metals Article
Effect of T6 Heat Treatment on the Microstructure and Hardness of Secondary AlSi9Cu3(Fe) Alloys Produced by Semi-Solid SEED Process Alberto Fabrizi , Stefano Capuzzi, Alessandro De Mori and Giulio Timelli * Department of Management and Engineering, University of Padova, Stradella S. Nicola, 3 I-36100 Vicenza, Italy;
[email protected] (A.F.);
[email protected] (S.C.);
[email protected] (A.D.M.) * Correspondence:
[email protected]; Tel.: +39-0444-998-769 Received: 3 September 2018; Accepted: 19 September 2018; Published: 23 September 2018
Abstract: The effect of the T6 heat treatment on the microstructure and hardness of a secondary semi-solid AlSi9 Cu3 (Fe) alloy have been investigated by using optical, scanning and transmission electron microscopy and hardness testing. The semi-solid alloy was produced using the swirled enthalpy equilibration device (SEED). The solution heat treatments were performed at 450, 470 and 490 ◦ C for 1 to 6 h followed by water quenching and artificial ageing at 160, 180 and 220 ◦ C for holding times ranging from 1 to 30 h. The microstructural investigations have revealed the spheroidization of the eutectic Si and the dissolution of the majority of Cu-rich compounds after all the solution heat treatments; moreover, the greater the solution temperature and time, the higher the hardness of the alloy. Unacceptable surface blistering has been observed for severe solution condition, 490 ◦ C for 3 and 6 h. The artificial ageing at 160 ◦ C for 24 h has led to the highest alloy strengthening thanks to the precipitation of β” and Q’ (or L) phases within the α-Al matrix. The hardening peaks at higher temperatures have been early achieved due to faster hardening kinetic; however, the lower number density of β” and Q’ (or L) phases and the presence of coarser θ’ precipitates result in a reduction of hardness values for peak aged condition at 180 and 220 ◦ C, respectively. Keywords: aluminium alloys; semi-solid; SEED; heat treatment; microstructure; mechanical properties
1. Introduction Aluminium-silicon casting alloys are widely used in the automotive industry due to their good castability, good corrosion resistance, high strength stiffness to weight ratio as well as recycling possibilities at low energy costs. Furthermore, mechanical properties of as-cast components can significantly be improved by suitable heat treatments. However, the application of a post-casting thermal treatment on high-pressure die-cast (HPDC) components is still a limited practice in industry. Due to the extreme turbulent flow of molten metal during filling of the die cavity, internal or sub-surface porosity, originated from entrapped gases, can usually occur in the diecastings [1]; consequently, during typical thermal treatment at elevated temperature (470÷500 ◦ C [2]), gas expansion can produce unacceptable surface blisters on the components [2,3]. Beside the use of vacuum-assisted HPDC to reduce the level of entrapped gas [4], semi-solid metal (SSM) processes have been considered a valid alternative for producing sound Al-Si-based light-weight parts with satisfactory response to heat treatments, especially for the automotive sector. Some examples are fuel rails, control arms, steering knuckles, engine brackets, or even engine blocks [5–7]. Generally, SSM processing is divided into two main categories such as rheocasting and thixocasting. The former involves the preparation of a SSM slurry directly from the molten metal,
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which is cooled and stirred up to the desired solid fraction, followed by forming process. Nowadays, this approach is more widely used than thixocasting where a specially prepared alloy (with globular or rosette-like α-Al grains) is reheated to the desired semi-solid forming temperature prior to casting. Contrary to the turbulent metal flow during die cavity filling in conventional HPDC, both the SSM casting approaches experience a laminar flow of semi-solid metal having a higher viscosity than the liquid state; consequently, the risk of gas entrapment is largely limited and blistering phenomena are prevented during further heat treatments or welding operations [8]. The typical thermal treatment for automotive components made by Al foundry alloy is T6 temper, which generally induces higher alloy strengthening. The T6 thermal cycle consists of a solution heat treatment followed by a water quenching and then an age hardening (or precipitation hardening). The solution heat treatment leads to the dissolution of intermetallic phases and the spheroidization of eutectic Si with a resulting improvement of alloy ductility [9,10]. The time for solution treatment is strongly dependent on the microstructural scale [10], ranging from few minutes up to several hours [11–13]. In general, too short solution treatment does not guarantee that all alloying elements are dissolved in the α-Al matrix and made available for further precipitation hardening; on the contrary, too long solution treatment shows economic limitations because it uses more energy and time than necessary. Age hardening at room temperature (natural ageing) or elevated temperature (artificial ageing, AA) increases the alloy strengthening because of the ultra-fine particles which precipitate from the supersaturated solid solution and act as obstacles to the dislocation movement. In the last years, an increasing attention from researchers has been paid on the heat treatment sequence of primary semi-solid Al-Si cast alloys such as A356, A357 alloys [14–16] in order to increase their mechanical properties. However, to date, only few studies have investigated the age hardening response of a secondary (or recycled) Al-Si alloy produced via SSM route [17] although secondary Al foundry alloys are the most used Al diecasting alloys in the automotive industry. In this work, the microstructure evolution and the precipitation hardening response of a rheocast secondary AlSi9 Cu3 (Fe) alloy used for automotive components were investigated as a function of different T6 heat treating parameters. 2. Materials and Methods In this study, automotive components (lower suspension arm) were manufactured with a secondary AlSi9 Cu3 (Fe) alloy (EN AC-46000) by the Swirled Enthalpy Equilibration Device (SEED), which operates by using a liquid-based method (rheocasting) for producing the Al slurry over a wide range of Al-Si alloy compositions (up to 17 wt.% Si) [18,19]. According to this process, the slurry is prepared by means of three stages [18]. Firstly, the molten metal of the desired composition and temperature is poured into a vessel whose thermal capacity is able to cool the alloy in order to obtain the desired solid fraction in the range between 30 and 45%. During this stage, the vessel is swirled to homogenize the temperature and the distribution of the solid phase which is generated at the walls of the vessel. The duration of this stage depends upon the dimensions of the container and the mass of the charge, even if an extremely close control of the mixing time is not necessary. Furthermore, Doutre et al. [18] demonstrated how the mixing intensity does not affect the process yield over the range of 150 to 300 rpm. In the second stage of the process, a valve at the bottom of the vessel can be opened and some of the remaining liquid phase is drained to produce a self-supporting semi-solid billet, which is depleted in the eutectic forming elements (Si, Mg, Fe) and enriched in the peritectic forming element (Ti) [18]. In the third stage, the vessel is inverted, and the billet is de-moulded and transferred to the forming unit. Information about the SEED method and processing parameters are reported in the Ref. [18,20] where the potential of the SEED process as a route of preparing high quality feedstock has been also tested to produce an automotive front upper control arm weighing approximately 1 kg.
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In the present work, the ingots were supplied by Raffineria Metalli Capra Spa (Brescia, Italy) while the demonstrators were produced at the CIE Automotive (Bilbao, Spain). The chemical composition of the alloy, measured on separately poured samples, is shown in Table 1. Table 1. Chemical composition of the experimental AlSi9 Cu3 (Fe) alloy (wt.%). Si
Fe
Cu
Mn
Mg
Ni
Zn
Cr
Sn
Ti
Al
8.820
0.691
2.122
0.23
0.32
0.065
0.90
0.05
0.01
0.04
bal.
The ingots were firstly melted at 670 ◦ C in a tower furnace of 800-kg capacity and then the molten metal was tapped into an electric resistance furnace where it was N-degassed with a rotary degasser (MTS Rotor Ø190, at 600 rpm, Foseco Española S.A., Izurtza, Spain) for 8 min. The metal was manually skimmed and held at 640 ◦ C. No grain refinement or chemical eutectic modification were performed. The melt was transferred by means of a ladle and poured into a steel vessel preheated at about 130 ◦ C to avoid thermal shock and to better control the amount of solid fraction. At this stage, in order to improve the distribution of the primary solid phase, the vessel was swirled for 115 s at a rate of 150 rpm with an eccentricity of the swirl motion. To achieve homogenisation of the temperature and structure, a delay time of 10 s was maintained before transferring the slurry to the shoot sleeve of a cold camber die casting machine. Then, the slurry was injected into the die cavity and the plunger velocity of the HPDC machine (Petransa, Alió, Tarragona, Spain) was lower than 1 ms−1 for the filling phase; once the die-cavity was filled, a pressure of ~115 MPa was applied to guarantee high-integrity castings. Detailed investigations and results regarding the final rheocasting quality are described in Ref. [21]. A calorimetry analysis of the as-rheocast alloy was preliminary performed with a Labsys Evo differential scanning calorimeter (DSC) (SETARAM Instrumentation, Caluire, France) to identify the solidification reaction sequence of the alloy. The samples in the form of 4 mm discs weighting approximately 30 mg were drawn from the castings. Samples were weighted before and after DSC scan and no significant mass increase was observed. The DSC test was carried out with an empty reference pan, as cooling (solidification) scan in a protective gas atmosphere of pure Argon. The sample was heated up to 800 ◦ C at 2 ◦ C/min and maintained at this temperature for 15 min before cooling to room temperature at 2 ◦ C/min. Specimens of about 30 × 10 × 15 mm3 size were drawn from the as-rheocast components and they were heat treated according to the following conditions: solution treatment inside an electric-resistance heated-air-circulating furnace, water quenching at room temperature and then artificial ageing. The solution treatments were performed at 450, 470 and 490 ◦ C for 1, 3 and 6 h, respectively. These process parameters were selected in order to minimize blistering and residual stresses in the castings after water quenching and to reduce the risk of localised incipient melting of Cu-rich compounds. Moreover, the experimental solution treatments were also chosen in accordance with the conventional solution cycles normally applied in the automotive industry, where shortening the total time is a key issue to increase the productivity and to reduce the manufacturing cost [11–13]. As mentioned before, after the solution heat treatment, the specimens were water quenched at room temperature. To avoid natural ageing prior to the artificial ageing, the specimens were stored at 4 ◦ C for 1 day; subsequently, the artificial ageing was carried out respectively at 160, 180 and 220 ◦ C for different ageing times ranging between 1 and 30 h. The temperature control during solution and ageing treatments was ±1 ◦ C. The heat-treated specimens were mechanically prepared by grinding with abrasive SiC papers and polishing with 6- and 3-µm diamond pastes and colloidal silica suspension. Optical microscope and quantitative metallographic analysis were used to determine the main characteristics of eutectic Si particles (size and aspect ratio) and the amount of undissolved Cu-rich phases. Before the quantification, the polished samples were chemically etched by using a water solution of HNO3 at 70 ◦ C to improve the contrast of Cu-rich phases. A FEI Quanta™ 250 scanning electron microscope
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quantification, the polished samples were chemically etched by using a water solution of HNO3 at 70 °C to improve the contrast of Cu-rich phases. A FEI Quanta™ 250 scanning electron microscope (SEM) (Thermo Fisher Scientific, Hillsboro, OR, USA) operating at 20 kV and coupled with an an energy energy dispersive dispersive spectrometer spectrometer (EDS) (EDS) (EDAX™, (EDAX™, AMETEK AMETEK BV, BV, Tilburg, The Netherlands) Netherlands) and and an electron back-scattered back-scattered diffraction diffraction (EBSD) (EBSD) (EDAX™, (EDAX™, AMETEK AMETEK BV, BV, Tilburg, The Netherlands) detector was used to study the nature and the distribution of the intermetallic phases. In order order to tobetter betterunderstand understandthe theeffects effects ageing treatment on alloy the alloy microstructure, of of thethe ageing treatment on the microstructure, peak ◦ peak conditions at 160, 180 Cand 220 °C were selected formicrostructural an extended microstructural ageingageing conditions at 160, 180 and 220 were selected for an extended characterization characterization using 2000 EXelectron II transmission electron (JEOL Ltd, by using a JEM™by2000 EXaIIJEM™ transmission microscope (TEM)microscope (JEOL Ltd,(TEM) Akishima, Japan) Akishima, Japan) operating at 200 kV. In particular, TEM observations supported by selected area operating at 200 kV. In particular, TEM observations supported by selected area electron diffraction electron diffraction (SAED) analyses were the carried to study the precipitation the α-Al (SAED) analyses were carried out to study phaseout precipitation in phase the α-Al phase. TEMinspecimens phase. TEM specimens were mechanically prepared practice. accordingFinally, to standard practice. the TEM were mechanically prepared according to standard the TEM discs Finally, (3 mm diameter, discsµm (3thickness) mm diameter, ~25 μm by thickness) were by ion milling up to precision the electron ~25 were processed ion milling up processed to the electron transparency using ion transparency using precision polishing system (PIPS, Model polishing system (PIPS, Modelion 691, Gatan, Pleasanton, CA, USA).691, Gatan, Pleasanton, CA, USA). In order order to evaluate evaluate the the strengthening strengthening response of the alloy as function function of the the solution solution and and ageing ageing treatments, treatments, Vickers Vickers hardness hardness measurements measurements were were performed performed on on each each sample sample by by applying applying aa load load of 60 kgf, in accordance accordance to to the the standard standard ASTM ASTM E92-17, E92-17, and and the the average averagehardness hardnessvalues valueswere werereported. reported. The hardness evolution of the alloy alloy allowed allowed to toeasily easilyindividuate individuatethe thepeak peakageing ageing(temperature, (temperature,time), time), that is the greatest material hardening obtained during ageing treatment.
Results and and Discussion Discussion 3. Results 3.1. As-Rheocast As-Rheocast Condition Condition 3.1. The microstructure as-rheocast AlSiAlSi9Cu3(Fe) exhibits the typical structure of rheocast 9 Cu3 (Fe) alloyalloy The microstructureofofthethe as-rheocast exhibits the typical structure of Al-Si alloy, characterized by globular or rosette-like primary α-Al crystals surrounded by the Al-Si rheocast Al-Si alloy, characterized by globular or rosette-like primary α-Al crystals surrounded by eutectic (Figure 1). The shape the α-Al is due is todue the to fragmentation of theofα-Al the Al-Sinetwork eutectic network (Figure 1). Theofshape of thecrystals α-Al crystals the fragmentation the dendrites due todue theto swirl motion during the SEED process. As reported in Ref.in [21], average size of α-Al dendrites the swirl motion during the SEED process. As reported Ref.the [21], the average the primary α-Al crystals, estimated by quantitative image analysis, results results in ~70 µm with aspect size of the primary α-Al crystals, estimated by quantitative image analysis, in ~70 μmanwith an ratio of 1.6. Moreover, the volume fraction of the primary α-Al phase is 55 ± 6%. aspect ratio of 1.6. Moreover, the volume fraction of the primary α-Al phase is 55 ± 6%. The eutectic eutectic Si Si particles particles generally generally show show fibrous fibrous morphology, morphology, as It is is worth worth The as illustrated illustrated in in Figure Figure 2. 2. It noticing that a larger amount of Al-Si eutectic is observed near to the casting surface (Figure 3): this noticing that a larger amount of Al-Si eutectic is observed near to the casting surface (Figure 3): this microstructural inhomogeneity is due to liquid segregation phenomena occurring during the SSM microstructural inhomogeneity is due to liquid segregation phenomena occurring during the SSM processing, as as described described in in Ref. processing, Ref. [22,23]. [22,23].
Figure Figure1. 1. Microstructure Microstructureof ofas-rheocast as-rheocast AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy.
Figure 4,4,the theDSC DSCcooling coolingcurve curve AlSi9 Cu3 (Fe) alloy is plotted. thermogram reveals In Figure ofof AlSi9Cu3(Fe) alloy is plotted. TheThe thermogram reveals the ◦ C, respectively, and according the liquidus and solidus temperatures approximately at 610 and 485 liquidus and solidus temperatures approximately at 610 and 485 °C, respectively, and according to to Ref. [24], sequence phaseprecipitation precipitationduring duringthe thesolidification solidificationpath pathcan canbe be identified identified as as Ref. [24], thethe sequence of of phase ◦ C), α-Al crystals (~578 ◦ C), eutectic Si and Fe-containing following: primary Fe-rich phase (~612 following: primary Fe-rich phase (~612 °C), α-Al crystals (~578 °C), eutectic Si phases (~567 ◦ C), Al2 Cu phase (~503 ◦ C) and formation of eutectic structures containing Al2 Cu and Al5 Mg8 Si2 Cu2 (~498 ◦ C).
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phases (~567 °C), Al22Cu phase (~503 °C) and formation of eutectic structures containing Al 22Cu5 and Metals 2018, 8, 750 of 18 Al55Mg88Si2i2Cu22 (~498 °C).
Figure Figure 2. 2. Eutectic Eutectic Si Si particles particles in in as-rheocast as-rheocast condition. condition.
Figure the casting. casting. Figure 3. 3. Distribution Distribution of of Al-Si Al-Si eutectic eutectic from from the the outer outer surface surface to to the the centre centre of of the
Figure Figure 4. 4. Differential Differential scanning scanning calorimetry calorimetry curve curve of of experimental experimental AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy alloy recorded recorded during during ◦ cooling at 2 °C/min; peak 1 = primary Fe-rich phase; peak 2 = α-Al crystals; peaks 3 and cooling at 2 C/min; peak 1 = primary Fe-rich phase; peak 2 = α-Al crystals; peaks 3 and 44 == eutectic eutectic Si Si and Fe-containing phases; peak 55 == Al Al222Cu peak 66 = = eutectic eutectic structures structures containing containing Al Al222Cu and Fe-containing phases; peak Cu phase; phase; peak Cu and and Al Al555Mg Mg888Si Si222Cu Cu22.2 .
Deeper SEM investigations investigations revealed revealed the intermetallic intermetallic phases phases distributed distributed solely solely in in the eutectic network (Figure 5); in addition, the EDS elemental maps seem to confirm the presence θAl-Si network (Figure 5); in addition, the EDS elemental maps seem to confirm the presence of θ-Alof2 Cu, Al-Si Al 22Cu, polygonal α-Al(Fe,Mn,Cr)Si (α-Fe),β-Al(Fe,Mn)Si acicular β-Al(Fe,Mn)Si (β-Fe) and Q-AlMgSiCu polygonal α-Al(Fe,Mn,Cr)Si (α-Fe), acicular (β-Fe) and Q-AlMgSiCu intermetallic intermetallic phases phases (Figure 6). (Figure 6).
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Figure 5.Scanning Scanningelectron electron microscopy microscopy micrograph micrograph of as-rheocast alloy. Figure 5. of as-rheocast alloy. Figure 5.5. Scanning micrographof ofas-rheocast as-rheocastalloy. Figure Scanningelectron electronmicroscopy microscopy micrograph
Figure 6. Composition mapsby by energydispersive dispersive spectroscopy related totothe micrograph in Figure 5 Figure 6. 6. Composition spectroscopyrelated relatedto themicrograph micrograph Figure Figure Compositionmaps maps byenergy energy dispersive spectroscopy the in in Figure 5 5 and showing the distributions of Mg, Si, Cu, Fe, Mn and Cr elements. Figure 6. Composition maps byofof energy dispersive spectroscopy related to the micrograph in Figure 5 and showing the distributions Mg, Cu, and Cr elements. elements. and showing the distributions Mg,Si, Si, Cu, Fe, Fe, Mn Mn and Cr and showing the distributions of Mg, Si, Cu, Fe, Mn and Cr elements.
ForFor a more ofthe theintermetallic intermetallic phases, EBSD analyses performed a moreaccurate accurateidentification identification of phases, EBSD analyses werewere performed on For a more accurate identification of the intermetallic phases, EBSD analyses were performed on the θ-Al 2Cu, α-Al(Fe,Mn,Cr)Si and β-Al(Fe,Mn)Si, and the and associated EBSD patterns well match the onthe the θ-Al α-Al(Fe,Mn,Cr)Si and β-Al(Fe,Mn)Si, the associated EBSD patterns well 2 Cu, For a more accurate identification of the intermetallic EBSD analyses were performed on θ-Al 2Cu, α-Al(Fe,Mn,Cr)Si and β-Al(Fe,Mn)Si, and thephases, associated EBSD patterns well match the nominal crystallography structures, i.e., tetragonal, cubic and monoclinic crystal systems, match the nominal crystallography structures, i.e., tetragonal, cubic and monoclinic crystal systems, thenominal θ-Al2Cu,crystallography α-Al(Fe,Mn,Cr)Si and β-Al(Fe,Mn)Si, and the associated EBSD patterns well systems, match the structures, i.e., tetragonal, cubic and monoclinic crystal respectively [25], as reported inFigure Figure 7. respectively [25], asas reported inin nominal crystallography structures, i.e., tetragonal, cubic and monoclinic crystal systems, respectively [25], reported Figure7.7. respectively [25], as reported in Figure 7.
Figure 7. Experimental patterns by electron back-scattered diffraction acquired from (a) Al2Cu, (b) αFigure Experimental patternsbybyelectron electronback-scattered back-scattereddiffraction diffractionacquired acquiredfrom from(a) (a)Al Al2Cu, Cu, (b) (b) αFigure 7. 7. Experimental patterns α-Fe 2 Fe and (c) β-Fe phases with superimposed tetragonal (space group, s.p. I4/mcm), cubic (s.p. Pm-3) Fe and (c) β-Fe phases with superimposed tetragonal (space group, s.p. I4/mcm), cubic (s.p. Pm-3) and (c) β-Fe phases with superimposed tetragonal (space group, s.p. I4/mcm), cubic (s.p. Pm-3) and and monoclinic (s.p. 2/m) pattern simulations, respectively. Figure 7. Experimental patterns by simulations, electron back-scattered and monoclinic (s.p. 2/m) pattern respectively.diffraction acquired from (a) Al2Cu, (b) αmonoclinic (s.p. 2/m) pattern simulations, respectively. Fe and (c) β-Fe phases with superimposed tetragonal (space group, s.p. I4/mcm), cubic (s.p. Pm-3) and monoclinic (s.p. 2/m) pattern simulations, respectively.
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3.2. As-Quenched Condition 3.2. As-Quenched Condition The specimens, which were solution heat treated at 450, 470 and 490 °C, showed essentially ◦ Themicrostructural specimens, which wereThe solution heat treated at 450, 470 and 490 showed essentially similar features. microstructure of the AlSi9Cu3(Fe) alloyC,solutioned at 470 °C similar microstructural features. The microstructure of the AlSi Cu (Fe) alloy solutioned 470 ◦ C for 3 for 6 h is shown as an example in Figure 8. As expected, the 9solution treatment has noatcoarsening 6 h is shown an example Figure 8. Asshow expected, the solution treatment has no coarsening effects on theasprimary α-Alincells, which a similar size to the as-rheocast condition (~70effects μm); on the primary α-Al cells, which show a similar size to the as-rheocast condition (~70 µm); while while the eutectic Si is considerably modified in size and morphology. The eutectic Si particles the eutectic is considerably modified in(Figure size and The eutectic particles fragment fragment andSiacquire a more round-shape 8b)morphology. due to the reduction of theSi total Al/Si interfacial and acquire a more (Figure 8b) due to thesmaller reduction of the total Al/Si interfacial energy energy during theround-shape solution treatment; moreover, eutectic Si particles in as-rheocast during the solution treatment; moreover, smaller eutectic Si particles in as-rheocast microstructure microstructure dissolve and favour the coarsening of coarser ones according to Ostwald ripening dissolve and[26]. favour the coarsening of coarser ones according to Ostwald ripening mechanism [26]. mechanism
Figure 8. Micrographs alloy solutioned solutioned at Figure 8. Micrographs at at (a) (a) low low and and (b) (b) higher higher magnifications magnifications of of AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy at ◦ 470 °C for 6 h. 470 C for 6 h.
The size and morphological morphological variations variations of of the eutectic eutectic Si particles as a function of solution treatments treatments are summarized summarized in in Table 2. The The eutectic eutectic Si Si particles particles fragment fragment and and spheroidize, spheroidize, mainly, mainly, during the the early earlystage stageofofthe thesolution solution treatment; then, they to coarsen during the subsequent treatment; then, they tendtend to coarsen during the subsequent stage stage of treatment, maintaining essentially the same shape. of treatment, maintaining essentially the same shape. Table 2. Size and morphology of the eutectic Si particles, and area fraction of Cu-rich compounds as Table 2. a function of the solution heat treatments. Alloy condition
Alloy Condition As-rheocast As-rheocast As-quenched
As-quenched As-quenched
As-quenched
As-quenched
Solutionheat Heattreatment Treatment Solution (◦ C)
Temperature(°C) Temperature
Time(h) (h) Time
--
--
450
450 470
470
490
1 13 36 1 6 3 16 31 3 66
Eutectic EutecticSi SiParticles particles (µm22)
Area Area (μm ) 2.4 ± 1.5 2.4 ± 1.5 1.1 ± 0.7 1.9 1.1±± 0.8 0.7 1.9 1.9±± 0.6 0.8 1.8 ± 0.6 1.9 ± 0.6 1.9 ± 0.7 1.9 1.8±± 0.8 0.6 1.6 0.7 1.9± ± 0.7 2.2 ± 0.9 1.9±± 0.8 0.8 2.2
Cu-Rich Compounds Cu-rich compounds
Aspect Ratio Aspect ratio
Area Fraction (%)(%) Area fraction
2.4 ± 1.3 2.4 ± 1.3 1.8 ± 0.5 2.01.8 ± 0.5 ± 0.5 2.02.0 ± 0.6 ± 0.5 2.1 ± 0.5 2.0 ± 0.6 1.9 ± 0.4 1.72.1 ± 0.4 ± 0.5 1.81.9 ± 0.4 ± 0.4 1.7 ± 0.5 1.7 ± 0.4 1.7 ± 0.4
2.1 ± 0.5 2.1 ± 0.5 1.0 ± 0.1 0.31.0 ± 0.1 ± 0.1 0.30.3 ± 0.1 ± 0.1 0.5 ± 0.2 0.3 ± 0.1 0.3 ± 0.1 0.20.5 ± 0.2 ± 0.2 0.20.3 ± 0.1 ± 0.1 0.1 ± 0.1 0.2 ± 0.2 0.1 ± 0.1
1 1.6 ± 0.7 1.8 ± 0.4 0.2 ± 0.1 As-quenched 490 3 2.2 ± 0.9 1.7 ± 0.5 0.1 ± 0.1 Moreover, fine polygonal and rod/platelet-shape particles appear within the α-Al matrix after 6 2.2 ± 0.8 1.7 ± 0.4 0.1 ± 0.1
solution treatments, as shown in Figure 8b. The EDS profile across the α-Al crystals suggests that the particles are Si-rich phase (Figure 9, red line). The bright field TEM micrograph correlated to Moreover, fine polygonal rod/platelet-shape appear within the α-Al matrixshape after SAED investigations confirmed and that most of the particlesparticles are Si crystals with irregular-polygonal solution treatments, as shown in Figure 8b. The EDS profile across the α-Al crystals suggests that the (Figure 10). particles areas-rheocast Si-rich phase (Figure the 9, red line). The field TEM micrograph correlated In the condition, Si content inbright the α-Al matrix, over its solubility leveltoatSAED room investigations confirmed that most of the particles are Si crystals with irregular-polygonal shape temperature, is almost homogeneously distributed; this is indicated by the Si profile across the width (Figure 10).
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In In the the as-rheocast as-rheocast condition, condition, the the Si Si content content in in the the α-Al α-Al matrix, matrix, over over its its solubility solubility level level at at room room Metals 2018, 8, 750 8 of 18 temperature, temperature, is is almost almost homogeneously homogeneously distributed; distributed; this this is is indicated indicated by by the the Si Si profile profile across across the the width width of of an an α-Al α-Al crystal crystal which which appears appears almost almost constant constant (Figure (Figure 9, 9, black black line). line). When When the the alloy alloy is is solution solution of antreated, α-Al crystal which Si appears (Figure 9, black line). When the alloy isTherefore, solution heat heat the move, clustering and clusters begin coarsening. the heat treated, the excess excess Si atoms atomsalmost move,constant clustering and the the clusters begin coarsening. Therefore, the treated, the excess Si atoms move, clustering and the clusters begin coarsening. Therefore, the final final Si particles, crossed by the EDS scan line, cause a corresponding rise of the Si level showing final Si particles, crossed by the EDS scan line, cause a corresponding rise of the Si level showing aa Si particles, by the EDS scan line, cause a corresponding rise ofprofiles the Si level showing aisfinal final “spiky” distribution profile. Moreover, by the in 9, thus final “spiky”crossed distribution profile. Moreover, bycomparing comparing the elemental elemental profiles inFigure Figure 9, ititis thus “spiky” distribution profile. of Moreover, by comparing the elemental profiles in Figure 9, it is leads thus clear clear the in phase after treatment to clear that that the precipitation precipitation of Si Si particles particles in the the α-Al α-Al phase after the the solution solution treatment leads to aa that the precipitation of Si particles in the α-Al phase after the solution treatment leads to a subsequent subsequent subsequent reduction reduction of of the the silicon silicon content content in in the the surrounding surrounding Al-matrix Al-matrix with with respect respect to to the the asasreduction of the silicon content in the surrounding Al-matrix with respect to the as-rheocast condition. rheocast rheocast condition. condition.
Figure α-Al Figure 9. 9. The The distribution distribution of of Si Si across across the the width width of of an an α-Al α-Al crystal crystal in in as-rheocast as-rheocast condition condition and and after after distribution of 66 hh of of solution solution treatment treatment at at 470 470◦°C. °C. C.
Figure electron microscopy (TEM) micrograph showing Si dispersoids Figure 10. 10. Transmission Transmission electron electron microscopy microscopy micrograph micrograph (TEM) (TEM) micrograph micrograph showing showing Si Si dispersoids dispersoids Figure 10. Transmission micrograph ◦ the α-Al matrix in the rheocast AlSi9Cu3(Fe) alloy after solution treatment at °C In withinthe theα-Al α-Almatrix matrixin inthe therheocast rheocast AlSi9Cu3(Fe) alloy after solution treatment at 470 470 °C for h.the In within AlSi after solution treatment at 470 C for 6for h.66Inh. 9 Cu3 (Fe) alloy the the area diffraction aa Si the inset, inset, the selected selected area electron electron diffraction (SAED) pattern acquired acquired from Si dispersoid dispersoid is inset, the selected area electron diffraction (SAED)(SAED) pattern pattern acquired from a Sifrom dispersoid is reported.is reported. reported.
Similar Si-containing dispersions were also found to form in an Al-7Si-Mg alloy during solution Similar Si-containing dispersions also found form in alloy solution treatments Moreover, Lasagni etwere al. [28] presence of nanometrical globular and Similar[27]. Si-containing dispersions were alsoshowed found to tothe form in an an Al-7Si-Mg Al-7Si-Mg alloy during during solution treatments Moreover, et the of globular and platelet-like Si precipitates in an Al-1.7Si alloy after solution treatment with a controlled cooling treatments [27]. [27]. Moreover, Lasagni Lasagni et al. al. [28] [28] showed showed the presence presence of nanometrical nanometrical globular and platelet-like Si in Al-1.7Si alloy solution treatment with controlled rate rate (20 K/min) and how the size of Si particles increases reducing theaacooling ratecooling (3 K/min) platelet-like Si precipitates precipitates in an an Al-1.7Si alloy after after solutionby treatment with controlled cooling rate (20 K/min) and how the size particles increases reducing the cooling rate K/min) at at expense their The precipitation of by Si the α-Al matrix solution (20the K/min) andof how thedensity. size of of Si Si particles increases by particles reducingin the cooling rate (3 (3 after K/min) at the the expense their density. The of Si particles in α-Al after treatment be ascribed to the excess in the supersaturated achievedtreatment after the expense of ofcould theirthus density. The precipitation precipitation ofSi Siatoms particles in the the α-Al matrix matrixmatrix after solution solution treatment could be rheocasting could thus thus process. be ascribed ascribed to to the the excess excess Si Si atoms atoms in in the the supersaturated supersaturated matrix matrix achieved achieved after after the the rheocasting process. The dissolution rheocasting process. of Cu-rich compounds was also investigated, as it is generally considered a key-factor for the strengthening enhancement. The fraction of undissolved Cu-rich phases in the as-quenched specimens is reported in Table 2. In the as-rheocast condition, the fraction of Cu-rich
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The dissolution of Cu-rich compounds was also investigated, as it is generally considered a keyfactor for 8, the asMetals 2018, 750strengthening enhancement. The fraction of undissolved Cu-rich phases in the 9 of 18 quenched specimens is reported in Table 2. In the as-rheocast condition, the fraction of Cu-rich compounds is about 2% and it drastically decreases in the earlier stages of solution treatment, more compounds about 2% and it drastically decreases in the by earlier stages the of solution rapidly as theissolution temperature increases. In addition, extending solutiontreatment, time up tomore 6 h, rapidly as the solution temperature increases. In addition, by extending the solution time up to 6 h, the Cu-rich compounds are almost completely dissolved; very few blocky θ-Al2Cu particles remain, the Cu-rich compounds are almost completely few blockyAl θ-Al particles remain, 2 Cu being more complex to dissolve if compared to dissolved; θ-phase invery binary eutectic + Al 2Cu structures and being more complex to dissolve if compared to θ-phase in binary eutectic Al + Al Cu structures and Q-Al5Mg8Si2Cu2 phase, as also resulted from the precipitation sequence in Figure 4.2As a consequence Q-Al Mg Si Cu phase, as also resulted from the precipitation sequence in Figure 4. As a consequence 5 8 dissolution, 2 2 of θ-phase the Cu level in the α-Al solid solution increases as it clearly appears by of θ-phase dissolution, the Cu level the α-Al solid increases as it clearly appears comparing the EDS profiles across theinα-Al crystals in solution the as-rheocast condition and after 6 h by of comparing the EDS profiles across the α-Al crystals in the as-rheocast condition and after 6 hZn of solution treatment at 470 °C (Figure 11a). Similar behaviour was also found for the Mg and ◦ solution treatment atby 470 (Figure 11a). profiles Similarinbehaviour wasrespectively. also found for the Mg and Zn elements, as reported theCEDS elemental Figure 11b,c, elements, as reported by the EDS elemental profiles in Figure 11b,c, respectively. Finally, the different solution treatments do not affect the α-Fe and β-Fe intermetallics because Finally, different solution dothe notDSC affectthermograph the α-Fe andin β-Fe intermetallics because of of their high the thermal stability, as ittreatments results from Figure 4. their high thermal stability, as it results from the DSC thermograph in Figure 4.
Figure 11. The distributions of (a) Cu, (b) Mg and (c) Zn across the width of α-Al crystals in as-rheocast Figure 11. The distributions of (a) Cu, (b) Mg and (c) Zn across the width of α-Al crystals in as-rheocast condition and after 6 h of solution treatment at 470 ◦ C. condition and after 6 h of solution treatment at 470 °C.
The Vickers hardness of the as-quenched AlSi9 Cu3 (Fe) alloy as a function of the solution treatment The Vickers hardness of the as-quenched AlSi9Cu3(Fe) alloy as a function of the solution conditions is illustrated in Figure 12; the hardness value of as-rheocast alloy has also been reported in treatment conditions is illustrated in Figure 12; the hardness value of as-rheocast alloy has also been the graph. reported in the graph. For the solution conditions considered in this work, it results that the hardness of the alloy increases by increasing the solution temperature and time; therefore, the optimal strengthening response of the alloy is achieved after thermal treatment at the most severe condition. It is worth noting that the solution treatments carried out at 450 ◦ C led to lower hardness values with respect to as-rheocast condition. In particular, the material softening is larger in the earlier stage
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For the solution conditions considered in this work, it results that the hardness of the alloy increases by increasing the solution temperature and time; therefore, the optimal strengthening response the alloy is achieved after thermal treatment at the most severe condition. Metals 2018,of 8, 750 10 of 18 It is worth noting that the solution treatments carried out at 450 °C led to lower hardness values with respect to as-rheocast condition. In particular, the material softening is larger in the earlier stage of the treatment, then the hardness slightly improves by prolonging the holding time and finally maintains at a steady value after 3 h. Similar behaviour was found for a thixocast 319 alloy solutioned at 500 ◦°C C [29]. The The initial initial softening softening of of the the alloy alloy may may be be due due to the eutectic Si particles coarsening, coarsening, which is characterized by the increase of the inter-particle distance of the eutectic Si [30]; on the other side, the subsequent strengthening can be associated with the increasing content of Cu, Mg and Si in solid solution. The The hardness hardness evolution evolution approaches approaches the the as-rheocast as-rheocast value, value, and and itit stabilizes stabilizes after after 44 h, h, indicating how the α-Al matrix is saturated. α-Al
Figure time. Figure 12. 12. Hardness Hardness of of rheocast rheocast AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy as function of solution temperature and time. ◦ C for 3 and 6 As mentioned, the alloy solutioned solutioned at at the most severe conditions conditions (i.e., (i.e., 490 490 °C for 3 and 6 h) exhibits higher hardness values; however, blistering phenomena were occasionally observed on the specimens’ thethe aforementioned heat treatment conditions. Blisters are generally specimens’ surface surfaceafter after aforementioned heat treatment conditions. Blisters are associated generally to the expansion of gas porosity, includes from physical gas from lubricant associated to the expansion of gaswhich porosity, whichgas includes gas fromentrapment, physical entrapment, gas from decomposition, and gas initially dissolved in the molten metal (e.g., hydrogen). The discrimination of lubricant decomposition, and gas initially dissolved in the molten metal (e.g., hydrogen). The each contribution to the total porosity duetotal to gas entrapment notentrapment possible in is this However, discrimination of each contribution to the porosity due toisgas notwork. possible in this the gas porosity due to dissolved hydrogen is not common in pressure casting processes, where the work. However, the gas porosity due to dissolved hydrogen is not common in pressure casting formation of gas pores is mainly attributed to the decomposition of coolant and die lubrication [31]. processes, where the formation of gas pores is mainly attributed to the decomposition of coolant and In the present die lubrication [31].study, the subsequent ageing treatments were applied to the alloy solutioned at 470 ◦In C for 6 h which achieved high hardness values and, concurrently, showed a free-defects surface. the present study, the subsequent ageing treatments were applied to the alloy solutioned at
470 °C for 6 h which achieved high hardness values and, concurrently, showed a free-defects surface. 3.3. Ageing Condition 3.3. Ageing Condition 3.3.1. Hardness Testing
age hardening 3.3.1.The Hardness Testing behaviour at 160, 180 and 220 ◦ C is shown in Figure 13. The results indicate that the maximum hardness value is achieved at 160 ◦ C for 24 h. By temperature increasing, the peak The age hardening behaviour at 160, 180 and 220 °C is shown in Figure 13. The results indicate hardness is reached for shorter times because of faster precipitation kinetics, i.e., 5 h at 180 ◦ C and that the maximum hardness value is achieved at 160 °C for 24 h. By temperature increasing, the peak 2 h at 220 ◦ C. It is well known that the precipitation rate during heat treatment is related to the hardness is reached for shorter times because of faster precipitation kinetics, i.e., 5 h at 180 °C and 2 diffusion of the solute elements in the α-Al matrix and this strictly depends on the temperature [32]. h at 220 °C. It is well known that the precipitation rate during heat treatment is related to the diffusion As described in the literature [15], the experimental peak value progressively decreases as the ageing of the solute elements in the α-Al matrix and this strictly depends on the temperature [32]. As temperature increases: 117 ± 3 HV at 160 ◦ C, 103 ± 3 HV at 180 ◦ C and 95 ± 2 HV at 220 ◦ C. It is described in the literature [15], the experimental peak value progressively decreases as the ageing worth noting that the hardness in the earlier stages of ageing is slightly lower when compared to the temperature increases: 117 ± 3 HV at 160 °C, 103 ± 3 HV at 180 °C and 95 ± 2 HV at 220 °C. It is worth as-quenched condition after the solution treatment at 470 ◦ C for 6 h; for these relative short times, the mechanical response appreciably decreases by decreasing the ageing temperature. As reported by Maksimovic et al. [33], this trend can be ascribed to the dissolution of GP zones formed during previous solution heat treatment.
noting that the hardness in the earlier stages of ageing is slightly lower when compared to the asquenched condition after the solution treatment at 470 °C for 6 h; for these relative short times, the mechanical response appreciably decreases by decreasing the ageing temperature. As reported by Maksimovic et al. [33], this trend can be ascribed to the dissolution of GP zones formed during Metals 2018, 8, 750 11 of 18 previous solution heat treatment.
Figure 13. 13. Age Age hardening curves of rheocast AlSi9 Cu3 (Fe) solutioned h. hardening curves of rheocast AlSi9Cu3(Fe) alloyalloy solutioned at 470at °C470 for 6◦ Ch. for The6asThe as-quenched condition after the solution treatment is also indicated. quenched condition after the solution treatment is also indicated.
3.3.2. Precipitation Microstructure 3.3.2. Precipitation Microstructure The artificial ageing treatments induce no visible effects at microscale in the alloy. Therefore, high The artificial ageing treatments induce no visible effects at microscale in the alloy. Therefore, magnification bright field TEM observations were carried out for the peak ageing conditions at 160, high magnification bright field TEM observations were carried out for the peak ageing conditions at ◦ C, in order to investigate the nature of strengthening phases responsible for the greatest 180 and 220 160, 180 and 220 °C, in order to investigate the nature of strengthening phases responsible for the material hardening. Where not differently specified in the text, TEM micrographs and SAED patterns greatest material hardening. Where not differently specified in the text, TEM micrographs and SAED have beenhave acquired the zone axis Alto zone have aaxis higher contrast between the precipitates patterns beenclose acquired closeAl the to have a higher contrast between and the the matrix. precipitates and the matrix. In In an an Al-Si-Cu-Mg Al-Si-Cu-Mg system, system, the the precipitation precipitation sequence sequence is is normally normally considered considered quite quite complex complex because of the occurrence of a great variety of structures that can range from binary to complex because of the occurrence of a great variety of structures that can range from binary to complex quaternary thethe strengthening mechanism of Al-Si-Cu-Mg quaternary phases. phases. Therefore, Therefore,aalarge largediscrepancy discrepancyabout about strengthening mechanism of Al-Si-Cualloys after solution and artificial ageing treatments emerges from the literature; consequently, Mg alloys after solution and artificial ageing treatments emerges from the literature; consequently, more investigations are are requested requested to to consolidate the understanding of the the age more experimental experimental investigations consolidate the understanding of age hardening mechanism. hardening mechanism. In In order order to to support support the the TEM TEM findings, findings, several several works works were were examined examined concerning concerning the the ageing ageing response of hypoeutectic cast Al-Si-Cu(-Mg) alloys and wrought Al-Mg-Si(-Cu) and Al-Cu-Mg(-Si) response of hypoeutectic cast Al-Si-Cu(-Mg) alloys and wrought Al-Mg-Si(-Cu) and Al-Cu-Mg(-Si) alloys. Although the the experimental experimental alloy alloy in in this this work work is is aa foundry foundry Al-Si-based Al-Si-based alloy, alloy, the alloys. Although the α-Al α-Al matrix matrix results to be Mgand Cu-enriched due to the dissolution of the Mgand Cu-rich intermetallic phases results to be Mg- and Cu-enriched due to the dissolution of the Mg- and Cu-rich intermetallic phases after Moreover, the the Si Si content content in in the the solid after the the solution solution heat heat treatment, treatment, as as illustrated illustrated in in Figure Figure 11. 11. Moreover, solid solution α-Al matrix is further reduced due the precipitation of Si particles, as previously solution α-Al matrix is further reduced due the precipitation of Si particles, as previously reported reported (Figures 10). Therefore, Therefore, the the age age hardening hardening mechanism mechanism in in wrought wrought Al-Mg-Si Al-Mg-Si or or Al-Cu-Mg Al-Cu-Mg based based (Figures 99 and and 10). alloys has to be also considered in order to better understand the precipitation mechanism in the alloys has to be also considered in order to better understand the precipitation mechanism in the rheocast AlSi Cu (Fe) alloy. 9 3 rheocast AlSi9Cu3(Fe) alloy. ◦ C for 24 h (Figure 14a) reveals a large The TEM investigation The TEM investigation of of the the peak peak ageing ageing condition condition at at 160 160 °C for 24 h (Figure 14a) reveals a large amount black dotsdots with awith square-like shape and an average sideaverage length ofside ~4 nm, homogeneously amount ofoffine fine black a square-like shape and an length of ~4 nm, distributed in the α-Al matrix. Several authors [34–36] have reported the presence of very similar homogeneously distributed in the α-Al matrix. Several authors [34–36] have reported the presence of precipitates Al-Mg-Si-(Cu) after ageingalloys treatments they treatments have been identified metastable very similarin precipitates inalloys Al-Mg-Si-(Cu) after and ageing and theyas have been coherent precipitation analysis ofthrough SAED patterns and of high-resolution identifiedβ” as phase metastable coherent through β’’ phasethe precipitation the analysis SAED patternsTEM and micrographs. The appearance of nanometric dots with similar size and shape were also observed in an high-resolution TEM micrographs. The appearance of nanometric dots with similar size and shape aged AlSi7Cu3Mg alloy by Li et al. [35], who stated that the square-shaped particles correspond to the cross sections of needle-like β” precipitates. Some needle-like precipitates that lie along the directions of the α-Al matrix are visible in Figure 14a, confirming the typical orientation relationships
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were also observed in an aged AlSi7Cu3Mg alloy by Li et al. [35], who stated that the square-shaped Metals 2018, 8, 750 12 of 18 particles correspond to the cross sections of needle-like β’’ precipitates. Some needle-like precipitates that lie along the directions of the α-Al matrix are visible in Figure 14a, confirming the typical orientation between β’’ mean and α-Al phases [37]. Theβ”mean length along of needle-like between β” relationships and α-Al phases [37]. The length of needle-like precipitates the β’’ Al precipitates along the Al directions directions was measured about 30 nm. was measured about 30 nm. Beside black dot dot particles, particles, some some precipitates precipitates appear appear with with aa rectangular rectangular cross-section cross-section with with Beside the the black 2 2 mean 5757 nm . By tilting the the sample closeclose to Al direction, these precipitates show a lathmean size size of of22×× nm . By tilting sample to Al direction, these precipitates show shaped morphology (Figure(Figure 14b) with dimension of ~7 nm. et al. that the a lath-shaped morphology 14b)awith a dimension of ~7Linm. Li [35] et al.observed [35] observed thatlaththe shaped precipitates in an aged AlSi7Cu3Mg alloy are quaternary Al-Mg-Si-Cu phase and their lath-shaped precipitates in an aged AlSi7Cu3Mg alloy are quaternary Al-Mg-Si-Cu phase and their chemical the Q’ Q’ phase. phase. The was also also chemical composition composition and and crystal crystal structure structure fit fit the The presence presence of of Q’ Q’ phase phase was observed in an aged AlSi9Cu3.5Mg alloy and the corresponding SAED pattern seems to be close to observed in an aged AlSi9 Cu3 .5Mg alloy and the corresponding SAED pattern seems to be close to the the experimental one obtained in the present work (see inset in Figure 14a). experimental one obtained in the present work (see inset in Figure 14a). It is worth worthmentioning mentioning that, beside Q’ phase, lath-like L precipitates with rectangular cross It is that, beside Q’ phase, lath-like L precipitates with rectangular cross section section and longer dimension lying along the Al direction, can also precipitate in aged Al-Mgand longer dimension lying along the Al direction, can also precipitate in aged Al-Mg-Si-Cu as Si-Cu[38,39]. as wellMoreover, [38,39]. Moreover, to the Lthe phase, the section Q’-lathshas section has the at an well contrary contrary to the L phase, Q’-laths the long sidelong at anside angle of ◦ with angle about ~11° with to respect to theAl Al directions or, equivalently, Q’ precipitates haveAl{510} Al about of ~11 respect the directions or, equivalently, Q’ precipitates have {510} habit habit planes with the Al matrix [36,40–42]. In an aged Al-Mg-Si alloy with Cu excess, it is found that planes with the Al matrix [36,40–42]. In an aged Al-Mg-Si alloy with Cu excess, it is found that the the decomposition of the supersaturated solid solution (SSSS)takes takesplace placeby byaa complex complex precipitation precipitation decomposition of the supersaturated solid solution (SSSS) sequence [43]: sequence [43]:
α(SSSS)→ →atomic atomicclusters clusters→ →GP GPzones zones→→β”,L β”,L→→β’,Q’ β’,Q’ α(SSSS) →→ QQ
(1) (1)
where where the the L L phase phase is is believed believed to to be be aa Q’ Q’ precursor. precursor.
◦ C, Figure 14. TEM TEM micrographs micrographs of of rheocast rheocast AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy alloy after peak ageing ageing treatment treatment at at 160 160 °C, acquired (a) close to to Al direction and and the the corresponding corresponding SAED SAED pattern pattern (inset), (inset), and and (b) after Al direction specimen tilting to AlAldirection. direction.
◦ C for 5 h is shown in Figure 15. peak aged aged condition condition at at 180 180 °C The microstructure for the peak for 5 h is shown in Figure 15. Parallelogram-like dots dots and and fine needles, associated associated to to Q’ and β’’ β” phases, are still the predominant Parallelogram-like precipitates in the alloy. alloy. While the dimensions of the precipitates seem comparable to those observed C for 24 h, after ageing at 160 ◦°C h, the the precipitate precipitate volume volume fraction fraction appears appears lower. lower. The SAED pattern (Figure 15) shows some different reflections compared to those observed in the diffraction pattern of the peak-aged condition C: due to the the complexity complexity of the the chemical chemical composition composition of the examined condition at 160 ◦°C: alloy, the alloy, the precipitation precipitation of of an an extra extra phase phase cannot cannot be be excluded. excluded. C for 2 h is reported in Figure TEM micrograph of peak aged alloy at 220 ◦°C Figure 16. If compared to previous peak aged conditions, fine parallelogram-like precipitates are not revealed. The significant ◦ C seems in agreement of β’’ β”fraction fractiondue duetotothe the temperature increase from to 220 reduction of temperature increase from 180180 to 220 °C seems in agreement with with al.who [44] found who found that β” precipitates not easily transform intostable more phases stable phases YangYang et al. et [44] that β’’ precipitates do notdo easily transform into more (semi◦ (semi-coherent Q’ phases) at for 175prolonged C for prolonged time, the transformation process coherent β’ and β’ Q’and phases) at 175 °C time, while thewhile transformation process becomes becomes more rapid at higher temperature (200 ◦ C). In the microstructure, fine precipitates with
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more rapid rapid at at higher higher temperature temperature (200 (200 °C). °C). In In the the microstructure, microstructure, fine fine precipitates precipitates with with narrow narrow more narrow rectangular morphology and {100} habit plane are also present and identified as Q’ phase. Al rectangular morphology morphology and and {100} {100}Al Al habit plane are also present and identified as Q’ phase. The rectangular habit plane are also present and identified as Q’ phase. The ◦C The dimensions of these precipitates seem comparable to those observed after ageing at 160 and 180 dimensions of of these these precipitates precipitates seem seem comparable comparable to to those those observed observed after after ageing ageing at at 160 160 and and 180 180 °C °C dimensions for 24 24 and and 55 h, h, respectively. respectively. for for 24 and 5 h, respectively.
Figure 15. 15. Microstructure Microstructure of rheocast rheocast AlSi9Cu3(Fe) treatment at 180 ◦°C; °C; SAED Microstructure of rheocast AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy after peak ageing treatment C; SAED Figure pattern in the inset. pattern in the inset.
Long needle-like needle-like precipitates length up up to to ~150 ~150 nm nm and and oriented orientedalong along Al Al direction Long precipitates with with aaa length length up to ~150 nm and oriented along Al direction ◦ C for 2 h. By tilting the TEM specimen close to direction, are also observed after ageing at 220 ageing at at 220 220 °C °C for for 2 h. By tilting the TEM specimen specimen close close to to Al Al direction, Aldirection, are also observed after ageing longer needles transform into platelet-shape precipitates (Figure 16b) with an average size about the longer needles transform into platelet-shape precipitates (Figure 16b) with an average size of the longer needles transform into platelet-shape precipitates (Figure 16b) with an averageofsize of 3 nm23. These 2. These 4.5 × 10 can be identified as θ’ phase that forms fine platelets (thickness ~1 nm) lying about 4.5 × 10 nm can be identified as θ’ phase that forms fine platelets (thickness ~1 nm) 3 2 about 4.5 × 10 nm . These can be identified as θ’ phase that forms fine platelets (thickness ~1 nm) on {100} planes, as largely in literature for Al-Si-Cu-Mg alloys [35,45,46]. lying onAl{100} {100} Al planes, as documented largely documented in literature for Al-Si-Cu-Mg alloys Moreover, [35,45,46]. lying on Al planes, as largely documented in literature for Al-Si-Cu-Mg alloys [35,45,46]. the SAED pattern (Figure 16a) shows reflections at {110} positions and sharp streaks parallel to the Moreover, the SAED pattern (Figure 16a) shows reflections at {110} Al positions and sharp streaks Al Moreover, the SAED pattern (Figure 16a) shows reflections at {110}Al positions and sharp streaks direction: by comparing it with the SAED pattern analyses of thermal treated Al-Cu and parallel to the Al direction: by comparing it with the SAED pattern analyses of thermal treated Al to the Al direction: by comparing it with the SAED pattern analyses of thermal treated parallel Al-Si-Cu–based alloys reported inreported literaturein it can be stated the experimental diffraction Al-Cu and and Al-Si-Cu–based Al-Si-Cu–based alloys reported in[47,48], literature [47,48], it can canthat be stated stated that the the experimental experimental Al-Cu alloys literature [47,48], it be that pattern supports that the platelets are consistent with θ’ phase. diffraction pattern supports that the platelets are consistent with θ’ phase. diffraction pattern supports that the platelets are consistent with θ’ phase.
Figure 16. 16. TEM TEM micrographs of rheocast AlSi9Cu3(Fe) alloy after peak ageing treatment at 220 °C, ◦ C, Figure TEM micrographs micrographs of of rheocast rheocast AlSi9Cu3(Fe) AlSi9 Cu3 (Fe) alloy alloy after after peak peak ageing ageing treatment treatment at at 220 220 °C, acquired (a) close to Al direction and the corresponding SAED pattern (inset) and (b) after acquired Al acquired (a) close to direction and and the the corresponding corresponding SAED SAED pattern pattern (inset) (inset) and and (b) (b) after Al direction specimen tilting tilting to to Al direction. specimen direction. Al direction. Al
The evolution evolution of of the the volume volume fraction fraction of of precipitates precipitates during during ageing ageing treatment treatment can can be be generally generally The estimated according to the Johnson-Mehl-Avrami-Kolmogorov (JMAK) relationship [49]: estimated according to the Johnson-Mehl-Avrami-Kolmogorov (JMAK) relationship [49]:
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The evolution of the volume fraction of precipitates during ageing treatment can be generally estimated according to the Johnson-Mehl-Avrami-Kolmogorov (JMAK) relationship [49]: fV = fm (1 − exp(−ktn )
(2)
where fV is the precipitated volume fraction at a given aging time, t, fm is the maximum volume fraction, k is the overall rate constant that depends on the nucleation and growth rates of the precipitates, and n is a constant usually known as Avrami’s exponent. The volume fraction of the different phases precipitated at the peak aged conditions and observed during TEM investigations was estimated by isothermal phase transformation calculation using JMatPro® commercial software (Version 7.0, Sente Software Ltd., Guildford, UK) [50], where no phase competition was considered. The values of the precipitation kinetics used in the JMAK equation at the ageing temperatures of 160, 180 and 220 ◦ C are listed in Table 3. The volume fractions of the β”, Q’ and θ’ phases at the different peak ageing conditions are shown in Figure 17a. Table 3. Values of n, k and fm used to calculate the volume fraction of the precipitates by means of Equation (2). Ageing Temperature (◦ C)
Phase
n
k
fm (%)
160
β” Q0
1.10 1.16
3.4 × 10−5 6.5 × 10−6
0.56 0.72
180
β” Q0
1.09 1.12
8.1 × 10−5 1.7 × 10−5
0.55 0.79
220
θ0 Q0
1.79 1.09
9.8 × 10−7 8.1 × 10−5
1.93 0.79
Using Equation (2) and the mean dimensions of the particles measured on the TEM micrographs, the average number density, NV , of the different types of precipitate per unit volume can be calculated at the peak ageing conditions as follows NV = fV /(w × l × t)
(3)
where w, l and t are the average width, length and thickness of β”, Q’ and θ’ precipitates previously described as needle-, lath- and plate-like particles. The calculated number density of precipitates is shown in Figure 17b. Figure 17 evidences how the greatest volume fraction as well as the highest number density of β” and Q’ phases is obtained by ageing at 160 ◦ C for 24 h. On the other side, the highest volume fraction of θ’ phase is attained after 2 h ageing at 220 ◦ C. However, increasing ageing temperature up to 220 ◦ C leads to the coarsening phenomenon of precipitates as defined by Ostwald ripening theory in which fine precipitates are incorporated into larger ones to reduce the surface energy per volume of system. This can explain the reduced number density of θ’ phase at 220 ◦ C for 2 h (Figure 17b). Increasing the number density of precipitates per unit volume leads to shortening the distance between fine precipitates. Thus, the hardness of the material is increased accordingly as a function of the different peak ageing condition, i.e., from 95 ± 2 HV (220 ◦ C for 2 h) up to 117 ± 3 HV (160 ◦ C for 24 h). According to the Orowan mechanism, the increased strength of Al matrix can be estimated through the following equation [51] ∆σOrowan = 0.84MGb/λ
(4)
where M is the Taylor factor, G is the shear modulus, b is the Burgers vector and λ is the inter-particle spacing. Then, the Orowan strength is inversely proportional to the inter-particle distance.
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According to the Friedel mechanism, the mechanical properties can be also attributed to the ability of the strengthening precipitates to withstand shearing by gliding dislocations [52]. In general, the precipitation hardening progressively increases with the coarsening of the precipitates until they are shared by dislocations. With further increase of the size of precipitates, the dislocations bypass the precipitates via the Orowan mechanism, i.e., by bowing. When the fraction of precipitates by shearing and bowing mechanisms is equal, the alloy reaches the peak Metals 2018, 8, x FORpassed PEER REVIEW 15 of 18 hardness. Some works reported how β” precipitates are generally sheared by dislocations [53], while theθ’θ’phase phaseisisconsidered consideredaashear-resistant shear-resistantphase, phase,i.e., i.e.,by-pass by-passby bydislocations dislocations[54]. [54].On Onthe theother otherside, side, the the Q’ precipitates interact withwith dislocation both by shearing and bowing depending the precipitatescan can interact dislocation both by shearing and mechanisms, bowing mechanisms, on the particle size [53]. size [53]. depending on the particle very difficult difficult here here to to reliably reliably ascertain ascertain the the exact exact interaction interaction mechanism mechanism between betweenthe the ItIt isis very dislocationsand andthe theprecipitates precipitatesatatthe thedifferent differentpeak peakageing ageingconditions. conditions.AAtransition transitionfrom fromdislocation dislocation dislocations shearing to to dislocation forfor thethe different phases (β”, Q’ and andθ’) forand the different shearing dislocation looping loopingcould couldoccur occur different phases (β’’, Q’θ’) and for the size andsize morphology of precipitates. different and morphology of precipitates. However, the the interaction interaction between between precipitates precipitates and and dislocations, dislocations, that that isis the the precipitation precipitation However, strengtheningofofthe thealloy, alloy,isisdetermined determinedby byseveral severalfactors factorsincluding includingbut butnot notlimited limitedtotosize, size,volume volume strengthening fraction,and andnumber numberdensity densityofofprecipitates. precipitates.The Thecrystal crystalstructure structureofofthe theprecipitates, precipitates,asaswell wellasasthe the fraction, coherency in the α-Al matrix, can significantly affect the dislocation/particle interaction mechanisms. coherency in the α-Al matrix, can significantly affect the dislocation/particle interaction mechanisms. reportedhow howβ’’ β”and andθ’θ’phases phasesshow showaacoherent coherentinterface interfacewith withα-Al α-Al[55], [55],while whileQ’ Q’phase phasecan canform form ItItisisreported coherentor orsemi-coherent semi-coherentinterface interface[56]. [56]. aacoherent
(a)
(b)
Figure Figure17. 17.(a) (a)Calculated Calculatedvolume volumefraction fractionand and(b) (b)number numberdensity densityper perunit unitvolume volumeofofthe thephases phases observed observedby byTEM TEMinvestigation investigationatatthe thedifferent differentpeak peakageing ageingconditions. conditions.
Conclusions 4.4.Conclusions thiswork, work, the the effects effects of treatments onon thethe microstructure and InInthis of different differentsolution solutionand andageing ageingheat heat treatments microstructure hardness of secondary rheocast AlSi Cu (Fe) alloy have been investigated. The following conclusions 9 AlSi9Cu3(Fe) 3 and hardness of secondary rheocast alloy have been investigated. The following can be drawn. conclusions can be drawn. •• ••
•
•
◦ Cfor Solution Solutiontreatment treatmentatat470 470°C for66hhresults resultsininbeing beingaagood goodcompromise compromiseininterms termsofofhardening hardening level leveland andsurface surfacequality qualityas asno noblisters blistersappear appearon onthe thecasting castingsurface. surface. ◦ Cleads The Thesolution solutionheat heattreatment treatmentininthe thetemperature temperaturerange rangebetween between450 450and and490 490°C leadstotothe the fragmentation of of eutectic Si particles andand the dissolution of coarse Mg and fragmentationand andspheroidization spheroidization eutectic Si particles the dissolution of coarse Mg Cu-rich phases, which contributes to enhance the the solute content in in thethe α-Al and Cu-rich phases, which contributes to enhance solute content α-Alsolid solidsolution. solution. Moreover, fine polygonal and rod/platelet-shape Si particles appear within the α-Al matrix after solution treatment. The best strengthening degree is achieved after artificial ageing at 160 °C for 24 h; however, by increasing the ageing temperature, the age hardening peaks result for shorter times. On the other hand, the hardness values progressively decrease. The co-existing of highest volume fraction and number density of fine needle-shaped β’’ and Q’
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•
•
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Moreover, fine polygonal and rod/platelet-shape Si particles appear within the α-Al matrix after solution treatment. The best strengthening degree is achieved after artificial ageing at 160 ◦ C for 24 h; however, by increasing the ageing temperature, the age hardening peaks result for shorter times. On the other hand, the hardness values progressively decrease. The co-existing of highest volume fraction and number density of fine needle-shaped β” and Q’ (or L) precipitates is responsible for the maximum strengthening level of the rheocast and heat treated AlSi9 Cu3 (Fe) alloy. Upon increasing the ageing temperature, fine precipitates are incorporated into larger ones leading to a lower number density of strengthening phases; this results in a reduction of the hardness. At higher ageing temperature, such as 220 ◦ C, the formation of large θ’ platelets explains the incremental softening of the alloy with a decrease of about 19% respect to the artificial ageing at 160 ◦ C for 24 h.
Author Contributions: All the authors actively participated to the research: A.F. and S.C. worked on the microstructural and mechanical characterization; G.T. followed the casting trials; A.D.M. carried out the DSC investigation; A.F. and G.T. discussed the results; A.F. wrote the draft and G.T. revised and improved it. Funding: This research was funded by Raffineria Metalli Capra Spa (Brescia, Italy). Acknowledgments: The authors would like to acknowledge the skilful contribution of CIE Automotive S.A. (Bilbao, Spain) for rheocasting experiments. Conflicts of Interest: The authors declare no conflict of interest.
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