+ Al5Cu2Mg8Si6). Therefore, the temperature at the second step of solution treatment
34
could be increased up to the onset temperature of peak (2) (~793 K (520
)) to quickly
dissolve the remaining Cu-rich intermetallics. The temperature of the second solution
treatment step should be lower than 793 K (520
) in order to avoid incipient melting of
(α-Al + Al2Cu + Si) eutectic phase.6
Temperature (K)
793
813
833
853
873
580
600
5
773
(3)
4
500 ºC 6h
3
500 ºC 8h
2
500 ºC 10h
(2)
1
(1)
0
Heat flow (mw/mg) endo
As-cast
500
520
540
560
Temperature (ºC)
Figure 2-17 DSC curves of AlSiCuMg alloy solution treated at 773K for different times (reprinted with
permission from Elsevier)6
Dissolution of Cu phases (e.g. Al2Cu/Al5Cu2Mg8Si6) which increases the Cu content in the
α-Al matrix is one of the major purposes of the SHT. In order to determine the efficiency of
a specific SHT, Sjolander et al.192 and Han et al.199,
205
proposed to measure the Cu
distribution in the -Al matrix by means of line scans in electron probe microanalysis
(EPMA). For instance, the solutionising time of Al–8Si–3Cu (wt-%) alloy with different
SDAS (10, 25, 50 μm) was studied by Sjolander et al.192 Figure 2-18 illustrates the
concentration of Cu in the dendrite arms for various specimens with different solution time
(0, 10, 60, 180, 360, 600 min). Homogenisation in the dendrite arms occurred very fast
(within 10 and 60 min), but the concentration of Cu was strongly dependent on the
microstructure and solutionising time. For the finest microstructure (SDAS of 10 m), 10
min of solutionising time seemed to be enough; but for the very coarse microstructure
(SDAS of 50 m), even 10 h of solutionising time at 768 K (495
) was not sufficient.192
35
b) 4
c) 4
3
3
3
2
wt% Cu
4
wt% Cu
wt% Cu
a)
2
1
1
0
-20
0
0
20
Distance from dendrite center (μm)
Solutionizing time
(min.)
SHT-600
SHT-360
SHT-180
SHT-60
SHT-30
SHT-10
As-Cast
2
1
-10
-5
0
5
10
Distance from dendrite center (μm)
0
-4
-2
0
2
4
Distance from dendrite center (μm)
Figure 2-18 Cu concentration measured across dendrite arms in different solutionising times at 768 K (495 )
for various samples: (a) SDAS 50 μm, (b) SDAS 25 μm and (c) SDAS 10 μm, (reprinted with permission
from Elsevier)192
2.6.1.2.
Quenching
The purpose of quenching is to maintain the solid solution by cooling rapidly to a low
temperature in order to prevent the diffusion of the elements. As a result, solute atoms, as
well as a significant fraction of thermal vacancies, are effectively frozen inside the material.
This causes the concentration of solute atoms to be greater than the equilibrium level and a
thermodynamically unstable supersaturated solid solution is created.195, 206, 207
In order to avoid premature precipitation, which could severely deteriorate the mechanical
properties, cooling rate should be fast enough. For aluminium alloys, the usual quenching
media are both cold (below 303 K (30
100
)) and hot water (between 338 and 373 K (65 and
)). During quenching by cold water, the water temperature should not be increased
by more than 10 K. Furthermore, the transfer time period of specimens from the furnace to
the quench media must be short so as to pass quickly enough through the critical
temperature range where very rapid precipitation can occur.207-209 However, it should be
taken into account that very fast quenching might cause distortion and residual thermal
stresses.209
2.6.1.3.
Aging
During aging of Al alloys, solid solution strengthening gradually disappears and the
coherent structure of Guinier–Preston (GP) zones leads to an intense strain field in the
surrounding area.210, 211 The mechanisms contributing to increase the yield strength by the
motion of dislocations through precipitates may include chemical, stacking fault, modulus,
coherency and order strengthening.212, 213 These mechanisms were thoroughly reviewed by
36
Ardell.214 Aging is performed by holding the supersaturated solid solution at temperatures
below the solvus line to form a fine distribution of precipitates from a supersaturated solid
solution (see Figure 2-16). The thermodynamically unstable supersaturated solid solution
will reach equilibrium conditions by aging at room temperature (natural aging) or with a
precipitation heat treatment (artificial aging). Time and temperature are the two main
parameters of aging which affect the strengthening mechanisms. Higher aging temperature
accelerates the aging process by increasing nucleation and growth rates.189, 213, 215
Several investigations have been carried out to understand the effect of underaging, peak
aging and overaging on: hardness,216-218 tensile strength,189,
behaviour,181,
219
207, 217
crack propagation
TMF behaviour220 and cyclic stress–strain response of AlSi(Cu,Mg)
alloys.221 The sequence of precipitation of -Al2Cu begins by GP zones, which are
thermodynamically the least stable but kinetically the most favoured phases: -Al → GP
zones (plate-like) →  (plate-like) →  (plate-like) →  (Al2Cu). GP zones and  are
fully coherent with the α-Al matrix,  particles can be either coherent or semi-coherent,
while  particles are incoherent.222-224 GP zones with 3–5 nm diameters consisting of
localised concentrations of Cu atoms have been observed in specimens aged at 373 K (100
) for 2.5 h.225 The required aging time at 373 K (100
) was reported to be at least 1000
h to obtain a microstructure where plate-shaped Cu-rich particles (GP zones)
predominate.188, 225 However, some authors have stated that GP zones undergo dissolution
at temperatures higher than 373 K (100
403 K (130
(150
);91, 195 the presence of GP zones after aging at
) for 16 h226 and the coexistence of ‘GP zones and ’ after aging at 423 K
) for 3.5 h225 have also been reported. The peak strength is influenced by the
amount, size and site density of  and  phases.211 According to reports,188, 225 the reason
for softening with overaging in 319-type Al alloys can be attributed to the coarsening of the
 phase. The transformation of  to  occurs only when aging at 523 K (or higher) and for
time periods greater than 1000 h.188, 225
Two different combinations of precipitates have been observed in peak-aged condition of
the AlSiCuMg alloy system: (1) precipitation of  (based on Mg2Si) and/or  and (2)
precipitation of Q and/or , where the  phase only appears for a high concentration of
Cu (1 wt-%).195, 227, 228 In several studies,227, 228 no -Al2Cu phase has been reported during
37
artificial aging of AlSiCuMg alloys when the Cu content was less than 1.0 wt-%. Figure
2-19 illustrates the DSC curves of as quenched and aged Al–7Si–3Cu–0.4Mg (wt-%) alloy
with a 10 K min−1 heating rate. Formation and dissolution temperature of GP zones, Q
phase,  phase and  phase were found to be at about 303–493 K (30- 220
(220- 270
), 543–633 K (270- 360
) and 633–733 K (360- 460
), 493–543 K
).121 It is worth
mentioning that the temperature at which a given peak occurs increases with increasing
scan rate.121, 229 An exothermic peak corresponding to GP zone formation was only detected
for the as quenched specimen. In the alloy with some impurities (e.g. 0.6 wt-% Fe and 0.5
wt-% Mn), GP zones could not be detected at all; instead, the precipitation of  phase
appeared at earlier stages.121
Temperature, °C
100
0
300
200
400
500
600
G
+d PZ
iss fo
ol rm
λ'
fo utio atio
rm
n n
.+
di
θ'
ss
fo
ol
rm
ut
.+
io
n
di
ss
ol
ut
io
n
+ θ fo
di rm
ss
ol .
ut
io
n
Heat flow
Exotherm
Al-7Si-3Cu-0.4Mg (wt%) aged at 160°C
Endotherm
As-Quenched
Aged- 10 2min
Aged- 10 3min
Aged- 10 4min
Heating rate: 10°C/min.
300
400
500
600
700
800
Temperature, K
Figure 2-19 DSC curve of Al7Si3Cu0.4Mg alloy, solution-treated 10hrs@773K, water-quenched and aged for
different times at 443K, (reprinted with permission from American Foundry Society)121
The aging time-period to reach peak strength is longer for AlSiCu alloy than for
AlSiCu(Mg) alloy.121 The required time period to obtain peak strength in AlSiCu(Mg) alloy
varies from 30 h up to 120 h and even longer at a lower temperature (433 K (160
230, 231
)).121,
The addition of Mg accelerates and intensifies the precipitation-hardening process of
AlSiCu alloys.121,
215, 230
Kang and co-workers121 reported that not only was the peak
hardness obtained for AlSiCu alloy lower than that obtained for AlSiCuMg alloy, but also
the aging time required to reach peak hardness for the former was ten times longer than for
the latter. On the other hand, Wang et al.227 stated that Cu addition to AlSiMg alloy not
only increases the age hardenability, but also extends the time (from about 700 min to 3000
min) required to reach the peak hardness.
38
The large discrepancy between the thermal expansion coefficients of -Al matrix (23.5 
10−6 K−1) and Si particles (9.6  10−6 K−1) generates a lot of dislocations during quenching
around the Si particles and makes these locations become a preferential site of nucleation
for the  phase.121, 191, 223 On the contrary, the Q phase can nucleate at locations of lower
surface energy since this phase is assumed to have a better coherency (semi-coherency)
with -Al. Therefore, Q can precipitate on a dislocation located anywhere in the matrix
giving a more homogeneous distribution of these precipitates. The lengths of diffusion are
reduced and then less time is required to reach peak hardness.121, 232 This could explain the
higher age hardening rate of AlSiCuMg alloy relative to AlSiCu alloy. Nevertheless, it has
been reported that at elevated temperature the Cu-containing – phase can be much
more stable than the Mg-containing – (Al5Cu2Mg8Si6) and – (Mg2Si) phases.121, 233;
also, S–S (Al2CuMg) phase has been reported to be more stable than – (Mg2Si)
phase.234
In addition to the presence of β-Mg2Si and Q-Al5Cu2Mg8Si6 phases in an aluminium 319
alloy, S-(Al2CuMg) phase was also identified by Medrano et al.191 The authors stated that
β-Mg2Si and S-(Al2CuMg) phases probably precipitated during solidification, and still
remained undissolved after solution treatment. According to Mondolfo,123 in AlCuMg alloy
with the ratio of Cu to Mg between 4:1 and 8:1, the aging agent would be both the Al2Cu
phase and Al2CuMg phase. In the case of AlSiCuMg alloy with high silicon content, the S(Al2CuMg) phase is not usually found, but it can be seen in small amounts owing to
compositional heterogeneities.200 Nevertheless, the presence of S-(Al2CuMg) phase in
AlSiCuMg alloys has been observed by some authors.191, 215, 225 Ma et al.187 pointed out the
presence of Al2Cu and Al2CuMg phase in Al–11Si–2.5Cu–0.4Mg (wt-%) alloy. Reif et
al.231,
235
likewise reported the presence of S-(Al2CuMg) phase with increasing Mg
addition to AlSiCu alloy.
It is worth noting that increasing the Mg level beyond 0.3 wt-% in 319-type Al alloys does
not significantly change the alloy strength,236, 237 but it can considerably reduce the ductility
of the alloys. In 356-type Al alloys, increasing the Mg content up to 0.5 wt-% enhances the
strength, while further increasing Mg content can have a negative effect on the strength of
the alloys.195 Wang et al.149 reported that the fatigue lifetime of A357 alloy (with 0.7 wt-%
39
Mg) was lower than that of A356 (with 0.4 wt-% Mg). In alloys with high Mg content, a
large fraction of the -Al8FeMg3Si6 phase could be formed which is stable during the
solution treatment.195, 237
The small precipitates/zones which are cut by the dislocations in motion lead to a
maximum yield stress once the dislocations pass through them. This causes the local work
hardening to be small and the plastic deformation to be restricted on a few active slip
planes, which would probably be very deleterious to fatigue lifetime.226 On the other hand,
for large particle size/interspacing, bypassing particles by dislocations results in rapid work
hardening and the plastic strains are distributed throughout the specimen. However,
because of weak strengthening of these precipitates, the yield stress is not high enough.
Fine226 stated that ‘the interesting possibility is to have a dispersion of two kinds of second
phase particles, small closely spaced particles to give high yield stress plus large particles
to distribute the plastic deformation throughout the material’.
Therefore regarding the operating condition, the aluminium alloys might be used after peak
strengthening with metastable microstructure (T6) or after overaging with equilibrium
microstructure (T7). For engine components which are exposed to TMF, T7 condition seems
to be more appropriate than T6, since:
(a) T6 condition can cause localised deformation;226
(b) prolonged exposure at service temperature leads to higher thermal growthiv in T6
condition.176,182 The thermal growth of W319 Al alloy was 0.045% and 0.006%,
respectively, in T6 and T7 conditions;188 and
(c) T7 shows more stable microstructure and higher TMF lifetime than T6.43
Dispersion hardening
Trying to improve the elevated temperature strengths of aluminium alloys has involved
continuing efforts for more than three decades.43,
226
Before going further, it could be
worthwhile to consider the reason for successfully engineered Ni-based superalloys being
mechanically stable at high temperatures (exceeding 0.75Tm).238 The interesting mechanical
iv
Dimensional change induced by solid phase transformation.
40
properties of Ni-based superalloys at elevated temperatures can be mainly related to the
presence of very large volume fractions of fine -Ni3(Al,Ti) precipitate with L12 structure,
which is coherent–coplanar and moderately ductile. 80, 226, 239 The term ‘coherent–coplanar’
means the precipitate/matrix interfacial energy is very low and the tendency for
coarsening/coalescence of the precipitate is very small. To develop an effective highstrength high-temperature Al alloy, it can be useful to remember the characteristics of this
precipitate in Ni superalloys.240
Softening of the precipitation hardened Al alloys (e.g. AlSiCu) is the major problem at
elevated temperatures because of the dissolution/coarsening of the metastable precipitates.
A high-strength high-temperature Al based alloy must have a distribution of fine
precipitates/dispersed phases, which must be thermodynamically stable, coherent–coplanar
and ductile.37, 119, 226 A low solid solubility as well as limited diffusivity of the solutes in αAl at the intended service temperature, which is essential to retard volume diffusion,
controls the rate of dissolution and coarsening of the precipitated phases.226,
240, 241
Moreover, the larger the interfacial energy, the higher the driving force for
coarsening/coalescence of the precipitates (Ostwald ripening). Therefore, the required
driving force for coarsening can be very small in the coherent–coplanar precipitates.223
Zedalis242 stated that the coarsening rate of the tetragonal Al3Zr dispersed phase (D023 with
semi-coherent interface) is 16 times higher than that of the cubic modified one (L12 with
coherent interface). Furthermore, the coherency of the precipitate/matrix interface
magnifies the strengthening efficiency of the dispersed phase. Accordingly, precipitated
phases with a similar crystal structure and a low lattice parameter mismatch with the α-Al
solid solution are preferred.226, 240, 242
Among the transition elements, only the first element of the third group (i.e. Sc) exhibits a
high symmetry L12 trialuminide (Al3Sc) structure which is an ordered fcc lattice of the
Cu3Au type of structure.240,
243
Group 4 (Ti, Zr, Hf) and group 5 (V, Nb, Ta) elements
crystallise with the body-centred tetragonal D022 and D023 (Al3M) structures, as shown
graphically in Figure 2-20. The brittle low symmetry tetragonal structure (of Al3M; M = Ti,
Zr, Hf, V, Nb) can be transformed to the cubic structure (L12) by alloying.240 Furthermore,
it has been stated244-247 that the precipitation sequence in the aging treatment of
supersaturated Al–Ti, Al–Zr and Al–Hf solid solutions occurs initially by the formation of
41
a metastable cubic L12 (Al3M) phase. The overall sequence of precipitation in Al–Zr and
Al–V systems has been reported248 to be: (supersaturated solid solution)  (cubic spheres
and rod L12)  (tetragonal plates D023/D022). Long term exposure (hundreds of hours) at
high enough temperatures (450°C) is required to transform these metastable phases to the
equilibrium
tetragonal
(Al3M)
structure.
In
other
words,
these
phases
are
thermodynamically metastable (Gibbs free energies of formation of the tetragonal (D023)
and cubic phase (L12) of Al3Zr are −40.75 kJ mol−1 and −38.35 kJ mol−1, respectively249),
but kinetically stable at elevated temperatures even close to 673 K (400
), because of the
extremely slow diffusion rate of these transition elements in α-Al. Moreover, some alloying
elements can reduce much more the rate of this transformation. Zedalis242 stated that
‘addition of V to Al–Zr alloy led to a reduction of the precipitate-matrix mismatch for both
phases, and also retarded both coarsening as well as the cubic to tetragonal transformation’.
Litynska250 wrote that the addition of 0.2% Zr to Al–1Mg–0.6Si–1Cu–0.4Sc (wt-%)
retarded the coarsening of Al3Sc phase and restricted the size of Al3(Sc,Zr) precipitates to
about 20–40 nm, which were fully coherent with the matrix. The retardation of coarsening
of Al3V phase by Zr addition was also confirmed by Fine et al.248
To have a coherent/coplanar interface, dispersoid phases with small lattice parameter
mismatch are preferred. For the transition elements Hf, Zr, Sc, Nb, Ti, V and Ta the lattice
parameter mismatches between the precipitate (pure binary Al3M (L12) trialuminides) and
α-Al matrix at room temperature are 0.04%, 0.75%, 1.32%, 1.49%, 2.04%, 4.44% and
5.26% respectively.240, 251
Al
M
C
C
L12
D022
D023
Figure 2-20 a) L1 , (b) D0 , and (c) D0 crystal structures, (reprinted with permission from Elsevier)252
42
Diffusivity, D (m2/s )
10-11 Al, 660°C
10-12
10-13
10-14 Al, 400°C
10-15
10-16 Al, 300°C
10-17
10-18
10-19
10-20
10-21
10-22
10-23
10-24
10-25
10-26
10-27
10-28
660°C
400°C
300°C
Sc Ti V Cr Mn Fe Co Ni Cu Zn Ga Ge
Figure 2-21 Calculated diffusivities for different solute elements at 573K (300 ), 673 K (400
660 (T of Al), (reprinted with permission from Carl Hanser Verlag) 240
), and 933 K Because of the low volume fraction of the dispersoid phases in Al alloys, the precipitates
should be very small and resistant to coarsening. Therefore, for an alloy subjected to
prolonged exposure at elevated temperatures, slow diffusion kinetics is required to maintain
strength. Figure 2-21 compares the calculated diffusivities of different solute elements in αAl at three different temperatures (i.e. 573, 673 and 933 K (300, 400 and 660
)). It has
been reported that elements belonging to the same group might be assumed to show similar
diffusion kinetics in α-Al (e.g. DZr  DTi ).240
Al
Al
Of the transition elements, Zr seems to be one of the most promising for the design of
lightweight high-strength high-temperature Al alloys.233, 240 Al3Zr phase not only impedes
the dislocation motion but also refines the microstructure of Al alloys. With the addition of
0.15 wt-% Zr to Al–2% Cu alloy, the columnar grain structure changed to equiaxed
structure.253 Fasoyinu et al.254 studied the effect of Zr, Sc and a combination of both on
grain refinement of 356 alloy; effective concentration ranges of Zr and Sc of 0.37–0.69 and
0.39–0.75 (wt-%), respectively, are required to achieve a considerable grain refinement.
Nevertheless, the phase and microstructure evolution of different Al based alloys (binary
AlZr or multicomponent AlSiCuMgZr alloys) in the presence of this element has been
keenly disputed.
Mahmudi and co-workers255, 256 investigated the effects of 0.15 wt-% Zr addition on the
mechanical properties of A319 Al alloy. The hardness and wear resistance of the A319+Zr
alloy were improved by 10% and 60%, respectively, compared to the A319 alloy, which
were ascribed to the presence of the Al3Zr phase. Garat et al.37 observed the presence of
43
fine, semi-coherent ternary (Al–Zr–Si) dispersoids in the α-Al dendrites of (A356+Zr)
alloy, which were formed during solution heat treatment above 773 K (500
). They
observed no binary Al3Zr phase in the microstructure. Ozbakir257 also reported that with
0.15 wt-% Zr addition to A356 alloys, the eutectic ternary ε-(Al–Si–Zr) phase was formed
instead of the peritectic binary Al3Zr phase. Prasad258 observed the presence of both Al–Zr–
Si and Al3Zr phases. Iveland259 reported the presence of rod-shaped AlSiZr and AlSiZrTi
precipitates in the heat treated microstructure of A356 alloy containing Zr and Ti. Recently,
the presence of relatively coarse Al3Zr particles (diameters ~ 600 nm) in (A356+Zr) as cast
alloy was reported by Baradarani et al.260 After solution treatment, very fine Al3Zr particles
were observed in the microstructure, which led to the conclusion that either the Al3Zr
particles were not completely dissolved during solution heat treatment or the particles reprecipitated after dissolution. Baradarani et al.260 and Srinivasan et al.249 stated that the
dissolution–precipitation mechanism was promoted by the motion of grain boundaries,
which activates dissolution ahead of the advancing boundary and precipitation behind.
Recent developments in Al–Si alloys and applications in engine components
The Al alloys that are usually used for the fabrication of engine cylinder heads can be
classified into two main categories:37, 47, 261
 aluminium alloys containing 5–9 wt-% of Si, 3–4 wt-% of Cu (generally, treated to
temper T5 or T7) (AlSiCu alloys, such as A319); and
 aluminium alloys containing 7–10 wt-% of silicon and 0.25–0.45 wt-% of magnesium
(generally, treated to temper T6 or T7) (AlSiMg alloys, such as A356).
The secondary alloys based on the 319-type Al alloy, with iron contents between 0.5 and
1% and moderately high contents of other impurities (e.g. zinc, lead), are particularly used
in gasoline engine cylinder heads with fairly low service temperature and pressure. Primary
alloys, based on the 319- and 356-type Al alloys with an iron content of less than 0.3%, are
generally used for highly stressed (diesel engine) cylinder heads. Owing to limited contents
of impurity elements (e.g. Fe, Zn), the primary alloys are more expensive than the standard
secondary alloys. Aluminium alloys based on the 356-type alloy present high ductility and
acceptable strength at ambient temperature. However, their strength significantly decreases
44
above 473 K (200°C). In contrast, the alloys based on the 319-type alloy exhibit higher
yield/creep strength above 473 K (200°C), but present lower ductility.37, 42, 47, 261
In the last decade, several investigations have been carried out as regards the trade-off
between various properties (tensile strength, ductility, creep resistance and fatigue
resistance) of these two large families of aluminium alloys. Four Al–Si based alloys
containing different Cu, Mg and Fe contents were studied by Chuimert et al.42 The alloys
are commonly used by the industry to produce cylinder heads. The results are summarised
as follows:
(1) Al–5Si–3Cu–0.25Mg–0.7Fe (wt-%)
(2) Al–5Si–3Cu–0.25Mg–0.7Fe–1Zn (wt-%)
(3) Al–5Si–3Cu–0.25Mg–0.15Fe (wt-%)
(4) Al–7Si–0.3Mg–0.15Fe (wt-%)
untreated
untreated
T7
T6
→ high strength, low ductility
→ high strength, low ductility
→ high strength, good ductility
→ low strength, extreme ductility
In conditions similar to those encountered in service, the TMF lifetimes of the third and
fourth alloys (with 0.15 wt-% iron content) were up to ~5 times greater than those of the
first and second alloys (untreated alloy with 0.7 wt-% iron content).
Jonason262 investigated thermal fatigue resistance of four different Al–Si alloys (i.e. Al–
8Si–3Cu–0.3Mg–0.7Fe/T5, Al–7Si–3Cu–0.3Mg–0.2Fe/T5, Al–7Si–3Cu–0.3Mg–0.2Fe/T6,
Al–9Si–0.3Mg–0.2Fe/T6 (wt-%)) by cyclically heating and cooling the intervalve seat area
between 313 and 503 K. The Al–9Si–Mg (wt-%)/T6 alloy was found to be the most fracture
resistant alloy with significant tendencies to plastic deformation, the excellent fracture
resistance being attributed to the higher ductility of the alloy. The Al–7Si–3Cu (wt-%)/T6
and Al–7Si–3Cu (wt-%)/T5 alloys were the second and third most fracture resistant alloys,
respectively. Gundlach et al.41 reported very interesting results on TMF resistance of fifteen
different AlSi based alloys (319 and 356 Al alloys) fabricated by seven different foundries.
Testing was done on 78 samples by imposing thermal cycles between 339 K (66
561 K (288
) and
) under axial constraint. The average number of cycles to failure ranged from
162 to 1286 cycles. The lowest and highest fatigue lifetime belonged to the 319 Al alloys.
The authors stated that ‘two unmodified 319 alloys had the lowest TMF lifetime; while two
of the most highly modified 319 alloys displayed the highest TMF resistance’. Also, the
overall TMF lifetime of 356 Al alloys, which was between 228 and 644 cycles, was lower
than that of 319 Al alloys. During the thermal stress cycle, the stress–temperature diagram
displayed a thermal stress hysteresis loop. In thermal cycling up to 477 K (204
), the
45
amount of thermal stress hysteresis was comparable in both 319 and 356 alloys; however,
at higher thermal cycling temperature, 356 alloys displayed considerably larger thermal
stress hysteresis. Superior elevated temperature strength and resistance to overaging of 319
alloys caused less plastic deformation with further benefit of narrowing of the thermal
stress hysteresis loop. The elevated temperature strength of 319 alloys was ascribed to the
presence of Cu-bearing phases.
Feikus47 investigated the addition of 0.5 and 1 wt-% Cu to an Al–8Si–0.3Mg–0.1Fe (wt-%)
alloy for manufacturing engine cylinder heads. No significant improvement in the room
temperature yield strength of the alloys containing Cu was observed after conventional T6
treatment. The tensile strength and creep resistance of the alloys containing Cu were
significantly improved in the temperature range of 423–473 K (150- 200
). A minor
reduction in elongation was also reported. The effect of Cu addition on the coefficient of
thermal expansion and thermal conductivity was negligible. It is interesting to note that the
mechanical properties of both Cu-containing alloys (0.5 and 1 wt-%) were almost
comparable. Subsequently, the impact of Ni (0.5 wt-%) and Mn (0.3 wt-%) on Al–7Si–
0.4Cu–0.4Mg–0.4Fe (wt-%) alloy was extensively studied by Heusler et al.84 The casting
process and the solidification rate were simultaneously investigated. The addition of Ni
improved the creep strength of the alloy; however, it had a rather small effect on the tensile
strength at elevated temperatures. The fatigue strength of the Ni-containing alloy was
approximately 20% higher than that of the AlSiMg alloy. It is important to note that when
the casting process and the cooling conditions were not optimised, the improvement of
mechanical properties by alloy optimisation remained marginal.
Lee et al.263 studied the impact of Al3M (M = Ti, V, Zr) precipitates in AlSiCuMg alloy.
They stated that these dispersoid phases enhanced the high temperature mechanical
properties by effectively blocking the movement of dislocations. Thereafter, Laslaz and
Garat261 investigated the tensile and creep properties at ambient temperature, 523 and 573
K (250 and 300
) of three different alloys (A, B and C) having the following chemical
compositions: A, Al–7Si–0.4Mg–0.15Fe–0.15Ti; B, alloy A + 0.5Cu; and C, alloy A +
0.5Cu + 0.15Zr. The addition of copper to alloy A, which represents alloy B, led to an
improvement in the yield strength and UTS at both ambient and elevated temperatures,
without affecting the elongation. The addition of zirconium to alloy B, which gives alloy C,
46
significantly increased the creep resistance, the deformation under constant load being
reduced by 75%. This was attributed to the precipitation of fine thermally stable AlSiZr(Ti)
dispersoids (1 μm). However, Zr addition had almost no influence on the tensile
properties. They also studied the effect of Mn and Mg additions in alloy C. The high
temperature (~523 K (250
)) mechanical strength improved with increasing Mn content
from 0.1 to 0.3% and with increasing Mg content from 0.3 to 0.5%. They preferred not
adding Ni in the alloy to avoid problems in recycling and to maintain the ductility of the
part. To further improve the mechanical strength and creep resistance at elevated
temperatures (503–653 K (230- 380
)), Laslaz233 investigated the effect of excluding Mg,
and, instead, adding vanadium as another peritectic element. The results are presented in
Table 2-6. These results confirm that tensile properties at 523 and 573 K (250 and 300
)
of the alloys without Mg (alloys 7–9) are better than those of the alloys containing Mg
(alloys 1, 2). At 573 K (300
), the yield strength of the alloys without Mg (alloys 7–9)
exceeds 50 MPa, while the yield strength of the alloys containing Mg (alloys 1–6) is below
50 MPa. The exclusion of Mg makes the aging sequence change from ,  binary phase
(based on Mg2Si) and ,  quaternary phase (based on Al5Cu2Mg8Si6) to ,  (Al2Cu). It
was found that ,  (Al2Cu) phases can be more stable at high temperatures than , 
(Mg2Si) and ,  (Al5Cu2Mg8Si6).233
Moreover, the elimination of Mg and the phase Q (Al5Cu2Mg8Si6), which invariably
reduces the melting point, allows one to increase the solution treatment temperature from T
≤ 773 K to 788–798 K (from T ≤ 500
to 505–525
). The possibility of higher
solutionising temperature has several advantages: greater homogenisation of copper phases,
better modification of Si particles and more complete precipitation of zirconium dispersoid
phases.195,
205, 233
Garat et al.37 confirmed the positive effect of Mg exclusion and the
presence of dispersoid phases on the tensile properties and creep strength.
Nevertheless, Garat264, 265 stated subsequently that adding a small amount of Mg (0.1–0.2
wt-%) to AlSiCu-type alloys is required to improve the LCF strength and room temperature
tensile strength. Adding Mg and V together also had a synergic effect on creep strength (at
573 K (300
)).264, 265 More recently, Iveland259 compared the creep resistance and LCF
behaviour of A356, (A356 + 0.5Cu), A319 and (A356 + 0.5Cu + 0.5Hf) alloys. They
47
observed the presence of ribbon- or nanobelt-like hafnium compound in the -Al matrix
which is a unique microstructure. LCF strength of (A356 + 0.5Cu + 0.5Hf) alloy was the
best, and A319 alloy showed better LCF strength than the rest. This discovery certainly
opens interesting possibilities for niche applications, but not for the high volume
automotive market because of the prohibitive cost of hafnium.
Table 2-6: Chemical composition, mechanical strength and creep properties of Al-Si alloys233
Chemical composition (wt.%)
Alloy
No.
Si
Fe
Cu
Mg
Mn
Zr
Mechanical properties
V
Ti
250 Rm
1
5 0.15 3.1
0.30
0.10 111
2
5 0.15 3.1
0.30
0.14 0.25 0.10
3
7 0.15
0.30
0.10 61
4
7 0.15
0.30
0.12 0.14 0.15 0.10 62
5
7 0.15 0.5
0.38
0.10 73
6
7 0.15 0.5
0.38
0.14
0.10 68
7
5 0.15 4.1 <0.05 0.15 0.14 0.25 0.14 126
8
7 0.15 3.0 <0.05 0.20 0.14 0.25 0.14 100
9
7 0.15 2.4 <0.05 0.19 0.14 0.25 0.14 94
* R : UTS (MPa), R . : Yield Strength (MPa), A: Elongation (%)
Creep properties
300 .
A
92
55
56
66
63
16
35
35
35
35
16
33
37
103
80
75
Rm
.
A
62
43
43
44
45
72
64
60
47
40
41
40
42
63
54
51
30
34
34
35
35
23
34
44
(σ . ) (MPa)
σ (250 ) σ (300 )
60
61
39
40
39
41
53
-
Summary
An increasing social demand for a reduction in fuel consumption and gas emissions calls
for the urgent substitution of cast iron with lighter metals (e.g. Al–Si alloys) in the
production of engine components. Al–Si alloy cylinder heads are already used for engines
with lower firing pressure and temperature peaks, such as gasoline engines. On the other
hand, higher service temperatures and stress amplitudes, which are required to improve
engine performance, might cause fatigue failure in Al–Si alloy cylinder heads. The thin
walls adjacent to the water ducts in the valve bridge of cylinder heads are the most critical
locations for TMF crack initiation.
To reach the optimum pressure and temperature levels desired to ensure efficient
functioning of cylinder heads without the need to develop new materials, the existing
capabilities of Al–Si based alloys have to be improved by optimisation of either production
process or chemical composition. The fatigue lifetime of Al–Si alloys is more affected by
the actual casting processes than by alloy chemistry. This is evident in defect-limited
48
26
28
22
24
22
22
32
-
specimens, where the initiation of fatigue cracks is greatly delayed. The most detrimental
defects of cylinder heads are porosity and inclusions. Thus, measures must be taken to
fortify Al–Si alloys and minimise the above-mentioned defects which accelerate cracking.
To this end, dispersion and precipitation hardening are the major processes adopted in
strengthening Al–Si hypoeutectic alloys. Some transition elements, which can be
precipitated as fine, stable, coherent particles, can significantly improve the TMF
performance. In addition, heat treatment processes play a vital role in microstructural
modification and mechanical properties. The lamellar morphology of brittle Si particles can
be modified to fibrous form by suitable solution heat treatments. A 20 K increase in the
temperature of the solution treatment (from 753 to 773 K (480- 500
)) significantly
enhanced the strength of hypoeutectic Al–Si alloys containing Cu and Mg. For those Al–Si
alloys containing high Cu and Mg content, the duration and temperature of the solution heat
treatment are still debated, and a unique combination of time and temperature might have to
be determined for every single chemical composition.
49
50
Chapter 3
Materials and methods.
This section presents the experimental procedures used to study the effect of alloying
elements (Cu, Mg, and Fe) on the solidification processing and microstructure evolution of
hypoeutectic Al-7(wt.%)Si alloy. Al-Si alloys containing different Cu, Mg, and Fe content
were evaluated. Thermodynamic simulation was carried out to predict the precipitation
sequence and mass fraction of the solidified phases. Ring mould test was utilized to
evaluate the hot tearing susceptibility of the studied alloys. Different characterization
methods were used to determine the precipitations and intermetallic phases. The following
sections describe the procedures with further details.
Alloy making and melting:
3.1.1. Alloy making and melting procedures to evaluate hot tearing susceptibility
About 2 Kg of the as-received 1050 Al-alloys (with chemical composition of 99.84Al,
0.055Si, 0.093Fe, 0.0019Cu) was melted in a clay-graphite crucible, by means of an
electrical resistance furnace. Controlled amounts of Si-containing master-alloys and pure
Fe, Mg and Cu were added to the melt (at ~740
± 2 ) to reach the chemical composition
of the defined alloys. The melt was mechanically stirred after each time of alloying element
addition. To reduce hydrogen concentration in the melt, degassing was carried out for 15
min by bubbling
gas through a lance. After degassing, the melt surface was carefully
skimmed to eliminate the oxide layer and then was kept under argon protective atmosphere
to avoid oxidation. Samples were taken before and after the trials to determine the chemical
composition. The chemical composition of the alloys was analyzed by flame atomic
absorption spectroscopy (AAS). The average chemical compositions are presented in Table
3-1. All chemical compositions are given in weight percent (wt.-%) unless otherwise stated.
The casting method, the procedure developed to index the hot tearing sensitivity and
quantifying microporosity content are thoroughly explained in the next chapter (chapter 4).
Table 3-1: chemical composition of the alloys used to evaluate hot tearing susceptibility (wt.%)
Alloy No.
Si
Cu
Mg
Fe
Al SDAS (μm)
#1 R (reference, A356) 7.1 0.01 0.32 0.13 Bal.
14∓1
#2 RC0.5
7.08 0.54 0.31 0.14 Bal.
14∓2
#3 RC0.5F0.7
7.18 0.51 0.37 0.79 Bal.
15 ∓2
#4 RC3
6.89 3.16 0.32 0.12 Bal.
15∓2
#5 RC3F0.7
6.98 3.1 0.33 0.77 Bal.
16∓2
#6 RC3(M0)
6.91 3.3 0.01 0.16 Bal.
16∓2
#7 RC3F0.7(M0)
6.94 3.08 0.01 0.74 Bal.
16∓2
“R” indicates the reference alloy. The symbols “C”, “F” and “M” represent the Cu, Fe and Mg
elements; the number after each symbol presents the concentration of the respective element.
Alloys #6 and #7 contain zero Mg which were characterized by “(M0)”.
3.1.2. Alloy making and melting procedures for microstructure evolution
The as-received A356.2 Al alloys in the form of 12.5 kg ingots (with chemical composition
of Al, 7Si, 0.12Fe, 0.37Mg), were cut into small pieces, cleaned (with ethanol to remove
excess chips and oil from sectioning prior to melting) and dried. They were used to
elaborate the alloys containing Mg. For the alloys without Mg, the as-received 1050 Alalloys (with chemical composition of 99.84Al, 0.055Si, 0.093Fe, 0.0019Cu), were used.
Then, about 150g of the as-received Al alloy was melted in a graphite crucible (3.5cm
52
diameter and 10.5cm long), by means of an electrical resistance furnace. Controlled
amounts of pure Si, Fe, Mg and Cu were added to the melt (at ~735 ± 2 ) to reach the
desired chemical composition. The melt was mechanically stirred by means of a boronnitride rod after each addition and the melt surface was skimmed to eliminate the oxide
layer prior to sampling. The holding time varied from 25 to 35 minutes.
To have a uniform chemical composition, sampling from the melt was carried out with
Pyrex tubes filled with the help of a propipette (see Figure 3-1). Tubes with 5 mm inside
diameter and 2 mm wall thickness were used in this purpose. One side of the tube was
attached to the propipette, and the other side was preheated first by fire flame and later by
immersion in the melt for ~5 seconds. After solidification and cooling, the tubes were
broken, and the Al alloy bars were extracted. Considering the small size of the bars and the
absence of hot spot along them, this method seems to be very effective to reduce macrosegregation. The chemical composition of the alloys, which was analyzed by atomic
absorption spectroscopy (AAS), is presented in Table 3-2.
The sampling by Pyrex tubes, which reduces the segregation, can be ideal for chemical
analysis. However, the microconstituents of the specimens prepared by this method were
too fine, which sometimes made difficult to identify the phases. Therefore, the rest of the
prepared melt was poured in a room temperature permanent mould (i.e. a cast-iron plate). A
cooling rate of ~1.15 Ks-1 was recorded during the solidification. Therefore, the specimens
prepared with the permanent mould were more appropriate for non-ambiguous phase
identification. The specimens prepared with the permanent mould were verified by (optical
and electron) microscopy and by DSC to have the same signatures as the specimens
sampled with the Pyrex tubes; the permanent mould specimens showed the same
microconstituents and the same DSC results (number of peaks and the temperature
corresponding to each peak) as the specimens sampled with the Pyrex tubes.
The sampling by the Pyrex tubes was used for phase identification at the beginning of the
project. However, since there was difficulty in phase identification, the specimens prepared
by the permanent mould were replaced to this end. In the methodology-section of each of
the following chapters, it is clarified which kind of the prepared specimens was used for the
microstructural characterization.
53
Figure 3-1: Pyrex tubes and propipette used in sampling
Table 3-2: chemical composition of the alloys (wt.%) used for microstructure evolution
Alloy No.
Si
Cu
Mg
Fe
Al
A356.0
Ref.
6.5-7.5
0.2
0.25-0.45
0.2
Bal.
#1
RC0.5
7.08
0.54
0.30
0.12
Bal.
#2
RC0.5F0.7
7.18
0.51
0.31
0.79
Bal.
#3
RC1.5
6.98
1.5
0.3
0.10
Bal.
#4
RC3
6. 9
3.38
0.35
0.12
Bal.
#5
RC3F0.7
6.98
3.1
0.33
0.77
Bal.
#6
RC3(M0)
6.9
3.3
0
0.13
Bal.
#7
RC3F0.7(M0)
6.94
3.08
0
0.74
Bal.
#8
RC1M0.4
6.81
1.05
0.39
0.08
Bal.
#9
RC1M0.8
6.82
0.99
0.78
0.06
Bal.
#10
RC1.6M0.4
6.77
1.64
0.38
0.06
Bal.
#11
RC1.6M0.8
6. 86
1.63
0.78
0.06
Bal.
“R” indicates the reference alloy. The symbols “C”, “F” and “M” represent the Cu, Fe and Mg
elements; the number after each symbol presents the concentration of the respective element. Alloys
#6 and #7 contain zero Mg which were characterized by “(M0)”.
Thermodynamic Prediction:
A comprehensive study of the thermodynamic evaluation of Al-7Si alloy with addition of
the elements was carried out using the Thermo-Calc software (with TTAl7 database). This
thermodynamic simulation can identify the phases, the transition temperatures and the
reactions that occur during the (equilibrium/ non-equilibrium) solidification interval of the
alloys at the defined compositions.
A computational algorithm developed by Larouche
1
was used to calculate the phase
precipitations and their mass fraction during solidification. The algorithm is based on the
54
assumptions of equilibrium at the solid/liquid interface, uniform composition in the liquid,
and mobility of each element in the primary phase with back diffusion model. This
algorithm was linked to the Thermo-Calc software package, and database TTAL7 was used
to calculate the thermodynamic variables. Computations based on this algorithm will be
referred below as the multiphase back diffusion (MBD) model.
Atomic absorption spectroscopy
The chemical analysis of all major elements was carried out by atomic absorption
spectroscopy (AAS). The specimens were cut and about 150 mg solid sample was placed in
a 100 ml polyethylene digestion bottle with a binary acid mixture (2ml concentrated HF (4849%) and 25ml of diluted HCl (10 vol. %) with 50 ml of distilled water). The bottle was
shaken overnight at 60 to entirely dissolve the sample. Subsequently, the solution was analyzed by atomic absorption spectroscopy (AAS).
Microstructural Analysis:
For microstructural investigations, the specimens were cut and cold-mounted in an epoxy
resin-hardener mixture. The specimens were then subjected to grinding and polishing
procedures to produce a mirror-like surface. Generally, the grinding procedures were
performed in successive steps using silicon carbide (SiC) abrasive papers in a sequence of
180, 320, 400, 600, and 1200 grits sizes under water spray. The starting grit size depended
on the condition of the initial surface. If the specimens had been cut by band saw, 180- grit
paper was first used. In the case of the specimens cut by diamond blade, 600-grit paper was
first used. Prior to polishing, the specimens were held in an ultrasonic bath for ~4 minutes
to remove any excess particles and dirt. Polishing were carried out using diamondsuspension, which contained a diamond particle size of 6 μm, as the first step of the
polishing process; it was followed by further polishing through the application of the same
suspension containing a smaller diamond particle size of 1 μm. The final polishing was
carried out using a colloidal silica suspension, having a particle size of 0.05 μm. Distilled
water was used as lubricant throughout the final polishing stage. After each polishing step,
the same ultrasonic bath treatment was applied.
55
The secondary phases were identified by means of scanning electron microscopy (SEMJEOL 840A) and electron probe microanalysis (EPMA-CAMECA SX100); moreover, they
were further studied by means of an optical microscope (OM-NIKON EPIPHOT) equipped
with a CLEMEX image analysis program (CLEMEX VISION PROFESSIONEL).
The SDAS, which is the linear distance between two secondary aluminum dendrites (arms),
was determined via the mean linear intercept (MLI) method. As illustrated in Figure 3-2,
the SDAS was identified as the ratio of length segment to the number of dendrite arms
266
.
Eight (and/or more) primary dendrites containing at least 5 secondary arms were considered
to measure the average value of SDAS in one sample.
Figure 3-2: SDAS mesurement of the specimens
Differential Scanning Calorimetry (DSC):
The specimens for DSC analysis were sectioned from the bars, using a low speed cutter
with a diamond blade so as not to cause any additional heat or stress on the samples.
Subsequently, the specimens were grounded to reach a desired weight (20-30 mg). DSC
tests were carried out on a power compensated Perkin-Elmer Diamond DSC under
protective argon atmosphere and using alumina crucibles in both reference and sample
pans.
To study the sequence of precipitation occurred during solidification, DSC device was
programmed as following: each sample was heated from room temperature to 450 at a
scanning rate of 100 /min, and then heated from 450
to 680
at a scanning rate of
10 /min; afterwards, held at this temperature for 1 minute to be completely homogenous.
The sample was later cooled down to 450 at a scanning rate of 5 or 10
56
/min (mentioned
in the corresponding text/ caption of figure) and finally cooled down from 450
to room
temperature with a scanning rate of 100 /min. To evaluate the efficiency of solution heat
treatment, the same heating program applicable for the as-cast specimens was applied on
the as-quenched specimens.
Heat Treatment:
Solution heat treatment (SHT) was conducted in an electric resistance furnace. The
temperature of (first step) SHT was always lower than the melting point (Tm) of the last
solidified melt, which has already been determined by means of DSC analysis. In some
cases, the second step of SHT was applied at higher temperatures (TSHT>Tm; e.g. 530 ). A
K-type thermocouple was used to monitor the TSHT. After holding the specimens in the
intended time/temperature of SHT, they were then water-quenched to room temperature (in
less than 4 seconds) to assure maximum solute-saturation. The cooling curve during
quenching, which was recorded by means of a Data-Logger (OM-DAQPRO-5300), is
presented at Figure 3-3.
RC3(M0)
Temperature (ºC)
500
400
300
200
100
0
3
5
7
9
11
13
15
17
Time (Second)
19
Figure 3-3: cooling curve of the alloy RC3(M0) during quenching, recorded by Data-Logger (OM-DAQPRO5300).
57
58
Chapter 4
.
“Hot Tearing Susceptibility of Al-Si Based Foundry Alloys Containing Various Cu,
Mg and Fe Content”
Résumé:
Les défauts de solidification (tendance à la fissuration à chaud et microporosité) des
alliages de fonderie Al-Si contenant différentes teneurs en Cu, Mg et Fe ont été étudiés à
l'aide d’un moule annulaire permanent. Une nouvelle méthode d’indexation semiquantitative, nommée la sensibilité à la fissuration à chaud (HTS), a été définie afin de
refléter le volume de fissures générées dans les échantillons ayant subies une déchirure à
chaud. En augmentant la teneur en Cu et en Fe des alliages, la valeur de HTS et la fraction
surfacique de porosité ont été augmentées. Les microstructures des alliages ont été
minutieusement étudiées pour comprendre l'effet des éléments sur les défauts. Une
présence accrue de la phase β-Al5FeSi a augmenté le niveau de microporosité en bloquant
physiquement l’alimentation en métal dans les poches liquides restantes. L'augmentation de
la concentration en Cu (de ~0,5 à 3%) dans les alliages a augmenté le niveau de
microporosité aussi.
Des calculs thermodynamiques ont également été utilisés dans l’analyse des
microstructures obtenues. La sensibilité à la fissuration à chaud (HCS) des alliages a été
simulée avec l'indice de fissuration à chaud de Katgerman. La température critique (Tcr)
utilisée dans l'indice théorique (HCS) correspond au moment où 2% du volume
interdendritique est occupé par des particules de phase secondaire. La corrélation obtenue
entre les résultats expérimentaux (HTS) et les résultats simulés (HCS) est excellente. Un
nouvel indice (βR) a été introduit par redéfinition de l'indice de fissuration à chaud (CSC =
Δtv / Δtr) initialement proposé par Clyne et Davies. βR représente le ratio de contraction de
solidification se produisant pendant la période de temps vulnérable (Δtv) et pendant la
période de temps de relaxation de la contrainte (∆tr). Les corrélations de βR avec le
pourcentage de surface de porosité et avec le HTS étaient excellentes toutes les deux.
Abstract:
Solidification defects (hot tearing tendency and microporosity) of seven different Al-Si
based foundry alloys were studied by means of the ring mould test. The sensitivity of the
alloys to hot tearing (HTS) was ranked by developing a new semi-quantitative index
through which the volume of the generated cracks in the torn specimens is compared
together. None of the studied alloys were susceptible to hot tearing at higher mould
temperature (>250
; at lower mould temperature, tendency to hot tearing was found to
increase with increasing Cu and Fe contents. Microstructure analysis illustrated that βAl5FeSi phase enhances microshrinkage porosity by physically impeding metal feeding.
The increase of Cu concentration from ~0.5 to ~3% in the Al-7Si foundry alloys increased
the level of shrinkage microporosity as well.
The HTS results presents a very good correlation with the results simulated by the
Katgerman’s hot tearing index (HCS). The critical temperature Tcr used in the HCS index
presumed the temperature at which 2% of the interdendritic volume is occupied by
secondary phase particles. Moreover, a new index βR based on the Clyne and Davies index
was introduced which reflects the ratio of solidification shrinkage in the vulnerable time
period (∆tv) and in the stress relief time period (∆tr). The correlations of βR with porosity
area% and βR with HTS were both excellent.
Introduction:
As a common and severe casting defect, hot-tearing occurs during solidification above the
non-equilibrium solidus of metals. It is generally caused by the thermal stress produced by
the restraining of solidification contraction. As the accumulated thermally induced stress
exceeds the strength of the mushy zone, and liquid feeding is insufficient to compensate the
solidification shrinkage in the vulnerable temperature range, hot tears can be generated
60
267,
268
. They are frequently observed near a hot spot region, where heat transfer is insufficient.
Their fracture surface is generally intergranular and have a dendritic morphology 267, 269.
Various theoretical models have been developed to evaluate the hot tearing tendency. Clyne
and Davies270, as pioneers, proposed that the hot tearing susceptibility could be
characterized by the ratio of the vulnerable time period tv to the time period available for
stress relief tr. This ratio was called the Cracking Susceptibility Coefficient (CSC) and
defined as follow: CSC= ∆tv/∆tr, where ∆tv is the vulnerable time period for tears to
propagate (critical time interval for interdendritic separation), and ∆tr is the time available
for stress relaxation processes (e.g. liquid/mass feeding). The authors pointed out that the
vulnerable region (∆tv) belongs to the solidification interval through which the fraction
solid (fs) is in between 0.9 and 0.99 and liquid flow is restricted to narrow interdendritic
channels. The time for stress relaxation (∆tr) is limited to 0.9 and 0.6 solid fraction (fs)
through which the permeability is supposed to heal the possible incipient tears. The
aforementioned fraction solid range was subsequently modified by Katgerman
271
.
According to Katgerman theory, the vulnerable time period (∆tv) is limited to a region with
a solid fraction in between 0.99 and
;
corresponds to the critical point after which
the system transits from a regime with adequate liquid feeding to a regime with inadequate
liquid feeding. Based on the Feurer’s criterion272, Katgerman
271
reported that the critical
point is attained when the velocity of volume shrinkage is equal to the maximum
volumetric flow rate per unit volume. The time period of stress relaxation (∆tr) is limited to
a region between
and
which correspond to dendrite coherency point. According
to Katgerman theory, the CSC was re-defined as: CSC= (t0.99-tcr)/( tcr-tcoh).
To evaluate the hot tearing tendency of different alloys, various methods such as ring
mould test
47, 269, 273-275
, the cold finger test276,
277
and Constrained Rod Casting (CRC)
mould278-280 were utilized. The ring mould test is a simple and widely used technique 47, 269,
273-275
in which, a rigid core resists the solidification contraction and induces tensile stress
onto the solidifying alloy.
Process parameters (e.g. mould temperature), mould design (e.g. presence/lack of hot spot)
and chemical composition of the alloy are the major factors to influence the hot tearing
susceptibility
267, 281
. While a lot of papers were published on hot tearing, few researches
61
reported the impact of mould temperature on the occurrence of hot tearing. Though it is
generally accepted that higher mould temperature improve permeability and liquid feeding;
which in turn, alleviate hot tearing susceptibility267,
268
. According to Li267, tears can be
generated at any mould temperature; but a higher mould temperature increases the crack
onset temperature and extends the propagation time which help to heal the crack with the
remaining liquid. Solidification interval and micro/macro-structure parameters (e.g. eutectic
fraction, and micro-constituents size and morphology) are the other main features which are
strongly affected by chemical composition. Longer solidification interval, which elongates
the vulnerable temperature range, increases hot tearing susceptibility.
Two categories of Al-Si based alloys, viz. 319- and 356- type alloys, are widely used in
automotive application (e.g. engine components) due to the low density, high thermal
conductivity and excellent mechanical properties. These Al alloys are prone for casting
defects (mainly shrinkage porosity) which significantly influence their quality
characteristics. Nevertheless, owing to their high Si content, the susceptibility of these
alloys to the defects is significantly lower than other Al alloys (e.g. AlZn, AlCu, AlZnMg).
It has been reported that the overall shrinkage during solidification process of pure Al, AlCu binary alloys and Al-7%Si alloys are respectively ~8.14%, ~8.4%, and ~4.5% 3‐5. Nonetheless, impurities and alloying elements can strongly affect their hot tearing
resistance and microporosity content. Few contributions can be found in the literature
concerning the hot tearing sensitivity of the Al-Si foundry alloys. Paray et al. 282 studied the
effect of strontium content and dissolved hydrogen concentrations on hot tearing
susceptibility of 319-type Al alloys; they reported beneficial effect of strontium addition
and higher hydrogen level in reducing hot tearing tendency. Bozorgi et al.
283
studied the
hot tearing tendency of five different AlSi7MgCu-alloys with varying Mg and Cu contents.
They stated that increasing Cu content enhanced hot tearing tendency, but increasing Mg
content had beneficial effect on hot tearing resistance.
Edward et al.4 and Cáceres et al.157 pointed out that increasing Cu concentration
significantly enhances volume fraction of porosity. Mackay et al.
3, 5, 284
investigated the
effect of Si and Cu content on soundness of cast structure. They stated that Al-9Si-1Cu
alloy had the lowest level of porosity and Al-7Si-4Cu alloy had the highest level. They
concluded that higher volume fraction of primary α-Al dendrites, lower volume of Al-Si
62
eutectic phase within larger freezing range, and higher volume fraction of the Cu and Mg
containing post-Al-Si eutectic phases were the main reason of higher porosity level in the
Al-7Si-4Cu alloy. Numerous researches can be found in studying the effect of Fe content
on porosity of AlSi alloys
122, 285-287
. The increased porosity level is associated with the
enhanced volume fraction of β-Al5FeSi platelets, which physically block the interdendritic
flow channels 127, 285, 287, 288.
The main purpose of this research is to evaluate the effect of Cu, Mg and Fe contents on
casting defects of the Al-7Si alloy. The sensitivity of the alloys to hot tearing (HTS) was
ranked by developing a new semi-quantitative index. The effect of the elements (Cu, Mg
and Fe) on area percentage of porosity was evaluated. Microstructures of the alloys were
studied to understand the effect of the elements on the defects. The hot cracking
susceptibility (HCS) of the alloys was simulated by the Katgerman’s hot tearing index.
Moreover, a new index (βR) was introduced based on the Clyne and Davies index to reflect
the ratio of solidification shrinkage in the vulnerable time period (∆tv) and in the stress
relief time period (∆tr). HSC and βR, both, were simulated by multiphase back diffusion
model developed by Larouche1.
Materials and Method:
The alloy making and melting procedures were described in preceding chapter (section
3.1.1).
The average chemical compositions and the secondary dendrite arm spacing
(SDAS) of the 7 alloys investigated are presented in Table 4-1.
Table 4-1: chemical composition (wt.%) and SDAS of the alloys
Alloy No.
Si
Cu
Mg
Fe
Al SDAS
#1 R (reference, A356) 7.1 0.01 0.32 0.13 Bal. 14∓1
#2 RC0.5
7.08 0.54 0.31 0.14 Bal. 14∓2
#3 RC0.5F0.7
7.18 0.51 0.37 0.79 Bal. 15 ∓2
#4 RC3
6.89 3.16 0.32 0.12 Bal. 15∓2
#5 RC3F0.7
6.98 3.1 0.33 0.77 Bal. 16∓2
#6 RC3(M0)
6.91 3.3 0.01 0.16 Bal. 16∓2
#7 RC3F0.7(M0)
6.94 3.08 0.01 0.74 Bal. 16∓2
“R” indicates the reference alloy. The symbols “C”, “F” and “M” represent the Cu, Fe and
Mg elements; the number after each symbol presents the concentration of the respective
element. Alloys #6 and #7 contain zero Mg which were characterized by “(M0)”.
Hot tearing experimentations were performed using a cast iron ring mould presented in
Figure 4-1. This mould had an enlarged section to generate a hot spot and to facilitate the
63
pouring of the melt. In order to reduce the friction between the mould wall and the melt and
to homogenize the heat transfer rates, a coating of boron-nitride (BN) was applied in the
mould cavity. This was done prior each series of test by cleaning and preheating the mould
up to 150
before applying the BN coating. The melt was poured at 735 when the target temperatures at two locations in the mould surface were met: near the core center and at the periphery of the mould. Figure 4-1: Schematic view of the ring mould used for the investigation of hot tearing tendency
Two kinds of experiments were made. The first series was made by casting rings using the
same mould temperature for every alloy. The temperature of the mould was set to ~260°C
resulting in no hot tearing in any alloys. The goal was to evaluate the microporosity
generated under similar casting conditions, giving an indication about the propensity of
each alloy to generate microscale type of defects. A cooling rate of ~6.5 Ks-1 was recorded
during solidification. The second series was made to compare the propensity of each alloy
to generate macroscale defects like hot tears. Since only one ring was cast per pouring, it
was not possible to cast the 7 alloys at the same mould temperature without losing some
discrimination power about the hot tearing susceptibility. If the mould was too hot (260°C
for instance), none of the alloys experienced hot tearing; if the mould was too cold, there
was difficulty to completely fill the mould cavity by liquid in some alloys (e.g. RC3 &
RC3F0.7). Therefore, it was decided to work at different mould temperatures for the
studied alloys (one defined temperature for each alloy). The strategy was to decrease
64
steadily the mould temperature down to the point where hot tearing occurred but not at the
price of incomplete filling. The mould temperatures for each alloy giving consistent results
are presented in Table 4-2. These temperatures alone give an indication of the hot tearing
susceptibility of the alloys. For example, it was found that with a mould temperature of
220 , a severe hot tearing was obtained in the RC3F0.7 alloy, while for the RC0.5F0.7
alloy, it was necessary to reduce the mould temperature down to 140
to produce a hot
tear.
Table 4-2: Mould temperature of the alloys
Alloy Number
RC3F0.7
RC3F0.7(M0)
RC3
RC3(M0)
RC0.5F0.7
RC0.5
R
Mould Temperature
220
220
180
180
140
120
100
4.2.1. Hot tearing indexation:
Classification of hot tearing susceptibility was performed by using a semi-quantitative
indexation method largely inspired from the one proposed by Paray et al.
282
. The index
called Hot Tearing Sensitivity (HTS) rates the severity of the tears obtained according to
this formula:
N
HTS   X  Y  Z
(1)
n 1
where X, Y and Z increases respectively according to the length of the tear in the
circumferential direction, the gap width across the tear and the tear depth. The summation
was made over all specimens cast with each alloy. In this study, at least five trials were
carried out for each alloy. Table 4-3 presents the different rating numbers dedicated to each
parameter (X, Y, Z) and Figure 4-2 shows representative examples of the torn specimens.
As defined above, one can say that HTS is a number representing the severity of the defect.
65
Table 4-3: Crack size parameters for hot tearing index
Category: Arc length
Rate Number (X)
No tear
light crack (C≤2/3T)
Severe crack (2/3T≤ C )
Category: Gap width
0
1
2
Rate Number (Y)
No tear
Small opening
Medium opening
Large opening
0
1
2
3
Category: Tear depth
Rate Number (Z)
No tear
surface crack
The crack penetrate up to 0.5T
The crack penetrate more than 0.5T
Complete fracture
C: crack length, T: Thickness of sample
0
1
2
3
4
30 mm
Schematic view of the cracked ring
2 mm
alloy R0.5Cu (X=1, Y=1, Z=1)
2 mm
2 mm
Alloy RC3 (X=2, Y=2, Z=3)
Alloy RC3F0.7(M0) (X=2, Y=3, Z=4)
Figure 4-2: macrographs to illustrate the different severity levels of hot tearing in the alloys
66
4.2.2. Samples preparation and characterization
Samples for microstructural examination were cut close to the hot-spot regions, mounted,
ground and polished using standard procedure. The polished sections were then studied to
identify the morphology and distribution of second phase particles around the tear surface.
The dendrite arm spacing (SDAS) was calculated by the standard linear intercept method.
The volume fraction of porosity of the alloys was evaluated close to the hot-spot regions
using the standard metallographic procedure. The surface fraction of porosity was
quantified by means of an optical microscope and Clemex image analyzer, assuming that
volume and surface fractions are equal. Image analysis of porosity was done on 4 different
cross sections near the hot spot and oriented perpendicularly to the radius of the ring, each
having an area of 5.8 mm2. The global mean value and standard deviation were calculated
with these 4 measurements.
To study the tear surface, the specimens were sectioned from the rings containing (hot) tear.
The rings with a small hot tearing level, i.e. the incompletely broken rings, were thoroughly
broken to subsequently evaluate the crack surface by SEM/EDX.
4.2.3. Thermodynamic Prediction:
The most important factor on hot tearing is the chemical composition affecting
solidification interval and the amount of liquid phase present at different solidification
stages. Therefore, to understand the variation of HTS, the solidification interval of the
alloys were investigated by means of the multiphase back diffusion (MBD) model1.
Experimental results and discussion
4.3.1. Microstructural constituents
Figure 4-3 illustrates the optical micrograph obtained on the 4 studied alloys. The
microstructure of alloy RC0.5 (and RC0.5F0.7) contains α-Al, Si, Cu-bearing intermetallics
(Q-Al Mg Si Cu
& θ-Al Cu) and Fe-bearing intermetallics (β‐Al FeSi & π-
Al Mg FeSi ). The alloy RC3 (and RC3F0.7) had the same microstructure as alloy RC0.5,
except π-Al Mg FeSi , which was not observed in its microstructure. In the alloys
RC3(M0) and RC3F0.7(M0), the phases are limited to α-Al, Si, β‐Al FeSi and θ-Al Cu.
67
Higher Fe content promotes increasing the volume fraction and length of the β-Al FeSi
phase. While there were only a few small β-Al FeSi phase in alloys containing less Fe (e.g. RC3), a large number of β-phases were found in the microstructure of the alloys containing
high level of iron (e.g. RC3F0.7).
Figure 4-4 displays the DSC cooling curves of the alloys RC3 and RC3F0.7. The peak at
~584
in the alloy RC3F0.7 corresponds to the formation temperature of β-Al5FeSi
phase (T β
)7. In the alloy RC3, the peak corresponds to T
initiates at lower
temperature. Merging with the peak correlated to Al+Si eutectic reaction, make impossible
to specify an exact temperature to T
in alloy RC3. However, as a whole, increasing Fe
content of the alloys (from ~0.12 to ~0.75%) enhanced T
and changed the reaction-
type of β-(Al FeSi) phase from post-eutectic to pre-eutectic7. Enhancement of the T
provides more time available for lengthwise growth and facilitates the diffusivity of Fe
atoms which considerably accelerate coarsening.
The predicted solidification temperature and the mass fraction of major phases in the alloys
RC3 and RC3F0.7 are compared in Figure 4-5. Silicon, in both alloys, is the predominant
secondary phase. The system containing less Fe (RC3) is composed of Cu-containing
phases (θ‐Al Cu & Q as the main intermetallics and a small quantity (~0.4%) of β(Al FeSi) phase. But in the system containing high Fe content (RC3F0.7), both β-(Al FeSi)
and Cu-containing phases (θ‐Al Cu & Q are the major intermetallics. Moreover, it can be
seen that in the system containing high Fe content, β-phase begins to solidify along with αAl and before the Al-Si eutectic reaction (as a pre-eutectic phase). N-Al7Cu2Fe phase was
rejected in the calculation due to negligible volume fraction. There is a good correlation
between the predicted and experimentally observed intermetallics, which indicates that the
thermodynamic calculations using the present databases can be used to predict the
microstructural evolution.
68
Al Cu
Al FeSi
Al Mg FeSi
Al FeSi
Al Mg Si Cu
Al Mg Si Cu
Al Cu
Si
RC0.5
RC3
Al FeSi
Al Cu
Al FeSi
Al Cu
Si
Si
RC3(M0)
RC3F0.7(M0)
Figure 4-3: As-cast microstructures of the four alloys studied.
T
T
qDSC (mW/mg)
Liquidus
2.7
1.7
RC3F0.7
RC3
0.7
-0.3
485
505
525
545
565
585
T (˚C)
605
Figure 4-4: DSC cooling curves of the alloys RC3 and RC3F0.7, with a scanning rate of 10 K/min..
69
FCC
Al5FeSi
Al5Cu2Mg8Si6
1
RC3
Mass Fraction of Phases
Mass Fraction of Phases
1
Si
Al2Cu
0.1
0.01
FCC
Al5FeSi
Al5Cu2Mg8Si6
Si
Al2Cu
RC3F0.7
0.1
0.01
0.001
0.001
0.0001
0.0001
505
555
T (˚C)
605
505
555
T (˚C)
605
Figure 4-5: evolution of the main intermetallic phases (calculated by MBD model) in alloys: (a) RC3, (b)
RC3F0.7.
4.3.2. Characterization of microporosity
The area percentages of microporosity near the hot spot of the alloys are presented in
Figure 4-6. These seven alloys can be divided into 3 categories depending of their
combined Cu and Fe contents and the level of microporosity produced. The first category
comprises the alloys R and RC0.5, both having the lowest combined amount of Cu and
Fe. These alloys showed the lowest amount of microporosity. Notice that the amount of
porosity in the alloy R was negligible. The alloys RC0.5F0.7, RC3(M0) and RC3 belong
to the second category, which is characterized by a slightly higher amount of
microporosity. Finally, the third category comprises alloys RC3F0.7(M0) and RC3F0.7,
both having the highest Cu and Fe content (~3%Cu & ~0.75%Fe). This category is
characterized by the highest amount of microporosity. These results indicate that the
amount of microporosity increases with the combined contents of Cu and Fe. This is in
agreement with the findings reported in literature
4, 157
. Figure 4-7 presents the
microstructure of specimens taken from each category. It is clear that the micropores are
of the type shrinkage porosity.
70
1
Area percentage of porosity
0.8
0.6
0.4
0.2
0
600 µm
Figure 4-6: microporosity content in the alloys.
RC0.5
RC3
RC3F0.7(M0)
Figure 4-7: Microstructure of 3 alloys showing the micropores formed near the hot spot. The amount of
porosity in these microstructure are in ascending order from left to right and can be representative of the 3
categories of alloys.
4.3.3. Hot tearing sensitivity
The HTS index of the alloys, which was calculated with formula (1), is plotted in Figure
4-8. None of the studied alloys were susceptible to hot tearing at higher mould temperature
( 250 ), but at lower mould temperature, the alloys containing high Cu (and Fe) content
71
were more vulnerable to hot tearing. A356 as reference alloy (R), which is the least prone
to hot tearing, is also included for comparison. It was found that the alloys containing high
Cu (~3%) and Fe (~0.7%) content (RC3F0.7 and RC3F0.7(M0)) are the most susceptible
alloys for hot tearing, and the alloy RC0.5 is the less susceptible (after A356-alloy). As a
whole, the HTS results showed that increasing the content of Cu and Fe in Al-Si alloys
reduce the hot tearing resistance; but adding Mg to the (319-type) alloys seems to have
negligible influence on hot tearing.
100
HTS
80
60
40
20
0
Figure 4-8: Hot tearing index (HTS) of the studied alloys
4.3.4. Hot tear surface analyses
Micrographs of typical crack surfaces of the alloys are presented in Figure 4-9. Silicon
particles were found on the tear surface of all alloys. The presence of Si particles in the tear
surface indicates that hot tearing was initiated at a temperature lower than the onset
temperature of the Al-Si eutectic reaction. In the tear surfaces of the alloys containing high
Fe content (~0.7%Fe), large β-Al5FeSi platelets were identified in the intergranular regions
(Figure 4-9 a and d). Optical micrograph across the tear region of the alloys (RC3(M0) and
RC3F0.7(M0)), which are presented in Figure 4-10, confirm the results of the fractographic
analysis. It can be assumed that the enhancement of HTS in the alloys containing high Fe
content (RC3F0.7) is linked to the increased occurrence of lamellar β-Al5FeSi phase; β72
Al5FeSi phase promotes the shrinkage porosity during solidification by physically blocking
the metal feeding, as shown in Figure 4-11.
In the tear surface of the alloy RC0.5F0.7, along with the presence of β-Al5FeSi and Si
particles, a layer of eutectic phases partially covered the tear area; the presence of these
eutectic phases can be attributed to the liquid feeding that occurred in a vain attempt to heal
the crack. Worth to mention that in the RC0.5 alloy, the tear was too superficial to perform
a fractographic analysis.
Si
Si
a) RC3F0.7
b) RC3
Si
Si
c) RC3(M0)
d) RC0.5F0.7
Figure 4-9: SEM micrographs of the hot tear section in the alloys.
73
a) RC3(M0)
b) RC3F0.7(M0)
Figure 4-10: Optical micrograph across the tear region of the alloys a) RC3(M0) and b) RC3F0.7(M0), solidarrow: β-Al5FeSi phase, dash-arrow: Si-particles.
Figure 4-11: Physically blocking the metal feeding by β-Al5 FeSi phase
4.3.5. Prediction Hot Tearing Susceptibility:
One of the objectives of this work was to provide a quantitative manner to evaluate how the
hot tearing susceptibility of Al-Si foundry alloys is affected by addition of the elements. In
a previous work done by Kamga et al. 278, a similar goal was put forward and an index was
proposed to explain the hot tearing susceptibility variation of Al-4.5%Cu alloys having
different amounts of Fe and Si. Taking as predominant the influence that Fe secondary
phases can have on the healing process, they rewrote the hot tearing index (HCS) of
Katgerman 271 as below:
HCS=
74
.
(2)
where Tcoh is the dendrite coherency temperature, Tcr is the temperature below which liquid
after feeding is inadequate and T0.01 is the temperature at which the volume fraction liquid
is equal to 0.01. The key parameter in this equation is Tcr, since the time allowed to the
healing process depends strongly on that. A higher value of Tcr means that liquid feeding is
stopped sooner during solidification for a given family of alloy. Kamga et al.278 obtained a
very good correlation between the value given by Eq. (2) and the index characterising the
severity of defects by defining Tcr as the temperature at which a given portion (ocr) of the
interdendritic volume is occupied by secondary phases:
 1- gl - gpp 
 = οcr
1g
pp


Tcr = temperature at which: 
(3)
where gl and gpp are, respectively, the volume fraction of the liquid and the primary phase.
The secondary phases responsible of the increase of hot tearing susceptibility were
basically the Fe intermetallic phases, which impeded the liquid feeding. The HCS index can
be calculated with the multiphase back diffusion (MBD) model1 and the value obtained for
each alloy is plotted against the measured index (HTS) in Figure 4-12. The correlation
between HCS and HTS is not very sensitive to values of ocr in the range 0.02-0.05, but for
the alloys investigated, the best correlation was obtained with ocr=0.02.
3.0
R² = 0.8849
2.0
HCS
R
RC0.5
RC0.5F0.7
1.0
RC3(M0)
RC3
RC3F0.7(M0)
RC3F0.7
0.0
0
20
40
60
80
100
HTS
Figure 4-12: Plot of HTS versus HCS assuming ocr= 0.02 and fraction liquid =0.78 at the dendrite coherency
point.
75
Notice that the fraction liquid at the dendrite coherency point was chosen based on
experimental results obtained by Veldman et al.289 on different aluminium alloys containing
7%Si and up to 4%Cu using a rheological method. The solidification paths were calculated
based on a constant cooling rate of 6.5 Ks-1 with the composition and SDAS given in Table
4-1. The plot presented in Figure 4-12 shows a clear trend between the calculated and the
measured index, but since the latter (HTS) is semi-quantitative, one can say that there is
maybe a subjective factor in the definition of HTS. It is why porosity was measured near
the hot spot of the specimens cast with the same mould temperature. Since porosity was
clearly related to shrinkage, the amount of this defect could be related to the calculated
(HCS) index. Figure 4-13 presents the plot of the porosity area% measured vs. the HCS
calculated as above. The correlation is not very good indicating that there is probably
something missing in the analysis.
R
Porosity area %
0.8
RC0.5
RcC0.5F0.7
0.6
RC3(M0)
RC3
0.4
RC3F0.7(M0)
RC3F0.7
0.2
0.0
0.0
0.5
1.0
1.5
2.0
HCS
2.5
3.0
Figure 4-13: Plot of porosity area versus HCS calculated with cr = 0.02 and fraction liquid =0.78 at the
dendrite coherency point.
Since porosity is generated by solidification shrinkage, perhaps one has to include
shrinkage in the definition of the hot tearing index. According to Clyne and Davies
theory270, the hot tearing susceptibility could be evaluated by the ratio of the vulnerable
time period tv to the time period available for stress relief tr (CSC= ∆tv/∆tr). Similarly,
one can write:
CSC 
76
 v  r

  R  r
 r  v
 v
(4)
where ∆εv and ∆εr are the solidification shrinkage occurred during the vulnerable time
period and the stress relief time period, respectively. The parameters
 v and  r
are the
average strain rates associated to these periods. Since the ratio βR=(∆εv/∆εr is related to
solidification shrinkage, one can expect to find a good correlation between this parameter
and the consequence of shrinkage, namely porosity area%. Figure 4-14 presents the
correlation obtained between the βR and the area% of porosity.
1.0
Porosity area %
R
R² = 0.8935
RC0.5
0.8
RcC0.5F0.7
0.6
RC3(M0)
RC3
0.4
RC3F0.7(M0)
RC3F0.7
0.2
0.0
0.5
0.7
0.9
1.1
1.3
1.5
βR
1.7
1.9
2.1
Figure 4-14: Plot of porosity area versus R calculated with cr = 0.02 and fraction liquid =0.78 at the dendrite
coherency point.
Clearly, porosity is strongly related to the ratio βR=(∆εv/∆εr . The calculation of shrinkage
deformations
v and r is detailed in the appendix. The ratio of strain rates in the
vulnerable and the relaxation regimes are not expected to vary significantly from one
alloy to the other even though the strain rates alone may vary somewhat. In fact,
 r
 v and
vary nearly at the same pace according to the applied cooling rate, so their ratio
remains almost constant. The ratio βR=(∆εv/∆εr
is consequently a key parameter
describing the severity of defects. If the correlation between
R
and porosity area % is
excellent, it is also good with the experimental index HTS as this is shown in Figure 4-15.
77
100
R² = 0.807
356
80
356Cu
356CuFe
319
HTS
60
319Mg
319Fe
40
319MgFe
20
0
0.5
1
1.5
R
2
Figure 4-15: Plot of HTS versus R calculated with cr = 0.02 and fraction liquid =0.78 at the dendrite
coherency point.
Solidification shrinkage and liquid feedability at the later stages of solidification process
are the major parameters that influence hot tearing tendency. These parameters are a
function of the chemical composition of the alloys, the solidification thermal conditions,
and process parameters (e.g. mould temperature). In this work, the analysis of hot tearing
severity and porosity content in the Al-Si based (356- and 319-type) foundry alloys
emphasize the importance of the after feeding critical temperature Tcr and of R, defined as
the ratio of shrinkage deformations occurring during the relaxation and vulnerable
solidification regimes. The importance of Tcr is clearly consensual in the literature but the
acceptance of the ratio R is more empirical since it comes from the index proposed by
Clyne and Davies270.
Conclusion:
The alloy RC0.5, with the lowest combined amount of Cu and Fe, presented the minimum
porosity area % (after A356 as reference alloy). The alloys RC3F0.7(M0) and RC3F0.7,
with the highest combined amount of Cu and Fe, experienced the maximum area% of
porosity. These results imply the direct correlation of microporosity with the Cu and Fe
contents of the alloys.
78
The hot tearing susceptibility of the alloys was evaluated experimentally by using a new
semi-quantitative indexation method called hot tearing sensitivity (HTS); which was
defined to reflect the volume of generated cracks in the torn specimens. Based on this
indexation, the studied alloys as the commercial foundry alloys are all resistant to hot
tearing; none of them were susceptible to hot tearing at higher mould temperature (>250 ).
At lower mould temperature, the alloys with the highest combined amount of Cu and Fe
(RC3F0.7(M0) & RC3F0.7) were the most prone to hot tearing, and the alloy containing
lowest Cu and Fe content (RC0.5) was the most resistant to hot tearing. Microstructure
analysis illustrated that the enhancement of hot tearing sensitivity by increasing Fe content
can be linked to an increased occurrence of lamellar β-
phase, which physically
block metal feeding. In order to better understand the effect of Cu (and Fe) on HTS and on
porosity area %, computational thermodynamic was done.
The multiphase back diffusion model1 was utilized to simulate the theoretical hot tearing
index (HCS) proposed by Katgerman271 for the alloys. A very good correlation was
obtained between the experimental hot tearing index (HTS) and the theoretical index
(HCS). Nevertheless, since the values of HTS were semi-quantitative, the HCS results were
compared with the area% of porosity of the alloys, as well. The correlation between HCS
and the area% of porosity was not very good, which implies the effect of another parameter
on area% of porosity of the alloys. Therefore, a new index (βR) was introduced, which
represents the ratio of solidification shrinkage (∆εv/∆εr) occurring during the vulnerable
time period (∆εv) and during the stress relief time period (∆εr). βR was strongly influenced
by the Cu and Fe contents of the alloys; the alloys with the highest combined amount of Cu
and Fe (RC3F0.7(M0) & RC3F0.7) illustrated the maximum βR values. Moreover, an
excellent correlation was found between βR and porosity area%; the correlation between βR
and HTS was also very good. These correlations indicate how the chemistry (Cu and Fe
content) of the alloys affect the HTS and the porosity area% by altering the ratio of
solidification shrinkage (∆εv/∆εr).
79
80
Chapter 5
.
“Evolution of Intermetallic Phases in Multicomponent Al-Si Foundry Alloys
Containing Different Cu, Mg and Fe Content”
Résumé:
L’effet de Cu et du Fe et les paramètres de traitement thermique de mise en solution (SHT)
sur l'évolution de la microstructure ont été étudiées. Les microstructures à l'état brut de
coulée et à l'état de traitement thermique de mise en solution ont été évaluées par
microscopie optique/électronique et par l’analyse calorimétrique différentielle à balayage
(DSC). Les évolutions de la microstructure ont été vérifiées par les calculs
thermodynamiques.
Les résultats (prévus et expérimentaux) ont démontré que la solubilité/stabilité de la phase
Q-Al5Cu2Mg8Si6 a été fortement influencée par la teneur en Cu. Par exemple, pour des
alliages d'aluminium à base de A356 et contenant de faible teneur en Cu (par exemple,
1,5%), le pic de DSC correspondant à la phase Q a disparu après 5 heures de traitement
thermique de mise en solution; cependant, dans les alliages contenant des teneurs élevées
en Cu (par exemple 3,4%), le pic de DSC a persisté à rester même après 20 heures de
traitement thermique de mise en solution. En outre, dans les alliages d'aluminium A356
contenant des teneurs élevées en Cu et en Fe, la durée du traitement thermique de mise en
solution conduit le Cu dissous à être graduellement perdu au profit de la phase NAl7Cu2Fe.
Abstract:
In this paper, the influence of Cu, Mg and Fe content on the microstructure evolution of AlSi based alloys has been studied. Initially, the as-cast microstructure of four Al-Si alloys
containing different Cu, Mg and Fe content was studied using differential scanning
calorimetry (DSC), optical and electron microscopy. Subsequently, the effect of different
solution heat treatments (SHT) on the microstructure evolution of the alloys was evaluated.
The microstructure evolutions after SHT were verified by thermodynamic calculations. The
results demonstrated that the dissolution of Q-Al Cu Mg Si phase was strongly dependent
on the Cu content of the alloy. That is, in 356 Al alloys containing low Cu content (e.g.
1.5%), the DSC peak corresponding to Q-phase disappeared after a SHT of 5 hours at 502
(935 F). However, in the alloys containing high Cu content (e.g. 3.4%), the peak was
still remaining even after 20 hours of SHT. In addition, the study also illustrated that in 356
Al alloys containing high Cu and Fe contents, longer solution treatment led the dissolved
Cu to be gradually lost to the N- Al7Cu2Fe phase.
Introduction
Excellent castability, better thermal conductivity and high strength to weight ratio make AlSi hypoeutectic alloys a suitable alternative for cast iron in the fabrication of engine
components (e.g. cylinder-heads)
42, 47, 84
. Hypoeutectic Al-Si alloys containing Cu and/or
Mg (e.g. 319 and 356) have been widely used in the automotive industry. The large eutectic
phases (θ-Al2Cu and Q-Al5Cu2Mg8Si6) that appear during solidification are generally
dissolved by applying an appropriate solution treatment, and are re-precipitated as fine
evenly distributed metastable phases to strengthen the alloys 37, 211.
The temperature and holding time period are the critical parameters of SHT. Lower
temperature/holding time-period might not be sufficient to dissolve the Cu- bearing
intermetallic phases. Higher SHT temperature (TSHT) can lead to incipient melting, which
deteriorates the mechanical properties due to void formation. Longer SHT not only
enhances the production costs, but can also lead the dissolved elements to be wasted on
other phases.
82
The temperatures at which the eutectic (θ and Q) phases can be melted while heating, are
required for the optimization of the SHT. Thermal analysis of the Al-Si-Cu-Mg alloys
6, 7
illustrated an endothermic peak occurring at ~507C (944F), which corresponded to Qphase. Melting of the θ-phase has been reported to start at about 525C (977F) 6, 7.
Solution heat treatment of Al-Si-Cu-Mg alloys is generally restricted to ~500C (932F) 196,
290
. Sokolowski et al. 197, 198 reported that the single step SHT of the Al-7Si-3.7Cu-0.23Mg
wt.%
alloy, which must be less than 500C (932F), is neither able to maximize the
dissolution of Cu rich intermetallic phases, nor is able to homogenize the microstructure
and modify the Si particle. As a result, they proposed a two-step SHT (e.g. 8 hours
@495C+ 2 hours @515C); by which the Cu bearing phases that solidified at the lowest
temperature could be dissolved at the first step
6, 10
. The second step of SHT helps the
dissolution of the remaining Cu-bearing intermetallic phase and further homogenisation of
the microstructure 6, 198.
Nevertheless, some authors reported fairly sluggish dissolution rates, or even the stability
of Q- Al Cu Mg Si phase at ~500C (932F) when the magnesium content is sufficiently
high 200; some others reported that Q-phase is stable during SHT due to its complex nature
199, 205
. Computational thermodynamic is useful to understand the stability/dissolution of Q-
phase at the corresponding SHT temperature. If the stability of Q-phase is as high as the
temperature of the first solution treatment step, the second step must be ignored. For alloys
in which Q-phase can be dissolved at the first step, the second step of SHT could further
homogenise the microstructure.
To minimize/eliminate un-dissolved Q-phase, Yan et.al.14 proposed that: TQ<TH<(TS-10C);
where TQ is the formation temperature of Q-phase, TH is the SHT temperature and TS is the
equilibrium solidus temperature. To satisfy this criterion, the alloy composition (mainly the
Cu and Mg contents) should be selected so that the formation temperature of Q-phase (TQ)
is lower than the equilibrium solidus temperature (TS) 14.
Solute atoms can be wasted to other phases during solution treatment. The presence of NAl7Cu2Fe phase has never been reported in the as-cast microstructure of Al-Si-Cu-Mg
alloys
11, 13, 287
; nevertheless, the transformation of β-Al5FeSi phase to N-Al7Cu2Fe phase
has been observed after SHT in few studies 11, 13, 192.
83
The major purpose of this work is to determine the evolution of Q-Al5Cu2Mg8Si6, θ-Al2Cu
and N-Al7Cu2Fe phases in Al-Si-Cu-Mg alloys. Four Al-Si alloys containing Cu, Mg and
Fe were investigated. The phases corresponding to each alloy were studied by DSC, optical
and electron microscopy. The solution heat treatments parameters were optimized to
maximize the dissolution of θ-Al2Cu, π-Al8Mg3FeSi6 and Q-Al5Cu2Mg8Si6 phases while
minimizing the loss of Cu into N-Al7Cu2Fe phase.
Experimental procedure
The alloy making and melting procedures were described in chapter 3 (section 3.1.2). The
specimens all were prepared by means of Pyrex tubes. The chemical composition and the
secondary dendrite arm spacing (SDAS) of the alloys are also given in Table 5-1.
Table 5-1: Chemical composition of the Al alloys (wt.%)
Alloy No.
A356
Reference (R)
#1
RC0.5
#2
RC1.5
#3
RC3
#4
RC3F0.7
Si
6.5-7.5
7.08
6.98
6. 90
6.98
Cu
0.20
0.54
1.5
3.38
3.1
Mg
0.25-0.45
0.30
0.30
0.35
0.33
Fe
0.20
0.12
0.10
0.12
0.77
SDAS (μm)
-12± 2
12± 2
13± 1
13± 1
Solution heat treatment was conducted in an electric resistance furnace. The temperature of
the first step of SHT was ~5C (9F) lower than the (non-equilibrium) solidus determined by
differential scanning calorimetry (DSC). For some alloys, the second step of SHT was
applied at a higher temperature. After SHT, the specimens were quenched in water to
assure maximum solute saturation. The specimens, which were solution heat treated at
different times/temperatures, were finally evaluated by means of DSC and electron probe
microanalysis (EPMA).
Samples for microstructural examination were mounted, ground and polished using
standard procedure. The polished sections were then studied with an optical microscope,
scanning
electron
microscopy
and
electron
probe microanalysis. Moreover, a
comprehensive study of the thermodynamic evaluation of Al-7Si alloys containing different
Cu, Mg and Fe content was carried out with the Thermo-Calc software using the TTAL7
database .
84
Results and discussion
5.3.1. As-cast microstructure
The as-cast microstructures of alloys (RC0.5) and (RC3) are presented in Figure 5-1. The
microstructure of alloy (RC0.5) was composed of soft α-Al dendrites, eutectic Si particles,
θ-Al2Cu phase, Q-phase and intermetallic Fe-containing phases (π- and β-phase). The
micro-constituents of alloy (RC1.5) were similar to alloy (RC0.5); but the higher Cu
content of this alloy promoted larger volume fraction of Cu bearing intermetallic phases (θ
and Q). The micro-constituents of alloy (RC3) were α-Al dendrites, eutectic Si particles, θAl2Cu phase, Q-phase and β-Al5FeSi phase. Since the chemical compositions of alloy
(RC3F0.7) and alloy (RC3) are similar; the same micro-constituents were observed in these
alloys. However, due to the higher Fe content, the size and distribution of the iron bearing
intermetallic phase (β-phase) was considerably larger in alloy (RC3F0.7).
Figure 5-2 illustrates the heating portion of DSC curves obtained with the set of 4 alloys in
their as-cast condition. The DSC curves were shifted vertically to avoid overlap. A welldefined peak (peak I) corresponding to the (non-equilibrium) solidus temperature of the
alloys can be seen in the DSC curves except for alloy (RC0.5).
Peak I, II and IV are well known peaks which correspond to the following reactions,
respectively:
Peak I: α-Al + Si + Al2Cu + Al5Cu2Mg8Si6 → Liquid
Peak II: α-Al + Al2Cu + Si → Liquid
Peak III, which appeared in alloys (356-3Cu) and (356Fe-3Cu), correlated with the reaction
below:
Peak III: α-Al+ N-Al7Cu2Fe + Si → Liquid + β-Al5FeSi
Peak IV: α-Al + Mg Si + π-Al8Mg3FeSi6 + Si → Liquid
85
Figure 5-1: As-cast microstructures of: a) alloy (RC0.5), b) alloy (RC3).
1.4
RC3F0.7
RC3
RC1.5
RC0.5
qDSC (mW/mg)
1.2
1
III
0.8
I
II
0.6
0.4
IV
0.2
500
520
540
560
T (˚C)
Figure 5-2: Heating DSC curves of the alloys in as-cast condition.
5.3.2. Microstructure of the solution treated specimens
5.3.2.1.
Alloy (RC0.5)
Figure 5-3(a) illustrates the heating DSC curves of alloy (RC0.5); one in the as-cast
condition and the other after SHT. According to ASTM B-917, six to twelve hours of SHT
at ~538C (1000F) is suitable for the A356 Al alloy
292
. As discussed earlier, Q- and θ‐
phases started to melt at around 507 and 525C (944 and 977F), respectively. The peaks (I
and II) corresponding to these phases in the alloy (RC0.5), disappeared during the heating
process in DSC. Nevertheless, two hours SHT at 502C (935F) was applied to insure the
entire dissolution of aforementioned phases (θ and Q). As shown in Figure 5-3(b), the π-
86
phase was still remaining after the first step of SHT. Subsequently, the second step of SHT
was conducted at 540C (1004F) for 8 hours. The second step helps further homogenization
of the microstructure and a more complete modification of the Si particles. Moreover, by
applying the second step of SHT, peak IV disappeared.
qDSC (mW/mg)
a)
RC0.5-SHT
RC0.5-AsCast
1
b)
0.9
π
0.8
0.7
IV
0.6
500
530
T (˚C)
560
Figure 5-3: a) Comparison of DSC curves of alloy (RC0.5) in as-cast and after solution treatment
(2h@502C+8h@540C), b) remaining of π-phase after the first step of solution treatment (2h@502C).
5.3.2.2.
Alloys (RC1.5) and (RC3)
Solution heat treatment of alloys (RC1.5) and (RC3) was conducted at 502C (935F). Figure
5-4 (a) illustrates the DSC curves of the alloy (RC1.5) in as-cast condition and after
different time-periods of SHT. There were no detectable peaks (I, II and III) corresponding
to the eutectic phases (Q, θ and π after the treatment. Worth to note that small amounts of
Q- and -phase could be found in the microstructure after 5 hours of SHT.
Figure 5-4 (b) illustrates the DSC curves of alloy (RC3) in as-cast condition and after SHT.
The area corresponding to peak I decreased with SHT, which indicates gradual dissolution
of the polynary eutectic phases (θ, Si and Q). Note that the area corresponding to peak II
was also decreased during the first step of SHT and almost disappeared after 10 hours of
SHT. Figure 5-5 illustrates the remnants of undissolved Q- and θ- phases after 8 hours of
SHT; particles of -phase were very tiny and dispersed in this microstructure. Figure 5-6,
which illustrates the EPMA results of alloy (RC3) after 15 hours of SHT, presents the
undissolved Q-phase in the microstructure. Even after 20 hours SHT at 502C (935F), a tiny
peak I was still observed.
87
AsCast
5h@502C
10h@502C
a)
qDSC(mW/mg)
0.8
1.7
qDSC (mW/mg)
0.9
0.7
0.6
I
II
IV
0.5
As-Cast
5h@502C
10h@502C
15h@502C
20h@502C
b)
1.5
1.3
1.1
I
0.9
II
0.7
0.4
0.5
0.3
500
550 T(°C)
0.3
500
550
T (˚C)
Figure 5-4: DSC curves of a) alloy (RC1.5) and b) alloy (RC3) after different solution treatment times at 502C (935F).
5.3.3. Time period of solution treatment
In alloy (RC1.5), the Q-phase can be dissolved at temperatures higher than 470C (878F).
Higher Cu content enhances the stability of Q-phase, such as that in alloy (RC3), where it is
stable up to 501C (933F). When the temperature of stability of the Q-phase is close to the
SHT, a long solution treatment can be required to dissolve it.
Longer SHT can cause coarsening of the microstructure and the loss of solute Cu to the NAl7Cu2Fe phase. The latter case might not be appreciable in Al alloys containing low
amounts of β-phase (e.g. alloy RC3); but in Al alloys with large volume fraction of β-phase
(e.g. alloy RC3F0.7), this might result in a significant loss of Cu in the primary α-Al phase.
Figure 5-7 illustrates the calculated mass fraction of phases at equilibrium for alloys (RC3)
and (RC3F0.7). One can see that a significant amount of N-Al7Cu2Fe phase is produced at
the expense of the -phase as the temperature decreased, thus reducing the Cu content in
the primary phase. This loss in Cu is particularly important when the alloy contains a high
level of Fe, like in alloy (RC3F0.7).
Worth to mention that the amount of N-phase tends to decrease at equilibrium as the
temperature increases; so applying a SHT at a higher temperature should help to reduce the
volume fraction of N-phase.
88
Figure 5-5: Remnant of θ‐ and Q- phase in alloy (RC3) after 8 hours of solution treatment.
To further evaluate the effect of SHT on the precipitation of N-phase, the alloy (RC3F0.7)
was solution treated at 502C (935F) for different time periods (5, 10 and 20 hours) and
evaluated by EPMA. Figure 5-8 illustrates the area fraction of intermetallic phases
containing Cu (i.e. N-, θ- and Q-phases) in alloy (RC3F0.7), as measured with EPMA
elemental mappings. Five hours of SHT reduced the area fraction to a minimum because of
the dissolution of θ- and Q-phases. Longer SHT increased the area fraction of Cu
containing phases due to the precipitation of N-phase. Moreover, this alloy was evaluated
by DSC in as-cast and solution treated conditions. The area under DSC curves
corresponding to peaks II and III, which were integrated after plotting a straight line, are
shown in Figure 5-8. By increasing the SHT time-period, the area under DSC curve
increased. This implies an increasing volume fraction of N-phase as SHT proceeds.
89
Figure 5-6: Remnant of Q-phase in alloy (RC3) after 15 hours of solution treatment at 502C (935F).
a) Al-7Si-3.5Cu-0.35Mg-0.15Fe
Si
Liq
0.06
Al
0.05
0.04
θ: Al2Cu
Q: Al5Cu2Mg8Si6
β: Al5FeSi
Si: Silicon
N: Al7Cu2M
Liq: Liquid
Al: α-FCC
θ
0.03
0.02
b) Al-7Si-3.5Cu-0.35Mg-0.75Fe
0.07
Si
Mass Fraction
Mass Fraction
0.07
0.06
0.05
0.04
0.03
β
Al
θ
0.01
0
300
400
0.01
β
Q
500
600
700
T (°C)
θ: Al2Cu
Q: Al5Cu2Mg8Si6
β: Al5FeSi
Si: Silicon
N: Al7Cu2M
Liq: Liquid
Al: α-FCC
N
0.02
N
Liq
Q
0
300
400
500
T (°C)
600
700
Figure 5-7: Calculated mass fraction of phases vs. temperature (in equilibrium state) for Al-7Si-3.5Cu0.35Mg containing a) 0.15 and b) 0.75 wt % Fe.
90
Area under peak II and III in DSC curves (mW/mg)
5.5
Area fraction (%) of Cu phases
2
3.5
1
Area fraction (%) of Cu phases
Area under DSC curve (mW/mg)
3
1.5
0
5
10
15
20
Time period (hour) of solution treatment at 502C (935.6F)
Figure 5-8: Area fraction of intermetallic phases containing Cu (EPMA mappings), and the area under DSC
curves (mW/mg) corresponding to peaks II and III in alloy (RC3F0.7).
5.3.4. High temperature solution heat treatment
To evaluate the effect of a high temperature SHT on the microstructure, alloys (RC3) and
(RC3F0.7) were solution heat treated during 5 hours at 535C (995F). The presence of
massive eutectic (θ and Q) phases nearby polygonal Si particles is the major characteristic
of incipient melting (see Figure 5-9). Figure 5-10 illustrates that all of the peaks observed
in the as-cast condition exist with more or less the same energy after 5 hours of solution
treatment at 535C (995F). Therefore, the micro-constituents which were locally melted after
SHT did not diffuse into Al matrix; instead, they were re-precipitated upon quenching with
a massive form.
91
siv
(θ e eu
+Q te
) ctic
M
as
Po
ly
go
na
l
Si
Si
Figure 5-9: Incipient melting after five hours of solution treatment of alloys (RC3F0.7) at 535C (995F). RC3- AsCast
RC3-5h@535C
RC3F0.7- AsCast
1.8
qDSC (mW/mg)
RC3F0.7-5h@535C
III
I
II
0.8
500
510
520
530
540
550
T (˚C)
Figure 5-10: DSC curves of alloys (RC3) and (RC3F0.7) in as-cast condition and after solution treatment
(5h@535C).
5.3.5. Stability of Q-phase
The Q-phase was easily dissolved during heating in alloy (RC0.5); but in alloy number
(RC3), it was still stable even after a few hours of solution heat treatment. It seems that the
stability of Q-phase is strongly dependent on the chemical composition of the alloy.
To evaluate the effect of different elements on the stability of Q-phase, the dissolution
temperature of Q-phase for Al-Si alloys containing various Si, Fe, Cu and Mg content was
calculated with Thermo-Calc (see Figure 5-11). The results demonstrate that the stability
92
of Q-phase is independent from the Si and Fe content. Nevertheless, it confirms the strong
influence of Mg and Cu content on the stability of Q-phase; for instance, by increasing Cu
content from 1 to 3.5% in Al-7Si-0.35Mg-0.25Fe, the stability enhances from 453 to 503C (847 to 937F).
Figure 5-12 presents the calculated isothermal section of Al-7Si-xCu-xMg-0.15Fe at 500C (932F). The highlighted area illustrates the regions wherein Q-phase would be stable with a
SHT at 500C (932F). For some chemical compositions, the stability of Q-phase could be
even higher than the equilibrium solidus; however, the SHT must be limited to nonequilibrium solidus
189, 293
. Mohamed et al.
194
reported that for an Al-6.6Si-3.2Cu alloy
containing ~0.3%Mg, a two-step SHT (8h@505C+2h@520C) caused better mechanical
properties, but for the alloys containing ~0.6% Mg, a single step SHT (8h@505) was
recommended. For the Al-Si-Cu alloys containing lower Mg content (~0.3%), the Q-phase
can be dissolved at T 485C (905F); but for the alloys containing 0.6% Mg content, the Qphase is stable up to 517C (962F). The volume fraction of Q-phase in as-cast condition and
after the SHT (8hrs@490C+4hrs@500C) was reported to be 2.11 and 2.19%, respectively,
in an Al-7Si-3.5Cu-0.6Mg alloy
205
. This indicates how Q-phase is stable in Al-Si-Cu
system when the Mg content is 0.6% and above. The effect of Mg content on the stability
of Mg-bearing intermetallics (e.g. Q and π) is explained with further details in the chapter
(7).
93
Al-xSi-3.5Cu-0.35Mg-0.25Fe (x=5-10)
Al-7Si-xCu-0.35Mg-0.25Fe (x=0.8-5)
Al-7si-3.5Cu-xMg-0.25Fe (x=0.15-0.8)
Al-7Si-3.5Cu-0.35Mg-xFe (x=0.15-1)
Dissolution T of Q (°C)
Mg
520
Fe
Si
Cu
500
480
460
440
420
400
(wt. %)
Figure 5-11: Variation of the dissolution temperature of Q-phase with chemical composition (calculated with
Thermo-Calc).
0.9
Al
+M
+S
i
π+
l+S
i
l+S
+A
β+
A
+β
Q+
%)
(M
ass
θ: Al2Cu
Q: Al5Cu2Mg8Si6
M: Mg2Si
β: Al5FeSi
π:Al8FeMg3Si6
Q
θ+
Mg
π+Al+Si
π+β+Al+Si
0.4
i
0.5
+S
Al
0.6
π+
0.7
Q+
0.8
i
0.3
θ+Al+Si
0.2
0.1
0
1
2
3
θ+β+Al+Si
4
5
6
Cu (Mass %)
7
8
9
Figure 5-12: Phase field distribution of Al-7Si-xCu-xMg-0.15Fe at 500C (932F), calculated by ThermoCalc).
94
Conclusion
1. The stability of Q-phase is strongly dependent on Cu content. In alloys containing
low Cu content (e.g. alloy RC1.5), the peak I corresponding to Q-phase disappeared
after 5 hours of solution treatment; but in alloy (RC3), it remained even after 20 hours
of solution treatment.
2. The transformation of β-Al5FeSi to N-Al7Cu2Fe during solution heat treatment leads
some part of the dissolved Cu in Al matrix to be wasted. The amount of Cu not
available to strengthen the primary phase increases with the volume fraction of βAl5FeSi.
3. A long solution heat treatment may promote the dissolution of Q-phase, but on the
other hand, it also promotes the growth of N-Al7Cu2Fe phase. In a solution heat
treatment of Al-Si-Cu-Mg alloys containing high Fe content, a compromise between
the growth of N-phase and dissolution of Q-phase is required.
4. The presence of massive eutectic (θ and Q) phases nearby polygonal Si particles was
the major characteristic of specimens having experienced incipient melting.
95
96
Chapter 6
.
“Assessment of Post-Eutectic Reactions in Multicomponent Al-Si Foundry Alloys
Containing Cu, Mg and Fe”
Résumé:
L’effet de la composition chimique des alliages Al-Si (Cu, Mg et Fe contenu) et des
paramètres de traitement thermique de mise en solution (SHT) sur l'évolution de la
microstructure ont été minutieusement étudiées. Les microstructures à l'état brut de coulée
et à l'état de traitement thermique de mise en solution ont été évaluées par microscopie
optique/électronique pour étudier les réactions post eutectiques. L’analyse calorimétrique
différentielle à balayage (DSC) a été utilisée pour examiner les transformations de phase
survenant au cours du processus de chauffage et de refroidissement. Les calculs
thermodynamiques ont été effectués pour évaluer la formation de la phase à l'état
d'équilibre et hors-équilibre. La phase Q-Al5Cu2Mg8Si6 a été solidifié soit à la même
température ou plus tôt que la phase θ-Al2Cu en fonction de la teneur en Cu de l'alliage.
Deux morphologies des microconstituants Al-Cu ont été observées dans la microstructure
de coulée: soit le format eutectique et le format bloc. Puisque le microconstituant en forme
de bloc contenait toujours une certaine teneur en Fe, il est appelé ci-après AlCuFe
intermétallique. Bien que l’AlCuFe-intermétallique a été à peine observée dans la
microstructure de coulée, la réaction de l'α-Al avec la phase β-Al5FeSi est à l’origine de la
formation de la phase N-Al7Cu2Fe au cours du traitement thermique de mise en solution.
L'effet de la teneur en Cu sur la température de solidification de la phase π-Al8Mg3FeSi6 a
également été étudié.
Abstract:
Post-eutectic reactions occurring in Al-Si hypoeutectic alloys containing different
proportions of Cu, Mg and Fe were thoroughly investigated in this work. As-cast
microstructures were initially studied by optical and electron microscopy to investigate the
microconstituents of each alloy. Differential scanning calorimetry (DSC) was then used to
examine the phase transformations occurring during the heating and cooling processes.
Thermodynamic calculations were carried out to assess the phase formation in equilibrium
and in non-equilibrium conditions. The Q-Al5Cu2Mg8Si6 phase was predicted to precipitate
from the liquid phase, either at the same temperature or earlier than the θ-Al2Cu phase
depending on the Cu content of the alloy. Two morphologies of Al-Cu intermetallics were
found in the as-cast microstructure: eutectic-like and bloc-like morphologies. Since the
block-like morphology contained some Fe content, it is entitled hereafter AlCuFe
intermetallic. However, the AlCuFe- intermetallic was barely observed in the as-cast
microstructure, the reaction of α-Al with the β-Al5FeSi phase caused the formation of the
N-Al7Cu2Fe phase during solution heat treatment. Thermodynamic calculations and
microstructure analysis helped to determine the DSC peak corresponding to the melting
temperature of the N- Al7Cu2Fe phase. The effect of Cu content on the solidification
temperature of π-Al8Mg3FeSi6 is also discussed.
Introduction
Al-Si based foundry alloys have become a suitable alternative for cast iron in the
fabrication of engine components (e.g. cylinder-heads) in recent years. Better thermal
conductivity and high strength to weight ratio are two main advantages of the Al-Si
hypoeutectic alloys. The major Al-Si alloys in the fabrication of engine components can be
classified into two main categories: 1) Al-7Si-3Cu alloys (e.g. as A319) containing Mg (<
0.4 wt.%)
85
87, 264, 265
; and Al-7Si-0.3Mg alloys containing Cu (e.g. A356+0.5Cu-wt.%) 47, 84,
. Copper and Mg play a vital role in the strengthening of Al-Si alloys even though Mg
addition could have a negative effect on high temperature mechanical properties of 319type Al alloys
37, 233
. However, the addition of Mg is required to improve the mechanical
properties at room temperature
98
264, 265
. The Al–Si alloys can be used as primary Al alloys
(Fe < 0.2 wt.%) or secondary Al alloys (with Fe content up to 1 wt.%) 85, 122. Further details
of the effect of these elements (Cu, Mg and Fe) on cast Al-Si alloys have been thoroughly
reviewed in a recent publication (chapter 2) 294.
To maximize the efficiency of strengthening, the as-solidified large eutectic phases (e.g. θAl2Cu and Q-Al5Cu2Mg8Si6) must be dissolved and re-precipitated by applying the
appropriate heat treatment 37, 195, 211. The temperature(s) and reaction(s) of the post-eutectic
phases during the last stage of solidification are critical parameters in the optimization of
the solution heat treatment (SHT). In Al-Si alloys containing Cu and Mg, the last solidified
eutectic reaction normally involves the θ‐ and Q-phases and was reported to occur at ~780 K 507
6, 7. However, the effect of chemical composition (e.g. Cu and Mg content) on
the precipitation/melting temperature of the θ‐ and Q-phase is not clear because of the
complexity of the system. There is some controversy regarding the precipitation of the θphase in the literature about the temperature where this phase appears for the first time
during solidification. Mulazimoglu et al.8 reported that the precipitation of the θ-phase
occurs at ~822 K (549
10
in 319.2 foundry alloy. This was neither confirmed by Samuel 9,
nor by other authors6, 7, via solidification or reheating experiments. Instead, it has been
reported that the θ-phase can grow with two distinct morphologies
9, 15, 295
; eutectic-like
morphology (with Cu concentration of ~28 wt.%) and block-like morphology (with Cu
concentration of ~40 wt.%). The DSC heating curves obtained on the 319 Al alloy
indicated two endothermic peaks (during heating), one at ~793 K (520
at ~806 K (533
and another one
, which were respectively ascribed to the melting of the eutectic-like and
block-like θ- Al2Cu phase
9, 10, 15, 293
. It is worth noting that, during solidification, the
occurrence of a peak ascribed to the formation of the block-like θ-Al2Cu phase
(corresponding to the peak at 806 K (533
in heating) has never been reported.
Because iron is a common impurity in aluminium alloys and is almost insoluble in the
primary phase, a variety of iron-bearing intermetallic phases can be found in the
microstructure. In Al-Si-Mg-Fe foundry alloys, the iron intermetallic phases: β-Al5FeSi and
π-Al8FeMg3Si6, are frequently observed. The latter could be entirely/partially dissolved
during SHT. Therefore, the precipitation/dissolution temperature of this phase and the
effect of alloying elements can play a vital role in the optimization of the SHT. The
precipitation/dissolution temperature of the phase π-Al8FeMg3Si6, has been reported to be
99
about 827 K (554
in Al-Si-Mg/Cu alloys 7, 140. The effect of Mg on this temperature has
been thoroughly investigated
7, 140, 270, 296
, but there is a dearth of information pertaining to
the influence of Cu on the precipitation/dissolution of this phase.
Addition of copper brings other intermetallic phases to the microstructure if the kinetics
conditions are favourable. Indeed, the presence of the Al7Cu2Fe phase has never been
reported in the as-cast microstructure of Al-Si-Cu-Mg alloys, but it has been observed in
the solution heat treated condition
11, 13, 195
. In a DSC analysis made on the solution heat
treated Al-7Si-3Cu-0.3Mg-0.8Fe alloy, the peak occurring at 795 K (522
during heating
was supposed to be caused by the following reaction (chapter 5) 297:
(α-Al+ N-Al7Cu2Fe+ Si → Liquid + β-Al5FeSi).
The Al7Cu2Fe phase, which is sometimes entitled as β(FeCu) or N-phase, has a broad
composition range from 29 to 39 wt.% Cu and from 12 to 20 wt.% Fe
90, 124
. The presence
of AlCuFe-intermetallic in the solution treated specimens of Al-Si foundry alloys has only
been reported in a few studies
11-14
, but the detail of the phase transformation, its effect on
thermal analysis and the effect of chemical composition has never been studied.
The major purpose of this work was to study the effect of Cu, Fe and Mg content on posteutectic reactions occurring in Al-Si foundry alloys. This study was undertaken to elucidate
the reactions involved during solidification and reheating; the latter giving some indications
of what happens during the SHT. Seven different Al-7Si based alloys containing various
Cu, Fe and Mg content were investigated. The alloys were initially studied by DSC, optical
and electron microscopy. Particular attention was paid to observe the products formed by
the transformations occurring at the beginning of the reheating cycle.
Moreover, a comprehensive study of the thermodynamic prediction of the microstructure
evolution in Al-7Si alloys containing different Cu, Mg and Fe content was carried out with
the Thermo-Calc software
298
using the TTAL7 database . the multiphase back diffusion
(MBD) model1 was used to calculate the phase precipitations and their mass fraction during
solidification.
100
Experimental Procedure
The alloy making and melting procedures were detailed in chapter 3 (section 3.1.2). The
average chemical compositions of the seven alloys investigated are presented in Table 6-1.
All micrographs taken from the specimens cast in the permanent mould and presented in
this paper were indicated in the figure captions. In all other cases, the specimens came from
the metal sampled with the Pyrex tubes.
Table 6-1: chemical composition of the alloys (wt.%)
Alloy No.
Si
Cu
Mg
Fe
Al
Ref. A356.0
6.5-7.5 0.2
0.25-0.45 0.2
Bal.
#1
RC0.5
7.08
0.54 0.30
0.12
Bal.
#2
RC0.5F0.7
7.18
0.51 0.31
0.79
Bal.
#3
RC1.5
6.98
1.5
0.3
0.10
Bal.
#4
RC3
6. 9
3.38 0.35
0.12
Bal.
#5
RC3F0.7
6.98
3.1
0.33
0.77
Bal.
#6
RC3(M0)
6.9
3.3
0
0.13
Bal.
#7
RC3F0.7(M0) 6.94
3.08 0
0.74
Bal.
“R” indicates the reference alloy. The symbols “C”, “F” and “M”
represent the Cu, Fe and Mg elements; the number after each symbol
presents the concentration of the respective element. Alloys #6 and #7
contain zero Mg which were characterized by “(M0)”.
The SHT was conducted in an electric resistance furnace. The temperature of the SHT was
~5 K lower than the solidus determined by differential scanning calorimetry (DSC). After
this treatment, the specimens were quenched in water to obtain maximum solute saturation.
The specimens, which were solution treated at different times/temperatures, were finally
evaluated by means of DSC and electron probe microanalysis (EPMA).
Samples for microstructural examination were mounted, ground and polished using
standard procedure. The polished sections were then studied with an optical microscope,
scanning electron microscopy and electron probe microanalysis. A scanning electron
microscope (SEM, JEOL JSM-6480LV) equipped with an electron backscattered
diffraction (EBSD) pattern acquisition camera and Channel 5 software
299
, were used to
confirm the crystallographic structure of iron-bearing intermetallics. Moreover, a
comprehensive study of the thermodynamic evaluation of Al-7Si alloys containing different
Cu, Mg and Fe content was carried out with the Thermo-Calc software and with the
multiphase back diffusion (BDM)1 model .
Results and discussion
101
6.3.1. Microstructure of the alloys
The as-cast microstructures of some alloys (RC0.5, RC3, RC3F0.7 and RC3(M0)) are
presented in Figure 1. The eutectic Si phase is dark gray in colour, but the intermetallic
particles are brighter; they are mostly concentrated in the interdendritic regions.
Microconstituents of the alloys containing 1.5 wt.% Cu or less (RC0.5, RC0.5F0.7 and
RC1.5) are very similar (Fig. 1a) and they are composed of α-Al dendrites, eutectic Si
particles, the θ-Al2Cu phase, Q-phase and Fe-containing intermetallic phases (π- and βphase). The major difference between the microstructure of the aforesaid alloys is the
varying volume fraction of Cu and Fe bearing intermetallic phases which are enhanced by
increasing Cu and Fe content. The microconstituents of alloy #4 (RC3) is comprised of αAl dendrites, eutectic Si particles, the θ-Al2Cu phase, Q-phase and β-Al5FeSi phase (Fig.
1b). Since the chemical composition of alloy #5 (RC3F0.7) is similar to that of alloy #4
(RC3) except for the Fe content, there was no difference in the microconstituents of these
alloys. Due to the higher Fe content, the size and distribution of the iron bearing
intermetallic phase (β-phase) was considerably larger in alloy #5-RC3F0.7 (Fig. 1c), but in
alloy #4 (RC3), it was hardly visible. The major microconstituents of alloy #6 and #7
(RC3(M0) and RC3F0.7(M0)) are the same (Fig. 1d), consisting of α- Al dendrites, eutectic
Si particles and the θ-Al2Cu and β-Al5FeSi phase.
The predicted mass fractions of the phases during the solidification process are presented in
Figure 6-2. The mass fraction of the post eutectic phases in alloys #2-RC0.5F0.7, #5RC3F0.7 and #7-RC3F0.7(M0) are similar to the one in alloys #1-RC0.5, #4-RC3 and #6RC3(M0) respectively; and therefore, their curves are not presented here. The θ-phase is
predicted to precipitate in all of the alloys; the Q-phase is present in all of the Mg
containing alloys (#1 to #5) and the π-phase is formed in the alloys containing 1.5 wt.% Cu
or less (#1 to #3). A small amount of the N-phase was predicted to precipitate in the alloys
containing high Cu content (#4 to #7).
102
Figure 6-1: as-cast microstructures of the experimental alloys: a) alloy#1 (RC0.5), b) alloy#4 (RC3), c)
alloy#5 (RC3F0.7), d) alloy#6 (RC3(M0)).
103
Figure 6-2: mass fraction of phases vs. temperature as calculated with the MBD model 1.
The predicted solidification temperatures of the post eutectic phases are presented in Table
6-2. In all cases, the θ-phase appears at the last stage of solidification. The solidification
temperature of the Q-phase is influenced by the Cu content. In the alloys containing high
Cu contents (RC3, RC3F0.7), the Q-phase is predicted to solidify along with the θ-phase
(i.e. at ~783 K 510
, but in the alloys containing lower Cu contents (RC0.5, RC0.5F0.7
and RC1.5), it solidifies earlier (i.e. at 810 and 799 K (537 and 526
in alloys RC0.5 and
RC1.5, respectively).
Table 6-2: the solidification temperatures K ( ) of the post eutectic phases predicted by the MBD model 1.
#3-RC1.5 #4-RC3,
#6-RC3(M0),
Alloy No. #1-RC0.5,
Phase
θ-Al2Cu
Q-Al5Mg8Cu2Si6
N-Al7Cu2Fe
π-Al8Mg3FeSi6
104
#2-RC0.5F0.7
783 (510)
810 (537)
-823 (550)
783 (510)
799 (526)
-805 (532)
#5-RC3F0.7
#7-RC3F0.7(M0)
783 (510)
783 (510)
791 (518)
--
797 (524)
-803 (530)
--
6.3.2. Thermal analysis of as-cast specimens
Figure 6-3 illustrates the DSC curves recorded during the heating of all the studied alloys in
as-cast condition. The DSC curves were shifted vertically to avoid overlap. Determination
of the solidus temperature (the temperature at which the last solidified eutectic is melted
while heating) by DSC analysis helps to specify the upper limit of the SHT temperature.
The non-equilibrium solidus temperatures of the alloys, calculated by the multiphase backdiffusion model, are also given in Figure 6-3. A well-defined peak corresponding to the
solidus temperature of the alloys, can be seen in the DSC curves except for the alloys
containing 0.5 wt.% Cu (#1-RC0.5, #2-RC0.5F0.7).
Peak I appeared at ~780 K (507
in alloys #3 to #5 (RC1.5, RC3 and RC3F0.7). This is a
critical temperature in the SHT of Al-Si-Cu-Mg alloys. According to the literature, this
peak corresponds to: (α-Al+ Si+ Al2Cu+ Al5Mg8Cu2Si6 ↔ liquid) 6, 7, 123. Peak II appeared
in the alloy containing 1.5 wt.% Cu or more (#3 to #7). This peak generally correlates with:
(α-Al+ Si+ Al2Cu ↔ liquid)
123, 141, 300
. However, according to the MBD model, the
temperature and sequence of the precipitation of the Q- and θ-phases are both affected by
the Cu and Mg contents. This will be elaborated on with more details in section (6.3.4).
Peak III appeared in the alloys containing 3 wt.% Cu or more (alloys #4 to #7, Figure 6-3a,
b). For the alloys containing low Fe content (#4-RC3 and #6-RC3(M0)), this peak was tiny
and masked by peak II, but in the alloys containing high Fe content (#5-RC3F0.7 and #7RC3F0.7(M0)) it was intense enough to be distinguished from peak II. The onset
temperature of this peak (III) varies with the Mg content of the alloys. In the Mg containing
alloys (RC3 and RC3F0.7), it appeared at ~(795 K) 522 ; but in the alloys free of Mg
(RC3(M0) and RC3F0.7(M0)), it occurred at ~805 K (532
. The predicted precipitation
temperatures of the N-phase, as illustrated in Figure 6-2 and listed in Table 6-2, are close to
the aforementioned DSC temperature. According to Samuel et al.
9, 10, 15
, peak III
corresponds to the melting of the blocky θ-Al2Cu phase in the alloy Al-7Si-3Cu. Further
analysis, which was carried out to correlate the appropriate reaction(s) to this peak, is
presented in the next section (6.3.3).
Peak IV was observed in the alloys containing 1.5 wt.% Cu and/or less (RC0.5, RC0.5F0.7
and RC1.5, Figure 6-3c, d). It appeared at ~817 K (544
in alloys #1 and #2 (RC0.5,
105
RC0.5F0.7). This peak can be assigned to the reaction: (α-Al+ Mg2Si+ π‐Al8Mg3FeSi6+ Si
↔ liquid)
7, 140, 296
. The predicted mass fraction of the Mg2Si phase was negligible and no
evidence of the Mg2Si phase in the microstructure was detected. But the π- Al8Mg3FeSi6
phase was easily observed in the as-cast microstructure of the alloys containing 1.5 wt.%
Cu or less (RC0.5, RC0.5F0.7 and RC1.5).
The cooling DSC curves of alloys #3 to #5 (RC1.5, RC3 and RC3F0.7) are presented in
Figure 6-4. The temperatures and the numbers of peaks in the cooling regime were different
compared to those identified in the heating regime. The peaks occurred at slightly lower
temperatures than those occurring during heating. Some of the peaks, which were seen in
the heating curves, were merged together and/or disappeared. In alloys #4 and #5 (RC3 and
RC3F0.7)), peak I and II occurred at almost the same temperature (at ~774 K (501 )),
however, in alloy #3, the peaks occurred at 768 K (495 ) and 785 K (512
Peak III was not seen in the cooling curves of the alloys.
106
, respectively.
Figure 6-3: DSC heating curves of the as-cast specimens with scanning rate of 10K/min. The solidus
temperatures (Ts) given above were calculated with the MBD model 1.
107
Figure 6-4: DSC cooling curves of alloys #3 to #5 (RC1.5, RC3 and RC3F0.7) with a scanning rate of 5
K/min. The starting temperature of the DSC cooling tests was 933 K (660 .
6.3.3. The N-phase
Two morphologies of Al-Cu microconstituents (i.e. eutectic-like and block-like), which are
shown in Figure 5, were observed in the as-cast microstructure. The concentration of Cu in
the AlCu eutectic-like microconstituent was between 30 to 38 wt.%. In the block-like
microconstituent, the concentration of Cu was between 38 to 45 wt.%. The block-like
microconstituent usually contained some Mg, Si and significant Fe content; the content of
Fe varied from 1 to 12.5 wt.%. In some cases, the stoichiometry of the block-like
microconstituent (Al6Cu2Fe0.7Si0.3), was close to the N-phase (Al7Cu2Fe). The Cu
concentration in the block-like microconstituent is comparable with the result reported by
Samuel15 for their blocky θ-phase, however, the presence of Fe in the blocky
microconstituent has never been reported in the literature. Since the block-like
microconstituent always contained some Fe content inside, hereafter, in this paper it will be
called AlCuFe-intermetallic.
108
Figure 6-5: morphology of the AlCu-eutectic and block-like AlCuFe- intermetallic phases in alloy RC3
(prepared with the permanent mould).
As mentioned earlier, Samuel et al. correlated peak III to the melting of the “blocky θphase” 9, 10, 15. To clarify which phases melt during the reactions producing peak II and peak
III, two as-cast specimens of alloy #7 (RC3F0.7(M0)) were placed in the DSC sample pan
and were respectively heated up to 800 K (527
, the temperature just after peak II) and up
to 810 K (537 , the temperature just after peak III). In both cases, they were rapidly
cooled after the reaction. To track the evolution of the microstructure, the micrographs of
the specimen before (in as-cast condition) and after the heat treatment were compared at the
same location. As shown in Figure 6-6, by heating the specimen up to 800 K (527
, the
AlCu eutectic microconstituent either melted or disappeared (transformed/ dissolved),
while AlCuFe- intermetallic were easily found in the microstructure. It is worth noting that
these microconstituents (AlCu eutectic and AlCuFe- intermetallic) shown in the optical
micrographs were verified with EPMA. The micrographs presented in Figure 6-7 show that
by heating the specimen up to 810 K (537
, the AlCu eutectic microconstituent and the
AlCuFe- intermetallic were both almost completely melted or disappeared. Therefore, peak
II seems to correspond to the melting of the AlCu eutectic microconstituent and peak III to
the melting of the AlCuFe- intermetallic.
109
Figure 6-6: a) alloy #7 (RC3F0.7(M0)) in as-cast condition; b) alloy #7 heated up to 800 k (527 , just beyond
peak II) and rapidly cooled (a and b were taken at the same location).
Figure 6-7: a) alloy #7 (RC3F0.7(M0)) in as-cast condition; b) alloy #7 heated up to 810 K (537 , just
beyond peak III) and rapidly cooled (a and b were taken at the same location).
The SHT of the alloys containing high Cu and Fe content (RC3F0.7 and RC3F0.7(M0))
helped to correlate peak III to the melting of the AlCuFe- intermetallic with more
confidence. The predicted mass fraction of the N-phase in alloys #5 and #7 (RC3F0.7 and
RC3F0.7(M0)) at solidus temperature is negligible (~0.08 and ~0.14% respectively), but in
the equilibrium condition it can be significantly enhanced (up to ~5%). Figure 8 presents
the EPMA elemental mapping of the as-cast microstructure of alloy #5 (RC3F0.7). This
mapping confirms that the area fraction of the AlCuFe- intermetallics is negligible, as
predicted by the MBD model. In order to evaluate the AlCuFe- intermetallic content in the
equilibrium state, alloys #5 and #7 (RC3F0.7 and RC3F0.7(M0)) were solution heat treated
at 775 K (502 ) for different time periods. Figure 6-9 illustrates the EPMA results for
alloy #5 (RC3F0.7) after a 15 min. SHT. As illustrated, the area fraction of the AlCuFeintermetallics has been considerably increased even after such a short time period of SHT.
110
The area fraction of the AlCuFe- intermetallics was correlated with the SHT time period, in
such a way that by increasing the time period, the area fraction was significantly enhanced.
Figure 6-10 illustrates the elemental mapping of alloy #5 (RC3F0.7) after a 20 hour SHT at
775 K (502 ). As shown in this figure, almost all of the Q and -phases in the as-cast
microstructure were dissolved, while the area fraction of the AlCuFe-intermetallics (mostly
the N-phase) was significantly enhanced. The DSC curves of alloy #5 (RC3F0.7) in the ascast and solutionized conditions, which are illustrated in Figure 6-11, confirm the EPMA
results, in such a way that peaks I and II got smaller by increasing the SHT time period and
peak III got enlarged. After 10 hours of the SHT, peaks I and II almost disappeared and
peak III got much larger. The conformity of the EPMA results with the DSC results implies
that peak III occurring in the heating regime, corresponds to the melting of the AlCuFeintermetallics (mostly the N-phase). The peak ascribed to the formation of the AlCuFeintermetallic particles during solidification was too shallow to be seen, likely because of the
very low mass fraction of the AlCuFe- intermetallics which were formed (see Figure 6-4).
Figure 6-8: elemental mapping of alloy #5 (RC3F0.7) in as-cast condition.
111
Figure 6-9: elemental mapping of alloy #5 (RC3F0.7) solution heat treated 15 min. at 775 K (502
quenched.
and
Figure 6-10: elemental mapping of alloy #5 (RC3F0.7) solution heat treated 20 hours at 502 .
Figure 6-11: DSC curves of alloy #5 (RC3F0.7) in as-cast and solution heat treated conditions; scanning rate
10 K/min.
112
Few authors have mentioned the appearance of peak III in the alloys Al-7Si-3Cu 9, 10, 15. All
of the aforementioned references stated that this peak corresponds to the melting of the
blocky θ-Al2Cu phase. Zolotorevsky et al.
90
reported the occurrence of the following
peritectic reaction in Al-Cu-Fe-Si alloy systems: (α-Al+ N-Al7Cu2Fe+ Si ↔ liquid+ βAl5FeSi). However, due to the low time available during non-equilibrium solidification, it
seems that the peritectic reaction cannot be completed during solidification. Therefore, in
the as-cast condition, the AlCuFe- intermetallic did not generally meet the stoichiometry of
the N-Al7Cu2Fe phase. One can assume however, that the AlCuFe- intermetallic was a
predecessor of the N-Al7Cu2Fe phase. It was only after applying the appropriate SHT, that
the Al, Cu and Fe contents in the AlCuFe- intermetallic generally reached to be 52.34, 33.9
and 13.21 wt.% (7.1, 2 and 0.9 at.%), respectively, meeting the stoichiometry of the NAl7Cu2Fe phase.
In order to confirm the crystallographic structure of the N-phase, the solution heat treated
specimens (8 hours at 775 K (502
) of alloy #5 and #7 (RC3F0.7 and RC3F0.7(M0))
were verified by EBSD analysis. The EBSD patterns and simulation results of the NAl7Cu2Fe phase are shown in Figure 6-12. Figure 6-12(b) is the indexed experimental
EBSD patterns for the N-Al7Cu2Fe phase and Figure 6-12(c) is the simulation results
calculated by the Channel 5 software. In EBSD analysis, the accuracy of the solution
provided by the software is presented by the mean angular deviation (MAD) between the
experimental and calculated patterns; a smaller MAD value indicates a closer match
between the experimental and simulated Kikuchi bands. For an accurate solution, the MAD
value must be lower than 0.7
301, 302
. As illustrated in this figure, the MAD value is 0.2,
which confirms the accuracy of the solutions obtained for the N-Al7Cu2Fe phase.
As shown in Figure 6-13(a), the N-Al7Cu2Fe phase can hardly be distinguished from the βAl5FeSi under an optical microscope, but they were easily differentiated by SEM as shown
in Figure 6-13(b). These two figures demonstrate that the solid state transformation of the
β-Al5FeSi to N-Al7Cu2Fe phase starts from the interface and extends inward to the βAl5FeSi phase. It is worth mentioning that the vast majority of the N-phase particles,
formed during the SHT, came from the transformation of the β-Al5FeSi particles rather than
the transformation of the AlCuFe- intermetallic formed during solidification. The volume
fraction of the latter was negligible in the alloys investigated.
113
Figure 6-12: phase morphology, EBSD pattern and simulation results for the N-Al7Cu2Fe phase in alloy #5RC3F0.7 (MAD=0.2).
Figure 6-13: N-Al7Cu2Fe phase under a) optical microscope and b) SEM.
In order to validate the actual reaction producing peak III, one must identify the product(s)
of the incipient melting of the N-phase. Therefore, initially the specimens were solution
heat treated for 10 hours at 775 K (502
to have a sufficient volume fraction of the N-
phase in the microstructure. Figure 6-14(a) illustrates the microstructure of alloy #7
(RC3F0.7(M0)) after the heat treatment. Subsequently, the specimens were heated up to the
temperature just beyond peak III (i.e. ~800 K (527
in RC3F0.7 and ~810 K (537
in
RC3F0.7(M0)) and then rapidly quenched. Figure 6-14(b) illustrates the microstructure of
alloy #7 (RC3F0.7(M0)), at the same location as Figure 6-14(a), after applying the second
step of SHT (i.e. 10 min. at ~810 K (537
). The circled areas in Figure 6-14(a) indicate
the presence of the N-phase after the first step SHT. The same locations in Figure 6-14(b)
are composed of AlFeSi and AlCu intermetallics. The N-phase therefore experienced
incipient melting and was supposedly substituted by AlFeSi intermetallics (mostly βAl5FeSi); the AlCu intermetallic being precipitated from the liquid phase upon cooling
from 537°C. Notice that according to equilibrium computations made with Thermo-Calc, β-Al5FeSi is more stable than the N-phase at 537°C for a system having the composition of
114
the alloy #7 (RC3F0.7(M0)). The situation is reversed at a lower temperature, where the Nphase becomes more stable than the β-Al5FeSi phase. This explains why the N-phase grows
at the expense of the β-Al5FeSi phase when the specimen is reheated.
Figure 6-15 compares the microstructure of alloy #5 (RC3F0.7) after (a) the first step of the
SHT (i.e. 10 hours at 775 K (502
), and (b) after the second step of the SHT (i.e. 10 min.
at 800 K (527 )). As shown, almost all of the areas containing the N-phase experienced
incipient melting and the products are AlFeSi intermetallic; porosity and other phases were
difficult to identify. Incipient melting seems to start at the interface of the α-Al and Nphase, to extend these phases inward. Thus, peak III can be correlated to the following
reaction through which the N-phase along with α-Al are transformed to liquid and βAl5FeSi: (α-Al+ N-Al7Cu2Fe+ Si → Liquid + β-Al5FeSi).
Figure 6-14: microstructure of alloy #7 (RC3F0.7(M0)) a) 10 hours solution treated at 775 K (502 ), b) 10
hours solution treated at 775 K (502 ), quenched and 10 min. solution treated at 810 K (537 ); (a and b
were taken at the same location); (solid black arrows are AlCuFe-, dotted red arrows are AlFeSi- and dashed
arrows are AlCu- intermetallics).
115
Figure 6-15: microstructure of alloy #5 (RC3F0.7) a) 10h solution treated at 775 K (502 ), b) 10h solution
treated at 775 K (502 ), quenched and 10 min. solution treated at 803 K ((530 )); (a and b were taken at
the same location).
6.3.4. Sequence of the θ- and Q-phases transformation in heating/cooling processes
According to the literature, the precipitation/melting temperature of the θ-phase during the
cooling/heating process of Al-Si hypoeutectic alloys containing Cu and Mg, normally
occurs some degrees (~15 K above the precipitation/melting temperature of the Q-phase 6,
7, 303
. However, according to the results obtained with the MBD model, which are
schematically illustrated in Figure 6-16, the temperature and sequence of the precipitation
of the Q- and θ-phases, are both strongly influenced by Cu and Mg contents. During
solidification, the precipitation temperature of the θ-phase increases with increasing Cu
content and decreasing Mg content, while the precipitation temperature of the Q-phases
decreases with increasing Cu content and decreasing Mg content. Similar results have
recently been reported by Yan et al. 14. In alloys #4 and #5 (RC3 and RC3F0.7), the Q- and
θ-phases are both predicted to precipitate at almost the same temperature (at ~783 K 510
), but in alloy #3 (RC1.5), the model predicts that the Q-phase should precipitate (at
~799 K 526
) some degrees above the onset temperature of the θ-phase, which
precipitates at ~783 k 510
.
Figure 6-16: effects of Cu and Mg contents on the precipitation temperature of the Q- and θ-phases; predicted
by the MBD model.
As shown in the cooling DSC curves of the alloys (Figure 6-4), peaks I and II were merged
together in alloys #4 and #5 (RC3 and RC3F0.7); but in the alloy #3 (RC1.5), these two
peaks were clearly appearing at two different temperatures. Similar DSC results for alloys
116
with chemical compositions comparable to alloys #3 (RC1.5) and #5 (RC3F0.7) have been
reported by Mrówka-Nowotnik et al.
303
and Martinez et al.
141
, respectively. Therefore,
unlike the heating DSC curves, there is a good consistency between the number of peaks
observed in the cooling DSC curves and the number of peaks predicted by the MBD model.
It seems that, as predicted by this model, the θ-phase and Q-phase in alloy #4 (RC3)
precipitate at almost the same temperature, but in alloy #3 (RC1.5), the Q-phase should
precipitate earlier than the θ-phase.
To validate the results given by the MBD model and to also verify the reacting phases
corresponding to peaks I and II in the heating process, specimens of alloys #3 (RC1.5) and
#4 (RC3) were heated in DSC up to 787 K (514 , the temperature just beyond peak I);
subsequently, they were rapidly quenched. As shown in Figure 6-17, both the θ- and Qphases experienced localised melting in alloy #4 (RC3) after the heat treatment. However,
in alloy #3 (RC1.5), as illustrated in Figure 6-18, only the θ-phase was locally melted; the
Q-phase was only partially dissolved by the solid state transformation and some remained
after the heat treatment. Therefore, the results indicate that, as predicted by the MBD
model, peak I corresponds to the melting of the θ- and Q‐phases in alloy #4 (RC3), but in
alloy #3 (RC1.5) it corresponds to the melting of the θ-phase alone. The melting of these
phases occurs when they react with the aluminium primary phase, as this is predicted by the
reverse eutectic reactions if local equilibrium is reached. However, it is not clear how to
define the “local equilibrium”, since the reactions likely start at interfaces, so the size and
composition of the reacting system are difficult to establish. As shown in Figure 6-17, the
AlCuFe- intermetallic also precipitated in the microstructure of alloy #4 (RC3) after
heating up to 787 k (514
. By heating a specimen of alloy #4 (RC3) up to 803 K (530 , right after peak II/III), all
secondary phases containing Cu (the Q-phase, θ-phase and AlCuFe- intermetallic)
experienced localised melting. Figure 6-19 compares the evolution of the microstructure in
this alloy before and after heating up to 803 K (530 ). It is worth mentioning that the
phase identifications were all validated with EPMA.
117
Figure 6-17: a) alloy #4 (RC3) in the as-cast condition. b) alloy #4 heated up to 787 K (514 , just beyond
peak I) and quenched (a and b were taken at the same location); (prepared with the permanent mould).
Figure 6-18:a) alloy #3 (RC1.5) in the as-cast condition. b) alloy #3 heated up to 787 k (514 , just beyond
peak I) and quenched (a and b were taken at the same location); (prepared with the permanent mould).
Figure 6-19: a) alloy #4 (RC3) in the as-cast condition. b) alloy #4 heated up to 803 K (530 , just beyond
peak III) and quenched (a and b were taken at the same location); (prepared with the permanent mould).
118
6.3.5. Effect of Cu content on the post-eutectic phases
The DSC cooling curves of the 356 Al alloys containing 0.5, 1.5 and 3 wt.% Cu (RC0.5,
RC1.5 and RC3) are compared in Figure 6-20. Peaks I and II, which were not seen in the
alloys containing 0.5 wt.% Cu (RC0.5 and RC0.5F0.7), appeared in the alloy containing 1.5
wt.% Cu (RC1.5) or more. The area corresponding to these peaks (I and II) in alloy #3
(RC1.5) is smaller than alloy #4 (RC3), which implies the lower volume fraction of the
eutectic (θ‐ and Q-) phases.
Though the DSC peaks I and II did not appear with the alloys containing 0.5 wt.% Cu
(RC0.5 and RC0.5F0.7, Figure 6-20 and Figure 6-3d), the phases (Q- and θ-) corresponding
to these peaks (I and II) were observed in the as-cast microstructure (Figure 1a). Moreover,
these phases (Q- and θ-) were predicted in all of the Mg containing alloys (#1 to #5) as
illustrated in Figure 6-2. The discrepancy between the DSC results, the predicted and
observed microstructures, could be due to the low volume fraction of the phases in the
alloys containing 0.5 wt.% Cu content (RC0.5 and RC0.5F0.7), which were not detected by
DSC.
Another major difference between the microstructures of alloys RC0.5, RC1.5 and RC3,
was the presence of the π-phase in alloys RC0.5, RC1.5, which corresponds to peak IV in
the DSC curve. Peak IV has been reported to occur at ~827 K (554
for the precipitation
of the π-phase in Al-Si-Mg/Cu alloys 7, 140. According to our DSC results, this peak started
at about 815 K (542
in alloy RC0.5, while in alloy RC1.5, it appeared approximately at
805 K (532 , Figure 6-20). Therefore, the precipitation/dissolution temperature of this
phase seems to be affected by the Cu content.
119
Figure 6-20: DSC cooling curves of alloys #1 (RC0.5), #3 (RC1.5) and #4 (RC3), with a scanning rate of 5
K/min.
Figure 21 illustrates the effect of Cu content on the predicted mass fraction and temperature
formation of the π-phase, which was calculated with the MBD model. As illustrated in this
figure, the mass fraction of the π-phase in an Al-7Si-xCu-0.35Mg-0.15Fe alloy was reduced
from ~0.25% to ~0.05% by increasing the Cu content from 0.5 wt.% Cu to 2 wt.% Cu,
respectively. Moreover, the precipitation temperature of the π-phase decreased from ~825
to 800 K (552 to 527
. Figure 6-20 and Figure 21 show that there is a good agreement
between the precipitation temperature of the π- phase measured by DSC analysis and
predicted by the MBD model.
Figure 6-21: evolution of mass fraction and temperature formation of the π-phase with Cu content (in Al-7SixCu-0.35Mg-0.15Fe), predicted by the MBD1.
Conclusion
1. It was found that the microconstituent called the “block-like θ-Al2Cu phase” is in
fact an AlCuFe- intermetallic compound containing a significant amount of Fe.
120
2. Though the AlCuFe- intermetallic was hardly found in the as-cast microstructure,
the reaction of α-Al with the β-Al5FeSi phase causes the formation of the NAl7Cu2Fe phase during the heating (>723 K (450
sufficiently high amount of Cu (e.g.
) of alloys containing a
3 wt.%).
3. By heating the Al-7Si alloy containing 3 wt.% Cu in DSC, two peaks appeared at
~794 and 805 K ~521 and 532
; these peaks were correlated to the melting of
the AlCu-eutectic and AlCuFe- intermetallic, respectively.
In the Al-7Si-3Cu alloy containing Mg, the DSC peak corresponds to the melting of
the AlCuFe- intermetallic appeared at 795 K (522
. The results are in good
agreement with the results predicted by the multiphase back-diffusion model.
4. The area fraction of the N-phase was significantly enhanced by increasing the time
period of the solution heat treatment. By reheating the solution treated specimen to
810 K (537
for the Al-7Si-3Cu-0.75Fe alloy, the N-phase was replaced by β-
Al5FeSi and other solid phases.
5. According
to
the
multiphase
back-diffusion
model,
the
solidification
sequence/temperatures of θ- and Q-phases are strongly affected by Cu and Mg
content. This has been confirmed by the thermal DSC analysis and metallographic
assessment. 6. In Al-7Si-0.3Mg-xCu alloys, the precipitation/dissolution temperature of the πphase was influenced by the Cu content. The DSC peak corresponding to the πphase during cooling occurred at ~817 K 544
occurred at ~808 K (535
with 0.5 wt.% Cu, while it
with 1.5 wt.% Cu. These results are in agreement with
the multiphase back-diffusion model.
121
122
Chapter 7
.
“Solubility/ Stability of Cu/Mg Bearing Intermetallics in Al-Si Foundry Alloys
Containing Different Cu and Mg Content”
Résumé:
Quatre alliages hypoeutectiques Al-Si contenant diverses teneurs en Cu (1 et 1,6 wt.%) et
Mg (0,4 et 0,8 wt.%) ont été étudiés afin d'évaluer avec plus de détails l'évolution des
intermétalliques contenant ces éléments Les fractions de phases contenant du Cu/Mg ont
été quantifiées avant et après le de traitement thermique de mise en solution (SHT) pour
évaluer la solubilité/stabilité des phases. Deux intermétalliques contenants du Mg (QAl5Cu2Mg8Si6, π- Al8FeMg3Si6) ayant une couleur grise sous le microscope optique ont été
discriminés avec l’aide d’attaques chimiques. En outre, les concentrations des éléments
(Cu, Mg et Si) dans la phase α-Al ont été analysées. Les résultats ont montré que, dans les
alliages contenant ~ 0,4% de Mg, la phase Q-Al5Cu2Mg8Si6 s’est dissous après le
traitement thermique de 6 heures à 505 ; mais dans les alliages contenant ~ 0,8% de Mg,
il était insoluble / partiellement soluble. Par ailleurs, après le traitement thermique à 505 ,
la phase Mg2Si a été partiellement substituée par la phase Q. L’application d’un traitement
thermique à des températures élevées (par exemple 525 ) a provoqué la fusion localisée
des intermétalliques contenant du Cu (Q et θ) dans les alliages contenant une haute teneur
en Mg (0.8 wt.%). L'application de traitement thermique en deux étapes (6h@ 505
+
8h@ 525 ) dans les alliages contenant ~ 0,4% de Mg, a contribué à dissoudre davantage le
reste des intermétalliques contenant du Mg et a en outre modifié la microstructure, mais
dans les alliages contenant ~ 0.8% de Mg, il a provoqué la fusion partielle de la phase Q. Il
y avait un bon accord entre les résultats expérimentaux et les résultats prévus par ThermoCalc. Pour réduire/éliminer les intermétalliques contenant du Cu / Mg non-dissous, la
solubilité des éléments (Cu et Mg) à la température de traitement thermique applicable doit
être prise en compte.
Abstract:
Evolutions of the Cu/Mg bearing intermetallics were thoroughly investigated in four Al-Si
hypoeutectic alloys containing various Cu (1 and 1.6 wt.%) and Mg (0.4 and 0.8wt.%)
contents. The area fractions of Cu/Mg bearing phases before and after solution heat
treatment (SHT) were quantified to evaluate the solubility/stability of the phases. Two Mgbearing intermetallics (Q-Al5Cu2Mg8Si6, π- Al8FeMg3Si6), which appear as gray colour
under optical microscope, were discriminated by the developed etchants. Moreover, the
concentrations of the elements (Cu, Mg and Si) in α-Al were analysed. The results
illustrated that in the alloys containing ~0.4%Mg, Q-Al5Cu2Mg8Si6 phase got dissolved
after 6 hours of SHT at 505 ; but in the alloys containing ~0.8%Mg, it was insoluble/ partially soluble. Furthermore, after SHT at 505 , Mg2Si was partially substituted by Qphase. Applying SHT at high temperatures (e.g. 525 ) caused localized melting of the
remaining Cu bearing intermetallics (Q and θ phases) in the alloys containing high Mg
content (0.8 wt.%). Applying a two-steps SHT (6h@505 +8h@525 ) in the alloys
containing ~0.4%Mg, helped to further dissolve the remaining Mg bearing intermetallics
and further modified the microstructure, but in the alloys containing ~0.8%Mg, it caused
partial melting of Q-phase.
Thermodynamic calculations were carried out to assess the phase formation in equilibrium
and in non-equilibrium conditions. There was a good agreement between the experimental
results and the predicted results. To minimize/eliminate the un-dissolved Cu/Mg bearing
intermetallics, the solubility of the elements (Cu and Mg) at the applicable SHT
temperature must be taken into account.
Introduction:
In the last decades, Al-Si based foundry alloys have been increasingly used in the
automotive industry mainly in the fabrication of engine components. High strength to
124
weight ratio, high thermal conductivity and excellent castability are the major advantages
of
the
Al–Si
hypoeutectic
alloys.
Nevertheless,
the
increase
of
operation
temperature/pressure of the engines necessitates strengthening of the Al–Si alloys.
Magnesium and Cu are the major/principle alloying element(s) of the commercial Al-Si
based foundry alloys due to their appreciable solubility and strengthening effects. The large
eutectic Cu/Mg bearing phases (θ-Al2Cu, Q-Al5Cu2Mg8Si6 and π-Al8FeMg3Si6), which
appear at the last stages of solidification, can get dissolved by applying an appropriate
solution heat treatment (SHT) and re-precipitated as fine evenly distributed metastable
phases to strengthen the alloys. However, they may be insoluble/ partially soluble
depending of the alloys chemistry (e.g. high Mg/Cu content and fraction) and the SHT
parameters used (time and temperature)294. If the large eutectic Cu/Mg bearing phases do
not dissolve during SHT, the hardening effect of the Cu/Mg elements will be reduced and
the ductility of the alloys will also suffer14. In order to minimize/eliminate un-dissolved
Cu/Mg bearing intermetallics, the alloy chemistry must be optimized.
The applicable SHT temperature (TSHT) is generally restricted by the non-equilibrium
solidus, which is the melting point of the last solidified phases (Tmp). Al-Si alloys
containing both Cu and Mg are generally limited to Tmp ~507 ; but for the alloys that
contain Cu and/or Mg individually, Tmp can be much higher10,
297, 304, 305
. The higher
applicable TSHT not only accelerate the dissolution rate of the Cu/Mg bearing intermetallics
but also further modify the microstructure (e.g. Si particles) of the alloys294. Another
strategy is to apply a two steps SHT: the temperature of the 1st SHT step (~500 ) is limited
by Tmp to avoid incipient melting of the Cu containing phases (θ-Al2Cu and QAl5Cu2Mg8Si6); after dissolution of the Cu bearing phases, the 2nd SHT step is applied at a
higher temperature (e.g. between 520 and 540
depending of the alloy chemistry) to
further dissolve the Mg bearing intermetallics and to further modify the microstructure192,
197, 294
.
Nevertheless, there is a controversy in literature about the stability of the Cu containing
phases. For instance, in Al-Si-Cu-Mg alloys, some researchers98, 199, 200, 205 indicated that Qphase is insoluble at ~500 , but others6, 297, 306 stated the complete/partial dissolution of the
Q phases. Lasa et al.300 even reported an increase in Q-content after SHT of the Al-13Si-
125
1.4Cu-1.3Mg-0.1Fe (wt.%) alloy; they mentioned that the dissolved Mg2Si was
transformed to Q-phase. Alfonso et al.306 reported very sluggish dissolution rate of θ-phase
(still visible after 72h SHT at 480 ) in Al-6Si-3Cu-0.6Mg (wt.%), but Moustafa et al.98
stated almost complete dissolution of θ-phase in Al-11Si-2.6Cu (wt.%) after 8h SHT at
500 . The present authors297 and Yan et al.14, reported that the solubility/ stability of the
Q-phase is strongly affected by the Cu-Mg content of the alloys. For the alloys containing
low Cu-Mg content, the Q-phase can be entirely soluble; but for the alloys containing high
Cu-Mg content, Q-phase can be insoluble/ partially soluble294, 297.
The microconstituents of an alloy, their corresponding volume fraction/ solidification
temperatures are strongly influenced by the Cu and Mg content. Samuel et al.9 reported that
the presence of very small Mg content (0.06 wt.%) leads to precipitation of Q, π and Mg2Si
phases in Al-6Si-4Cu-0.5Fe wt.% (319-type Al alloy). It has been reported that a Mg level
beyond 0.3 wt.% in 319-type Al alloys does not affect considerably the alloy strength,
while it can reduce significantly the alloy ductility307, 308. The presence of a large volume
fraction of insoluble Mg-bearing intermetallics was responsible for this reduced ductility.
Therefore, the alloy chemistry (mainly Cu and Mg content) must be optimized to minimize/
eliminate the insoluble/ partially soluble intemetallics.
To reduce the un-dissolved Q-phase in Al-9Si-0.1Fe(%) alloys containing Cu and Mg, Yan
et.al.14 suggested that: TQ < TH < (TS ˗ 10 ); where TQ is the precipitation temperature of
Q-phase, TH is the solution heat treatment temperature and TS is the equilibrium solidus
temperature. To satisfy this criteria, the preferred Mg and Cu content and their relations
were suggested to be: (Cu + 10·Mg) = 5.25 (wt.%), 0.5 < Cu < 2 wt.% and
0.27 < Mg < 0.53 wt.%. The lower and upper limits of this criterion were proposed to be
TQ < (TS ˗ 15 ) and TQ < (TS ˗ 5 ), respectively. To satisfy the lower and the upper limits
of the criterion, the Mg and Cu relations must be: 4.7 < (Cu + 10·Mg) < 5.8 (wt.%).
Studying the evolution of the Cu/Mg bearing intermetallics can be helpful to optimize the
alloy chemistry and the SHT process. The most common Cu/Mg bearing intermetallics
which frequently appear in the Al-Si based foundry alloys are as follow: θ-Al2Cu and QAl5Cu2Mg8Si6, π and Mg2Si. In optical microscope (OM), θ-phase with yellow colour
appears as the brightest phase and the Mg2Si phase with Chinese script morphology appears
126
as the darkest phase. These two phases are easily discriminated from the other
microconstituents of the Al-Si based alloys under OM. But, the Mg bearing intermetallic
(Q- and π-) phases both appear as light gray with more/less similar morphology, which
makes impossible to be differentiated under OM. Therefore, these phases must be either
discriminated (under OM) by means of an appropriate etchant or by means of electron
microscopy (SEM/EPMA). In the literature, HNO3 was used to differentiate the Cu based
phases (e.g. Q) by which the Cu phases change to dark grey 3, 309-312, and H2SO4 was used to
discriminate the Fe bearing intermetallics 310-312; but the details of the procedures were not
reported.
The major purpose of this work was to elucidate the effect of Cu and Mg content on the
solubility/ stability of Cu/Mg bearing intermetallics in Al-Si foundry alloys. Four Al-Si
foundry alloys containing various Cu and Mg contents were studied. To quantify the area
fraction of the phases, two etchants were developed to discriminate the Cu/Mg bearing
intermetallics under optical microscope. The maximum soluble Cu-Mg contents in Al-Si
foundry alloys at the applicable solution treatment temperature (TSHT) were investigated.
The evolutions of the Cu/Mg bearing intermetallics were thoroughly studied in the as-cast
and solution heat treated condition. This experimental work was paralleled by a
comprehensive study of the thermodynamic prediction of the microstructure evolution in
Al-7Si alloys containing different Cu and Mg content. These predictions were carried out
with the Thermo-Calc software
298
using the TTAL7 database
291
. A multiphase back
diffusion (MBD) model1 was used to calculate the phase precipitation sequence and the
mass fraction of microconstituents during solidification.
Materials and methods
The alloy making and melting procedures were detailed in chapter 3 (section 3.1.2). The
specimens prepared by the Pyrex tubes were only used for chemical analysis, and the
specimens cast in the permanent mould were used for microstructure characterization. The
average chemical compositions of the seven alloys investigated are presented in Table 7-1.
127
Table 7-1: chemical composition of the alloys (wt.%)
Alloy No.
Si
Cu
Mg
Fe
Al
Cu/Mg
Ref. A356.0
6.5-7.5
0.2
0.25-0.45
0.2
Bal.
-#1
RC1M0.4
6.81
1.05
0.39
0.08 Bal.
~3
#2
RC1M0.8
6.82
0.99
0.78
0.06 Bal.
~1
#3
RC1.6M0.4
6.77
1.64
0.38
0.06 Bal.
~4
#4
RC1.6M0.8
6. 86
1.63
0.78
0.06 Bal.
~2
“R” indicates the reference alloy. The symbols “C” and “M” represent the Cu
and Mg elements; the number after each symbol presents the concentration of
the respective element.
Samples for microstructural examination were sectioned from the bar, mounted, ground and
polished using standard procedure. The polished sections were then evaluated with an
optical microscope (OM-NIKON EPIPHOT) and with electron probe microanalysis
(EPMA-CAMECA SX100) equipped with a wavelength dispersive spectrometer (WDS).
In order to enhance the contrast between the Mg bearing intermetallics (Q and π) under
optical microscope (OM), two different solutions were developed: 1st etchant was (3 ml
HNO3 + 100 ml H2O), and 2nd etchant was (3 ml HNO3 + 1 ml HCl + 100 ml H2O); the
required time period for the etching process with both aforesaid solutions was ~15 min. To
validate the accuracy of the results, the phases differentiated under OM were validated by
EPMA. Quantitative metallography was carried out by the image processing ImageJ
software on three different polished specimens (for each individual alloy); a minimum of
36 fields (at least twelve fields per specimen) each with ~13958 μm2 surface area were
analysed per alloy at a magnification of 400X. The samples were scanned in a regular and
systematic manner. The reported mean value and standard deviation for each alloy were
calculated with the measurements made on these three sections. To validate the measured
area fraction of each Mg bearing intermetallic by OM, the same coordinates of three
micrographs already taken by OM were analysed by X-ray elemental mapping (with
EPMA) and the area fraction of each phase was verified. It is worth to mention that the
quantified area fraction of the intermetallics was assumed to be equal to their volume
fraction.
The solution heat treatment (SHT) was conducted in an electric resistance furnace. The
temperature of the solution treatment (TSHT) was ~505 . For some specimens, the 2nd step
of SHT was applied at higher temperatures (e.g. at ~525 ). The total time period of SHT
128
was 14 hours. For the SHT with two steps, the time period of the 1st step (505 ) was 6
hours which was continued by 8 hours SHT at ~525 . After the 14 hours SHT, the
specimens were quenched in water to obtain the maximum solute saturation.
Line-scans were conducted across dendrite arms using WDS to measure the element
concentrations (i.e. Cu, Mg and Si) into the α-Al matrix. At least 8 dendrite arms, three
points over each dendrite, were scanned per specimen. The line-scans were carefully taken
from areas free and fairly away from the other particles. A conventional vickers
microhardness tester (MATSUZAWA- (MMT-X7A)) was used to measure the hardness of
the α-Al matrix; the indentation load of 50 gram-force and a 15 second loading time were
used. The average of at least 8 measurements was reported as the microhardness value. The
indentations were always pointed in the α-Al matrix fairly away from the other particles.
Moreover, a comprehensive study of the thermodynamic evaluation of Al-7Si alloys
containing different Cu, Mg and Fe content was carried out with the Thermo-Calc software
and with the multiphase back diffusion (BDM)1 model .
Results and Discussion
7.3.1. Characterizing the microconstituents under OM:
Figure 7-1 (a) presents the as-cast microstructures of the alloy #2 (RC1M0.8). As shown,
all the microconstituents are individually discriminated under OM except the two Mg
containing intermetallics: Q-Al5Cu2Mg8Si6 and π-Al8FeMg3Si6. The eutectic Si phase is
dark gray in colour, θ-phase with yellow colour appears as the brightest phase, and the
Mg2Si phase with Chinese script morphology is the darkest phase in the microstructure.
The β-Al5FeSi phase with gray colour can be recognized by its platelet (needle like)
morphology. Nevertheless, the Mg containing phases (Q and π), which both appear in light
gray, require an appropriate solution to be differentiated.
Figure 7-1 (b) and (c) show the same microstructure after being treated by the 1st (HNO3)
and by the 2nd (HNO3+HCl) etchants, respectively. As shown in Figure 7-1 (b), after the
treatment with the 1st etchant (HNO3), Q-phase changed to dark colour (almost the same
colour as Mg2Si), however π‐phase altered slightly. After etching with HNO3, the 129
specimens were polished to remove the effect the etchant and were consequently etched
with the 2nd solution (HNO3+HCl) to compare the results. As shown in Figure 7-1 (c), Qphase remained almost intact but π-phase became slightly darker. Noteworthy that Mg2Si
remained with its own original dark colour after treated by the two aforesaid solutions.
Figure 7-1: microstruture of alloy #2-RC1M0.8 (SHTed 14h @ 505 : a) microstruture before tretment with
the etchant, b) after treatment with (HNO3), c) after treatment by (HNO3+HCl); the micrographs were taken at
the same coordinate.
In order to confirm the results, the etched specimens were verified by EMPA. Figure 7-2 (a
and b) respectively illustrate the microstructure of alloy #4 (RC1.6M0.8) before and after
treatment by (HNO3) under OM. The elemental mapping of the specimen at the same
location is presented in Figure 7-2 (c-d). As shown, there is an excellent agreement between
the results of OM and EPMA to distinguish the phases. Moreover, Figure 7-3 (a and b)
respectively illustrate the microstructure of alloy #4 (RC1.6M0.8) before and after
treatment with the 2nd etchant (HNO3+HCl) under OM. The elemental mappings of the
specimen are also presented in Figure 7-3 (c-d). As shown, the discrimination of the
130
microconstituents (mainly Mg-bearing intermetallics) under OM is sufficiently good to
provide accurate image analysis of phase fractions, removing the obligation to obtain
EPMA mappings for phase identification.
Since the treatment of the specimens by the 1st solution (HNO3) made a better contrast
between the Q and π-phases, it is preferable to use this solution to quantify the area fraction
of the phases. However, by this method, as shown in Figure 7-4, Q and Mg2Si phases both
appear as dark colour in the microstructure.
The area fraction of the phases (Q and π), which was initially counted by the OM, was
verified by EPMA as well. The appearance of a phase with a range of colors under EPMA
reduces the accuracy of the image analysis in which the phase is selected by colorthreshold. For instance, to measure the area fraction of Q-phase, the mapping of Mg
element (already presented in Figure 7-2 d) was considered. By changing the value of hue
from (134, 168) to (134, 169) in the threshold-color section of ImageJ, the measured area
fraction of Q-phase was enhanced from ~2.7 to 10.6%; the area selected in the image
processing are compared in Figure 7-5(c and d). This large imprecision is mainly due to the
presence of solute Mg in α-Al matrix which changed the color of the matrix to light blue
(almost the same color as Q-phase). However, by manually selecting/masking Q-phase (the
white area in Figure 7-5 b), the corresponding area fraction of Q-phase was ~3.1% which is
in accordance with the OM results (~3.2%).
131
Figure 7-2: as-cast microstructure of alloy #4-RC1.6M0.8: a) before tretment with the etchant, b) etched with
HNO3 c-e) EPMA elemental mapping of the Cu, Mg and Fe elements; the dashed lines in EPMA micrographs
correspond to the same coordinate of the OM micrographs.
132
Figure 7-3: as-cast microstructure of alloy #4-RC1.6M0.8: a) before tretment with the etchant, b) etched with
HNO3+HCl c-e) EPMA elemental mapping of the Cu, Mg and Fe elements; the dashed lines in EPMA
micrographs correspond to the same coordinate of the OM micrographs.
Figure 7-4: SHTed (6h@505 ) microstructure of alloy #2-RC1M.08 a) before treatment with the etchant,
and b) after being etched with HNO3; the micrographs were taken at the same coordinate; gray colour of Qphase was changed to dark colour (like Mg2Si) after being etched.
133
(a)
(b)
(c)
(d)
Figure 7-5: a EPMA elemental mapping of the Mg element correspond to the dashed area in Figure 7-2-d, b)
the area correspond to Q-phase (white area) was manually masked, c) the hue in threshold color of ImageJ 134, 168 and the counted area fraction is 2.9% d the hue 134, 169 and the counted area fraction is 10.7%. 7.3.2. Stoichiometry of the phases after etching:
The solutions used here in etching process seem attacked inside of the Cu/Mg bearing
phases and changed the stoichiometry of some phases (in particular Mg2Si). For example,
the stoichiometry of the Mg2Si phase in as-cast condition, which was checked by EPMA,
was Mg1.9Si. But after the etching processes, the phase (Mg2Si) generally appeared like
porosity and the Mg element was almost eliminated; in some area, the remaining Si (which
was not dissolved by the solution) was detected by EPMA. Moreover as presented in the
Table 7-2, the stoichiometry of the Q and π-phases were slightly altered.
Table 7-2: stoichiometry of Q- and π-phases measured before and after etching.
From literature
as-cast condition
After etching
HNO3
Q-Al5Cu2Mg8Si6 Al5Cu2Mg8Si6.5 Al6Cu2Mg7.5Si6.5
π-Al8Mg3Si6Fe Al9.5Mg3.5Si5Fe Al10.2Mg3.5Si5.5Fe
(HNO3+HCl)
Q-Al5Cu2Mg8Si6 Al5Cu2Mg8Si6.5
Al6Cu2Mg7Si6
π-Al8Mg3Si6Fe Al9.5Mg3.5Si5Fe Al13.5Mg3.5Si7Fe
7.3.3. Effect of Cu/Mg content of the alloys on evolution of as-cast microstructure
The predicted mass fractions of the phases formed during the solidification process of Al7Si-0.07Fe-1.6Cu-xMg alloy are presented in Figure 7-6 with respect of the Mg content.
The dashed-vertical-green lines correspond to the chemistry of the alloys #3-RC1.6M0.4
and #4-RC1.6M0.8. By increasing the Mg content of the alloy, the mass fraction of θ-phase
gradually reduces, but the mass fraction of the Mg bearing intermetallics (Q, π and Mg2Si)
134
considerably enhances. By increasing Mg content from 0.1 to 0.8%, β-Al5FeSi is gradually
substituted with π‐phase so that the mass fraction of π-phase enhances (from 0%) to
~0.5% (and β-Al5FeSi decreases from 0.23 to 0.04%).
1.6% Cu
#3
#4
Phase fraction, Mass%
0.012
Ɵ-Al2Cu
Q
β-AlFeSi
π
Mg2Si
0.006
0
0
0.2
0.4
Mg content, wt%.
0.6
0.8
Figure 7-6: The effect of Mg content on phase fractions in Al-7Si-1Cu-0.07Fe-xMg (wt.%); the dashedvertical-red lines correspond to the chemistry of the #3-RC1.6M0.4 and #4-RC1.6M0.8 (the results were
predicted by MBD1).
The as-cast microstructures of the alloys are presented in Figure 7-7. The microstructure of
the alloys containing ~0.4% Mg (RC1M0.4 & RC1.6M0.4) were composed of α-Al
dendrites, eutectic Si phase, θ-Al2Cu particles, Q-phase and Fe-containing intermetallics (πand β-phase); the β-phase was barely found in the microstructure. By increasing the Mg
content to 0.8% (i.e. in alloys #2-RC1M0.8 & #4-RC1.6M0.8), β-phase was replaced by πphase and Mg2Si appeared in the microstructure. These are all in excellent agreement with
the predicted results.
135
Figure 7-7: as-cast microstructures of the experimental alloys: a) alloy#1 (RC1M0.4), b) alloy#2 (RC1M0.8),
4 c) alloy#3 (RC1.6M0.4), d) alloy#4 (RC1.6M0.8).
The quantified area fractions of the phases in as-cast condition are presented in Figure 7-8;
the predicted volume fraction of the phases by MBD[26] are also included for comparison.
As stated earlier and shown in Figure 7-4, Q- and Mg2Si-phases both appeared in etched
microstructure with more/less the same colour (dark); therefore, these phases both are
counted together (Q+Mg2Si) in image analysis. Moreover, due to less contrast in colour
between the Fe-bearing intermetallics (i.e. π and β), these two phases were counted
together, as well.
The mass fraction of Fe-bearing intermetallics (β-Al5FeSi and π-phases) was enhanced
from 0.3% (in alloy #3: RC1.6M0.4) to 0.7 % (in #4: RC1M0.8). As shown in Figure 7-8,
there is good correlation between the predicted results by MBD and the experimental
results. The same scenario has been reported by Wang et al.[27] for A356 (Al-7Si-0.4Mg0.09%Fe) and A357 (Al-7Si-0.7Mg-0.09%Fe) alloys; where, the volume fraction of Fe-
136
bearing intermetallics enhanced from ~0.5% in A356 alloys to 1.6% in A357 alloys. They
stated that the Fe-bearing intermetallics, which were almost exclusively β-Al5FeSi in
A356, changed dominantly to large π-phase along with a small proportion of β-Al5FeSi in
A357[27].
Area/Vol. fraction of phases
(Q+Mg2Si)-Experiment
(Q+Mg2Si)- MBD prediction
(π+β)- Experiment
1
(π+β)- MBD Prediction
0.5
0
1
2
alloy No.: #2
#4
3
#1
Cu/Mg (wt.%)
4
#3
Figure 7-8: the quantified area fractions and predicted volume fraction by MBD1 of the phases Q Mg2Si and π β in as‐cast condition vs. ratio of Cu/Mg.
7.3.4. Effect of Cu/Mg content on maximum applicable SHT temperature
For the Al-Si alloys with the chemical compositions more/less similar to the chemistry of
the studied alloys, Ammar et al. [28] recommended ~525 as the solution heat treatment
temperature; solution heat treatment at this temperature not only did not deteriorate the
microstructure but also it improved the mechanical properties [28]. Therefore, SHT was
initially performed at ~525
for the studied alloys. Figure 7-9 illustrates the in-situ
micrographs of alloy #2 (RC1M0.8) in as-cast condition and after being SHTed for 5 hours
at 525 . As shown, π and Mg2Si phases, both were still stable/ partially soluble at 525 ;
but the Cu containing phases (Q and θ) were both melted, instead of getting dissolved.
Consequently, in next step, the SHT was applied in two steps: 1th step: (6h@to505 ) + 2nd
step: (8h@525 ).
137
Figure 7-9: microstructure of alloy #2 (RC1M0.8), a) in as-cast condition, b) after 5 hours SHT at 525 ; the
micrographs were taken at the same coordinate; Cu-bearing intermetallics were melted but Mg2Si and πphases were remained almost intact.
Figure 7-10 present the in-situ micrographs of the alloys in as-cast condition, after the 1st
step of SHT (6h@505 ) and after the 2nd step of SHT (8h@525 ). In-situ microstructure
analysis of the alloys helped to better understand the evolution of the Cu/Mg bearing
intermetallics after each SHT step. After the 1st step, the Si‐particles were spheroidized and the θ-particles got completely dissolved in all of the studied alloys. Depending the
Cu/Mg content of the alloys, the Mg containing phases (Q, π and Mg2Si) were soluble/
insoluble. For instance, in alloy #3-RC1.6M0.4 (Figure 7-10 d, e & f) almost all of the Qphase disappeared after the 1st step; but in alloy #2-RC1M0.8 (Figure 7-10 a, b & c) and
#4-RC1.6M0.8 (Figure 7-10 g, h & i), it was insoluble/ partially soluble. The stability of
the Cu-bearing intermetallics (e.g. Q) in the alloys containing 0.8% Mg restricts the SHT
temperature to 505
and prevents applying the 2nd SHT step (i.e. at 525 ). As shown in
Figure 7-10 (g, h & i), the remained Q-phase was melted after the 2nd step SHT. This is in
agreement with the incipient melting of Cu-bearing intermetallics (Q-phase and
undissolved θ-phase) after SHT at 520
in 319 type Al alloys reported by Han et al.[4, 13].
Therefore, though two step SHT can be applied for the alloys containing 0.4%Mg (#1RC1M0.4 & #3-RC1.6M0.4), only one step SHT (TSHT at 505°C) is recommended for the
alloys containing 0.8%Mg (#2-RC1M0.8 & #4-RC1.6M0.8).
138
Figure 7-10: evolution of as-cast microstructure (a, d & g) after applying 1st step SHT (6h@505 ) (b, e & h)
and after 2nd step SHT (8h@525 ) (c, f & i); a-c: alloy #2 (RC1M0.8), d-f: #3 (RC1.6M0.4) and g-i: alloy #4
(RC1.6M0.8); the micrographs of each alloy were taken at the same coordinate.
7.3.5. Microstructure evolution and age hardening after SHT at 505 :
The specimens were all SHTed at 505
(for 14 hours) and etched with HNO3 to reveal the
different phases. The measured area fractions of the Mg-bearing intermetallics (Q+Mg2Si)
and Fe-bearing (π+β) intermetallics in the as-cast condition and after SHT are provided in
Figure 7-11 with respect of the ratio of Cu/Mg. The predicted volume fractions of the
phases at 505 are also shown for comparison. In the alloys containing 0.4% Mg
(Cu/Mg> 2.6, alloys #1 & #3), Q phases got dissolved almost completely after the SHT;
however, π-phase was partially dissolved. For the alloys containing 0.8%Mg (Cu/Mg<2.6,
alloys #2 and #4), all of the Mg/Fe bearing phases (Q+Mg2Si & π β) were insoluble/
139
partially soluble. As one can see, there is good agreement between experimental and
predicted results.
The concentration of the elements (Cu, Mg and Si) in α-Al after the SHT (14h@505 )
was evaluated to verify the solutes in the matrix. The results are presented in Figure 7-12
where each value represents the average of at least 24 readings, carefully pointed away
from the other particles. Moreover, the predicted equilibrium concentrations of the
elements (Cu, Mg and Si) in α-Al at 505
are also provided for comparison. As shown,
there is satisfactory agreement between the predicted and the experimental results. These
concentration of the elements in α-Al confirm that in the alloys containing 0.4% Mg
(Cu/Mg> 2.6, alloys #1 & #3), the majority of the Mg containing phases were dissolved
after the SHT since the concentration of Mg in α-Al reached ~0.35%. However, for the
alloys containing 0.8% Mg (Cu/Mg<2.6, alloys #2 and #4), the concentration of Mg in α-Al
reached up to ~0.43% which implicitly indicate that the majority of the Mg containing
phases were still remained un-dissolved.
Worth to note that by applying SHT at 505 , the Mg2Si phase was partially transformed to
Q-phase in alloy #4 (RC1.6M0.8) after SHT. This solid state phase transformation can be
clearly established from Figure 7-13. The same phenomenon has been previously observed
in a hyper-eutectic Al-13Si-1.4Cu-1.3Mg-0.1Fe (wt.%) alloy after 5 hours SHT at 500
by
Lasa et al.300. They stated that Mg2Si dissolved, and simultaneously the Q-phase nucleated
on the Mg2Si particles. This is also in agreement with the predicted results according which
Q-phase is always more stable than Mg2Si for a system having the chemistry as the alloy #4
(RC1.6M0.8).
140
Vol./area fraction of phases (π+β)
Vol./area fraction of phases (Q+Mg2Si)
AsCast-Experiment
14hSHT@505˚C
1
b) Vol. Fraction of π+β
1
a) Vol. Fraction of Q+Mg2Si
Prediction@505˚C
0.5
AsCast-Experiment
14hSHT@505˚C
Prediction@505˚C
0.5
0
0
1
2
alloy No.:#2
3
#4
#1
Cu/Mg (wt.%)
1
4
#3
2
alloy No.: #2
3
#4
#1
Cu/Mg (wt.%)
4
#3
Figure 7-11: Phase fractions (in as-cast condition, after SHT (14h@505 ) and predicted@505 ) in the
alloys vs. ratio of Cu/Mg a) Q+Mg2Si b) π β . Cu%-14h.SHT@505˚C
0.005
1.2
Mg%-14h.SHT@505˚C
Cu%-Prediction@505˚C
Si%-14h.SHT@505˚C"
Si%-Prediction@505˚C
wt.% of dissolved element ( Si)
Mg%-prediction@505˚C
wt.% of dissolved element ( Mg)
wt.% of dissolved element ( Cu)
0.025
0.02
0.004
0.015
0.01
0.7
0.003
1
2
alloy No.: #2 #4 #1
3
4
Cu/Mg (wt.%)
#3
0.95
1
2
3
4
Cu/Mg (wt.%)
alloy No.: #2 #4 #1
#3
1
alloy No.: #2
2
3
4
Cu/Mg (wt.%)
#4
Figure 7-12: Concentration of the elements [after SHT (14h@505 ) and predicted@505
matrix vs. ratio of Cu/Mg a) Cu b) Mg c Si.
#1
#3
in the α-Al
Figure 7-13: microstructure of alloy #4 (RC1.6M0.8); a) in as-cast condition, b) after 1st step SHT
(6h@505 ); the micrographs were taken at the same coordinate; the phase inside the solid black line
represents Mg2Si phase which shrunk in favour of Q-phase after applying the SHT, Q- and π-phases were
remained almost intact.
141
The age-hardening curves of the SHTed specimens, which were aged at 180
for different
time periods, are plotted in Figure 7-14. Increasing Cu content from ~1% (#1-RC1M0.4 &
#2-RC1M0.8) to ~1.6% (#3-RC1.6M0.4 & #4-RC1.6M0.8) considerably enhanced the
hardness (of α-Al matrix) after the SHT, the aged and the over-aged conditions; but
increasing the Mg content from ~0.4 (alloys #1-RC1M0.4 & #3-RC1.6M0.4) to ~0.8
(alloys #2-RC1M0.8
& #4-RC1.6M0.8) did not appreciably affect the hardness. The
former can be due to better dissolution of Cu bearing intermetallics (mainly θ) which
enhances the chance of precipitation of metastable particles while aging; and the latter can
stand for the insoluble large Mg bearing intermetallics by which the Mg element loses the
chance of precipitation as fine metastable particles while aging. Transmission electron
microscopy analysis is required to evaluate with further details the microstructure evolution
before/after aging process.
Hv
130
110
#4 (RC1.6M0.8)
#3 (RC1.6M0.4)
90
#2 (RC1M0.8)
#1 (RC1M0.4)
70
1
10
100
1000
10000
Time (min.)
Figure 7-14: Hardness vs. aging time relation; (specimens were SHTed (14h@505 ), quenched and aged at
180 .
7.3.6. Effect of high temperature SHT on dissolution of intermetallics
In the alloys containing 0.4%Mg (#1-RC1M0.4 & #3- RC1.6M0.4), solubility of Cu
bearing intermetallics (Q- and θ phases) at 505 allows to apply the 2nd SHT step at higher
temperatures. Solution heat treatment at higher temperatures helped to dissolve the
142
remaining Mg bearing intermetallics and to further modify the Si particles. Figure 7-15
compares the concentration of the elements (Si, Mg & Cu) in α-Al after different SHT
processes. As shown, the concentration of Mg and Si in α-Al was enhanced with increasing
the SHT temperature; however the Cu concentration was remained almost constant since
the Cu bearing intermetallics (θ and Q) are (thermodynamically) soluble at TSHT>483 .
There is excellent consistency between the experimental and predicted results, which are
both compared in this figure. In addition, Figure 7-16 illustrates the effect of the SHT
processes on the hardness (of α-Al matrix) in alloy #3(RC1.6M0.4). As presented, the
hardness (of α-Al matrix) is enhanced by increasing the SHT temperature, which implies
the presence of more solutes in α-Al.
Si- ThermoCalc
Cu- ThermoCalc
Mg- ThermoCalc
0.005
0.022
0.017
0.004
0.012
0.007
Mass % of Mg in α-Al
Mass % of Si and Cu in α-Al
Si-Exp
Cu-Exp
Mg-Exp
0.003
480
490
500
510
Temp (˚C)
520
530
Figure 7-15: concentration of the elements (prediction & experiment) in the α-Al matrix of alloy #3
(RC1.6M0.4) vs. the different SHT processes [(14h@490 , (14h@505 , (6h@505+8h@520 and 6h@505+8h@530
.
143
95
Hv
sample #3 (0.4Mg-1.6Cu)
90
85
80
480
490
500
510
Temp. (˚C)
520
530
Figure 7-16: Hardness of alloy #3-RC1.6M0.4 vs. the different SHT processes [(14h@490
(6h@505+8h@520 and 6h@505+8h@530
.
, (14h@505
,
General discussion
The predicted equilibrium concentration of Mg in α-Al for the alloys Al-7Si-yCu-xMg0.1Fe at two different temperatures (505 & 530 ) is presented in Figure 7-17. With
increasing Mg content of the alloys up to 0.33% at 505 , the equilibrium concentration of
Mg in α-Al
is linearly enhanced up to ~0.36% (section (a) of the curve);
subsequently, by increasing Mg content up to 0.8% (section (b-d)),
remains
either constant or enhances slightly. In section (a) of the curves, Mg is entirely soluble in αAl; ones crossing the section (b), Mg-bearing intermetallic(s) appear(s) and α-Al passes its
maximum solid solubility. As shown, the same scenario is repeated at 530 , except the maximum solubility of Mg in α-Al which reaches to ~0.41%. Worth to note that
the effect of Cu content (0.5 to 5%) of the Al-Si based alloys on the maximum solubility of
Mg in α-Al is negligible (
~ 0.33 to 0.36, respectively).
The predicted concentrations of Mg in α-Al (
) are compared with the
experimental data for different chemistries in Table 7-3; as shown, there is excellent
agreement between the experimental and the predicted results. Accordingly, one can
conclude that in the Al-Si based alloys for which the maximum applicable SHT
temperature (
) is limited to 505 , the optimum Mg content of the alloys is ~0.33%;
and in the alloys for which the SHT can be applied at higher temperatures, the Mg content
of the alloy can be increased ( e.g. to 0.41% for SHT at ~530 ). This is in agreement with
144
the results reported in literature; Samuel et al.308 recently studied the effect of Mg content
(~ 0.00 and 0.3 and 0.6%) in tensile strength properties of 319–type Al alloys; ~0.3% Mg
content reported as the optimum content. Figure 7-17: maximum dissolved Mg (mass %) in α-Al vs. Mg content of Al-7Si-yCu-xMg-0.1Fe alloy
(y=0.5, 1 & 1.5; x= 0.00 to 0.8%) at two different temperatures (505 and 530 ). The microconstituents in
sections: a=(Al+ Si+ β-Al5FeSi), b=( Al+ Si+ β-Al5FeSi+ π-Al8FeMg3Si6), c=( Al+ Si+ π-Al8FeMg3Si6); for
Cu= 0.5 and 1%: d=( Al+ Si+ π-Al8FeMg3Si6+ Mg2Si) and for Cu= 1.5: d=( Al+ Si+ π-Al8FeMg3Si6+ QAl5Cu2Mg8Si6).
Table 7-3: concentrations of Mg element in α‐Al after different SHT conditions in the studied alloys.
Alloy No.
#1-RC1M0.4
#2-RC1M0.8
#3-RC1.6M0.4
#4-RC1.6M0.8
SHT condition
14h@ 505
0.352 (0.362*)
0.434 (0.445)
0.368 (0.372)
0.431 (0.426)
6h@505+8h@ 520
--0.399 (0.404)
-6h@505+8h@ 530
--0.405 (0.403)
-* The predicted equilibrium values at the corresponding TSHT (505, 520 & 530 ) are listed in the parenthesis.
7.4.1. Stability of the Cu/Mg bearing intermetallics:
As mentioned earlier, there is always controversy in literature about the fact that the Cu/Mg
bearing intermetallics are soluble/ insoluble6,
11, 200, 294, 306
. To evaluate the capacity to
dissolve a phase, two major factors must be considered: a) whether the phase is
thermodynamically stable at the maximum applicable SHT temperature (
) and b)
whether the dissolution kinetics of the phase is rapid enough during the time period of the
SHT (tSHT). Therefore, initially the stability of phases must be evaluated at equilibrium at
. If they are unstable, then tSHT must be long enough to obtain complete dissolution.
Thermodynamically stable: Figure 7-18(a, b & c) respectively illustrate the (predicted)
equilibrium precipitation temperature for θ, Q and π-phases (in Al-7Si-yCu-xMg-0.1Fe)
145
over which the corresponding Cu/Mg bearing intermetallic is thermodynamically unstable
(TTS). If TTS of a phase is higher than the applicable TSHT, the corresponding phase cannot
be dissolved totally. The horizontal (red) plan was plotted to discriminate the chemistries
for which the TTS is less than 505
is always less than 505
(505 :
at 1st step of SHT). TTS for θ-phase (T )
except for the alloys containing Cu >3.9%; it is why θ-phase is
generally known as a dissolving phase in Al-Si based alloys. TTS for Q-phase (T ) varies
strongly with the Cu and Mg content. For the Al-Si-Mg alloys containing 1%Cu (or less),
T
is less than 505
(regardless of Mg content); but for the alloys containing more Cu,
T
can be higher/lower than 505
depending on the Mg content of the alloys. TTS for π-
phase (T ) is mostly higher than 505
unless the Mg content of the alloy is less than 0.33
%; it is why π-phase is generally known as an insoluble phase. Consequently, to dissolve
the Cu/Mg bearing intermetallics without experience melting, the alloy composition
(mainly the Cu and Mg contents) should be selected so that T
and T
are lower than the
applicable TSHT (~505 ).
Taking into account the criterion of the maximum solubility of Mg in α-Al (at the
applicable TSHT) can be very helpful to design a chemistry (of Al-Si-Cu-Mg alloys) for
which the Cu/Mg bearing intermetallics can all be dissolved. The maximum soluble Mg
content of the AlSiCuMg alloys vs. the applicable TSHT are presented in Figure 7-19. For
instance, the maximum soluble Mg content of the Al-7Si-1.5%Cu alloy at 530
is 0.41 %;
for this chemistry (Al-7Si-1.5Cu-0.41Mg-0.1Fe), the Cu bearing intermetallics (θ- & Qphases) are all thermodynamically unstable at TSHT≥477
(T =476
& T =392 ) and
the remaining Mg bearing intermetallics (i.e. π-phase) can be dissolved at TSHT≥528
(T =527 ). As elaborated earlier in the introduction, Yan et al.14 also suggested a
criterion to optimize the Cu and Mg contents of the Al-Si alloys [4.7<(Cu+10Mg)<5.8
(wt.%)] by which the Q-phase are entirely soluble for the designed chemistries.
Nevertheless the stability of the other Mg bearing intermetallics (e.g. π-phase) was not
considered in this criterion and there are chemistries for which the phase(s) cannot be
dissolved.
146
Figure 7-18: the green plan corresponds to TTS of each phase (the predicted equilibrium precipitation
temperature for θ, Q and π-phases in Al-7Si-yCu-xMg-0.1Fe), the red-plan corresponds to the
(=505 , and the vertical (blue & red) lines correspond to the chemistry of the studied alloys.
Al-7Si-4Cu-0.1Fe-xMg
Max. Mg content of the alloys soluble at
TSHT (wt.%)
Al-7Si-1.5Cu-0.1Fe-xMg
Al-7Si-0.5Cu-0.1Fe-xMg
0.41
0.35
0.29
490
500
510
520
530
540
SHT Temp. (˚C)
550
Figure 7-19: maximum Mg content of the Al-Si alloys soluble at the corresponding temperature; three Al-7Si0.1Fe-yCu-xMg (y=0.5-1.5 & 4%Cu) are compared.
The predicted results are in agreement with the empirical stability of the Cu/Mg bearing
intermetallics. In alloy #3-RC1.6M0.4, T =403
7-20, by applying SHT 6h @ 505
and T =483°C; as shown in Figure
, θ‐phase and the majority of Q‐phase were dissolved. Nevertheless, in alloy #4 (RC1.6M0.8), T =384°C and T =537
and
T =549 ; as shown in Figure 7-10, the θ-phase disappeared after the SHT, but the
majority of large eutectic Mg-bearing phases (mainly Q and π-phases) were remained
intact. Even applying 10 hours SHT at 505 did not considerably change these phases i.e. Q and π) in alloy #4 as displayed in Figure 7-21. Consequently, applying (the 2nd step)
SHT at higher temperatures (e.g. at 525
can cause partial melting of the remaining Cu-
bearing phases (i.e. Q-phase)10, 205, 308.
147
Figure 7-20: alloy #3 (RC1.6M0.4), a) in as-cast condition, b) after 1st step SHT (6h@505 ); the
micrographs were taken at the same coordinate; θ-phase and the majority of Q-phase were dissolved.
Figure 7-21: alloy #4(RC1.6M0.8), a) in as-cast condition, b) after 10 hours SHT at 505 ; the micrographs
were taken at the same coordinate; Q- and π- phases remained almost intact after SHT and Mg2Si was
transformed to Q-phase.
Kinetics of dissolution: The time period of the SHT process (tSHT) plays also a vital role to
dissolve the Cu/Mg bearing intermetallics (θ- & Q-phases). For the alloys having (T
&
T ) < 505 , tSHT must be long enough to dissolve the Cu bearing phases before applying
the 2nd SHT step. As mentioned earlier, θ and Q-phases are both unstable in alloy #3RC1.6M0.4, provided that the time period of the (1st step) SHT to be long enough to
dissolve them completely. Figure 7-22 illustrates microstructure of alloy #3-RC1.6M0.4
after a two-steps SHT (2 hours@505
+ 5hour@525 ); since the time period of 1st step
SHT (2 hours) was not long enough to entirely dissolve the Cu-bearing intermetallics (θ
and Q), the remaining of the Cu-bearing intermetallics (both θ and Q) experienced melting
in the 2nd step SHT (5h@525 ).
148
One of the reason of easier and faster dissolution of θ-phase is that
higher than the T
(~505 ) is much
in the studied alloys; therefore, kinetically there is sufficient driving
force for dissolution. But for the Mg bearing intermetallics (Q & π phases), TTS is either
very close to
(505 ), which necessitates much longer tSHT or is higher than the
in which the phases cannot be dissolved completely.
Figure 7-22: alloy #3 (RC1.6M0.4), a) in as-cast condition, b) after two step SHT (2 hours@505 +
5hour@525 ); the micrographs were taken at the same coordinate; as shown in dashed-dotted area, the Cubearing intermetallics (Q & θ) were melted after the SHT.
Conclusion:

Two etchants were developed to discriminate the Mg bearing intermetallics (Qand π-phases) under optical microscope. After treatment by the etchants, the
stoichiometry of some phases (mainly Mg2Si) was altered.

Mg2Si phase which appeared in as-cast microstructure of the alloys containing
0.8%Mg (#2-RC1M0.8 & #4-RC1.6M0.8) was partially transformed to Q-phase
in alloy #4 after SHT (6h@505 ).

The content of the elements (Cu and Mg) and their ratio (Cu/Mg) play a major
role in the /stability of Mg bearing intermetallics. In the alloys containing 0.4%Mg
(#1-RC1M0.4 & #3- RC1.6M0.4), Q phase could be dissolved; but in the alloys
containing 0.8%Mg (#2-RC1M0.8 & #4- RC1.6M0.8), the majority of Q (and
Mg2Si) phase remained after SHT for 14h @ 505 . π-phase, which was
dissolved in alloy #3-RC1.6M0.4, was partially dissolved in the other alloys.
149

By applying a single step SHT at 525 , most of the Cu-bearing intermetallics (θ
and Q) were locally melted; however, by applying the two-steps SHT
(6h@505 +8h@525 ), they were dissolved in the alloys containing 0.4% Mg
(#1-RC1M0.4 & #3- RC1.6M0.4) and experienced partial melting in the alloys
containing 0.8%Mg (#2-RC1M0.8 & #4- RC1.6M0.8).

In the alloys for which the applicable TSHT is limited to 505 , the Mg content is
recommended to be 0.33% (or less) to minimize/eliminate the undissolved Cu/Mg
bearing intermetallics. According to the thermodynamic prediction, the Mg
content can be enhanced up to 0.43% if the alloy can be SHTed at 540
experience melting.
150
without
Chapter 8
Perspective and general conclusions
This thesis was aimed to study the effect of Cu and Mg as alloying elements and Fe as
impurity on the Al-7 (wt. %) Si based foundry alloys. It was mainly focused on the
influence of the above mentioned elements on solidification defects, and on the evolution
of post eutectic phases during solidification/solution-treatment process. This chapter
summarizes the work taken place within the framework of this project, and the original
aspects of the work are highlighted. In addition, the main conclusions drawn from the work
are outlined. Consequently, some suggestions are provided for future researches to study
with further details the effect of alloying elements on microstructure evolution and
mechanical properties of the Al-Si based foundry alloys.
General conclusions
1. The alloys containing the highest combined Cu and Fe content (i.e. RC3F0.7(M0))
& RC3F0.7) experienced the maximum amount of microporosity and the alloys
with the lowest combined amount of Cu and Fe (RC0.5) presented the minimum
microporosity. Therefore, the amount of microporosity is correlated with the
combined amount of Cu and Fe.
2. A new semi-quantitative indexation method called hot tearing sensitivity (HTS) was
introduced to evaluate the hot tearing susceptibility of the alloys, which was defined
to reflect the volume of generated cracks in the torn specimens. The Al-7Si based
alloys (356Cu and 319-type alloys) are really resistant to hot tearing; at higher
mould temperature ( 250 ), the studied alloys all were resistant to hot tearing. By
reducing the mould temperature, the alloys containing high Cu and Fe content
(RC3F0.7(M0)) & RC3F0.7) were the most sensitive to hot tearing, and the alloys
containing less Cu and Fe content (e.g. RC0.5) were the most resistant to hot
tearing. The enhancement of the hot tearing sensitivity by increasing Fe content was
linked to an increased density/size of lamellar β-Al5FeSi phase, which impede
liquid feeding.
3. The theoretical hot tearing index (HCS) proposed by Katgerman271 was simulated in
the studied alloys using the multiphase back diffusion (MBD) model1. The
temperature, at which 2% of the interdendritic volume is occupied by secondary
phase particles was considered as the critical temperature (Tcr) used in this
theoretical index (HCS). A very good correlation was obtained between the
experimental hot tearing index (HTS) and the theoretical index (HCS).
4. By re-definition of the hot tearing index (CSC= ∆tv/∆tr) originally proposed by
Clyne and Davies270, a new index (βR) was introduced. βR express the ratio of
solidification shrinkage occurring during the vulnerable time period (∆tv) and during
the stress relief time period (∆tr). The alloys with the highest combined amount of
152
Cu and Fe (RC3F0.7(M0) & RC3F0.7) illustrated the maximum βR values, and the
alloy containing the lowest Cu and Fe content (e.g. RC0.5) displayed the minimum
βR value. The correlations of βR with the porosity area% and with the HTS both
were very good.
5. In order to identify the phases involved in the reaction producing each DSC peak, a
procedure of microstructure characterization was developed in which an as-cast
specimen is heated in DSC up to a temperature just beyond the desired peak and
then rapidly cooled. Comparing the evolution of the microconstituents at the same
location before and after the heat treatment helped to correlate the phases involved
during the reaction producing the DSC peak.
6. The microconstituent called “block-like θ-Al2Cu phase” contained a considerable
amount of Fe. This phase is in fact an AlCuFe intermetallic which is presumably a
predecessor of the stable N-Al7Cu2Fe phase. The area fraction of AlCuFeintermetallic was correlated with the time period of SHT. Tough it was negligible in
the as-cast microstructure, it was significantly enhanced after SHT. By heating
(>450 ) the alloys containing sufficiently high amount of Cu (e.g.
3%), α-Al
reacts with β-Al5FeSi and causes the formation of AlCuFe-intermetallic (mostly NAl7CuFe phase).
7. By heating Al-7Si alloy containing 3%Cu in DSC, two peaks appeared at ~521 and
~532.5 . Comparison of the microstructures in the as-cast condition and after
heating the specimen to 527 , showed melting of AlCu eutectic phase and
unmelted AlCuFe-intermetallic. By heating the as-cast specimens to 537 , both the
AlCu eutectic phase and AlCuFe-intermetallic experienced melting. Thus, the peaks
at 521 and 532.5 were correlated to melting of AlCu eutectic phase and AlCuFeintermetallic, respectively. By heating the solution heat treated specimen to 537
in
Al-7Si-3Cu-0.75Fe alloy, N-phase was replaced with β-Al5FeSi and other solid
phases once cooled again at room temperature.
153
a. In Al-7Si-3Cu alloy containing 0.3%Mg, the DSC peak corresponds to the
melting of AlCuFe- ntermetallic appeared at 522 . The results are in good
agreement with the results predicted by MBD model.
8. According to the MBD model, the solidification sequence/temperatures of θ- and Qphases are strongly influenced by Cu and Mg content. In Al-7Si-0.3Mg-3Cu alloy,
θ- and Q-phases both are predicted to solidify at almost the same temperature
(~510 ); but in Al-7Si-0.3Mg alloy containing 1.5%Cu, Q-phase and θ-phase are
respectively solidified at ~526 and 510 . Experimental results seem to be in
agreement with the predicted results:
a. In cooling process of Al-7Si-0.3Mg-1.5Cu, two DSC peaks appeared at
~512 and ~495 ; however, in Al-7Si-0.3Mg alloy containing 3%Cu, the
peaks were merged and appeared at ~499 .
b. In heating process of the Al-7Si-0.3Mg alloys containing 1.5 and 3%Cu, two
DSC peaks appeared at ~507 and ~519 . By heating Al-7Si-0.3Mg-3Cu
alloy to 514 , both θ‐ and Q‐phases were melted; but in the alloy containing 1.5%Cu, only θ‐phase was melted and Q‐phase persisted to remain. This supports the aforementioned MBD model results. 9. According to DSC results and the results predicted by MBD model, the
precipitation/dissolution temperature of π-phase was influenced by the Cu content.
The DSC peak corresponding to π-phase appeared at ~544
in alloys Al-7Si-
0.3Mg containing 0.5%Cu; by increasing Cu content to 1.5%, it solidified at
~535 . These results were confirmed by the MBD model.
10. The temperature and holding time period are the critical parameters of SHT.
a. Lower temperature/holding time might not be sufficient to dissolve the Cubearing intermetallic phases.
b. Higher solution treatment temperature can lead to incipient melting; of the
major characteristics of the specimens having experienced incipient melting
154
are the presence of massive eutectic (θ and Q) phases nearby polygonal Si
particles was.
c. Longer solution treatment not only enhances the production costs, but can
also lead the dissolved elements to be wasted on other phases.
11. Some part of the dissolved Cu in Al matrix is wasted to N-Al7Cu2Fe during SHT.
Longer time period of SHT can lead more dissolved Cu to be wasted. Moreover, the
amount of Cu not available to strengthen the primary phase increases with the
volume fraction of β-Al5FeSi.
12. The stability of Mg bearing intermetallics is strongly correlated with the Mg and Cu
content of the alloy.
a. In the alloys containing low Cu content (e.g. alloy RC1.5), the DSC peak
corresponding to Q-phase disappeared after 5 hours of SHT; but in the alloys
containing high Cu content (RC3), the peak was persisted to remain even after 20
hours of SHT.
b. In the alloys containing 0.4%Mg (RC1M0.4 & RC1.6M0.4), Q phase was
unstable; but in the alloys containing 0.8%Mg (RC1M0.8 & RC1.6M0.8), Q phase
could not be dissolved completely after SHT for 14h @ 505 . Moreover, the
majority of π-phase, which was dissolved in alloy RC1.6M0.4, was only partially
dissolved in the other alloys.
c. By applying a single step SHT at 525 , most of the Cu-bearing intermetallics (θ
and Q) were locally melted; however, by applying the two-steps SHT
(6h@505 +8h@525 ), they were dissolved in the alloys containing 0.4% Mg
(RC1M0.4 & RC1.6M0.4) and experienced partial melting in the alloys containing
0.8%Mg (RC1M0.8 & RC1.6M0.8).
13. The Mg bearing Q- and π-intermetallics both appear as light gray with more/less
similar morphology under optical microscope, which make them impossible to be
differentiated. Two etchants (HNO3 and HNO3+HCl) were developed to
155
discriminate them (Q & π) under optical microscope. The developed etchants
changed the stoichiometry of some phases (mainly Mg2Si).
14. Mg2Si phase was partially transformed to Q-phase while SHT (6h@505 ) in the
alloy RC1.6M0.8.
15. If the applicable SHT temperature is limited to 505 , the Mg content is
recommended to be 0.33% (or less) to minimize/eliminate the un-dissolved Cu/Mg
bearing intermetallics. If the alloy can be solution heat treated at higher temperature
without experience melting, the Mg content can be increased (for instance up to
0.43% if the alloy is solution heat treated at 540 ).
16. Higher Cu and Mg contents of the alloys produce large intermetallics (QAl5Cu2Mg8Si6, π- Al8FeMg3Si6) which cannot be dissolved completely. These large
and insoluble intermetallics have been reported to have two major negative impacts
on mechanical properties: 1) enhance stress concentration while in service which
can affect the strength and ductility of the alloys, 2) decrease the precipitation
hardening efficiency by wasting the hardening elements (Cu and Mg). Furthermore,
higher Cu content (e.g. ~3 wt.%), significantly enhanced the casting defects (hot
tearing susceptibility and porosity area %). Based on the computations and the
experimental results, the range of Cu and Mg is recommended to be 1-1.5 wt.% and
0.3-0.4 wt.%, respectively.
156
Recommendations for future works:
This project was mainly focused to investigate the effect of Cu, Fe and Mg on solidification
defects and microstructure evolution of Al-7Si alloys. Although the effect of the elements
on solidification defects and microstructure evolution was thoroughly studied, further
investigations and characterizations would be suggested to this project:
 To study the mechanism/sequence of precipitation of the post eutectic phases at the last
stage of solidification:

The chemical composition of the liquid right before the DSC peak of the post
eutectic phase is predicted and the specimen is prepared. Subsequently, the
microstructure evolution of the specimen is investigated.
 To study the effect of adding Mn on the precipitation of N-Al7Cu2Fe phase:

By adding Mn to the base alloy, α-Al15(Fe,Mn)3Si2 phase appears in the as-cast
microstructure at the expense of β-Al5FeSi.
 To simulate the dissolution mechanisms of the partially solvable post eutectic phases
(e.g. Q- Al5Cu2Mg8Si6 and π-Al8Mg3FeSi6 phases) with Dictra-ThermoCalc
 To study the effect of various Cu contents (e.g. 0.5, 1, 1.5, 2 and 3 wt. %) on the
mechanical properties (tensile strength, ductility and fatigue strength) of Al-7Si-0.35Mg
(wt. %):
 Optimize the heat treatment procedures (solution treatment and aging process) of
the selected alloys and evaluate the solubility of the post eutectic phases (e.g. π);
 Compare the mechanical properties (tensile, fatigue and creep strength) of the
alloys;
 To study the effect of Si content (between 5 to 9 %) on thermal analysis, microstructure
evolution and mechanical properties (tensile, fatigue and creep strength) of Al-Si based
alloys
 To compare mechanical properties of the optimized alloy(s) in the preceding steps with
the mechanical properties of the secondary Al-Si foundry alloys (containing more
impurities like Cu and Fe); the secondary Al-Si foundry alloys are more cost effective
than primary alloys.
 To study the effect of transition elements (e.g. Sc, Zr, Hf, Mo and Mn) on microstructure
evolution and mechanical properties of the optimized alloy(s) in the preceding step.
157
158
Chapter 9
Appendix (1): calculation of
R
Appendix
(ratio of solidification shrinkage)
In the vulnerable regime, the shrinkage deformation v occurs between the critical
temperature (Tcr) and the solidus (Tsol). This deformation is calculated as follow:
 v 
Tcr

Tsol
d 1  
1
 
 ln  sol 
 cr 
where  is the mass density of the alloy. The shrinkage deformation in the relaxation
regime occurs between the temperature of dendrite coherency (Tcoh) and the critical
temperature (Tcr). Similarly, one can write:
 
 r  ln  cr 
 coh 
The variables sol, cr et coh are respectively the mass density at Tsol, Tcr and Tcoh. Since
R  v r , one obtains:
  sol 
 cr 
 ln 

 cr 
 coh 
 r  ln 
The mass density of the alloy is calculated with the rule of mixture:
1



f


Where f is the mass fraction of phase  as calculated by the multiphase back diffusion
(MBD) model. For the liquid and primary solid phases, the density is adjusted according to
their composition by using these equations:
 liq   Alliq
M liq
M Al
 FCC   AlFCC
M FCC
M Al
 Alliq and  AlFCC are the densities of pure aluminum in respectively the liquid and
solid state. MAl is the molar mass of aluminum and M liq , M FCC are respectively the
Where
average molar masses of the liquid and primary solid phase, which are calculated via the
MBD model. The density of pure aluminum phases are supposed to vary with temperature
according to these equations:
Alliq (g/cm3 )  2.7658  3.935104 T
313
AlFCC (g/cm3 )  2.7233  6.2228105 T 1.23107 T 2
314
where T is in Kelvin.
The mass density of the secondary phases is given in Table 4. A mid-value of 3.45 g/cm³
was chosen for -Al5FeSi. These values were assumed constant in the calculations.
Table 9-1: mass density of the secondary phases in Al-Si based foundry alloys.
Phase
Density (g/cm3) Reference
315
Silicon
2.33
316
Al2Cu
4.35
1
Al5Cu2Mg8Si6
2.34
315
Mg2Si
1.99
317
3.3 – 3.6
-Al5FeSi
317
2.82
Al8FeMg3Si6
160
Appendix (2): Back diffusion model (BDM)
Solidification paths vary between two extreme conditions, 1- global equilibrium and 2- no
diffusion in solid phases (Scheil condition). In the equilibrium solidification path, fractions
of phases are computed based on the equilibrium phase diagram. It is well known that
equilibrium solidification conditions are rarely met; because, in this condition, very high
mass diffusivities are required to achieve close-to-equilibrium conditions in solidification
process. Figure 9-1 (a) presents the variation of chemistry in a cylindrical specimen at three
different steps (solid fraction= 0.25, 0.5 and 0.75) during equilibrium solidification. As the
temperature is lowered more solid forms and, provided cooling is slow enough to allow
extensive solid state diffusion, the solid and liquid will always be homogeneous with
compositions following the solidus and liquidus lines. The relative amounts of solid and
liquid at any temperature are simply given by the “lever rule”.
On the opposite side, we have the Scheil solidification path, where diffusion in the solid is
assumed to be zero. According to this assumption, the amount of liquid at a given
temperature is overestimated in the solidification interval. Figure 9-1 (b) illustrates the
variation of chemistry in a cylindrical specimen at three different steps (solid fraction=
0.25, 0.5 and 0.75) during Scheil solidification. The first solid forms when the cooled end
of the bar reaches the liquidus (T1 in Figure 9-1 (b)). This first solid will be purer than the
liquid from which it forms so that solute is reject into the liquid and raises its concentration.
The temperature of the interface must decrease below T1 before further solidification can
occur, and the next layer of solid will be slightly richer in solute than the first. As this
sequence of events continues the liquid becomes progressively richer in solute and
solidification takes place at progressively lower temperatures (Figure 9-1 (b)). Local
equilibrium can be assumed at the solid/liquid interface during solidification. However,
since there is no diffusion in the solid, the separate layers of solid retain their original
compositions. Thus the mean composition of the solid ( s) is always lower than the
composition at the solid/liquid interface, as shown by the dashed line in Figure 9-1 (b).
The liquid can become richer in solute after each step, and it may even reach a eutectic
composition, XE. Since the last liquid is rich of all solutes rejected by the primary solid
phase, there are a lot of possibilities for secondary phase’s formation, which have to
decrease the expected solidus temperature.
161
Taking into account the back-diffusion of solute in the primary solid phase, the actual
solidification paths are between these two extreme conditions; Figure 9-1 (c) presents the
evolution of the solid/liquid/interface during solidification under back diffusion condition.
It has been proposed in literature that using microsegregation models, we can calculate the
solute composition in the liquid in terms of the fraction of phases. The model of Brody and
Flemings [1] is the most widely used model; their proposed expression can be generalized
to equilibrium and Scheil conditions. For a multicomponent alloy having a nominal
composition CNOM for a given solute i, the Brody–Flemings expression proposed by Kurz
and Fisher [2] can be written as follow:
(
1
1
2
(1)
where Cl is the solute content in the liquid phase, ki is the equilibrium partition coefficient,
fs is the mass fraction of solid and i is the back-diffusion parameter taking values between
0 and 0.5, for respectively Scheil (no back-diffusion) and global equilibrium conditions; for
the values in between, a certain amount of solute diffuses in the primary solid phase, but
with a very slow rate to reach equilibrium. The amount of solute diffusing in the solid
phase depends on the mass diffusivity, the solidification time and a solidification
characteristic length.
The partition coefficient ki in Eq. (1) was presumed to be constant to obtain an integrable
form of the incremental mass conservation equation. As stated by Chang [3], this
assumption can be acceptable in binary alloys; but in the multicomponent alloys, it can
produce serious errors. The multiphase back diffusion model (BDM) presents a scheme to
resolve the problem Eq. (1), and links a thermodynamics computational tool like ThermoCalc [21] to a mass conservation equation to do the computation.
In BDM computation, the chemistry of the alloys, size of system (λ/2, λ: dendrite arm
spacing), solidification rate, are taken from experimental results. The geometry can be
assumed plate (G=1) or columnar (G = 2) depending the solidification conditions. All the
calculations, here in this thesis, were made with a temperature step of 0.25 K.
162
T
T1
T2 X
s
T3
T
XL
L
Solid fraction= 25%
Solid fraction= 50%
Solid fraction= 75%
TE
X
X
X
X1
X0
Xmax
XE
X1
X1
distance
X
distance
distance
(a)
(b)
(c)
Figure 9-1: Calculated composition profiles of a specimen obtained at 3 different solidification steps (solid
fractions: 0.25, 0.50 and 0.75), a) in equilibrium condition, b) in Scheil condition, c) in BDM condition.
163
164
Bibliography
1.
D. Larouche, Calphad, 2007, Vol.31(4), pp.490-504.
2.
D. Liu, H. V. Atkinson, and H. Jones, Acta Materialia 2005, Vol.53 pp.3807-3819.
3.
R. Mackay and J. H. Sokolowski, International Journal of Cast Metals Research, 2010 Vol.23(1),
pp.7-22.
4.
G. A. Edwards, G. K. Sigworth, C. H. Cáceres, D. H. StJohn, and J. Barresi, AFS Transactions, 1997,
Vol.105, pp.809-818.
5.
R. Mackay, J. H. Sokolowski, R. Hasenbusch, and W. J. Evans, AFS Transactions, 2003, Vol.111,
pp.229-240.
6.
G. Wang, X. Bian, W. Wang, and J. Zhang, Materials Letters, 2003, Vol.57(24-25), pp.4083-4087.
7.
L. Backerud, G.Chai, and J.Tamminen: 'Solidification Characteristics of Aluminum Alloys', 2-Foundry
Alloys, 115-150; 1990, USA, Skan Aluminum.
8.
M. H. Mulazimoglu, N. Tenekedjev, B. M. Closset, and J. E. Gruzleski, Cast Metals, 1993, Vol.6(1),
pp.16-28.
9.
E. Samuel, A. Samuel, and H. Doty, AFS Trans, 1996, Vol.104, pp.893-901.
10.
F. H. Samuel, Journal of Materials Science, 1998, Vol.33(9), pp.2283-2297.
11.
H. R. Ammar, A. M. Samuel, F. H. Samuel, G. K. Sigworth, and J. C. Lin, AFS Transactions, 2010,
Vol.118, pp.9-27.
12.
E. Cerri, E. Evangelista, S. Spigarelli, P. Cavaliere, and F. DeRiccardis, Materials Science and
Engineering A, 2000, Vol.284(1-2), pp.254-260.
13.
A. Lise Dons, Journal of Light Metals, 2001, Vol.1(2), pp.133-149.
14.
Xinyan Yan and Jen C. Lin: US Patent, 0105045 A1, May 2013.
15.
P. Ouellet, F. H. Samuel, D. Gloria, and S. Valtierra, International Journal of Cast Metals Research,
1997, Vol.10, pp.67-78.
16.
C. Halászi: 11th European Automotive Congress (EAEC), Budapest, Austria, 2007, GTE.
17.
L. Remy: 'Thermal-Mechanical Fatigue (Including Thermal Shock)', Centre des Materiaux, UMR
CNRS 7633, Ecole des Mines de Paris, France, 1999.
18.
W. Z. Zhuang and N. S. Swansson: 'Thermo-Mechanical Fatigue Life Prediction: A Critical Review',
DSTO Aeronautical and Maritime Research Laboratory, Melbourne Victoria Australia, 1998.
19.
R. Neu and H. Sehitoglu, Metallurgical and Materials Transactions A, 1989, Vol.20(9), pp.1755-1767.
20.
M. Riedler, H. Leitner, B. Prillhofer, G. Winter, and W. Eichlseder, Meccanica, 2007, Vol.42(1),
pp.47-59.
21.
R. Minichmayr, M. Riedler, G. Winter, H. Leitner, and W. Eichlseder, International Journal of
Fatigue, 2008, Vol.30(2), pp.298-304.
22.
X. Su, M. Zubeck, J. Lasecki, H. Sehitoglu, C. C. Engler-Pinto, C. Y. Tang, and J. E. Allison:
Thermomechanical Fatigue Behavior of Materials, ASTM International, W. Conshohocken, PA, 2003,
240-254.
23.
E. Wilfried, W. Gerhard, M. Robert, and R. Martin: Recent Trends in Processing and Degradation of
Aluminium Alloys, InTech, Montanuniversität Leoben, Austria, 2011, 329-346.
24.
D. Löhe, T. Beck, and K. H. Lang: Fifth International Conference on Low Cycle Fatigue, Berlin,
Germany, 2004, Deutscher Verband für Materialforschung und -prüfung, 161-175.
25.
R. Tichánek, M. Španiel, and M. Divi, Acta Polytechnica, 2005, Vol.45 (3), pp.43-48.
26.
M. Španiel and R. Tichánek, MECCA Journal of Middle European Construction and Design of Cars,
2004, Vol.2(2).
27.
R. Fuoco and M. F. Moreira, AFS Trans, 2009, Vol.117, pp.225-240.
28.
S. Thalmair, J.Thiele, A. Fischersworring-Bunk, R. Ehart, and M. Guillou, SAE Technical Paper,
2006, Vol.01-0541.
29.
M. H. Shojaefard, M. R. Ghaffarpour, A. R. Noorpoor, and S. Alizadehnia, Proceedings of the
Institution of Mechanical Engineers, Part D: Journal of Automobile Engineering, 2006 Vol.220(5),
pp.627-636.
30.
K. S. Lee, D. N. Assanis, J. Lee, and K. M. Chu, SAE Technical Paper, 1999, Vol.01-0973.
31.
M. Shalev, Y. Zvirin, and A. Stotter, International Journal of Mechanical Sciences, 1983, Vol.25(7),
pp.471-483.
32.
J. J. Thomas, L. Verger, A. Bignonnet, and E. Charkaluk, Fatigue & Fracture of Engineering
Materials & Structures, 2004, Vol.27(10), pp.887-895.
33.
R. Bertodo and T. J. Carter, journal of strain analysis 1971, Vol.6(1), pp.1-12.
34.
T. Takahashi, T. Nagayoshi, M. Kumano, and K. Sasaki, SAE Technical Paper, 2002, Vol.01-0585.
35.
P. M. Norris, K. L. Hoag, and W. Wepfe, Experimental Heat Transfer, 1994, Vol.7(1), pp.43 - 53.
36.
T. Takahashi, K. Sasaki, and M. Iida: 7th Asia Pacific Industrial Engineering and Management
Systems Conference, Bangkok, Thailand, 2006, 1849-1859.
37.
M. Garat and G. Laslaz, AFS Transactions, 2007, Vol.115, pp.89-96.
38.
V. Firouzdor, M. Rajabi, E. Nejati, and F. Khomamizadeh, Materials Science and Engineering: A,
2007, Vol.454-455, pp.528-535.
39.
M. M. Rahman, A. K. Ariffin, S. Abdullah, A. B. Rosli, and M. S. M. Sani: Proceedings of the
International Conference on Mechanical Engineering Dhaka, Bangladesh, 2007
40.
S. Trampert, T. Gocmez, and S. Pischinger, Journal of Engineering for Gas Turbines and Power,
2008, Vol.130(1), pp.6-10.
41.
R. B. Gundlach, B. Ross, A. Hetke, S. Valtierra, and J. F. Mojica, AFS Trans, 1994, Vol.102, pp.205223.
42.
R. Chuimert and M. Garat: 3rd International Symposium Aluminium + Automobile, Dusseldorf, Feb.
1988, 154-159.
43.
T. Takahashi and K. Sasaki, Procedia Engineering, 2010, Vol.2(1), pp.767-776.
44.
T. Takahashi and K. Sasaki, Society of Automotive Engineers, 1998 Vol.107(980688), pp.454–461.
45.
S. C. Lee and L. C. Weng, Metallurgical and Materials Transactions A, 1991, Vol.22(8), pp.18211831.
46.
G. E. Dieter: 'Mechanical Metallurgy'; 1988, McGraw-Hill, New York.
47.
F. J. Feikus, AFS Transactions, 1998, Vol.106, pp.225- 231.
166
48.
H. Sehitoglu, T. Smith, X. Qing, H. Maier, and J. Allison, Metallurgical and Materials Transactions
A, 2000, Vol.31(1), pp.139-151.
49.
J. Boileau and J. Allison, Metallurgical and Materials Transactions A, 2003, Vol.34(9), pp.1807-1820.
50.
J. M. Boileau, J. W. Zindel, L. A. Godlewski, J. E. Allison, and K. A. Kofeldt, SAE Technical Paper,
2007, Vol.01-1224.
51.
Y. Jang, Y. Jeong, C. Yoon, and S. Kim, Metallurgical and Materials Transactions A, 2009,
Vol.40(5), pp.1090-1099.
52.
A. Hany: 'Influence of metallurgical parameters on the mechanical properties and quality indices of AlSi-Cu-Mg and Al-Si-Mg casting alloys', PhD thesis, Universite du Quebec a Chicoutimi, Quebec,
Canada, 2010.
53.
K. Mollenhauer and H. Tschoeke, eds. 'Handbook of Diesel Engines', (ed.K. Mollenhauer, et al.), 3096; 2010, New York, Springer.
54.
J. Sullivan, R. Baker, B. Boyer, R. Hammerle, T. Kenney, L. Muniz, and T. Wallington,
Environmental Science & Technology, 2004, Vol.38(12), pp.3217-3223.
55.
T. Denton: 'Automobile Mechanical and Electrical Systems: Automotive Technology: Vehicle
Maintenance and Repair', 81-258; 2011, Waltham, MA, Elsevier.
56.
T. Smith, H. Sehitoglu, E. Fleury, H. Maier, and J. Allison, Metallurgical and Materials Transactions
A, 1999, Vol.30(1), pp.133-146.
57.
R. Molina, M. Leghissa, and L. Mastrogiacomo, Metallurgical Science and Technology, Teksid
Aluminum 2004, Vol.22(2), pp.2-8.
58.
W. L. Guesser: 'Compacted Graphite Iron – a new material for diesel engine cylinder blocks', Brazilian
MRS Meeting, Rio de Janeiro, 2003.
59.
B. S. Andersson, Tribology Series, 1991, Vol.18, pp.503-506.
60.
M. Priest and C. M. Taylor, Wear, 2000, Vol.241(2), pp.193-203.
61.
K. Funatani, K. Kurosawa, P. A. Fabiyi, and M. F. Puz, SAE Technical Paper, 1994(940852), pp.8996.
62.
V. D. N. Rao, D. M. Kabat, H. A. Cikanek, C. A. Fucinari, and G. Wuest, SAE Technical Paper, 1997,
Vol.970023, pp.99-124.
63.
A. Datta, J. D. Carpenter, R. D. Ott, and P. J. Blau, SAE Technical Paper, 2002, Vol.01-0490.
64.
M. James, J. M. Kihiu, G. O. Rading, and J. K. Kimotho: Sustainable Research and Innovation
Conference Proceedings, Vol. (3) 2011.
65.
S. Dawson, China Foundry, 2009, Vol.6(3), pp.241-246.
66.
M. Medraj and A. Parvez, Automotive 2007, Vol.45, pp.45-47.
67.
J. E. Allison and G. S. Cole, JOM, January 1993, Vol.45, pp.19-24.
68.
M. K. Kulekci, Int J Adv Manuf Technol, 2008, Vol.39, pp.851-865.
69.
F. H. Froes, D. Eliezer, and E. Aghion, JOM, 1998, Vol.50(9), pp.30-34.
70.
K. K. Chawla and N. Chawla: Automotive Composites, in John Wiley & Sons, G. Nicolais, 2012,
Encyclopedia of Composites, Hoboken, New Jersey.
71.
A. H. Musfirah and A. G. Jaharah, Journal of Applied Sciences Research, 2012, Vol.4(9), pp.48654875.
72.
A. A. Luo, JOM, February 2002, Vol.54(2), pp.42-48.
73.
G. S. Cole and A. M. Sherman, Materials Characterization, 1995, Vol.35(1), pp.3-9.
167
74.
B. Bronfin, E. Aghion, F. v. Buch, S. Schumann, and H. Friedrich: Magnesium: Proceedings of the 6th
International Conference Magnesium, Weinheim, 2004.
75.
B. Bronfin, M. Katsir, O. Bar-Yosef, F. Moll, and S.Schumann: TMS (The Minerals, Metals &
Materials Society), Magnesium Technology, 2005, 395-401.
76.
M. O. Pekguleryuz and A. A. Kaya: TMS (The Minerals, Metals & Materials Society), Magnesium
Technology, 2004, 281-287.
77.
C. H. Caceres, Metallurgical and Materials Transactions A, 2007, Vol.38(7), pp.1649-1662.
78.
S. Das, JOM, 2008, Vol.60(11), pp.63-69.
79.
V. Kevorkijan, Metalurgija - Journal of Metallurgy, 2002, Vol.8(3), pp.251-258.
80.
M. A. Meyers and K. K. Chawla: 'Mechanical Behavior of Materials'; 2009, Cambridge, Cambridge
University Press.
81.
W. L. Guesser, I. Masiero, E. Melleras, and C. Cabezas, Society of Automotive Engineers, 2004.
82.
R. Marquard and H. Sorger: CGI Design and Machining Workshop, Sintercast, PTW Darmstadt, Bad
Homburg, Germany, 1997.
83.
S. Thalmair, A. Fischersworring-Bunk, F. J. Klinkcnberg, K. H. Lang, and M. Frikha, Bulletin - Cercle
d'Études des Metaux Saint Etienne, 2008, Vol.17(17), pp.83-90.
84.
L. Heusler, F. J. Feikus, and M. O. Otte, AFS Transactions, 2001, Vol.109, pp.443-451.
85.
C. C. Engler-Pinto, J. V. Lasecki, J. M. Boileau, and J. E. Allison, SAE Technical Paper, 2004, Vol.011029, pp.459-464.
86.
Z. Li, A. M. Samuel, F. H. Samuel, C. Ravindran, S. Valtierra, and H. W. Doty, Materials Science and
Engineering A, 2004, Vol.367(1-2), pp.96-110.
87.
M. Tash, F. H. Samuel, F. Mucciardi, and H. W. Doty, Materials Science and Engineering: A, 2007,
Vol.443(1-2), pp.185-201.
88.
J. L. Jorstad, AFS Transactions, 2009, Vol.117, pp.241-249.
89.
Specification for Aluminum- Alloy Sand Castings, 2001,ASTM B 26/B 26M- 01, US.
90.
V. S. Zolotorevsky, N. A. Belov, and M. V. Glazoff: 'Casting Aluminum Alloy', 1-95; 2007, Elsevier
Science.
91.
J. Hatch: 'Aluminum: properties and physical metallurgy'; 1984, American Society for Metals.
92.
S. Shankar, Y. Riddle, and M. Makhlouf, Metallurgical and Materials Transactions A, 2004,
Vol.35(9), pp.3038-3043.
93.
M. M. Makhlouf and H. V. Guthy, Journal of Light Metals, 2001, Vol.1(4), pp.199-218.
94.
K. Gall, N. Yang, M. Horstemeyer, D. McDowell, and J. Fan, Metallurgical and Materials
Transactions A, 1999, Vol.30(12), pp.3079-3088.
95.
J. Luft, T. Beck, and D. Löhe: International Congress on Fracture, University of Karlsruhe, Germany,
2005.
96.
S. Khan and R. Elliott, Journal of Materials Science, 1996, Vol.31(14), pp.3731-3737.
97.
B. Cantor and K. O'Reilly, eds. 'Solidification and Casting', (ed.B. Cantor, et al.), 326-338; 2003,
Bristol and Philadelphia, Institute of Physics Publishing.
98.
M. A. Moustafa, F. H. Samuel, and H. W. Doty, Journal of Materials Science, 2003, Vol.38(22),
pp.4507-4522.
99.
T. S. Furlan and R. Fuoco, AFS Transactions, 2008 Vol.116, pp.281-297.
168
100. Shu-Zu Lu and A. Hellawell, Metallurgical and Materials Transactions A, 1987, Vol.18(10), pp.17211733.
101. N. Fat-Halla, Journal of Materials Science, 1989, Vol.24(7), pp.2488-2492.
102. W. Wang: 'Na, Sr and Sb interactions in Al-Si alloy melts and their effects on modification', PhD
thesis, McGill University, 1991.
103. S. Kumari, R. M. Pillai, and B. C. Pai, Journal of Alloys and Compounds, 2008, Vol.460(1-2), pp.472477.
104. S. Kumari, R. Pillai, B. Pai, K. Nogita, and A. Dahle, Metallurgical and Materials Transactions A,
2006, Vol.37(8), pp.2581-2587.
105. A. K. Dahle, K. Nogita, S. D. McDonald, C. Dinnis, and L. Lu, Materials Science and Engineering: A,
2005, Vol.413–414(0), pp.243-248.
106. W. Prukkanon, N. Srisukhumbowornchai, and C. Limmaneevichitr, Journal of Alloys and Compounds,
2009, Vol.477(1–2), pp.454-460.
107. W. Prukkanon, N. Srisukhumbowornchai, and C. Limmaneevichitr, Journal of Alloys and Compounds,
2009, Vol.487(1–2), pp.453-457.
108. K. Nogita, S. D. McDonald, and A. K. Dahle, Materials Transactions, 2004, Vol.45(2), pp.323-326.
109. H. V. Guthy: 'Evolution of the Eutectic Microstructure in Chemically Modified and Unmodified
Aluminum Silicon Alloys ', M.Sc. thesis, Worcester Polytechnic Institute 2002.
110. M. Timpel, N. Wanderka, R. Schlesiger, T. Yamamoto, N. Lazarev, D. Isheim, G. Schmitz, S.
Matsumura, and J. Banhart, Acta Materialia, 2012, Vol.60(9), pp.3920-3928.
111. O. Elsebaie: 'Effects of strontium-modification, iron-based intermetallics and aging conditions on the
impact toughness of Al-(6--11)%Si alloys', PhD thesis, Universite du Quebec a Chicoutimi,
Chicoutimi, Canada, 2010.
112. D. Emadi, J. E. Gruzleski, and J. M. Toguri, Metall. Trans. B, 1993, Vol.24B, pp.1-9.
113. C. Dinnis, A. Dahle, J. Taylor, and M. Otte, Metallurgical and Materials Transactions A, 2004,
Vol.35(11), pp.3531-3541.
114. J. E. Gruzleski and B. M. Closset: 'The treatment of liquid aluminum-silicon alloys'; 1990, Des Plaines,
Illinois, American Foundrymens Society.
115. D. A. Lados, D. Apelian, and A. M. d. Figueredo: Proceedings from the 2nd International Aluminum
Casting Technology Symposium, Columbus, OH, ASM International, 2002.
116. G. K. Sigworth, AFS Transactions, 2008 Vol.116, pp.115-139.
117. S. Hegde and K. Prabhu, Journal of Materials Science, 2008, Vol.43(9), pp.3009-3027.
118. J. G. Kaufman and E. L. Rooy: 'Aluminum Alloy Castings: Properties, Processes, and Applications';
2004, ASM International, Materials Park.
119. J. A. Lee: NASA Marshall Space Flight Center - Huntsville, AL, Personal communication, 1999.
120. Y. Haizhi, Journal of Materials Engineering and Performance, 2003, Vol.12 pp.288-297.
121. H. G. Kang, M. Kida, H. Miyahara, and K. Ogi, AFS Trans, 1999, Vol.107, pp.507–515.
122. J. A. Taylor: 35th Australian Foundry Institute National Conference, Adelaide , South Australia, 2004,
Australian Foundry Institute (AFI), 148-157.
123. L. F. Mondolfo: 'Aluminum Alloys: Structure and Properties'; 1976, Butterworths, London,
124. N. A. Belov, D. G. Eskin, and A. A. Aksenov: 'Multicomponent Phase Diagrams: Applications for
Commercial Aluminum Alloys'; 2005, Oxford: Elsevier.
169
125. D. Yang: 'Role of magnesium addition on the occurrence of incipient melting in experimental and
commercial aluminum-silicon-copper alloys and its influence on the alloy microstructure and tensile
properties', M.Sc. thesis, Universite du Quebec a Chicoutimi,Quebec, Canada, 2006.
126. M. O. Otte, S. D. McDonald, J. A. Taylor, D. H. St.John, and W. Schneider, Transactions of the
American Foundrymen's 1999, Vol.107, pp.471.
127. J. Taylor, G. Schaffer, and D. StJohn, Metallurgical and Materials Transactions A, 1999, Vol.30(6),
pp.1657-1662.
128. P. N. Crepeau, Transactions of the American Foundrymen's Society, 1995, Vol.103, pp.361.
129. S. Tabibian, E. Charkaluk, A. Constantinescu, A. Oudin, and F. Szmytka, Procedia Engineering, 2010,
Vol.2(1), pp.1145-1154.
130. H. Y. Kim, S. W. Han, and H. M. Lee, Materials Letters, 2006, Vol.60(15), pp.1880-1883.
131. L. Lu and A. Dahle, Metallurgical and Materials Transactions A, 2005, Vol.36(3), pp.819-835.
132. H. Y. Kim, T. Y. Park, S. W. Han, and H. M. Lee, Journal of Crystal Growth, 2006, Vol.291(1),
pp.207-211.
133. Y. Awano and Y. Shimizu, AFS Transactions, 1990, Vol.98, pp.889-895.
134. S. S. Kumari, R. M. Pillai, and B. C. Pai, Journal of Alloys and Compounds, 2008, Vol.453(1-2),
pp.167-173.
135. S. Seifeddine, S. Johansson, and I. L. Svensson, Materials Science and Engineering: A, 2008,
Vol.490(1–2), pp.385-390.
136. W. Shuncai, L. Chunzhi, and Y. Minggao, Materials Research Bulletin, 1989, Vol.24(10), pp.12671270.
137. J. Y. Hwang, H. W. Doty, and M. J. Kaufman, Philosophical Magazine, 2008, Vol.88(4), pp.607-619.
138. L. Arnberg, L. Backerud, and G. Chai: 'Solidification Characteristics of Aluminium Alloys', 3Dendrite Coherency; 1996, Des Plaines, IL, American Foundryman's Society.
139. R. I. MacKay and J. E. Gruzleski, International Journal of Cast Metals Research, 1998, Vol.10(5),
pp.255-265.
140. Q. G. Wang and C. J. Davidson, Journal of Materials Science, 2001, Vol.36(3), pp.739-750.
141. E. J. Martínez, M. A. Cisneros, S. Valtierra, and J. Lacaze, Scripta Materialia, 2005, Vol.52(6),
pp.439-443.
142. L. A. Dobrzański, R. Maniara, M. Krupiński, and J. H. Sokolowski, Journal of Achievements in
Materials and Manufacturing Engineering, 2007, Vol.24(2), pp.51-54.
143. F. Samuel, A. Samuel, P. Ouellet, and H. Doty, Metallurgical and Materials Transactions A, 1998,
Vol.29(12), pp.2871-2884.
144. A. Samuel, J. Gauthier, and F. Samuel, Metallurgical and Materials Transactions A, 1996, Vol.27(7),
pp.1785-1798.
145. M. Djurdjevic, T. Stockwell, and J. Sokolowski, Int J Cast Metals Res, 1999, Vol.12, pp.67-73.
146. Z. Li, A. M. Samuel, F. H. Samuel, C. Ravindran, and S. Valtierra, Journal of Materials Science, 2003,
Vol.38(6), pp.1203-1218.
147. M. F. Ibrahim, E. Samuel, A. M. A. Mohamed, A. M. Samuel, A. M. A. Al-Ahmari, and F. H. Samuel,
AFS Transactions, 2011, Vol.119, pp.149-170.
148. Q. Wang, Metallurgical and Materials Transactions A, 2004, Vol.35(9), pp.2707-2718.
149. Q. G. Wang, D. Apelian, and D. A. Lados, Journal of Light Metals, 2001, Vol.1(1), pp.85-97.
170
150. T. Beck, D. Löhe, J. Luft, and I. Henne, Materials Science and Engineering: A, 2007, Vol.468-470,
pp.184-192.
151. Q. G. Wang, D. Apelian, and D. A. Lados, Journal of Light Metals, 2001, Vol.1(1), pp.73-84.
152. R. Neu and H. Sehitoglu, Metallurgical and Materials Transactions A, 1989, Vol.20(9), pp.1769-1783.
153. C. C. Engler-Pinto Jr, H. Sehitoglu, H. J. Maier, and T. J. Foglesong, European Structural Integrity
Society, 2002, Vol.Volume 29, pp.3-13.
154. P. Mohanty, F. Samuel, and J. Gruzleski, Metallurgical and Materials Transactions A, 1993,
Vol.24(8), pp.1845-1856.
155. C. Chama, Journal of Materials Engineering and Performance, 1992, Vol.1(6), pp.773-779.
156. J. M. Zeng, Z. B. Xu, and J. He, Advanced Materials Research, 2008, Vol.51(93), pp.93-98.
157. C. H. Cáceres, M. B. Djurdjevic, T. J. Stockwell, and J. H. Sokolowski, Scripta Materialia, 1999,
Vol.40(5), pp.631-637.
158. S. G. Shabestari and H. Moemeni, Journal of Materials Processing Technology, 2004, Vol.153-154,
pp.193-198.
159. N. Roy, A. Samuel, and F. Samuel, Metallurgical and Materials Transactions A, 1996, Vol.27(2),
pp.415-429.
160. J. F. Major, AFS Transactions, 1997, Vol.105, pp.901-906.
161. H. Arami, R. Khalifehzadeh, M. Akbari, and F. Khomamizadeh, Materials Science and Engineering:
A, 2008, Vol.472(1-2), pp.107-114.
162. M. K. Surappa, E. Blank, and J. C. Jaquet, Scripta Metallurgica, 1986, Vol.20(9), pp.1281-1286.
163. Z. Ma: 'Effect of iron-intermetallics and porosity on tensile and impact properties of aluminum-siliconcopper and aluminum- silicon-magnesium cast alloys', PhD thesis, Universite du Quebec a Chicoutimi,
Quebec, Canada, 2002.
164. B. Skallerud, T. Iveland, and G. Härkegård, Engineering Fracture Mechanics, 1993, Vol.44(6),
pp.857-874.
165. M. E. Seniw, M. E. Fine, E. Y. Chen, M. Meshli, and J. Gray: Paul C. Paris International Symposium
‘Fatigue of Materials’, TMS-ASM Fall Meeting, 1997.
166. J. G. Conley, B. Moran, and J. Gray, SAE Technical Paper, 1998, Vol.980455.
167. A. Wickberg, G. Gustafsson, and L. E. Larsson, SAE Technical Paper, 1984(840121).
168. W. Chen, B. Zhang, T. Wu, D. Poirier, P. Sung, and Q.T. Fang: The 1st International Al Casting Tech.
Symposium, Rosemont, 1998.
169. W. Chen, B. Zhang, T. Wu, D. Poirier, P. Sung, and Q. T. Fang: Automotive Alloys II, TMS,
Warrendale, 1998, 99-113.
170. M. E. Seniw, J. G. Conley, and M. E. Fine, Materials Science and Engineering A, 2000, Vol.285(1-2),
pp.43-48.
171. M. F. Hafiz and T. Kobayashi, Scripta Metallurgica et Materialia, 1994, Vol.30(4), pp.475-480.
172. M. F. Hafiz and T. Kobayashi, Journal of Materials Science, 1996, Vol.31(23), pp.6195-6200.
173. F. T. Lee, J. F. Major, and F. H. Samuel, Fatigue Fract. Eng. Mater. Struct, 1995, Vol.18(3), pp.385396.
174. F. Lee, J. Major, and F. Samuel, Metallurgical and Materials Transactions A, 1995, Vol.26(6),
pp.1553-1570.
175. H. Toda, J. Katano, T. Kobayashi, T. Akahori, and M. Niinomi, Material Trans., 2005, Vol.46(1),
pp.111-117.
171
176. Y. Jang, S. Jin, Y. Jeong, and S. Kim, Metallurgical and Materials Transactions A, 2009, Vol.40(7),
pp.1579-1587.
177. X. Zhu, A.Shyam, J. W. Jones, H. Mayer, J. V.Lasecki, and J. E. Allison, International Journal of
Fatigue, 2006, Vol.28(11), pp.1566-1571.
178. C. Nyahumwa, N. Green, and J. Campbell, Metallurgical and Materials Transactions A, 2001,
Vol.32(2), pp.349-358.
179. Q. Wang, P. Crepeau, C. Davidson, and J. Griffiths, Metallurgical and Materials Transactions B,
2006, Vol.37(6), pp.887-895.
180. B. Zhang, D. Poirier, and W. Chen, Metallurgical and Materials Transactions A, 1999, Vol.30(10),
pp.2659-2666.
181. D. S. Jiang, L.-H. Chen, and T.-S. Liu, Materials Transactions, JIM, 2000, Vol.41(4), pp.499-506.
182. Y. Jang, S. Jin, Y. Jeong, and S. Kim, Metallurgical and Materials Transactions A, 2010,
Vol.41A(19), pp.19-21.
183. J. Campbell, Mater. Sci. Technol., 2006, Vol.22(2), pp.127-145.
184. J. Campbell, Metallurgical and Materials Transactions A, 2010 Vol.41(1), pp.18.
185. M. Angeloni: 'Fatigue Life Evaluation of A356 Aluminum Alloy Used for Engine Cylinder Head', PhD
thesis, Sao Paulo, 2011.
186. X. Pan, C. Lin, H. Brody, and J. Morral, Journal of Phase Equilibria and Diffusion, 2005, Vol.26(3),
pp.225-233.
187. Z. Ma, E. Samuel, A. M. A. Mohamed, A. M. Samuel, F. H. Samuel, and H. W. Doty, Materials &
Design, 2010, Vol.31(8), pp.3791-3803.
188. J. M. Boileau, C. A. Cloutier, L. A. Godlewski, P. A. R. Symansksi, C. Wolverton, and J. E. Allison,
SAE Technical Paper, 2003, Vol.01·0822.
189. J. Gauthier, P. Louchez, and F. Samuel, International Journal of Cast Metals Research, 1995,
Vol.8(2), pp.91-106.
190. S. Imurai, J. Kajornchaiyakul, C. Thanachayanont, J. T. H. Pearce, and T. Chairuangsri, Chiang Mai J.
Sci. , 2010, Vol.37(2), pp.269-281.
191. F. J. T. Medrano, S. Valtierra, J. E. Gruzleski, F. H. Samuel, and H. W. Doty, AFS Transactions, 2008,
Vol.116(02), pp.99-114.
192. E. Sjölander and S. Seifeddine, Materials & Design, 2010, Vol.31(Supplement 1), pp.S44-S49.
193. G. E. Totten and D. S. MacKenzie: 'Handbook of Aluminum, Volume 1, Physical Metallurgy and
Processes', 259-305; 2003, New York, CRC Press.
194. A. M. A. Mohamed, F. H. Samuel, and S. Al kahtani, Materials Science and Engineering: A, 2012,
Vol.543(0), pp.22-34.
195. E. Sjölander and S. Seifeddine, Journal of Materials Processing Technology, 2010, Vol.210(10),
pp.1249-1259.
196. R. Fuoco and E. R. Correa, AFS Transactions, 2002, Vol.110, pp.417-434
197. J. H. Sokolowski, X. C. Sun, G. Byczynski, D. O. Northwood, D. E. Penrod, R. Thomas, and A.
Esseltine, Journal of Materials Processing Technology, 1995, Vol.53(1-2), pp.385-392.
198. J. H. Sokolowski, M. B. Djurdjevic, C. A. Kierkus, and D. O. Northwood, Journal of Materials
Processing Technology, 2001, Vol.109(1-2), pp.174-180.
199. Y. M. Han, A. M. Samuel, F. H. Samuel, and H. W. Doty, International Journal of Cast Metals
Research, 2008, Vol.21 (5), pp. 387-393.
200. L. Lasa and J. M. Rodriguez-Ibabe, Materials Characterization, 2002, Vol.48(5), pp.371-378.
172
201. M. A. Azmah Hanim, S. Chang Chung, and O. Khang Chuan, Materials & Design, 2011, Vol.32(4),
pp.2334-2338.
202. H. d. l. Sablonnière and F. H. Samuel, International Journal of Cast Metals Research, 1996, Vol.9,
pp.151-225.
203. B. K. Prasad and T. K. Dan, Journal of Materials Science Letters, 1991, Vol.10(23), pp.1412-1414.
204. S. Shivkumar, J. S. Ricci, C. Keller, and D. Apelian, J. Heat Treat, 1990, Vol.8, pp.63-70.
205. Y. Han, A. Samuel, F. Samuel, S. Valtierra, and H. Doty, AFS Transactions, 2008, Vol.116, pp.79-90.
206. M. L. Newman: 'Modeling the behavior of a type-319 aluminum alloy during quenching', PhD thesis,
University of Illinois at Urbana-Champaign, United States 2002.
207. D. Zhang and L. Zheng, Metallurgical and Materials Transactions A, 1996, Vol.27(12), pp.3983-3991.
208. G. E. Totten and D. S. Mackenzie, Materials Science Forum, 2000, Vol.331-337, pp.589-594.
209. F. C. Campbell: 'Manufacturing Technology for Aerospace Structural Materials'; 2006 Oxford,
Elsevier.
210. S. Weakley-Bollin, W. Donlon, C. Wolverton, J. Allison, and J. Jones, Metallurgical and Materials
Transactions A, 2004, Vol.35(8), pp.2407-2418.
211. A. Wiengmoon, J. T. H. Pearce, T. Chairuangsri, S. Isoda, H. Saito, and H. Kurata, Micron, 2013,
Vol.45(0), pp.32-36.
212. M. Song, Materials Science and Engineering: A, 2007, Vol.443(1-2), pp.172-177.
213. M. I. N. Song, X. I. A. Li, and K. Chen, Metallurgical and Materials Transactions A, 2007, Vol.38(3),
pp.638-648.
214. A. Ardell, Metallurgical and Materials Transactions A, 1985, Vol.16(12), pp.2131-2165.
215. J. Gauthier, P. Louchez, and F. Samuel, International Journal of Cast Metals Research, 1995,
Vol.8(2), pp.107-114.
216. A. Zanada and G. Riontino, materials Science Forum, 2000, Vol.331-337, pp.229-234.
217. R. X. Li, R. D. Li, Y. H. Zhao, L. Z. He, C. X. Li, H. R. Guan, and Z. Q. Hu, Materials Letters, 2004,
Vol.58(15), pp.2096-2101.
218. Q. G. Wang and C. H. Cáceres, Materials Science and Engineering A, 1997, Vol.234-236, pp.106-109.
219. M. J. Caton, J. W. Jones, and J. E. Allison, Materials Science and Engineering A, 2001, Vol.314(1-2),
pp.81-85.
220. K. Moizumi, K. Mine, H. Tezuka, and T. Sato, Materials Science Forum, 2002, Vol.396 - 402,
pp.1371-1376.
221. S.-W. Han, K. Katsumata, S. Kumai, and A. Sato, Materials Science and Engineering: A, 2002,
Vol.337(1-2), pp.170-178.
222. M. S. Misra and K. J. Oswalt, AFS Transactions, 1982, Vol.90 pp.1-10.
223. D. A. Porter and K. E. Easterling: 'Phase Transformations in Metals and Alloys'; 1992, New York,
Chapman & Hall.
224. S. C. Wang and M. J. Starink, Int Mater Rev., 2005, Vol.50(4), pp.193-215.
225. R. Jahn, W. T. Donlon, and J. E. Allison, TMS (The Minerals, Metals & Materials Society), 1999,
pp.247-264.
226. M. E. FINE, Metallurgical Transactions A, 1975, Vol.6A, pp.625- 630.
227. G. Wang, Q. Sun, L. Feng, L. Hui, and C. Jing, Materials & Design, 2007, Vol.28(3), pp.1001-1005.
228. Y. J. Li, S. Brusethaug, and A. Olsen, Scripta Materialia, 2006, Vol.54(1), pp.99-103.
173
229. G. W. Smith, W. J. Baxter, and R. K. Mishra, Journal of Materials Science, 2000, Vol.35(15),
pp.3871-3880.
230. G. Wang, X. Bian, X. Liu, and J. Zhang, Journal of Materials Science, 2004, Vol.39(7), pp.2535-2537.
231. W. Reif, J. Dutkiewicz, R. Ciach, S. Yu, and J. Król, Materials Science and Engineering A, 1997,
Vol.234-236, pp.165-168.
232. D. J. Chakrabarti and D. E. Laughlin, Progress in Materials Science, 2004, Vol.49(3-4), pp.389-410.
233. G. Laslaz: US Patent, 0133949 Al, Jun. 22, 2006.
234. R. K. Mishra, A. K. Sachdev, and W. J. Baxter, AFS Transactions, 2004, Vol.112, pp.179-191.
235. W. Reif, S. Yu, J. Dutkiewicz, R. Ciach, and J. Król, Materials and Design, 1997, Vol.18(4-6),
pp.253-256.
236. F. Samuel, A. Samuel, and H. Liu, Journal of Materials Science, 1995, Vol.30(10), pp.2531-2540.
237. M. F. Ibrahim, E. Samuel, A. M. Samuel, A. M. A. Al-Ahmari, and F. H. Samuel, Materials & Design,
2011, Vol.32(4), pp.2130-2142
238. M. E. Fine: Dispersion Strengthened Aluminum Alloys, TMS, Warrendale, 1988, 103-121.
239. M. E. Fine: 'Alloy Design of Nanoscale Precipitation Strengthened Alloys: Design of a Heat Treatable
Aluminum Alloy Useful to 400° C', Northwestern University, Department of Materials Science and
Engineering, McCormick School of Engineering and Applied Science, 2006.
240. K. E. Knipling, D. C. Dunand, and D. N. Seidman, Zeitschrift für Metallkunde, 2006, Vol. 97(3),
pp.246-265.
241. J. R. Davis: 'Aluminum and aluminum alloys '; 1993, Materials Park OH, ASM Specialty Handbook.
242. M. Zedalis and M. Fine, Metallurgical and Materials Transactions A, 1986, Vol.17(12), pp.21872198.
243. J. Royset and N. Ryum, Int Mater Rev, 2005, Vol.50(1), pp.19-44.
244. J. F. Nie, A. Majumdar, and B. C. Muddle, Materials Science and Engineering: A, 1994, Vol.179180(Part 1), pp.619-624.
245. E. Nes, Acta Metallurgica, 1972, Vol.20(4), pp.499-506.
246. S. K. Pandey and C. Suryanarayana, Materials Science and Engineering: A, 1989, Vol.111, pp.181187.
247. N. Furushiro and S. Hori, Acta Metallurgica, 1985, Vol.33(5), pp.867-872.
248. L. M. Angers, Y. C. Chen, M. E. Fine, J. R. Weertman, and M. S. Zedalis: Aluminum Alloys: Their
Physical and Mechanical Properties, EMAS,Warley, 1986, 321-337.
249. D. Srinivasan and K. Chattopadhyay, Metallurgical and Materials Transactions A, 2005, Vol.36(2),
pp.311-320.
250. L. Lityńska-Dobrzyńska, Solid State Phenomena, 2007, Vol.130, pp.163-166.
251. K. E. Knipling: 'Development of a nanoscale precipitation-strengthened creep-resistant aluminum alloy
containing trialuminide precipitates', PhD thesis, Northwestern University, Illinois, 2006.
252. J. W. Martin: 'Concise Encyclopedia of the Structure of Materials', 235-247; 2007, Oxford, Elsevier.
253. E. Samuel, A. M. Samuel, F. H. Samuel, and H. W. Doty, AFS Transactions, 2010, Vol.118, pp.77-82.
254. F.A.Fasoyinu, D.Cousineau, P.Newcombe, T.Castles, and M.Sahoo, AFS Transactions, 2001, Vol.109,
pp.397-418.
255. P. Sepehrband, R. Mahmudi, and F. Khomamizadeh, Scripta Materialia, 2005, Vol.52(4), pp.253-257.
174
256. R. Mahmudi, P. Sepehrband, and H. M. Ghasemi, Materials Letters, 2006, Vol.60(21-22), pp.26062610.
257. E. Ozbakir: 'Development of Aluminum Alloys for Diesel-Engine Applications', M.Sc. thesis, McGill
University,Montréal, Québec, 2008.
258. P. Prasad: 'Characterization of new, cast, high temperature aluminum alloys for diesel engine
applications', M.Sc. thesis, University of Cincinnati, United States,Ohio, 2006.
259. T. Iveland, S. Brusethaug, P. As Holt, B. Barlas, D. Massinon, and P. Meyer: International Patent,
159169 Al, Dec. 22, 2011.
260. B. Baradarani and R. Raiszadeh, Materials & Design, 2011, Vol.32(2), pp.935-940.
261. G. Laslaz and M. Garat: US Patent, 0224145 Al, Oct. 13, 2005.
262. P. Jonason, AFS Transactions, 1992, Vol.100 pp.601-607.
263. J. A. Lee and P. S. Chen: US Patent, 0192627 Al, Oct. 16, 2003.
264. M. Garat: US Patent, 0126947 Al, Jun. 2, 2011.
265. M. Garat, International Journal of Metalcasting, 2011, Vol.5(3), pp.17-24.
266. J. Pavlovic-Krsti: 'Impact of casting parameters and chemical composition on the solidification
behaviour of Al-Si-Cu hypoeutectic alloy', PhD thesis, 2010.
267. S. Li: 'Hot Tearing in Cast Aluminum Alloys: Measures and Effects of Process Variables', PhD thesis,
WORCESTER POLYTECHNIC INSTITUTE, 2010.
268. A. Nabawy, A. Samuel, F. Samuel, and H. Doty, Journal of Materials Science, 2012, Vol.47(9),
pp.4146-4158.
269. J.-M. Drezet, JOURNAL DE PHYSIQUE IV, 1999, Vol.9, pp.53-62.
270. T. W. Clyne and G. J. Davies, The British Foundryman, 1981, Vol.74, pp.65-73.
271. L. Katgerman, JOM 1982, Vol.34 (2), pp.46-49.
272. U. Feurer, Quality Control of Engineering Alloys and the Role of Metals Science, 1977 Vol.Delft
University of Technology, pp.131-145.
273. D. Benny Karunakar, R. Naresh Rai, S. Patra, and G. Datta, The International Journal of Advanced
Manufacturing Technology, 2009, Vol.45(9), pp.851-858.
274. J. Deshpande: 'The Effect of Mechanical Mold Vibration on the Characteristics of Aluminum Alloys',
M.Sc. thesis, Worcester Polytechnic Institute 2006.
275. J. Deshpande and M. M. Makhlouf, AFS Transactions, 2008 Vol.116, pp.247-263.
276. D. Warrington and D. G. McCartney, Cast Metals, 1989, Vol.2, pp.134.
277. D. G. Eskin, Suyitno, and L. Katgerman, Progress in Materials Science, 2004, Vol.49(5), pp.629-711.
278. H. Kamguo Kamga, D. Larouche, M. Bournane, and A. Rahem, Materials Science and Engineering:
A, 2010, Vol.527(27-28), pp.7413-7423.
279. S. Lin, C. Aliravci, and M. O. Pekguleryuz, Metallurgical and Materials Transactions A, 2007,
Vol.38(5), pp.1056-1068.
280. G. CAO, C. ZHANG, H. CAO, Y. A. CHANG, and S. KOU, METALLURGICAL AND MATERIALS
TRANSACTIONS A, 2010, Vol.41A.
281. M. O. Pekguleryuz, S. Lin, E. Ozbakir, D. Temur, and C. Aliravci, International Journal of Cast
Metals Research, 2010, Vol.23(5).
282. F. Paray, B. Kulunkt, and J.E.Gruzleski, Int J Cast Metals Res, 2000, Vol.13, pp.147-159.
283. S. Bozorgi, K. Haberl, C. Kneissl, T. Pabel, and P. Schumacher: TMS, 2011.
175
284. R. Mackay, R. Hausenbusch, and J. Sokolowski, Materials Science Forum, 2007, Vol.539 - 543,
pp.392-397.
285. J. Taylor, G. Schaffer, and D. StJohn, Metallurgical and Materials Transactions A, 1999, Vol.30(6),
pp.1643-1650.
286. J. Taylor, G. Schaffer, and D. StJohn, Metallurgical and Materials Transactions A, 1999, Vol.30(6),
pp.1651-1655.
287. C. M. Dinnis, J. A. Taylor, and A. K. Dahle, Materials Science and Engineering: A, 2006, Vol.425(1–
2), pp.286-296.
288. W.Schneider, M. Otte, C. Dinnis, and J. Taylor, Casting Plant and Technology International, 2007.
289. N. Veldman, A. Dahle, D. StJohn, and L. Arnberg, Metallurgical and Materials Transactions A, 2001,
Vol.32(1), pp.147-155.
290. P. Ouellet and F. H. Samuel, Journal of Materials Science, 1999, Vol.34(19), pp.4671-4697.
291. TTAL7, TT Al-based Alloys Database, version 7.0, T. Ltd., Surrey Technology Center, Guildford, UK,
2010.
292. Standard Practice for Heat Treatment of Aluminum-Alloy Castings from All Processes, 2001,ASTM B
917/B 917M – 01, US.
293. P.-S. Wang, S.-L. Leea, J.-C. Lina, and M.-T. Jahn, Journal of Materials Research, 2000, Vol.15(09),
pp.2027-2035.
294. M. Javidani and D. Larouche, International Materials Reviews, 2014, Vol.59(3), pp.132-158.
295. M. B. Djurdjevic, W. Kasprzak, C. A. Kierkus, W. T. Kierkus, and J. H. Sokolowski: 105th Casting
Congress, Dallas, USA, 2001.
296. J. A. Taylor, D. H. St John, J. Barresi, and M. J. Couper, Mater. Sci. Forum, 2000, Vol.331-337,
pp.277-282.
297. M. Javidani, D. Larouche, and X. G. Chen, AFS Trans, 2014, Vol.122(14-056).
298. B. Sundman, B. Jansson, and J.-O. Andersson, Calphad, 1985, Vol.9, pp.153-190.
299. 'Oxford Instruments HKL, The HKL Channel 5 software', 2007 [viewed 2014]; Available from:
http://caf.ua.edu/wp-content/uploads/docs/JEOL-7000F-Oxford_Channel_5_User_Manual.pdf.
300. L. Lasa and J. M. Rodriguez-Ibabe, Journal of Materials Science, 2004, Vol.39(4), pp.1343-1355.
301. K. LIU, X. CAO, and X.-G. CHEN, Metallurgical and Materials Transactions A, 2012, Vol.43A,
pp.1097-1101.
302. K. Liu, X. Cao, and X. G. Chen, Metallurgical and Materials Transactions A, 2013, Vol.44(8),
pp.3494-3503.
303. G. Mrówka-Nowotnik and J. Sieniawski, Archives of Materials Science and Engineering, 2011,
Vol.47(2), pp.85-94.
304. M. Javidani, D. Larouche, and X. G. Chen, Metallurgical and Materials Transactions A, 2015,
Vol.(recently accepted).
305. D. L. Zhang, L. H. Zheng, and D. H. StJohn, Journal of Light Metals, 2002, Vol.2(1), pp.27-36.
306. I. Alfonso, C. Maldonado, G. Gonzalez, and A. Bedolla, Journal of Materials Science, 2006,
Vol.41(7), pp.1945-1952.
307. M. F. Ibrahim, E. Samuel, A. M. Samuel, A. M. A. Al-Ahmari, and F. H. Samuel, Materials &amp;
Design, 2011, Vol.32(7), pp.3900-3910.
308. Y. Han, A. M. Samuel, H. W. Doty, S. Valtierra, and F. H. Samuel, Materials & Design, 2014,
Vol.58(0), pp.426-438.
176
309. S. J. Mackay R, International Journal of Metalcasting, 2010, Vol.4(4), pp.33-50.
310. E. Tillová, M. Chalupová, and L. Hurtalová: 'Evolution of Phases in a Recycled Al-Si Cast Alloy
During Solution Treatment', in 'Scanning Electron Microscopy', (ed. V. Kazmiruk), 2012, InTech.
311. L. Hurtalova, E. Tillova, and M. Chalupova, Engineering Transactions, 2013, Vol.61(3), pp.197–218.
312. L. Hurtalova, E. Tillova, and M. Chalupova, The Archive of Mechanical Engineering, 2012,
Vol.LIX(4).
313. W. J. Coy and R. S. Mateer, Asm Transactions Quarterly, 1965, Vol.58(1), pp.99-&.
314. W. Gasior, Z. Moser, and J. Pstrus, Journal of Phase Equilibria, 1998, Vol.19(3), pp.234-238.
315. W. M. Haynes: 'CRC handbook of chemistry and physics : a ready-reference book of chemical and
physical data'; 2012, Boca Raton, Fla, CRC Press.
316. D. A. Porter and K. E. Easterling: 'Phase Transformations in Metals and Alloys'; 1992, CRC.
317. N. A. Belov, A. A. Aksenov, and D. G. Eskin: 'Iron in aluminum alloys : impurity and alloying
element'; 2002, London, Taylor & Francis.
177