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Role of Zr and Sc addition in controlling the microstructure and tensile properties of aluminum–copper based alloys
Materials and Design 88 (2015) 1134–1144
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Role of Zr and Sc addition in controlling the microstructure and tensile properties of aluminum–copper based alloys A.M. Samuel a, S.A. Alkahtani b, H.W. Doty c, F.H. Samuel a,⁎ a b c
Université du Québec à Chicoutimi, Québec, Canada Industrial Engineering Program, Mechanical Engineering Department, College of Engineering, Prince Sattam bin AbdulAziz University, Al Kharj, Saudi Arabia General Motors, Materials Engineering, 823 Joslyn Avenue, Pontiac, MI 48340, USA
a r t i c l e
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Article history: Received 23 June 2015 Received in revised form 11 September 2015 Accepted 15 September 2015 Available online 18 September 2015 Keywords: Aluminum alloys Zr and Sc addition Tensile properties Heat treatment Grain refining
a b s t r a c t The present work was carried out on eight alloys containing Al–2%Cu–1.5%Si–0.5%Fe–0.6%Mn-0.4%Mg–0.07%Ti, with small amounts of Zr, Sc or both. Tensile test bars were produced using low pressure die casting. All bars were solution heat-treated, followed by artificial aging. The as-cast microconstituents were determined using electron probe microanalyzer equipped with both electron dispersive x-ray and wavelength dispersive spectroscopic systems. The results show that in addition to the currently observed Fe-and Cu-based intermetallics, several Zr- and Sc-containing intermetallics were also formed during solidification (of the present alloys). Neither the Sc/Zr nor the Sc/Al ratio is constant, and both ratios may vary from one particle to another. The presence of Zr or Sc has a significant grain refining effect, equivalent to that obtained from the addition of a Ti-based grain refiner. The main influence of a Zr or Sc addition on the tensile properties of the base alloy appears during aging, and before the onset of softening. Although the addition of Zr or Sc improves the alloy tensile properties, Sc is less effective compared to Zr. Also, a Sc content beyond 0.15% results in no further improvement in the alloy strength. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Zirconium (Zr) is a transition element used in aluminum alloys to enhance their microstructure and mechanical properties. When Zr is added in small quantities, it forms tiny, coherent precipitates of L12ordered trialuminide (Al3Zr) through the decomposition of a supersaturated solid solution of aluminum [1,2]. The Al3Zr phase is distinguished by two different crystallographic structures: (i) the first is the stable tetragonal DO23 structure which is body-centered with eight atoms per unit cell, while (ii) the second is the coherent metastable Al3Zr phase with an L12-type structure which is simple cubic with four atoms per unit cell and having a slight lattice mismatch with α-Al [3,4]. The structure of the Al3Zr intermetallic phase is dependent on the level of Zr added to the alloy, as well as on the cooling rate. Furthermore, Al3Zr can dissolve up to 5% Sc and form a Al3(Zr1 − xScx) phase particle; when the Al3Zr phase dissolves the Sc, the lattice parameter of this phase is observed to decrease [5]. The Al3(Zr1 − xScx) phase may form upon the decomposition of the solid solution or as a primary phase during solidification. The primary phase type of Al3(Zr1 − xScx) has a greater efficiency as a grain refiner than the primary Al3Zr phase particle. Moreover, the primary Al3Zr
⁎ Corresponding author. E-mail address: [email protected] (F.H. Samuel).
http://dx.doi.org/10.1016/j.matdes.2015.09.090 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
phase particles may act as nucleation sites for Al3Sc [6]. As the Al3Sc phase nucleates on the surface of the Al3Zr phase particle, the grain refining capacity of the primary Al3Zr phase particles for α-Al shows improvement because of the small lattice misfit between the Al3Sc phase and α-Al [7,8]. The primary particles of Al3(ScxZr1 − x) have a greater refining efficiency than the primary particles of Al3 Sc or Al3Zr [9]. The Al 3Sc phase can also dissolve up to 5% of the titanium to produce the Al3(Sc1 − xTix); thus in the presence of titanium combined with zirconium Al 3(Sc1 − x − yZrxTiy) may be formed [10,11]. The primary phase particles of Al3(ScxZr1 − x) normally appear in two dimensions in the form of a star-like shape. Scandium reacts with Si and Cu to form two phases: (i) the VAlScSi phase, with a stoichiometric composition of AlSc2Si2 and (ii) the W-AlCuSc phase [12–14]. It is also reported that Sc tends to diminish the solubility of copper and silicon in the aluminum matrix. Seidman et al. [15] and Davydov et al. [16] studied the precipitation of Sc phases in Al–0.5%Sc alloy. The authors reported that Al3Sc precipitates may take place below 350 °C. These Al3 Sc precipitates strengthen the aluminum matrix by pinning the dislocation lines and resisting their motion. In a previous study [17], the authors reported on the effect of various alloying elements on the evolution of microstructure in an Al–2%Cu based alloy during solidification. The present work was undertaken to emphasize the role of Zr and Sc addition in controlling the
A.M. Samuel et al. / Materials and Design 88 (2015) 1134–1144 Table 1 List of the various 220 Al–2%Cu based alloys used in the present study and alloy chemistry in wt.%. Alloy code
Alloy chemistry
220A 220J 220K 220L 220M 220N 220O 220P
220 alloy: Al–2%Cu–1.32%Si–0.42%Mg–0.58%Fe–0.59%Mn–0.07%Ti 220A + 0.15%Ti + 0.30%Zr + 0.15%Sc 220A + 0.15%Zr 220A + 0.30%Zr 220A + 0.50%Zr 220A + 0.15%Sc 220A + 0.3%Sc 220A + 0.5%Sc
microstructure, by applying optical microscopy and electron probe microanalysis to identify and quantify the phases formed by the interactions between the different alloying elements Zr, Sc, Ti, as well as to study the effects of these intermetallic phases on the grain structure. The present investigation also aims at determining the optimum age hardening heat treatment required to maximize the tensile properties. 2. Experimental procedure Alloy 220 was received in the form of 12-kg ingots. The ingots were melted in a 60 kg-capacity crucible, using an electrical resistance
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furnace. The melting temperature was held at 740 ± 5 °C. The molten metal was degassed for 30 min using pure dry argon injected into the molten metal by means of a graphite rotary degassing impeller. The degassing time/speed was kept constant at 30 min/150 rpm. After degassing, the molten metal surface was skimmed carefully to remove all the oxide layers and dross before pouring the melt into the ingot molds. Samplings for chemical analysis were also taken simultaneously for each alloy melt composition prepared. The ingots thus prepared were later remelted in the furnace, and used for preparing melts for casting the LPDC (low pressure die casting) test bars. Tensile test bars were cast via LPDC. The permanent die and the filling system were placed over the furnace containing the molten alloy. The cavity is filled by forcing the molten metal (using pressurized gas at 0.3 to 1.5 bar) to rise into a ceramic tube which connects the die to the furnace. Once the die cavity is filled, the overpressure in the furnace is removed, and the residual molten metal in the tube flows again into the furnace. The various parts of the die are then separated, and the casting extracted. The tensile test bars were subsequently cut from their castings with dimensions corresponding to ASTM B-108 standard (gauge length 50 mm, diameter 12.5 mm, total length 200 mm). Table 1 lists the chemical composition of the used alloy. It should be mentioned here that all element contents are expressed in wt% throughout the article unless otherwise precised.
Fig. 1. Macrostructure of as cast alloys: (a) 220A—432 μm, (b) 220J—52 μm, (c) 220M — 80 μm, and (d) 220P — 70 μm.
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The as-cast test bars were subjected to heat treatment which consisted of solution heat treatment at 495 °C for 8 h (heating rate was about 2 °C/min.), followed by quenching in warm water, and then artificial aging at temperatures of 155, 180, 200, 220, 240 and 300 °C for 5 h at each temperature. The tensile testing was carried out using a servo-hydraulic mechanical testing machine at a strain rate of 1 × 10−4 s− 1. The average ultimate tensile strength (UTS), yield strength (YS) and percent elongation (%El) values were calculated over the five bars tested for each condition. All tests were carried out at room temperature. Grain size measurements were carried out for all of the alloys investigated. For these measurements, after the final grinding stage, the samples were etched using a solution made up of 12.5 g CrO3, 2.5 mL HF,
30 mL HCl, 40 mL HNO3 and 42.5 mL water. Grain size measurements were then carried out using an optical microscope and the line intercept method. Samples were sectioned from as-cast and heat-treated tensile-tested bars for metallographic examination. The samples were mounted in bakelite, ground, and polished to a fine finish, using standard polishing procedures. An electron probe microanalyzer, with electron dispersive x-ray spectrometry (EDS) and wavelength dispersive spectroscopic (WDS) facilitates, was used mainly to detect and analyze the chemical compositions of the various intermetallic phases which formed during solidification. Volume fraction and line-scan measurements were also carried out using the same set-up. X-ray images of various elements constituting the identified phases were also obtained to determine the
Fig. 2. Backscattered electron images of: (a) 220A, (b) 220J, (c) high magnification of (b), (d) 220M, and (e) 220P alloys.
A.M. Samuel et al. / Materials and Design 88 (2015) 1134–1144
distribution of these elements in the phases. In these cases the samples were cleaned in high purity ethanol in an ultrasonic agitator before being examined in the electron probe microanalyzer. In order to understand the role of precipitation, samples from the tensile tested bars were cut 1 cm away from the fracture surface, mechanically polished, followed by ion bombardment prior to examination using field emission scanning electron microscopy. 3. Results and discussion
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Table 3 Total volume fraction of intermetallics observed in the 220 as-cast alloy samples. Alloy
A
J
K
L
M
Average Std Dev.
3.98 0.42
5.97 0.6
4.7 0.42
5.1 0.48
6.42 0.62
Alloy
N
O
P
Average Std Dev.
4.21 0.4
5.21 0.63
4.87 0.59
3.1. Microstructural characterization The grain refining effect of these additives may be seen more clearly in the macrostructures of the etched alloy samples, as shown in Fig. 1. It is interesting to note the increase in the level of refinement on going from Alloys A to P as more grain refiners are added. The fineness of the macrostructure is reflected by the secondary dendrite arm spacing (SDAS), so that the finer the microstructure obtained, the smaller the average SDAS value exhibited by the solidified alloy sample. It is well known that SDAS has a strong influence on the mechanical properties of cast aluminum alloys [18,19]. Fig. 2(a) shows the morphology of the intermetallics observed in the base alloy after casting. The α-Fe (Al15(Mn,Fe,Cu)3Si2) is gray in color, while the Al2Cu-phase may appear either in block-like or eutectic form (Al + Al2Cu). The Q-phase, Al5Mg8 Cu2 Si6, may appear in the form of small particles growing out of the Al 2 Cu phase or alone. Based on the preceding, the Al–2%Cu base alloy possesses simple microstructural components consisting of three main intermetallics, namely Al15(Mn,Fe,Cu)3Si2, Al2Cu and Q-Al5Mg8Cu2Si6 phases, in addition to the α-Al matrix. The absence of β-Al5 SiFe phase in the microstructure may be attributed to the high Mn content in this alloy [20]. Table 2 lists the main intermetallics observed in the present alloys whereas Table 3 shows their volume fraction. The star-like phase, Al3(Sc1 − xZrx), is considered to be the most important of the Sc–Zr intermetallics because of its heightened efficiency in nucleating α-Al grains. This type of efficiency may be attributed to the L1 2 -crystal structure of the star-like phase which is characterized by its low mismatch with the α-Al lattice structure [21]. Fig. 2(b) and (c) show the backscattered image of a star phase particle observed in the 220J alloy. It consists of different layers distinguished
Table 2 Chemical composition of intermetallics observed in the present work. Alloy code
Composition
Suggested formula
220A
Al–72.3 at.%, Fe–9.5 at.%, Si–10.5 at.%, Cu–2.4 at.%, Mn 5.2 at.% Al–18.4 at.%, Mg–40.3 at.%, Si–31 at.%, Cu–9.2 at.% Al–68 at.%, Cu–30.6 at.%, Si–1.1 at.% Al–78.1 at.%, Zr–7.6 at.%, Si–1.5 at.%, Sc–12.1 at.% Al–78.2 at.%, Zr–10.3 at.%, Ti–3.6 at.%, Sc–7.4 at.%
Al15(Fe,Mn,Cu)3Si2
220A 220A 220 J 220 J 220 J 220 M 220 M 220 M 220 M 220 M 220P 220P 220P
Al–75 at.%, Sc–11.1 at.%, Zr–10 at.%, Ti–3.2 at.% Al–68.4 at.%, Zr–31.4 at.% Al–63 at.%, Zr–14.2 at.%, Ti–17 at.%, Si–5.8 at.% Al–75.5 at.%, Zr–20.8 at.%, Ti–2.4 at.% Al–67.5 at.%, Zr–23.1 at.%, Ti–8.7 at.% Al–63.1 at.%, Zr–14 at.%, Ti–16.9 at.%, Si–5.75 at.% Al–77 at.%, Sc–23 at.% Al–38 at.%, Mg–2.7 at.%, Sc–26.8 at.%, Si–32 at.% Al–19.8 at.%, Mg–3.4 at.%, Sc–39.5 at.%, Si–36.7 at.%
Al5Mg8Si6Cu2 Al2Cu Al3(Zr,Sc) (Sc/Zr = 1.6) Al3(Zr,Sc,Ti) (Sc/Zr = 0.72) (Sc/Zr = 1) Al3Zr (Al,Si)3(Zr,Ti) Al3(Zr,Ti) Al2(Zr,Ti) (Al,Si)2(Zr,Ti) Al3Sc AlSiSc (Si/Al = 1) AlSi2Sc2 (Si/Al = 2)
by the change in brightness across the particle caused the variation of the Zr and Sc concentrations within the star phase as shown in the colored maps illustrated in Fig. 3. Based on the WDS analysis of these particles (Table 3) the average composition of the star phase was determined to be Al3(Sc,Zr,Ti) with a Zr:Sc ratio ranging between 0.72 and 1.6, whereas Ti concentration is constant at 3.6 at.%. It has been suggested by Zakharov et al. [22] that as the alloy melt passes the liquidus line, crystals of Al 3Zr are primarily nucleated; these nuclei then grow by absorbing the zirconium from the surrounding liquid. As the Zr is depleted from these regions, the nuclei of the star phase, namely Al3Zr, continue to grow through the precipitation of Al3Sc on the surface of these nuclei. At this point, the Al3Zr starts to precipitate on the surface of Al3Sc as the scandium also begins to become progressively depleted from the surrounding liquid, and so on in sequence. The fluctuations in the precipitation sequences are considered to be the main reason for the layer formation throughout the star phase [14]. Toropova et al. [23] proposed that since the Al 3 Zr phase nucleates as the primary phase, the star phase grows predominantly through the precipitation of Al3Sc on the surface of the Al3Zr in the form of layers displaying different degrees of brightness. The variation in brightness from one layer to another is ascribed to the variation in the amount of Zr atoms replacing Sc atoms in the Al3Sc phase. This suggestion is in agreement with the element distribution shown in Fig. 3 which is also in agreement with what has already been suggested earlier by other researchers [12, 24]. It has been observed that in the case where Sc/Zr is higher than unity, the particles precipitate in two different forms: star-like and multi-faceted shape as shown in Fig. 4. The reason however is not known. Fig. 2(d) exemplifies the precipitation of Zr-rich phases in 220M alloy. It is evident from this figure that the particles are having different shapes, size and distribution. Table 3 lists the chemical composition of some of the Zr-rich phase particles. Srinivasan and Chattopadhyay [25] proposed that as a result of high zirconium content (0.5%), all Zr-containing alloys are located in the L + Al3Zr region of the Al– Zr phase diagram during the melting stage [26]. This statement indicates that the Al3Zr particles are not formed from the melt during solidification but may come directly from the Al–15%Zr master alloy added to the liquid melt. Since these particles do not dissolve in the melt, they provide favored nucleation sites for the growth of intermetallics from the melt during solidification. In addition to zirconium trialuminide, some ternary and quaternary intermetallics have been detected viz. (Al,Si)3 (Zr,Ti), Al3 (Zr,Ti), Al2(Zr,Ti) and (Al,Si)2(Zr,Ti). Based on the homogeneous distribution of Zr, Ti and Si within the phase, it may be suggested that AlZrSi and AlZrTiSi, are formed at the same temperature of the precipitation of Al3Zr phase (≈665 °C [27]). Fig. 2(e) shows the backscattered electron image obtained from 220P alloy revealing the precipitation of Sc-rich phase (white circles) mainly in the form of star-like particles. Table 3 reveals that in addition to scandium trialuminide, other compounds can also precipitate mainly V-AlSiSc phase [28] with two Sc/Al ratios. However, none of these phases contains Ti. Fig. 5 exhibits the presence of Al3Sc phase in the form of a fan-like shape. Based on the Al–Sc binary phase diagram at scandium level of 0.55%Sc, the primary particles of L12–Al3Sc are formed
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Fig. 3. (a) Backscattered image (CP) showing an Al3(Sc,Zr,Ti) phase particle in the 220 J alloy in the as cast condition and corresponding X-ray images of (b) Ti, (c) Zr and (d) Sc. Arrows indicate high Zr concentration within the star-like particles. The Sc/Zr ratio is ~unity.
Fig. 4. (a) Backscattered image showing an Al3(Sc,Zr,Ti) phase particle in the 220 J-alloy in the as cast condition and corresponding X-ray images of (b) Ti, (c) Zr, and (d) Sc. The Sc/Zr ratio is ~1.6.
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Fig. 5. (a) Backscattered image showing an Al3Sc phase particle in the 220P alloy in the as cast condition, and corresponding X-ray images of (b) Si, (c) Al, (d) Sc and (e) corresponding EDS spectrum.
in a eutectic reaction L + Al3Sc → (Al) (refined) at 665 °C [29], and act as nucleation sites for α-Al in the molten metal during solidification. The V-phase is a ternary intermetallic compound (AlScSi) formed through invariant reactions, and has a gray color and a stoichiometric composition of AlSc2Si2 [30,31]. Fig. 6 is an example of the distribution of Si, Al, and Sc in the V-phase observed in the 220P alloy — white arrow points to the presence of Al2Cu phase particle adjacent to the V-phase. It should
be mentioned here that W-AlSiCu phase has not been detected in the present investigation. 3.2. Tensile properties In this section the alloys have been divided into three groups: (i) 220A and 220J, (ii) 220K, 220L and 220M and (iii) 220N, 220O and
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Fig. 6. (a) Backscattered image showing an AlSi2Sc2-V phase particle in the 220P alloy in the as cast condition, and corresponding X-ray images of (b) Si, (c) Al and (d) Sc,
220P. The variation in tensile properties of these alloys in the as-cast and solution heat-treated (SHT) conditions is shown in Fig. 7. Although the addition of 0.5%Sc caused a significant reduction in the alloy original grain size (432 μm vs. 52 μm), the corresponding tensile values are noticeably lower than those obtained from 220 M alloy containing 0.5%Zr. Increasing the amount of added Sc from 0.15% to 0.5% also has no corresponding effect on the alloy mechanical performance. As expected solution heat treatment resulted in partial dissolution of the alloying elements and hence relatively better UTS levels compared to those obtained in the as-cast condition. The alloy UTS is mainly controlled by the microstructure and soundness of the casting. As shown in the microstructure section, the microstructure of the investigated alloys in the as-cast condition contains a large number of solidified intermetallic particles or phases (Table 3). These hard particles have a harmful effect on the alloy UTS and %El. Also their sharp edges act as crack initiators. During tensile testing, cracking commences by the breaking up of these intermetallics or by decohesion of the intermetallic particles from the surrounding matrix [32,33]. When the alloy is solution heat treated at sufficiently high temperature, as the solution heat treatment is applied, these intermetallics will break down and dissolve by losing their elements which then return to the matrix producing a supersaturated solid solution of αAl; this solid solution strengthens the matrix by increasing the resistance to dislocation movement during the deformation of tensile testing [34,35].
The effect of artificial aging in the temperature range of 155 °C–300 °C on the tensile properties of 220A and 220J alloys is shown in Fig. 8. As expected, increasing the aging temperature from 155 °C to 300 °C will cause significant reduction in the alloy UTS, as shown in Fig. 8(a), by nearly 30%. Maximum strength was achieved when the alloys were aged at 180 °C, representing peak-aging condition. From Figs. 7 and 8, it is evident that aging the 220A and 220J alloys at 180 °C for 5 h increased their UTS values by about 60% with respect to the as-cast condition and 35% compared to SHT values. Similar observations are made for the YS levels. The addition of Zr–Ti produces a refined non-dendritic structure that decreases the probability for porosity formation and increases the amount of dissolved Al2Cu after solution heat treatment, and thereby the level of Cu in solid solution of α-Al, which then allows for a larger amount of precipitation hardening to take place upon aging, and therefore improve the alloy strength. Additionally, dispersoid precipitates of Al3Zr and/or Al3(Zr1 − xTix) may act as nucleation sites for the hardening phases during the aging process, resulting in further improvement in strength [36–38] Although the equilibrium solid solubility of zirconium in aluminum is very low, the high binding energy results in substantial vacancy trapping in solution treated and quenched alloys which results in modification of precipitation of S′-phase in dilute Al– Cu–Mg alloys. In order to enhance the mechanical properties of an aluminum alloy, it is important to obtain a thermally stable microstructure and coarsening-resistant dispersoids. This may be achieved by adding
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Fig. 7. Tensile properties of the present alloys in the as cast and solution heat treated conditions: (a) UTS and YS, (b) %El.
Zr, which has the smallest diffusion flux in aluminum of all the transition metals; the presence of Zr leads to the formation of fine dispersoids that resist coarsening at higher temperatures, which helps to improve/ maintain the mechanical properties [38,39]. Fig. 9 presents the tensile values of Zr-containing alloys. As can be seen, increasing the amount of added Zr resulted in increasing the alloy UTS values of all alloys. For example, the as cast value of the 220 M alloy increased from 280 MPa the as cast condition to 375 MPa after aging at 180 °C, representing an increase of about 35–40%, caused by simultaneous precipitation of Al2Cu and Al3Zr. In consistence with the data plotted in Fig. 8(a), the presence of Zr resulted in the extension of the aging temperature before the onset of softening, as shown in Fig. 9(a). Fig. 8(b) displays the variation in YS levels as a function of aging temperature. It can be seen that the YS values of the 220 M alloy are very close to those obtained for the corresponding UTS in the temperature range 155 °C–200 °C, followed thereafter by a fast drop to reach 100 MPa at 300 °C (almost 50% reduction compared to as-cast condition).The advantage of adding Zr in relatively small amounts, e.g. 0.15% (220K alloy), appears in the increase in the alloy ductility from 4.2% in the as-cast condition to 5.7% after aging at 300 °C. Applying the quality index equation proposed by Drouzy et al. [40] for Al–Si–Mg alloys, Q ðMPaÞ ¼ UTS ðMPaÞ þ 150 logð%ElÞ the quality index value of 220 L alloy is seen to vary from 313 MPa in the as cast condition to 395 MPa after aging at 180 °C (peak-aging) to 345 MPa after aging at 300 °C.
Fig. 8. Tensile properties of 220A and 220J alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
The work of Knipling et al. [41,42] on the precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc–0.1Zr (at.%) alloys during isochronal aging revealed that precipitation of Al3Sc(L12) commences between 200 °C and 250 °C in the Al–0.1Sc alloy, reaching a peak microhardness of 668 ± 20 MPa at 325 °C. In the Al–0.1Zr alloy, precipitation of Al3Zr (L12) commences between 350 °C and 375 °C, achieving a peak microhardness of 420 MPa at 425–450 °C. While the Al–0.1Sc–0.1Zr alloy achieves peak microhardness after isochronal aging to 400 °C, it overages after extended annealing at this temperature indicating that it is not suitable for extended use at 400 °C. A pronounced synergistic effect
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Fig. 10. Tensile properties of 220N, 220O, 220P alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
Fig. 9. Tensile properties of 220K, 220L, 220M alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
is observed when both Sc and Zr are present. Above 325 °C, Zr additions provide a secondary strength increase that is attributed to precipitation of Zr-enriched outer shells onto the Al3Sc precipitates, leading to a peak microhardness of 618 MPa at 400 °C for Al–0.06Sc–0.06Zr.
The variation in the tensile properties on the Sc-containing alloys as a function of aging temperature is presented in Fig. 10. Apparently, increasing the Sc content from 0.15 to 0.5% resulted in close UTS and YS values regardless of the aging temperature. It is evident from Fig. 10 that Sc addition is less effective in improving the alloy UTS, YS and %El values compared to that exhibited by Zr addition when the alloys are aged in the temperature range 155 °C–240 °C. Aging beyond 240 °C resulted in sharp decrease in the alloy strength regardless of the Zr or Sc content.
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Fig. 11. Precipitation of Al2Cu phase in 220A alloy as a function of heat treatment condition: (a) as cast, (b) SHT, (c) aged-180 °C, (d) aged 240 °C, (e) EDs spectrum corresponding to the arrowed particle in (d), (f) aged 300 °C.
3.3. Precipitation Fig. 11 presents the precipitation taking place in the base 220A alloy at different working conditions. Fig. 11(a) shows the presence of fine bright particles in the matrix of the alloy after casting, together with massive Al2Cu phase particles (white arrow). The variation in the precipitate size may be due to the cooling temperature range after solidification. After solutionizing at 495 °C for 8 h, the bulky particles dissolved almost completely in the aluminum matrix (Fig. 11(b)). However, some particles (less than 0.25 μm long) can easily be seen in the matrix. These particles have been identified as Si precipitates. Dense ultrafine spherical particles throughout
the entire matrix were the main feature obtained when the alloy was aged at 180 °C for 5 h (Fig. 11(c)) causing the observed increase in the alloy strength. Increasing the aging temperature to 240 °C resulted in a partial change in the morphology of the precipitated particles from spherical to elongated or rod-like shape as illustrated in Fig. 11(d). The white arrow in Fig. 11(d) points to undissolved Curich phase. The associated EDS spectrum (Fig. 11(e)) reveals that the undissolved particles could be Al 7Cu2 Fe phase [43]. Aging at 300 °C resulted in complete transformation of the spherical particles into rod shaped ones (Fig. 11(f)) indicating the precipitation of incoherent Al2 Cu phase and hence the observed decline in the alloy strength.
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4. Conclusions Based on the results obtained from the present study, the following conclusions may be drawn: 1. Addition of Zr or Sc in small amounts (about 0.15%) is sufficient to cause about 85% reduction in the alloy grain size with no hot tearing. 2. In addition to Fe- and Cu-based intermetallics, addition of Zr leads to precipitation of Al3Zr, (Al,Si)3(Zr,Ti), Al3(Zr,Ti), Al2(Zr,Ti) and (Al,Si)2(Zr,Ti) phases whereas the addition of Sc leads to the formation of Al3Sc, AlSiSc and AlSi2Sc2 phases during the course of solidification. 3. Combined addition of Zr and Sc causes the precipitation of Al3(Zr,Sc), Al3(Zr,Sc,Ti) and Al3(Sc,Zr,Ti) phases. In all cases neither the Sc/Zr nor Sc/Al ratio is constant. 4. Although the addition of Zr or Sc leads to significant improvement in the alloy strength, the effectiveness of Zr is greater than that obtained from Sc under the same condition (amount and heat treatment). 5. Addition of Sc beyond 0.15% has no further influence on the grain size or on the alloy strength. 6. The importance of additions of Zr or Sc appears not only in improving the alloy strength but also in extending the aging temperature before the onset of softening. 7. During aging up to 300 °C, copper aluminide is the main precipitating phase during aging, which influences the alloy strength. 8. Hardening due to precipitation of Al3Zr or Al3Sc phases has not been detected, indicating that most of the observed Zr- and Sc-based phases occurred only during the solidification process. Acknowledgments The authors would like to thank Ms. Amal Samuel for enhancing the images used in the present work. References [1] N.A. Belov, A.N. Alabin, D.G. Eskin, V.V. Istomin-Kastrovskii, Optimization of hardening of Al–Zr–Sc cast alloys, J. Mater. Sci. 41 (2006) 5890–5899. [2] D. Srinivasan, K. Chattopadhyay, Non-equilibrium transformations involving L12– Al3Zr in ternary Al–X–Zr alloys, Metall. Mater. Trans. A 36A (2005) 311–320. [3] N.A. Belov, D.G. Eskin, A.A. Aksenov, Multicomponent phase diagrams: application for commercial aluminum alloys, 2005 413. [4] L.S. Toropova, D.G. Eskin, M.L. Kharaktrova, T.V. Dobakina, Advanced aluminum alloys containing scandium: structure and properties, Gordon and Breach Science Publishers, Canada, 1998 175. [5] J. Royset, N. Ryum, Scandium in aluminum alloys, Int. Mater. Rev. 50 (2005) 19–44. [6] A.A. Rao, B.S. Murty, M. Chakraborty, Role of zirconium and impurities in grain refinement of aluminum with Al–Ti–B, Mater. Sci. Technol. 13 (1997) 769–777. [7] J.D. Robson, P.B. Prangnell, Dispersoid precipitation and process modelling in zirconium containing commercial aluminum alloys, Acta Mater. 49 (2001) 599–613. [8] J.D. Robson, Optimizing the homogenization of zirconium containing commercial aluminum alloys using a novel process model, Mater. Sci. Eng. A 338A (2002) 219–229. [9] M.S. Kaiser, M.K. Banerjee, Effect of ternary scandium and quaternary zirconium and titanium additions on the tensile and precipitation properties of binary cast Al–6 Mg alloys, Jordan J. Mech. Ind. Eng. 2 (2008) 93–99. [10] O.N. Senkov, M.R. Shagiev, S.V. Senkova, D.B. Miracle, Precipitation of Al3(Sc,Zr) particles in an Al–Zn–Mg–Cu–Sc–Zr alloy during conventional heat treatment and its effect on tensile properties, Acta Mater. 56 (2008) 3723–3738. [11] C.B. Fuller, J.L. Murray, D.N. Seidman, Temporal evolution of the nanostructure of Al(Sc,Zr) alloys: part I-chemical compositions of Al3(Sc1 − xZrx) precipitates, Acta Mater. 53 (2005) 5401–5413. [12] E.A. Marquis, D.N. Seidman, Nanoscale structural evolution of Al3Sc precipitates in Al(Sc) alloys, Acta Mater. 49 (2001) 1909–1919. [13] V.V. Zakharov, Effect of scandium on the structure and properties of aluminum alloys, Met. Sci. Heat Treat. 45 (2003) 246–253.
[14] K.B. Hyde, A.F. Norman, P.B. Prangnell, The effect of cooling rate on the morphology of primary Al3Sc intermetallic particles in Al–Sc alloy, Acta Mater. 49 (2001) 1327–1337. [15] D.N. Seidman, E.A. Marquis, D.C. Dunand, Precipitation strengthening at ambient and elevated temperatures of heat-treatable Al(Sc) alloys, Acta Mater. 50 (2002) 4021–4035. [16] V.G. Davydov, T.D. Rostova, V.I. Yelagin, Scientific principles of making an alloying addition of scandium to aluminum alloys, Mater. Sci. Eng. A 280A (2000) 30–36. [17] G.A. Zaki, A.M. Samuel, H.W. Doty, F.H. Samuel, Microstructural evolution during solidification of Al–Cu based alloys, Int. J. Mater. Sci. (2015) (in press - posted online on 10 August 2015). [18] E. Samuel, B. Golbahar, A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of grain refiner on the tensile and impact properties of Al–Si–Mg cast alloys, Mater. Des. 56 (2014) 468–479. [19] A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of grain refining and Srmodification interactions on the impact toughness of Al–Si–Mg cast alloys, Mater. Des. 56 (2014) 264–273. [20] Z. Ma, A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of Fe content on the fracture behaviour of Al–Si–Cu cast alloys, Mater. Des. 57 (2014) 366–373. [21] A.M. Nabawy, A.M. Samuel, F.H. Samuel, H.W. Doty, Influence of additions of Zr, Ti– B, Sr, and Si as well as of mold temperature on the hot-tearing susceptibility of an experimental Al–2% Cu–1%Si alloy, J. Mater. Sci. 47 (2012) 4146–4158. [22] H. Yong-dog, Z. Xin-ming, Effect of minor Sc and Zr on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Trans. Nonferrous Metals Soc. China 16 (2006) 1228–1235. [23] O.N. Senkova, M.R. Shagieva, S.V. Senkovaa, D.B. Miracleb, Precipitation of Al3(Sc,Zr) particles in an Al–Zn–Mg–Cu–Sc–Zr alloy during conventional solution heat treatment and its effect on tensile properties, Acta Mater. 56 (2008) 3723–3738. [24] E. Cloueta, A. Barbua, L. Laéb, G. Martinc, Precipitation kinetics of Al3Zr and Al3Sc in aluminum alloys modeled with cluster dynamics, Acta Mater. 53 (2005) 2313–2325. [25] D. Srinivasan, K. Chattopadhyay, Non-equilibrium transformations involving L12– Al3Zr in ternary Al–X–Zr alloys, Metall. Mater. Trans. A 36A (2005) 311–320. [26] B.M. Gable, G.J. Shiflet, E.A. Starke, Alloy development for the enhanced stability of Ω precipitates in Al–Cu–Mg–Ag alloys, Metall. Mater. Trans. A 37A (2006) 1091–1105. [27] B. Forbord, H. Hallem, K. Martbinsen, The effect of alloying elements on precipitation and recrystallisation in Al–Zr alloys, Proceedings of the 9th International Conference on Aluminum Alloys, Brisbane, Australia 2004, pp. 1179–1185. [28] H. Yong-dong, Z. Xin-ming, Effect of minor Sc and Zr on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Trans. Nonferrous Metals Soc. China 16 (2006) 1228–1235. [29] M.L. Kharakterova, D.G. Eskin, Precipitation hardening in ternary alloys of the Al–Sc– Cu and Al–Sc–Si systems, Acta Metall. Mater. 42 (1994) 2285–2290. [30] M.L. Kharakterova, D.G. Eskin, Effect of scandium and zirconium on structure and age hardening of the Al–Mg–Si alloys, Russ. Metall. (Metally) 1 (1997) 104–109. [31] H.G. Paris, T.H. Sanders, Assessment of scandium additions in aluminum alloy design, Proceedings of ICAA-6: Aluminum Alloys, vol. 1, Georgia Institute of Technology, Atlanta 1998, pp. 489–504. [32] M.F. Ibrahim, A.M. Samuel, F.H. Samuel, A preliminary study on optimizing the heat treatment of high strength Al–Cu–Mg–Zn alloys, Mater. Des. 57 (2014) 342–350. [33] O. Elsebaie, A.M. Samuel, F.H. Samuel, H.W. Doty, Impact toughness of Al–Si–Cu– Mg–Fe cast alloys: effects of minor additives and aging conditions, Mater. Des. 60 (2014) 496–509. [34] J.E. Hatch, Aluminum: properties and physical metallurgy, American Society for Metals, Materials Park, OH, 1984 135. [35] A.M. Samuel, J. Gauthier, F.H. Samuel, Microstructural aspects of the dissolution and melting of Al2Cu phase in Al–Si alloys during heat treatment, Metall. Mater. Trans. A 27A (1996) 1785–1798. [36] S. Qzbilen, H.M. Flower, Zirconium-Vacancy Binding and Its Influence on S'-Precipitation in an Al–Cu–Mg Alloy, Acta Metall. 37 (1989) 2993–3000. [37] G.E. Dieter, Mechanical Metallurgy, 3rd ed. McGraw-Hill, Boston, 1986. [38] R. Mahmudi, P. Sepehrband, H.M. Ghasemi, Improved properties of A319 aluminum casting alloy modified with Zr, Mater. Lett. 60 (2006) 2606–2610. [39] P. Sepehrband, R. Mahmudi, F. Khomamizadeh, Effect of Zr addition on the aging behavior of A319 aluminum cast alloy, Scr. Mater. 52 (2005) 253–257. [40] M. Drouzy, S. Jacob, M. Richard, Interpretation of tensile results by means of quality index and probable yield strength, AFS Int. Cast Met. J. 5 (1980) 43–50. [41] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc–0.1Zr (at.%) alloys during isochronal aging, Acta Mater. 58 (2010) 5184–5195. [42] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep resistant aluminum-based alloys — a review, Z. Metallkd. 97 (2006) 246–265. [43] Z. Li, A.M. Samuel, F.H. Samuel, Role of P and Fe on the precipitation of copper intermetallics in 319 alloys, J. Mater. Sci. Lett. 22 (2003) 585–587.
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Role of Zr and Sc addition in controlling the microstructure and tensile properties of aluminum–copper based alloys A.M. Samuel a, S.A. Alkahtani b, H.W. Doty c, F.H. Samuel a,⁎ a b c
Université du Québec à Chicoutimi, Québec, Canada Industrial Engineering Program, Mechanical Engineering Department, College of Engineering, Prince Sattam bin AbdulAziz University, Al Kharj, Saudi Arabia General Motors, Materials Engineering, 823 Joslyn Avenue, Pontiac, MI 48340, USA
a r t i c l e
i n f o
Article history: Received 23 June 2015 Received in revised form 11 September 2015 Accepted 15 September 2015 Available online 18 September 2015 Keywords: Aluminum alloys Zr and Sc addition Tensile properties Heat treatment Grain refining
a b s t r a c t The present work was carried out on eight alloys containing Al–2%Cu–1.5%Si–0.5%Fe–0.6%Mn-0.4%Mg–0.07%Ti, with small amounts of Zr, Sc or both. Tensile test bars were produced using low pressure die casting. All bars were solution heat-treated, followed by artificial aging. The as-cast microconstituents were determined using electron probe microanalyzer equipped with both electron dispersive x-ray and wavelength dispersive spectroscopic systems. The results show that in addition to the currently observed Fe-and Cu-based intermetallics, several Zr- and Sc-containing intermetallics were also formed during solidification (of the present alloys). Neither the Sc/Zr nor the Sc/Al ratio is constant, and both ratios may vary from one particle to another. The presence of Zr or Sc has a significant grain refining effect, equivalent to that obtained from the addition of a Ti-based grain refiner. The main influence of a Zr or Sc addition on the tensile properties of the base alloy appears during aging, and before the onset of softening. Although the addition of Zr or Sc improves the alloy tensile properties, Sc is less effective compared to Zr. Also, a Sc content beyond 0.15% results in no further improvement in the alloy strength. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Zirconium (Zr) is a transition element used in aluminum alloys to enhance their microstructure and mechanical properties. When Zr is added in small quantities, it forms tiny, coherent precipitates of L12ordered trialuminide (Al3Zr) through the decomposition of a supersaturated solid solution of aluminum [1,2]. The Al3Zr phase is distinguished by two different crystallographic structures: (i) the first is the stable tetragonal DO23 structure which is body-centered with eight atoms per unit cell, while (ii) the second is the coherent metastable Al3Zr phase with an L12-type structure which is simple cubic with four atoms per unit cell and having a slight lattice mismatch with α-Al [3,4]. The structure of the Al3Zr intermetallic phase is dependent on the level of Zr added to the alloy, as well as on the cooling rate. Furthermore, Al3Zr can dissolve up to 5% Sc and form a Al3(Zr1 − xScx) phase particle; when the Al3Zr phase dissolves the Sc, the lattice parameter of this phase is observed to decrease [5]. The Al3(Zr1 − xScx) phase may form upon the decomposition of the solid solution or as a primary phase during solidification. The primary phase type of Al3(Zr1 − xScx) has a greater efficiency as a grain refiner than the primary Al3Zr phase particle. Moreover, the primary Al3Zr
⁎ Corresponding author. E-mail address: [email protected] (F.H. Samuel).
http://dx.doi.org/10.1016/j.matdes.2015.09.090 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
phase particles may act as nucleation sites for Al3Sc [6]. As the Al3Sc phase nucleates on the surface of the Al3Zr phase particle, the grain refining capacity of the primary Al3Zr phase particles for α-Al shows improvement because of the small lattice misfit between the Al3Sc phase and α-Al [7,8]. The primary particles of Al3(ScxZr1 − x) have a greater refining efficiency than the primary particles of Al3 Sc or Al3Zr [9]. The Al 3Sc phase can also dissolve up to 5% of the titanium to produce the Al3(Sc1 − xTix); thus in the presence of titanium combined with zirconium Al 3(Sc1 − x − yZrxTiy) may be formed [10,11]. The primary phase particles of Al3(ScxZr1 − x) normally appear in two dimensions in the form of a star-like shape. Scandium reacts with Si and Cu to form two phases: (i) the VAlScSi phase, with a stoichiometric composition of AlSc2Si2 and (ii) the W-AlCuSc phase [12–14]. It is also reported that Sc tends to diminish the solubility of copper and silicon in the aluminum matrix. Seidman et al. [15] and Davydov et al. [16] studied the precipitation of Sc phases in Al–0.5%Sc alloy. The authors reported that Al3Sc precipitates may take place below 350 °C. These Al3 Sc precipitates strengthen the aluminum matrix by pinning the dislocation lines and resisting their motion. In a previous study [17], the authors reported on the effect of various alloying elements on the evolution of microstructure in an Al–2%Cu based alloy during solidification. The present work was undertaken to emphasize the role of Zr and Sc addition in controlling the
A.M. Samuel et al. / Materials and Design 88 (2015) 1134–1144 Table 1 List of the various 220 Al–2%Cu based alloys used in the present study and alloy chemistry in wt.%. Alloy code
Alloy chemistry
220A 220J 220K 220L 220M 220N 220O 220P
220 alloy: Al–2%Cu–1.32%Si–0.42%Mg–0.58%Fe–0.59%Mn–0.07%Ti 220A + 0.15%Ti + 0.30%Zr + 0.15%Sc 220A + 0.15%Zr 220A + 0.30%Zr 220A + 0.50%Zr 220A + 0.15%Sc 220A + 0.3%Sc 220A + 0.5%Sc
microstructure, by applying optical microscopy and electron probe microanalysis to identify and quantify the phases formed by the interactions between the different alloying elements Zr, Sc, Ti, as well as to study the effects of these intermetallic phases on the grain structure. The present investigation also aims at determining the optimum age hardening heat treatment required to maximize the tensile properties. 2. Experimental procedure Alloy 220 was received in the form of 12-kg ingots. The ingots were melted in a 60 kg-capacity crucible, using an electrical resistance
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furnace. The melting temperature was held at 740 ± 5 °C. The molten metal was degassed for 30 min using pure dry argon injected into the molten metal by means of a graphite rotary degassing impeller. The degassing time/speed was kept constant at 30 min/150 rpm. After degassing, the molten metal surface was skimmed carefully to remove all the oxide layers and dross before pouring the melt into the ingot molds. Samplings for chemical analysis were also taken simultaneously for each alloy melt composition prepared. The ingots thus prepared were later remelted in the furnace, and used for preparing melts for casting the LPDC (low pressure die casting) test bars. Tensile test bars were cast via LPDC. The permanent die and the filling system were placed over the furnace containing the molten alloy. The cavity is filled by forcing the molten metal (using pressurized gas at 0.3 to 1.5 bar) to rise into a ceramic tube which connects the die to the furnace. Once the die cavity is filled, the overpressure in the furnace is removed, and the residual molten metal in the tube flows again into the furnace. The various parts of the die are then separated, and the casting extracted. The tensile test bars were subsequently cut from their castings with dimensions corresponding to ASTM B-108 standard (gauge length 50 mm, diameter 12.5 mm, total length 200 mm). Table 1 lists the chemical composition of the used alloy. It should be mentioned here that all element contents are expressed in wt% throughout the article unless otherwise precised.
Fig. 1. Macrostructure of as cast alloys: (a) 220A—432 μm, (b) 220J—52 μm, (c) 220M — 80 μm, and (d) 220P — 70 μm.
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The as-cast test bars were subjected to heat treatment which consisted of solution heat treatment at 495 °C for 8 h (heating rate was about 2 °C/min.), followed by quenching in warm water, and then artificial aging at temperatures of 155, 180, 200, 220, 240 and 300 °C for 5 h at each temperature. The tensile testing was carried out using a servo-hydraulic mechanical testing machine at a strain rate of 1 × 10−4 s− 1. The average ultimate tensile strength (UTS), yield strength (YS) and percent elongation (%El) values were calculated over the five bars tested for each condition. All tests were carried out at room temperature. Grain size measurements were carried out for all of the alloys investigated. For these measurements, after the final grinding stage, the samples were etched using a solution made up of 12.5 g CrO3, 2.5 mL HF,
30 mL HCl, 40 mL HNO3 and 42.5 mL water. Grain size measurements were then carried out using an optical microscope and the line intercept method. Samples were sectioned from as-cast and heat-treated tensile-tested bars for metallographic examination. The samples were mounted in bakelite, ground, and polished to a fine finish, using standard polishing procedures. An electron probe microanalyzer, with electron dispersive x-ray spectrometry (EDS) and wavelength dispersive spectroscopic (WDS) facilitates, was used mainly to detect and analyze the chemical compositions of the various intermetallic phases which formed during solidification. Volume fraction and line-scan measurements were also carried out using the same set-up. X-ray images of various elements constituting the identified phases were also obtained to determine the
Fig. 2. Backscattered electron images of: (a) 220A, (b) 220J, (c) high magnification of (b), (d) 220M, and (e) 220P alloys.
A.M. Samuel et al. / Materials and Design 88 (2015) 1134–1144
distribution of these elements in the phases. In these cases the samples were cleaned in high purity ethanol in an ultrasonic agitator before being examined in the electron probe microanalyzer. In order to understand the role of precipitation, samples from the tensile tested bars were cut 1 cm away from the fracture surface, mechanically polished, followed by ion bombardment prior to examination using field emission scanning electron microscopy. 3. Results and discussion
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Table 3 Total volume fraction of intermetallics observed in the 220 as-cast alloy samples. Alloy
A
J
K
L
M
Average Std Dev.
3.98 0.42
5.97 0.6
4.7 0.42
5.1 0.48
6.42 0.62
Alloy
N
O
P
Average Std Dev.
4.21 0.4
5.21 0.63
4.87 0.59
3.1. Microstructural characterization The grain refining effect of these additives may be seen more clearly in the macrostructures of the etched alloy samples, as shown in Fig. 1. It is interesting to note the increase in the level of refinement on going from Alloys A to P as more grain refiners are added. The fineness of the macrostructure is reflected by the secondary dendrite arm spacing (SDAS), so that the finer the microstructure obtained, the smaller the average SDAS value exhibited by the solidified alloy sample. It is well known that SDAS has a strong influence on the mechanical properties of cast aluminum alloys [18,19]. Fig. 2(a) shows the morphology of the intermetallics observed in the base alloy after casting. The α-Fe (Al15(Mn,Fe,Cu)3Si2) is gray in color, while the Al2Cu-phase may appear either in block-like or eutectic form (Al + Al2Cu). The Q-phase, Al5Mg8 Cu2 Si6, may appear in the form of small particles growing out of the Al 2 Cu phase or alone. Based on the preceding, the Al–2%Cu base alloy possesses simple microstructural components consisting of three main intermetallics, namely Al15(Mn,Fe,Cu)3Si2, Al2Cu and Q-Al5Mg8Cu2Si6 phases, in addition to the α-Al matrix. The absence of β-Al5 SiFe phase in the microstructure may be attributed to the high Mn content in this alloy [20]. Table 2 lists the main intermetallics observed in the present alloys whereas Table 3 shows their volume fraction. The star-like phase, Al3(Sc1 − xZrx), is considered to be the most important of the Sc–Zr intermetallics because of its heightened efficiency in nucleating α-Al grains. This type of efficiency may be attributed to the L1 2 -crystal structure of the star-like phase which is characterized by its low mismatch with the α-Al lattice structure [21]. Fig. 2(b) and (c) show the backscattered image of a star phase particle observed in the 220J alloy. It consists of different layers distinguished
Table 2 Chemical composition of intermetallics observed in the present work. Alloy code
Composition
Suggested formula
220A
Al–72.3 at.%, Fe–9.5 at.%, Si–10.5 at.%, Cu–2.4 at.%, Mn 5.2 at.% Al–18.4 at.%, Mg–40.3 at.%, Si–31 at.%, Cu–9.2 at.% Al–68 at.%, Cu–30.6 at.%, Si–1.1 at.% Al–78.1 at.%, Zr–7.6 at.%, Si–1.5 at.%, Sc–12.1 at.% Al–78.2 at.%, Zr–10.3 at.%, Ti–3.6 at.%, Sc–7.4 at.%
Al15(Fe,Mn,Cu)3Si2
220A 220A 220 J 220 J 220 J 220 M 220 M 220 M 220 M 220 M 220P 220P 220P
Al–75 at.%, Sc–11.1 at.%, Zr–10 at.%, Ti–3.2 at.% Al–68.4 at.%, Zr–31.4 at.% Al–63 at.%, Zr–14.2 at.%, Ti–17 at.%, Si–5.8 at.% Al–75.5 at.%, Zr–20.8 at.%, Ti–2.4 at.% Al–67.5 at.%, Zr–23.1 at.%, Ti–8.7 at.% Al–63.1 at.%, Zr–14 at.%, Ti–16.9 at.%, Si–5.75 at.% Al–77 at.%, Sc–23 at.% Al–38 at.%, Mg–2.7 at.%, Sc–26.8 at.%, Si–32 at.% Al–19.8 at.%, Mg–3.4 at.%, Sc–39.5 at.%, Si–36.7 at.%
Al5Mg8Si6Cu2 Al2Cu Al3(Zr,Sc) (Sc/Zr = 1.6) Al3(Zr,Sc,Ti) (Sc/Zr = 0.72) (Sc/Zr = 1) Al3Zr (Al,Si)3(Zr,Ti) Al3(Zr,Ti) Al2(Zr,Ti) (Al,Si)2(Zr,Ti) Al3Sc AlSiSc (Si/Al = 1) AlSi2Sc2 (Si/Al = 2)
by the change in brightness across the particle caused the variation of the Zr and Sc concentrations within the star phase as shown in the colored maps illustrated in Fig. 3. Based on the WDS analysis of these particles (Table 3) the average composition of the star phase was determined to be Al3(Sc,Zr,Ti) with a Zr:Sc ratio ranging between 0.72 and 1.6, whereas Ti concentration is constant at 3.6 at.%. It has been suggested by Zakharov et al. [22] that as the alloy melt passes the liquidus line, crystals of Al 3Zr are primarily nucleated; these nuclei then grow by absorbing the zirconium from the surrounding liquid. As the Zr is depleted from these regions, the nuclei of the star phase, namely Al3Zr, continue to grow through the precipitation of Al3Sc on the surface of these nuclei. At this point, the Al3Zr starts to precipitate on the surface of Al3Sc as the scandium also begins to become progressively depleted from the surrounding liquid, and so on in sequence. The fluctuations in the precipitation sequences are considered to be the main reason for the layer formation throughout the star phase [14]. Toropova et al. [23] proposed that since the Al 3 Zr phase nucleates as the primary phase, the star phase grows predominantly through the precipitation of Al3Sc on the surface of the Al3Zr in the form of layers displaying different degrees of brightness. The variation in brightness from one layer to another is ascribed to the variation in the amount of Zr atoms replacing Sc atoms in the Al3Sc phase. This suggestion is in agreement with the element distribution shown in Fig. 3 which is also in agreement with what has already been suggested earlier by other researchers [12, 24]. It has been observed that in the case where Sc/Zr is higher than unity, the particles precipitate in two different forms: star-like and multi-faceted shape as shown in Fig. 4. The reason however is not known. Fig. 2(d) exemplifies the precipitation of Zr-rich phases in 220M alloy. It is evident from this figure that the particles are having different shapes, size and distribution. Table 3 lists the chemical composition of some of the Zr-rich phase particles. Srinivasan and Chattopadhyay [25] proposed that as a result of high zirconium content (0.5%), all Zr-containing alloys are located in the L + Al3Zr region of the Al– Zr phase diagram during the melting stage [26]. This statement indicates that the Al3Zr particles are not formed from the melt during solidification but may come directly from the Al–15%Zr master alloy added to the liquid melt. Since these particles do not dissolve in the melt, they provide favored nucleation sites for the growth of intermetallics from the melt during solidification. In addition to zirconium trialuminide, some ternary and quaternary intermetallics have been detected viz. (Al,Si)3 (Zr,Ti), Al3 (Zr,Ti), Al2(Zr,Ti) and (Al,Si)2(Zr,Ti). Based on the homogeneous distribution of Zr, Ti and Si within the phase, it may be suggested that AlZrSi and AlZrTiSi, are formed at the same temperature of the precipitation of Al3Zr phase (≈665 °C [27]). Fig. 2(e) shows the backscattered electron image obtained from 220P alloy revealing the precipitation of Sc-rich phase (white circles) mainly in the form of star-like particles. Table 3 reveals that in addition to scandium trialuminide, other compounds can also precipitate mainly V-AlSiSc phase [28] with two Sc/Al ratios. However, none of these phases contains Ti. Fig. 5 exhibits the presence of Al3Sc phase in the form of a fan-like shape. Based on the Al–Sc binary phase diagram at scandium level of 0.55%Sc, the primary particles of L12–Al3Sc are formed
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Fig. 3. (a) Backscattered image (CP) showing an Al3(Sc,Zr,Ti) phase particle in the 220 J alloy in the as cast condition and corresponding X-ray images of (b) Ti, (c) Zr and (d) Sc. Arrows indicate high Zr concentration within the star-like particles. The Sc/Zr ratio is ~unity.
Fig. 4. (a) Backscattered image showing an Al3(Sc,Zr,Ti) phase particle in the 220 J-alloy in the as cast condition and corresponding X-ray images of (b) Ti, (c) Zr, and (d) Sc. The Sc/Zr ratio is ~1.6.
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Fig. 5. (a) Backscattered image showing an Al3Sc phase particle in the 220P alloy in the as cast condition, and corresponding X-ray images of (b) Si, (c) Al, (d) Sc and (e) corresponding EDS spectrum.
in a eutectic reaction L + Al3Sc → (Al) (refined) at 665 °C [29], and act as nucleation sites for α-Al in the molten metal during solidification. The V-phase is a ternary intermetallic compound (AlScSi) formed through invariant reactions, and has a gray color and a stoichiometric composition of AlSc2Si2 [30,31]. Fig. 6 is an example of the distribution of Si, Al, and Sc in the V-phase observed in the 220P alloy — white arrow points to the presence of Al2Cu phase particle adjacent to the V-phase. It should
be mentioned here that W-AlSiCu phase has not been detected in the present investigation. 3.2. Tensile properties In this section the alloys have been divided into three groups: (i) 220A and 220J, (ii) 220K, 220L and 220M and (iii) 220N, 220O and
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Fig. 6. (a) Backscattered image showing an AlSi2Sc2-V phase particle in the 220P alloy in the as cast condition, and corresponding X-ray images of (b) Si, (c) Al and (d) Sc,
220P. The variation in tensile properties of these alloys in the as-cast and solution heat-treated (SHT) conditions is shown in Fig. 7. Although the addition of 0.5%Sc caused a significant reduction in the alloy original grain size (432 μm vs. 52 μm), the corresponding tensile values are noticeably lower than those obtained from 220 M alloy containing 0.5%Zr. Increasing the amount of added Sc from 0.15% to 0.5% also has no corresponding effect on the alloy mechanical performance. As expected solution heat treatment resulted in partial dissolution of the alloying elements and hence relatively better UTS levels compared to those obtained in the as-cast condition. The alloy UTS is mainly controlled by the microstructure and soundness of the casting. As shown in the microstructure section, the microstructure of the investigated alloys in the as-cast condition contains a large number of solidified intermetallic particles or phases (Table 3). These hard particles have a harmful effect on the alloy UTS and %El. Also their sharp edges act as crack initiators. During tensile testing, cracking commences by the breaking up of these intermetallics or by decohesion of the intermetallic particles from the surrounding matrix [32,33]. When the alloy is solution heat treated at sufficiently high temperature, as the solution heat treatment is applied, these intermetallics will break down and dissolve by losing their elements which then return to the matrix producing a supersaturated solid solution of αAl; this solid solution strengthens the matrix by increasing the resistance to dislocation movement during the deformation of tensile testing [34,35].
The effect of artificial aging in the temperature range of 155 °C–300 °C on the tensile properties of 220A and 220J alloys is shown in Fig. 8. As expected, increasing the aging temperature from 155 °C to 300 °C will cause significant reduction in the alloy UTS, as shown in Fig. 8(a), by nearly 30%. Maximum strength was achieved when the alloys were aged at 180 °C, representing peak-aging condition. From Figs. 7 and 8, it is evident that aging the 220A and 220J alloys at 180 °C for 5 h increased their UTS values by about 60% with respect to the as-cast condition and 35% compared to SHT values. Similar observations are made for the YS levels. The addition of Zr–Ti produces a refined non-dendritic structure that decreases the probability for porosity formation and increases the amount of dissolved Al2Cu after solution heat treatment, and thereby the level of Cu in solid solution of α-Al, which then allows for a larger amount of precipitation hardening to take place upon aging, and therefore improve the alloy strength. Additionally, dispersoid precipitates of Al3Zr and/or Al3(Zr1 − xTix) may act as nucleation sites for the hardening phases during the aging process, resulting in further improvement in strength [36–38] Although the equilibrium solid solubility of zirconium in aluminum is very low, the high binding energy results in substantial vacancy trapping in solution treated and quenched alloys which results in modification of precipitation of S′-phase in dilute Al– Cu–Mg alloys. In order to enhance the mechanical properties of an aluminum alloy, it is important to obtain a thermally stable microstructure and coarsening-resistant dispersoids. This may be achieved by adding
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Fig. 7. Tensile properties of the present alloys in the as cast and solution heat treated conditions: (a) UTS and YS, (b) %El.
Zr, which has the smallest diffusion flux in aluminum of all the transition metals; the presence of Zr leads to the formation of fine dispersoids that resist coarsening at higher temperatures, which helps to improve/ maintain the mechanical properties [38,39]. Fig. 9 presents the tensile values of Zr-containing alloys. As can be seen, increasing the amount of added Zr resulted in increasing the alloy UTS values of all alloys. For example, the as cast value of the 220 M alloy increased from 280 MPa the as cast condition to 375 MPa after aging at 180 °C, representing an increase of about 35–40%, caused by simultaneous precipitation of Al2Cu and Al3Zr. In consistence with the data plotted in Fig. 8(a), the presence of Zr resulted in the extension of the aging temperature before the onset of softening, as shown in Fig. 9(a). Fig. 8(b) displays the variation in YS levels as a function of aging temperature. It can be seen that the YS values of the 220 M alloy are very close to those obtained for the corresponding UTS in the temperature range 155 °C–200 °C, followed thereafter by a fast drop to reach 100 MPa at 300 °C (almost 50% reduction compared to as-cast condition).The advantage of adding Zr in relatively small amounts, e.g. 0.15% (220K alloy), appears in the increase in the alloy ductility from 4.2% in the as-cast condition to 5.7% after aging at 300 °C. Applying the quality index equation proposed by Drouzy et al. [40] for Al–Si–Mg alloys, Q ðMPaÞ ¼ UTS ðMPaÞ þ 150 logð%ElÞ the quality index value of 220 L alloy is seen to vary from 313 MPa in the as cast condition to 395 MPa after aging at 180 °C (peak-aging) to 345 MPa after aging at 300 °C.
Fig. 8. Tensile properties of 220A and 220J alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
The work of Knipling et al. [41,42] on the precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc–0.1Zr (at.%) alloys during isochronal aging revealed that precipitation of Al3Sc(L12) commences between 200 °C and 250 °C in the Al–0.1Sc alloy, reaching a peak microhardness of 668 ± 20 MPa at 325 °C. In the Al–0.1Zr alloy, precipitation of Al3Zr (L12) commences between 350 °C and 375 °C, achieving a peak microhardness of 420 MPa at 425–450 °C. While the Al–0.1Sc–0.1Zr alloy achieves peak microhardness after isochronal aging to 400 °C, it overages after extended annealing at this temperature indicating that it is not suitable for extended use at 400 °C. A pronounced synergistic effect
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Fig. 10. Tensile properties of 220N, 220O, 220P alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
Fig. 9. Tensile properties of 220K, 220L, 220M alloys in the aged conditions: (a) UTS, (b)YS, and (c) %El.
is observed when both Sc and Zr are present. Above 325 °C, Zr additions provide a secondary strength increase that is attributed to precipitation of Zr-enriched outer shells onto the Al3Sc precipitates, leading to a peak microhardness of 618 MPa at 400 °C for Al–0.06Sc–0.06Zr.
The variation in the tensile properties on the Sc-containing alloys as a function of aging temperature is presented in Fig. 10. Apparently, increasing the Sc content from 0.15 to 0.5% resulted in close UTS and YS values regardless of the aging temperature. It is evident from Fig. 10 that Sc addition is less effective in improving the alloy UTS, YS and %El values compared to that exhibited by Zr addition when the alloys are aged in the temperature range 155 °C–240 °C. Aging beyond 240 °C resulted in sharp decrease in the alloy strength regardless of the Zr or Sc content.
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Fig. 11. Precipitation of Al2Cu phase in 220A alloy as a function of heat treatment condition: (a) as cast, (b) SHT, (c) aged-180 °C, (d) aged 240 °C, (e) EDs spectrum corresponding to the arrowed particle in (d), (f) aged 300 °C.
3.3. Precipitation Fig. 11 presents the precipitation taking place in the base 220A alloy at different working conditions. Fig. 11(a) shows the presence of fine bright particles in the matrix of the alloy after casting, together with massive Al2Cu phase particles (white arrow). The variation in the precipitate size may be due to the cooling temperature range after solidification. After solutionizing at 495 °C for 8 h, the bulky particles dissolved almost completely in the aluminum matrix (Fig. 11(b)). However, some particles (less than 0.25 μm long) can easily be seen in the matrix. These particles have been identified as Si precipitates. Dense ultrafine spherical particles throughout
the entire matrix were the main feature obtained when the alloy was aged at 180 °C for 5 h (Fig. 11(c)) causing the observed increase in the alloy strength. Increasing the aging temperature to 240 °C resulted in a partial change in the morphology of the precipitated particles from spherical to elongated or rod-like shape as illustrated in Fig. 11(d). The white arrow in Fig. 11(d) points to undissolved Curich phase. The associated EDS spectrum (Fig. 11(e)) reveals that the undissolved particles could be Al 7Cu2 Fe phase [43]. Aging at 300 °C resulted in complete transformation of the spherical particles into rod shaped ones (Fig. 11(f)) indicating the precipitation of incoherent Al2 Cu phase and hence the observed decline in the alloy strength.
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4. Conclusions Based on the results obtained from the present study, the following conclusions may be drawn: 1. Addition of Zr or Sc in small amounts (about 0.15%) is sufficient to cause about 85% reduction in the alloy grain size with no hot tearing. 2. In addition to Fe- and Cu-based intermetallics, addition of Zr leads to precipitation of Al3Zr, (Al,Si)3(Zr,Ti), Al3(Zr,Ti), Al2(Zr,Ti) and (Al,Si)2(Zr,Ti) phases whereas the addition of Sc leads to the formation of Al3Sc, AlSiSc and AlSi2Sc2 phases during the course of solidification. 3. Combined addition of Zr and Sc causes the precipitation of Al3(Zr,Sc), Al3(Zr,Sc,Ti) and Al3(Sc,Zr,Ti) phases. In all cases neither the Sc/Zr nor Sc/Al ratio is constant. 4. Although the addition of Zr or Sc leads to significant improvement in the alloy strength, the effectiveness of Zr is greater than that obtained from Sc under the same condition (amount and heat treatment). 5. Addition of Sc beyond 0.15% has no further influence on the grain size or on the alloy strength. 6. The importance of additions of Zr or Sc appears not only in improving the alloy strength but also in extending the aging temperature before the onset of softening. 7. During aging up to 300 °C, copper aluminide is the main precipitating phase during aging, which influences the alloy strength. 8. Hardening due to precipitation of Al3Zr or Al3Sc phases has not been detected, indicating that most of the observed Zr- and Sc-based phases occurred only during the solidification process. Acknowledgments The authors would like to thank Ms. Amal Samuel for enhancing the images used in the present work. References [1] N.A. Belov, A.N. Alabin, D.G. Eskin, V.V. Istomin-Kastrovskii, Optimization of hardening of Al–Zr–Sc cast alloys, J. Mater. Sci. 41 (2006) 5890–5899. [2] D. Srinivasan, K. Chattopadhyay, Non-equilibrium transformations involving L12– Al3Zr in ternary Al–X–Zr alloys, Metall. Mater. Trans. A 36A (2005) 311–320. [3] N.A. Belov, D.G. Eskin, A.A. Aksenov, Multicomponent phase diagrams: application for commercial aluminum alloys, 2005 413. [4] L.S. Toropova, D.G. Eskin, M.L. Kharaktrova, T.V. Dobakina, Advanced aluminum alloys containing scandium: structure and properties, Gordon and Breach Science Publishers, Canada, 1998 175. [5] J. Royset, N. Ryum, Scandium in aluminum alloys, Int. Mater. Rev. 50 (2005) 19–44. [6] A.A. Rao, B.S. Murty, M. Chakraborty, Role of zirconium and impurities in grain refinement of aluminum with Al–Ti–B, Mater. Sci. Technol. 13 (1997) 769–777. [7] J.D. Robson, P.B. Prangnell, Dispersoid precipitation and process modelling in zirconium containing commercial aluminum alloys, Acta Mater. 49 (2001) 599–613. [8] J.D. Robson, Optimizing the homogenization of zirconium containing commercial aluminum alloys using a novel process model, Mater. Sci. Eng. A 338A (2002) 219–229. [9] M.S. Kaiser, M.K. Banerjee, Effect of ternary scandium and quaternary zirconium and titanium additions on the tensile and precipitation properties of binary cast Al–6 Mg alloys, Jordan J. Mech. Ind. Eng. 2 (2008) 93–99. [10] O.N. Senkov, M.R. Shagiev, S.V. Senkova, D.B. Miracle, Precipitation of Al3(Sc,Zr) particles in an Al–Zn–Mg–Cu–Sc–Zr alloy during conventional heat treatment and its effect on tensile properties, Acta Mater. 56 (2008) 3723–3738. [11] C.B. Fuller, J.L. Murray, D.N. Seidman, Temporal evolution of the nanostructure of Al(Sc,Zr) alloys: part I-chemical compositions of Al3(Sc1 − xZrx) precipitates, Acta Mater. 53 (2005) 5401–5413. [12] E.A. Marquis, D.N. Seidman, Nanoscale structural evolution of Al3Sc precipitates in Al(Sc) alloys, Acta Mater. 49 (2001) 1909–1919. [13] V.V. Zakharov, Effect of scandium on the structure and properties of aluminum alloys, Met. Sci. Heat Treat. 45 (2003) 246–253.
[14] K.B. Hyde, A.F. Norman, P.B. Prangnell, The effect of cooling rate on the morphology of primary Al3Sc intermetallic particles in Al–Sc alloy, Acta Mater. 49 (2001) 1327–1337. [15] D.N. Seidman, E.A. Marquis, D.C. Dunand, Precipitation strengthening at ambient and elevated temperatures of heat-treatable Al(Sc) alloys, Acta Mater. 50 (2002) 4021–4035. [16] V.G. Davydov, T.D. Rostova, V.I. Yelagin, Scientific principles of making an alloying addition of scandium to aluminum alloys, Mater. Sci. Eng. A 280A (2000) 30–36. [17] G.A. Zaki, A.M. Samuel, H.W. Doty, F.H. Samuel, Microstructural evolution during solidification of Al–Cu based alloys, Int. J. Mater. Sci. (2015) (in press - posted online on 10 August 2015). [18] E. Samuel, B. Golbahar, A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of grain refiner on the tensile and impact properties of Al–Si–Mg cast alloys, Mater. Des. 56 (2014) 468–479. [19] A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of grain refining and Srmodification interactions on the impact toughness of Al–Si–Mg cast alloys, Mater. Des. 56 (2014) 264–273. [20] Z. Ma, A.M. Samuel, H.W. Doty, S. Valtierra, F.H. Samuel, Effect of Fe content on the fracture behaviour of Al–Si–Cu cast alloys, Mater. Des. 57 (2014) 366–373. [21] A.M. Nabawy, A.M. Samuel, F.H. Samuel, H.W. Doty, Influence of additions of Zr, Ti– B, Sr, and Si as well as of mold temperature on the hot-tearing susceptibility of an experimental Al–2% Cu–1%Si alloy, J. Mater. Sci. 47 (2012) 4146–4158. [22] H. Yong-dog, Z. Xin-ming, Effect of minor Sc and Zr on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Trans. Nonferrous Metals Soc. China 16 (2006) 1228–1235. [23] O.N. Senkova, M.R. Shagieva, S.V. Senkovaa, D.B. Miracleb, Precipitation of Al3(Sc,Zr) particles in an Al–Zn–Mg–Cu–Sc–Zr alloy during conventional solution heat treatment and its effect on tensile properties, Acta Mater. 56 (2008) 3723–3738. [24] E. Cloueta, A. Barbua, L. Laéb, G. Martinc, Precipitation kinetics of Al3Zr and Al3Sc in aluminum alloys modeled with cluster dynamics, Acta Mater. 53 (2005) 2313–2325. [25] D. Srinivasan, K. Chattopadhyay, Non-equilibrium transformations involving L12– Al3Zr in ternary Al–X–Zr alloys, Metall. Mater. Trans. A 36A (2005) 311–320. [26] B.M. Gable, G.J. Shiflet, E.A. Starke, Alloy development for the enhanced stability of Ω precipitates in Al–Cu–Mg–Ag alloys, Metall. Mater. Trans. A 37A (2006) 1091–1105. [27] B. Forbord, H. Hallem, K. Martbinsen, The effect of alloying elements on precipitation and recrystallisation in Al–Zr alloys, Proceedings of the 9th International Conference on Aluminum Alloys, Brisbane, Australia 2004, pp. 1179–1185. [28] H. Yong-dong, Z. Xin-ming, Effect of minor Sc and Zr on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Trans. Nonferrous Metals Soc. China 16 (2006) 1228–1235. [29] M.L. Kharakterova, D.G. Eskin, Precipitation hardening in ternary alloys of the Al–Sc– Cu and Al–Sc–Si systems, Acta Metall. Mater. 42 (1994) 2285–2290. [30] M.L. Kharakterova, D.G. Eskin, Effect of scandium and zirconium on structure and age hardening of the Al–Mg–Si alloys, Russ. Metall. (Metally) 1 (1997) 104–109. [31] H.G. Paris, T.H. Sanders, Assessment of scandium additions in aluminum alloy design, Proceedings of ICAA-6: Aluminum Alloys, vol. 1, Georgia Institute of Technology, Atlanta 1998, pp. 489–504. [32] M.F. Ibrahim, A.M. Samuel, F.H. Samuel, A preliminary study on optimizing the heat treatment of high strength Al–Cu–Mg–Zn alloys, Mater. Des. 57 (2014) 342–350. [33] O. Elsebaie, A.M. Samuel, F.H. Samuel, H.W. Doty, Impact toughness of Al–Si–Cu– Mg–Fe cast alloys: effects of minor additives and aging conditions, Mater. Des. 60 (2014) 496–509. [34] J.E. Hatch, Aluminum: properties and physical metallurgy, American Society for Metals, Materials Park, OH, 1984 135. [35] A.M. Samuel, J. Gauthier, F.H. Samuel, Microstructural aspects of the dissolution and melting of Al2Cu phase in Al–Si alloys during heat treatment, Metall. Mater. Trans. A 27A (1996) 1785–1798. [36] S. Qzbilen, H.M. Flower, Zirconium-Vacancy Binding and Its Influence on S'-Precipitation in an Al–Cu–Mg Alloy, Acta Metall. 37 (1989) 2993–3000. [37] G.E. Dieter, Mechanical Metallurgy, 3rd ed. McGraw-Hill, Boston, 1986. [38] R. Mahmudi, P. Sepehrband, H.M. Ghasemi, Improved properties of A319 aluminum casting alloy modified with Zr, Mater. Lett. 60 (2006) 2606–2610. [39] P. Sepehrband, R. Mahmudi, F. Khomamizadeh, Effect of Zr addition on the aging behavior of A319 aluminum cast alloy, Scr. Mater. 52 (2005) 253–257. [40] M. Drouzy, S. Jacob, M. Richard, Interpretation of tensile results by means of quality index and probable yield strength, AFS Int. Cast Met. J. 5 (1980) 43–50. [41] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc–0.1Zr (at.%) alloys during isochronal aging, Acta Mater. 58 (2010) 5184–5195. [42] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep resistant aluminum-based alloys — a review, Z. Metallkd. 97 (2006) 246–265. [43] Z. Li, A.M. Samuel, F.H. Samuel, Role of P and Fe on the precipitation of copper intermetallics in 319 alloys, J. Mater. Sci. Lett. 22 (2003) 585–587.