optimization of solution treatment of cast al-7si-0.3mg and al-8si-3cu-0.5mg alloys

Optimization of Solution Treatment of Cast Al-7Si-0.3Mgand Al-8Si-3Cu-0.5Mg Alloys

EMMA SJOLANDER and SALEM SEIFEDDINE

The influence of solidification rate on the solution-treatment response has been investigated foran Al-7Si-0.3Mg alloy and an Al-8Si-3Cu-0.5Mg alloy. The concentrations of Mg, Cu, and Si inthe matrix after different solution-treatment times were measured using a wavelength dispersivespectrometer. All Mg dissolves into the matrix for the Al-Si-Mg alloy when solution treated at803 K (530 �C) because the p-Fe phase is unstable and transforms into short b-Fe plates whichrelease Mg. The Q-Al5Mg8Cu2Si6 phase do not dissolve completely at 768 K (495 �C) in the Al-Si-Cu-Mg alloy and the concentration in the matrix reached 0.22 to 0.25 wt pct Mg. Thedistance between p-Fe phases and Al2Cu phases was found to determine the solution-treatmenttime needed for dissolution and homogenization for the Al-Si-Mg alloy and Al-Si-Cu-Mg alloy,respectively. From the distance between the phases, a dimensionless diffusion time was calcu-lated which can be used to estimate the solution-treatment times needed for different coarse-nesses of the microstructure. A model was developed to describe the dissolution andhomogenization processes.

DOI: 10.1007/s11661-013-2141-9� The Minerals, Metals & Materials Society and ASM International 2013

I. INTRODUCTION

A solution treatment at a temperature close to theeutectic temperature is the first step of a T6 heattreatment. The purposes of the solution treatment ARE(a) to dissolve particles formed during solidification,containing Mg and Cu; (b) to homogenize alloyingelements in the matrix; and (c) to spheroidize theeutectic silicon particles. The solution-treatment processneeds to be optimized because too short a solutiontreatment means that not all the alloying elements addedwill be dissolved and made available for precipitationhardening, whereas too long a solution treatment meansconsumption of more energy than is necessary. Otherimportant factors include the influences of residual Mg-and Cu-containing particles on the elongation to frac-ture, as well as the spheroidization and coarsening of Siparticles. A successful solution treatment depends on theas-cast microstructure in combination with the solution-treatment parameters chosen. The alloys of interest inthis article are cast Al-Si alloys having Si concentrationsranging from 6 to 9 wt pct, Cu concentrations of up to4 wt pct, and Mg concentrations between 0.3 and 0.7 wtpct.

In Al-Si-Mg alloys, phases containing Mg which mayform during solidification, are the Mg2Si, and the p-Fephase (Al8Mg3FeSi6).

[1] The p-Fe phase has a Chinesescript or blocky morphology and is often formed on the

b-Fe plates (Al5FeSi).[1] The p-Fe phase was reported by

Taylor et al.[1] to be the only Mg-containing phaseformed in an Al-7Si-0.3Mg-0.12Fe alloy having asecondary dendrite arm spacing (SDAS) 40 lm. Whenthe Mg concentration of the alloy was increased, theMg2Si phase started to form and reached 0.2 vol pct for0.7 wt pct Mg, while the fractions of the b-Fe phase andp-Fe phase were unaffected.In Al-Si-Cu-Mg alloys, the Al2Cu phase and the Q

phase (Al5Mg8Cu2Si6) may form in addition to the b-Fe,p-Fe, and Mg2Si phases.

[2] The Al2Cu phase forms asblocky Al2Cu or as eutectic (Al-Al2Cu) particles oftenon the b-Fe plates or as small blocky particles on theeutectic Si particles.[3] The Q phase often forms on theAl2Cu phase during the final stages of solidification.[4]

For high Mg concentrations, the scale of the Q phaseincreases, and separate particles start to form close tothe Al2Cu phase.[4,5]

Phases formed during solidification have differentpropensities to dissolve or transform during solutiontreatment depending on the alloy concentrations and thetemperature that can be used. Al-Si-Mg alloys can besolution treated at high temperatures ranging from813 K to 823 K (540 �C to 550 �C).[6] The dissolution ofthe Mg2Si phase is a fast process due to the hightemperature and the high diffusion rate of Mg in Al.Rometsch et al.,[7] for example, report that a shortsolution treatment of 8 through 15 minutes at 813 K(540 �C) is enough to dissolve and homogenize an A356alloy with SDAS 40 lm, while solution treatment of50 minutes is needed for an A357 alloy because of itscoarser microstructure (SDAS 55 lm) and higher Mgconcentration. The p-Fe phase transforms into b-Fephase and Mg in solid solution during solution treat-ment at 813 K (540 �C) if the Mg concentration is low

EMMA SJOLANDER, Postdoctoral, and SALEM SEIFEDDINE,Assistant Professor, are with the Materials and Manufacturing -Casting, Department of Mechanical Engineering, School of Engineer-ing, Jonkoping University, Box 1026 551 11, Jonkoping, Sweden.Contact e-mail: [email protected]

Manuscript submitted March 1, 2011.

METALLURGICAL AND MATERIALS TRANSACTIONS A

(0.3 to 0.4 wt pct). On the other hand, if the Mgconcentration is high (0.6 to 0.7 wt pct), there will be notransformation, and the process may even bereversed.[1,8–10]

Al-Si-Cu-Mg alloys must be solution treated at alower temperature to avoid local melting of Cu-richphases. Melting of the Q phase is reported to take placeat 778 K (505 �C) for an A319 alloy with 0.5 wt pctMg.[11] The exact temperature that can be used withoutmelting depends on the alloy concentration, the solid-ification rate and the heating rate to the solution-treatment temperature, as shown by for example Samuelet al.[12] for a 380 alloy. The Q phase can be either stable,grow or dissolve during solution treatment dependingon the alloy composition and solution-treatment tem-perature. The Q phase is reported to be stable or todissolve very slowly for alloys having high Cu concen-tration (3.5 to 4.4 wt pct) and various Mg concentra-tions when solution treated at 773 K (500 �C).[13,14] Fora lower Cu concentration of 2.8 wt pct the Q phase wasshown to be stable at 753 K (480 �C), while it dissolvedat 778 K (505 �C).[15] Lower solution-treatment temper-atures and lower diffusivity of Cu in Al make longsolution-treatment times necessary for Al-Si-Cu-Mgalloys to achieve a high and homogenous concentrationof alloying elements in solid solution. Solution-treat-ment times of 8 through 10 hours at temperatures from763 K to 778 K (490 �C to 505 �C) are recommended inthe literature.[14,16–18]

Spheroidization of Si particles during the solution-treatment process has been carefully examined duringthe recent years[19] and will not be treated in this article.The focus of this investigation is analyzing the dissolu-tion of Cu- and Mg-rich particles and homogenizationof Cu, Mg, and Si in the matrix. Two alloys, an Al-7Si0.3Mg alloy and an Al-8Si-3Cu-0.5Mg alloy, werechosen as they are typical alloys used in the automotiveindustry. The aims of the investigation are (a) todetermine how the solidification rate influences thescale, type, fraction, morphology, and distribution ofthe phases formed during solidification as well as theconcentration gradient in the matrix; (b) to study thedissolution and transformation of phases during solu-tion treatment; and (c) to measure the alloying elementconcentration’s increase in the matrix during solutiontreatment. From this knowledge a model is developedthat can be used as an aid to optimize the solution-treatment process for a particular alloy and the as-castmicrostructure. The research and the models developedfor solution treatment are activities aimed towardincreasing our knowledge of the heat-treatment responseand for the development of a model that takes all stepsin a T6 heat treatment into account for Al-Si-Cu-Mgcasting alloys.

II. MATERIALS AND EXPERIMENTS

Two Al-Si alloys were cast having different Mg andCu concentrations, see Table I. The alloys were Srmodified using an Al-10Sr master alloy and grain-refined using an Al-5Ti-1B master alloy. Cylindricalrods (length 18 cm and diameter 1 cm) were cast in apreheated permanent mold. The rods were remelted withthe gradient solidification technique which gives sampleswith a low content of porosity defects because of thegood feeding. Further information about the gradientsolidification technique can be found elsewhere.[20]

Different solidification rates, obtained by means ofdifferent withdrawal speeds of the samples from thefurnace, were used to achieve different coarsenesses ofthe microstructure. Rods with SDAS of ~10, 25, and50 lm were produced to simulate high-pressure diecasting, gravity die casting, and sand casting. Solutiontreatment was conducted in an electrical furnace at atemperature of 768 K (495 �C) for the Al-Si-Cu-Mgalloy and at 803 K (530 �C) for the Al-Si-Mg alloy.Times from 10 minutes up to 10 hours were used. Thesamples were quenched in 323 K (50 �C) water. Thetime for heating the samples to the solution-treatmenttemperature (10 through 15 minutes) is excluded fromthe recorded time.The microstructures were studied using a scanning

electron microscope equipped with energy dispersivespectrometer (EDS) and wavelength dispersive spec-trometer (WDS). The area fractions of Mg2Si, Al2Cu, p-Fe, b-Fe, and Q particles were measured in the as-castconditions, and the evolution of the particles withsolution-treatment time was observed. Cu, Mg, and Siconcentrations were measured across dendrite arms forsamples solution treated for various times, using WDS.The acceleration voltage was set to 20 kV for Cu and10 kV for Mg and Si measurements, and pure elementswere used as standards. Three points were measuredover a single dendrite arm, and at least three dendriteswere measured for each sample. Dendrites that were farfrom Mg- and Cu-containing phases were chosen. Fe-containing phases were studied using EDS.A model was developed which was tested on concen-

tration measurements from the current article as well ason data from the literature for the Al-Si-Mg alloy. Thedistances between phases were measured on componentsand die-cast material to compare with the gradientsolidified material used in the current article. Thermo-Calc with the TTAL database version 3[21] was used toobtain the equilibrium solubility of Si, Cu, and Mg inthe a-Al phase at the solution-treatment temperatures.The measured Mg concentration for the Al-Si-Cu-Mgalloy was, however, lower than the equilibrium concen-tration, and measurements were done after a solution

Table I. Alloy Composition in Weight Percent

Si Cu Mg Fe Ti Al Sr (ppm)

Al-Si-Mg 7.1 0.0 0.32 0.11 0.13 bal. 351Al-Si-Cu-Mg 8.5 3.1 0.47 0.17 0.23 bal. 350


treatment for 40 hours at 748 K, 758 K, 768 K, and778 K (475 �C, 485 �C, 495 �C, and 505 �C) for thecoarsest microstructure to study this deviation further.

III. RESULTS

The SDAS, area fractions, and the distance, betweenMg- and Cu-containing particles for the different coarse-nesses of the microstructure are presented in Table II forthe Al-Si-Mg alloy and in Table III for the Al-Si-Cu-Mgalloy. The areas investigated were normalized to thecoarseness of themicrostructure andwere chosen to coverthe same numbers of dendrites on each frame.

A. Al-Si-Mg (As-cast)

Thep-Fephase is themainphase containingMg, and theMg2Si phase was only observed in the coarsestmicrostruc-ture, see Table II. The area fraction of the p-Fe phaseincreases,while thatof theb-Fephasedecreaseswithafinermicrostructure. Similar results, without Mg2Si phases andthe p-Fe phase (0.9 vol pct) dominant compared with theb-Fe phase (0.1 vol pct), have been reported by Tayloret al.[1] for an Al-7Si-0.3Mg-0.12Fe alloy having SDAS40 lm. The p-Fe phases are small and uniformly distrib-uted in the finest microstructure, while they are present aslarge scripts in the two coarser microstructures, seeFigure 1(a). In the coarsest microstructure, the p-Fe phaseis formed on the b-Fe phase, see Figure 1(c).

The as-cast concentration profiles of Mg and Si forthe three coarsenesses of the microstructure are pre-sented in Figures 2(a) and (b). The Mg concentration inthe coarsest microstructure is higher than the finer ones,

which is probably a result of back diffusion due to thelonger solidification and cooling time. Similar concen-trations of Mg in the center of dendrites in the as-castcondition have been reported in the literature. Pedersenand Arnberg[22] report 0.05 and 0.11 wt pct Mg for anAl-7Si-0.24Mg-0.11Fe alloy having SDAS of 13 lm and52 lm, respectively, and Rometsch et al.[9] report0.13 wt pct Mg for an Al-7Si-0.4Mg-0.13Fe alloy havingSDAS 40 lm.The as-cast Si concentration is low for the coarsest

microstructure, with the lowest concentration at thedendrite edge, while the finest microstructure has ahigher Si concentration with the highest value at thedendrite edge, see Figure 2(b). The anomalous segrega-tion profiles for a slow solidification rate was shown byDons et al.[23] to be due to diffusion of Si from the a-Almatrix on to the eutectic Si particles during cooling aftersolidification.

B. Al-Si-Cu-Mg (As-cast)

The microstructural features of the as-cast Al-Si-Cu-Mg alloy are given in Table III. The Al2Cu phase ispresent both as blocky Al2Cu, as eutectic (Al-Al2Cu),and as small particles in the Si eutectic, in accordancewith the observations by Samuel et al.[3] The Q phase ispresent for all three solidification rates. In the coarsestmicrostructure, it is present either between Al2Cuparticles or as large independent particles, see Fig-ure 3(c). This is also true for the intermediate micro-structure, but the fraction of large independent particlesis lower. The observation of two types of Q phaseshaving different morphologies is in agreement withresults for A319 alloys with Mg> 0.4 wt pct cast in

Table II. Microstructural Features for the As-cast Al-Si-Mg Alloy

Withdrawal Area Pct of Phases** Distance Between Particles�

Speed (mm/s) SDAS* (lm) Mg2Si (Pct) b-Fe (Pct) p-Fe (Pct) Mg2Si (lm) p-Fe (lm)

0.03 51 (7) 0.07 (0.14) 0.26 (0.19) 0.3 (0.5) 380 (310) 240 (150)0.3 28 (3) 0 (–) 0.05 (0.07) 0.6 (0.5) — 50 (20)3 10 (1) 0 (–) 0 (–) 0.7 (0.3) — 9 (4)

Average values and standard deviations within brackets.*Secondary dendrite arm spacing. Average of 40 measurements.**The area investigated depended on the SDAS. SDAS 51:1.0 mm2, SDAS 28:0.25 mm2 and SDAS 10:0.04 mm2.�Average of 50 measurements.

Table III. Microstructural Features for the As-cast Al-Si-Cu-Mg Alloy

Solidification Area Pct of Phases** Distance Between Particles��

Rate (mm/s) SDAS (lm)* Al2Cu (Pct) Fe Rich (Pct) Q (Pct) Al2Cu (lm) Q (lm)

0.03 49 (7) 1.5 (0.7) 0.4 (0.2) 0.4 (0.2) 200 (40) 160 (70)0.3 24 (3) 1.2 (0.4) 0.2 (0.1) 0.3 (0.2) 62 (15)3 9 (1) 2.1 (0.3) � � 14 (3)

Average values and standard deviations within brackets.*Secondary dendrite arm spacing. Average of 40 measurements.**The area investigated depended on the SDAS. SDAS 49:3.1 mm2, SDAS 24:0.77 mm2 and SDAS 9:0.12 mm2.�The particles were too small and close to each other in the SDAS 9 lm sample to be able to distinguish the different phases and therefore no

quantitative measurement could be made.��Average of 50 measurements.


permanent molds reported by Yang,[5] who observedthat the Q phase forms in two reactions, one ahead ofthe Al2Cu phase having a script-like morphology andanother at lower temperature in a complex eutectic withSi particles and Al2Cu particles.

The b-Fe plates is the only Fe-containing phasepresent in the coarsest microstructure. For the interme-diate microstructure, b-Fe plates are present as well asan Al-Fe-Si-Cu phase having a more compact morphol-ogy. The fraction of the Mg2Si phase is very low; noparticles were found for the coarsest and the finestmicrostructures, while a few were found for the inter-mediate microstructure. These observations are inagreement with Samuel et al.[4] who observed smallMg2Si phases were present on Si particles and b-Feplates in nonmodified Al-6Si-3.8Cu-(0.3-0.6)Mg alloys,while no Mg2Si phases were observed when about220 ppm Sr was added. The phases present in the finestmicrostructure are very difficult to distinguish as theyare very small and closely spaced, see Figure 3(a). TheAl2Cu phase and Q phase were identified as well as a Fe

phase, probably containing Cu. All phases are evenlydistributed for the finest microstructure.The as-cast concentrations of Mg and Si in den-

drites for the different coarsenesses of the microstruc-ture in the Al-Si-Cu-Mg alloy are similar to the onesfor the Al-Si-Mg alloy presented in Figures 2(a) and(b). The Cu concentration in the dendrites in Al-Si-Cu-Mg is weakly influenced by the solidification rate,with a slightly higher Cu concentration in the coarsermicrostructures, see Figure 2(c). The presence of Mgatoms does not seem to influence the as-cast Cuconcentration profiles as these are similar to the onesobtained for an Al-8Si-3.1Cu-0.12Fe alloy reportedearlier.[24]

C. Al-Si-Mg (Solution Treated)

All p-Fe phases transform into b-Fe and Mg in solidsolution after 10 minutes of solution treatment at 803 K(530 �C) for the finest microstructure, and a high Mgconcentration in the matrix is obtained, see Figure 4(a).

Fig. 1—Microstructures of the SDAS 28 lm sample: (a) As-cast showing the script form of the p-Fe phase, and (b) solution treated for 30 minat 803 K (530 �C) showing b-Fe phases formed as a result of the transformation. Microstructures of the SDAS 51 lm sample: (c) As-cast show-ing a p-Fe phase (gray) formed on a b-Fe plate (white); and (d) solution treated for 6 h at 803 K (530 �C) showing thin b-Fe phases as well assome remaining p-Fe phase.


At longer solution-treatment times, a small reduction inthe size of the b-Fe phases formed through the trans-formation is seen.

Fragmentation of the p-Fe phases is seen after10 minutes of solution treatment for the intermediatemicrostructure, and small b-Fe phases can be seen onthe periphery of the p-Fe particles. After 30 minutes,most of the p-Fe particles are completely transformed,now appearing as short b-Fe phases, see Figure 1(b).After 2 hours, no p-Fe phases are found. These obser-vations are consistent with the Mg concentrationmeasurements, where 0.3 wt pct Mg is obtained after30 minutes of solution treatment, see Figure 4(b).

For the coarsest microstructure, both p-Fe phases andMg2Si phases are still found after 30 minutes thetransformation of the p-Fe phase has started, and somesmall b-Fe plates are observed. After 3 hours of solutiontreatment, most of the p-Fe phases have transformed,and only a few p-Fe phases remain after 6 hours.Figure 1(d) shows b-Fe formed as a result of thetransformation of the p-Fe phase. These observationsare consistent with the concentration measurementswhere a slow increase in Mg concentration is seen with

solution-treatment time, see Figure 4(c). Three hoursare needed to reach a high Mg level around 0.25 wt pct.To obtain all Mg in solid solution, a longer solution-treatment time of around 10 hours is needed.A homogeneous Si concentration in the dendrite

ranging from 1.1 to 1.2 wt pct, close to the equilibriumvalue of 1.06 wt pct at 803 K (530 �C),[21] is reached,after 10 minutes for the finest microstructure and after30 minutes for the two coarser ones. The homogeniza-tion of Si is as fast as or faster than for Mg, andtherefore does not determine the time needed to achievecomplete dissolution and homogenization.

D. Al-Si-Cu-Mg (Solution Treated)

During solution heat treatment, there is a rapidincrease in Cu concentration for the finest microstruc-ture, and the maximum concentration is reached after30 minutes at 768 K (495 �C), see Figure 5(a). Most ofthe Al2Cu phase have dissolved after 10 minutes, andafter 1 hour, no Al2Cu phases can be found. Phasespresent after 1 hour are the Q phase and the Al-Si-Cu-Fe phase in the shape of plates and tiny worms,

Fig. 2—As-cast concentration profiles. (a) Mg, (b) Si for the Al-Si-Mg alloy, and (c) Cu for the Al-Si-Cu-Mg alloy.


respectively, see Figure 3(b). The area fraction of the Qphase is observed by visual inspection to decrease after3 hours (no measurement was done). The Mg concen-tration in the finest microstructure increases rapidlyfrom the as-cast value in the range from 0.06 to 0.09 wtpct to around 0.22 wt pct after 10 minutes The Mgconcentration remained at this level for solution-treat-ment times up to 1 hour, after which a small increase inMg concentration to 0.25 wt pct was obtained. Theincrease in Mg concentration after 3 hours confirms thevisual observation of a decrease in area pct of the Qphase after 3 hours, indicating that the equilibriumcondition at 768 K (495 �C) might not have yet beenreached, and it may be possible to obtain a higher Mgconcentration for longer solution-treatment times.

After 1 hour, a high and homogeneous Cu concen-tration is reached in the intermediate microstructure, seeFigure 5(b), and most of the Al2Cu phases havedissolved. Fe-containing phases are the compact Al-Si-Cu-Fe phase and b-Fe plates. Some of the b-Fe platescontain Cu and start to transform to an AlCuFe phase,while others remain as b-Fe phases. For longer solution-treatment times, the fraction of Cu-containing b-Fe

plates increases. A transformation of b-Fe plates intoAl7Cu2Fe has been reported earlier for Al-9Si-3Cu-(0.2to 0.5)Mg alloys when solution treated at 763 K(490 �C).[25] No changes in the morphology, composi-tion, or fraction of the Q phase are observed. The Mgconcentration increases rapidly for the intermediatemicrostructure and 0.22 wt pct Mg is dissolved in thematrix after 10 minutes No further increase in Mgconcentration is obtained for the longer solution-treat-ment times investigated, i.e., up to 6 hours.Much longer times are needed to achieve a high and

homogeneous Cu concentration for the coarsest micro-structure because of the greater distance between Cucontaining phases, see Table III. After 10 minutes, theCu atoms in the dendrite start to homogenize, and ahomogeneous concentration profile within the dendritesis obtained after 1 hour. The Cu concentration is still ata low level after 1 hour, and large script of the eutectic(Al-Al2Cu) phase exists. After 3 hours, a few of the b-Feplates start to fragment, and a few of the plates containCu, see Figure 3(d). Undissolved Al2Cu is still presentafter 10 hours, and the Cu concentration in the matrix isabout 2.7 wt pct, i.e., longer solution-treatment times

Fig. 3—Microstructures of the SDAS 9-lm sample: (a) As-cast and (b) solution treated for 1 h at 768 K (495 �C). Microstructures of the SDAS49 lm sample: (c) As-cast showing large Q phases, and (d) solution treated for 3 h at 768 K (495 �C) showing fragmentation of the b-Fe plate.The Q phase is easily seen as the surrounding Al2Cu phase has started to dissolve.


are needed to achieve complete dissolution and homog-enization. The Mg concentration in the matrix for thecoarsest microstructure is shown in Figure 5(d).Homogenization of Mg within the dendrite is reachedafter 10 minutes The concentration increases furtherand reaches about 0.23 wt pct Mg after 3 hours, where itremains stable for the rest of the times investigated, i.e.,up to 10 hours. No decrease in the Q phase has beenobserved for the times investigated, which is consistentwith the Mg concentration measurements.

The dissolution and homogenization of Si occur rapidly,as was already seen for the Al-Si-Mg alloy. A homoge-neous Si concentration in the dendrite ranging from 0.8 to0.9 wt pct, close to the equilibrium value of 0.82 wt pct at768 K (495 �C),[21] is reached after 10 minutes, 1, and3 hours for the finest, intermediate, and coarsest micro-structures, respectively. The longer times needed to reachthe Si equilibriumconcentration for theAl-Si-Cu-Mgalloyare because of the slower diffusion rate at the lowersolution-treatment temperature, but the initial and equi-librium Si concentration also have an impact.

IV. DISCUSSION

A dimensionless diffusion time, Dt=r2s , which isindependent of the coarseness of the microstructure, isintroduced to study the diffusion kinetics. In theexpression, D is the diffusion coefficient at the solu-tion-treatment temperature, t is the time for the solutiontreatment, and rs is the radius of the spherical diffusionfield. The radius of the spherical diffusion field is ameasure of the distance the Mg/Cu atoms have to moveto remove concentration differences and is calculatedfrom the distance between areas with Mg/Cu-richparticles, l, as rs ¼ l 3=4pð Þ1=3.Figure 6(a) shows the increase in Mg concentration in

the center of dendrites for the three coarsenesses of themicrostructure for the Al-Si-Mg alloy after varioussolution-treatment times at 803 K (530 �C). Using thedistance between p-Fe phases to calculate the radius ofthe spherical diffusion field and the dimensionlessdiffusion time made the three curves merge into one,see Figure 6(b), meaning that the difference in increase

Fig. 4—Mg concentrations measured over dendrite arms for various solution-treatment times at 803 K (530 �C) for (a) SDAS 10 lm, (b) SDAS28 lm, and (c) SDAS 51 lm, for the Al-Si-Mg alloy. The values are average measurements on at least three dendrites. The standard deviationsare 0.01 to 0.02 wt pct for the SDAS 10-lm samples, from 0.01 to 0.04 wt pct for the SDAS 28-lm samples, and from 0.01 to 0.06 wt pct forthe SDAS 51-lm samples.


Fig. 5—Cu and Mg concentrations measured over dendrite arms for various solution-treatment times at 768 K (495 �C) for (a) SDAS 9 lm, (b)SDAS 24 lm and (c, d) SDAS 49 lm, for the Al-Si-Cu-Mg alloy. The values are averages from at least three dendrites. The standard deviationsare 0.02 to 0.5 wt pct Cu for the SDAS 9-lm samples, 0.01 to 0.2 wt pct Cu for the SDAS 24-lm samples, and 0.06 to 0.5 wt pct Cu and 0.01to 0.03 wt pct Mg for the SDAS 49-lm samples.

Fig. 6—Mg concentration increase in the center of the dendrite for the Al-Si-Mg alloy as a function of (a) the solution-treatment time and (b)the dimensionless diffusion time, using the distance between p-Fe phases to calculate the radius of the diffusion field. The dotted line shows thecalculated curve.


of Mg concentration with time for the different coarse-nesses of the microstructure seen in Figure 6(a) is due tothe difference in diffusion distance. It should be men-tioned that the initial Mg concentrations are differentfor the different coarsenesses of the microstructure,although this does not seem to influence the concentra-tion increase for longer solution-treatment times, as theconcentration gradients within the dendrites homoge-nize quickly because of diffusion of Mg within the a-Alphase and not from distant p-Fe phases.

The increases in Cu and Mg concentrations in thecenter of dendrites for the Al-Si-Cu-Mg alloy aftervarious solution-treatment times are presented in Fig-ures 7(a) and 8(a), respectively. The diffusion kineticswere evaluated by the same procedure as for the Al-Si-Mg alloy, using a dimensionless diffusion time. Thedistances between Al2Cu phases given in Table III wasused to calculate the radius of the diffusion field for Cu,and the result is shown in Figure 7(b). The concentra-tion increase for the intermediate microstructure doesnot merge perfectly with those of the finest and coarsestmicrostructures. A better agreement was observed forthe Al-Si-Cu alloy in an earlier investigation,[24] but theagreement is still reasonably good for the Al-Si-Cu-Mgalloy. The Q phase is the main Mg-rich particle present.It is, however, not dissolving during solution treatmentat 768 K (495 �C) and does not contribute to theincrease in Mg concentration in the dendrites. Theobserved increase in Mg concentration mainly originatesfrom diffusion of Mg from the a-Al phase of the Al-Sieutectic which has a higher Mg concentration than thedendrite in the as-cast condition because of segregationduring solidification. SDAS is a measure of the distancebetween regions of a-Al rich in Mg, and using SDAS tocalculate the radius of the diffusion field for Mg makesthe Mg concentration curves merge into one, seeFigure 8(b).

The solution-treatment response of the Al-Si-Cu-Mgalloy is compared in Figure 9 with the response of anAl-8Si-3.1Cu-0.12Fe alloy,[24] cast and solution treated

in the same way as for the present article. The Cuconcentration is seen to increase a little faster for the Al-Si-Cu-Mg alloy for short solution-treatment times, seeFigure 9. For longer times, no difference between thetwo alloys is seen. The distribution of Al2Cu particles isslightly finer for the Al-Si-Cu alloy compared with theAl-Si-Cu-Mg alloy and could not be responsible for thefaster increase in Cu concentration at short times for theAl-Si-Cu-Mg alloy. Nor can a higher diffusivity of Cuwhen Mg is present in solid solution explain the fasterincrease, as it does not persist for longer solution times.It can therefore be concluded that diffusion of Cu atomsin the a-Al matrix is not strongly influenced by thepresence of Mg atoms.

V. MODEL

If the as-cast microstructure is known, either fromexperiments or from solidification calculations, then thetime needed for dissolution and homogenization can becalculated. Solution treatment of cast aluminum alloyshas earlier been successfully modeled.[7,25–27] The modelused in the current study is based on a model developedby Vermolen et al.[26,27] The as-cast concentrationprofile and the fraction and distance between Cu/Mg-rich particles are needed as input to the model.Modeling the as-cast concentration profile is a difficulttask as segregation profiles are present both on the scaleof dendrites and on the scale of grains. During solutiontreatment, the concentration gradients within the den-drites and surrounding Al-Si eutectic homogenize quiterapidly. To increase the concentration of Cu/Mg fur-ther, distant Cu/Mg-containing particles must dissolve,and the Cu/Mg atoms released must diffuse a longdistance to remove the concentration gradients. The aimof the model is to estimate the solution-treatment timesneeded for different coarsenesses of the microstructure.The concentration gradients within the dendrites thenbecome less important as they homogenize rapidly, see

Fig. 7—Cu concentration’s increase in the center of the dendrite for the Al-Si-Cu-Mg alloy as a function of (a) the solution-treatment time; and(b) the dimensionless diffusion time, using the distance between Al2Cu phases to calculate the radius of the diffusion field. The modeled resultsfor the original geometry and the modified geometry are shown.


Figures 4 and 5. In this model, Scheil segregation is usedto calculate the concentration gradient between distantCu/Mg-rich particles, while the local concentrationgradients within the dendrites are neglected. The Scheilcalculation gives the fraction of Cu/Mg-rich particles,and the distances between them are obtained frommeasurements.

A spherical diffusion geometry with a sphericalparticle in the center, see Figure 10(a), was used as itgave the best fit to the experimental results. Thecompositions of the Cu/Mg-rich phases are constantduring the dissolution process, and Cu/Mg diffuses fromthe outer layer of the particles into the matrix. Theconcentrations in the particles, Cparticle, are assigned totheir stoichiometric compositions. Local equilibrium atthe particle–matrix interface is assumed, and the con-centration, Ci, is assumed to be the solubility limit ofCu/Mg in the matrix at the solution-treatment temper-ature. This is calculated using ThermoCalc.[21] Thisboundary condition at the particle–matrix interface is

relaxed when the particle has dissolved. The model isbased on diffusion, and Ficks Second law is solved withthe finite difference method. Diffusion is calculated in afixed system, and then the movement of the particle–matrix interface is calculated using mass balance. Theradius of the particle is reduced as the solute concen-tration in the matrix increases. The diffusion coefficientsfor Cu and Mg used in the model are derived from areview of diffusion coefficients.[28]

The model is able to predict the concentrationincrease of Mg in the Al-Si-Mg alloy well, see Fig-ure 6(b). The maximum Mg concentration predicted bythe model, 0.34 wt pct, is a little higher than themeasured 0.31 wt pct. The deviation could either be dueto Mg bound in particles or the uncertainty in thedetermination of the Mg concentration of the alloyusing optical emission spectroscopy or the Mg concen-tration of the matrix using WDS.The model was also applied on data for Al-Si-Mg

alloys available from the literature. Three datasets wereanalyzed treating Sr-modified Al-7Si-(0.26 to 0.62)Mgalloys having SDAS in the range from 23 to 55 lmsolution treated at 813 K (540 �C).[9,29] The distancebetween Mg-rich particles was not known, only theSDAS. Using the SDAS to calculate the radius of thediffusion field gave a good fit to the experimental results.The distance between Mg-rich particles was estimatedfrom the SDAS with the help of data from the currentinvestigation and was then used to calculate the radiusof the diffusion field which resulted in a much too slowincrease in Mg concentration in the matrix. The distancebetween Mg-rich phases was measured for die -castsamples of the two alloys used in the current investiga-tion, as well as for two components, and it was possibleto exclude a more homogenous distribution of Mg-richphases for die-cast samples compared with gradient-castsamples as a possible explanation for the faster concen-tration increase reported on in the literature for die-castsamples. Further studies are needed to understand thisdifference.

Fig. 8—Mg concentration’s increase in the center of the dendrite for the Al-Si-Cu-Mg alloy as a function of (a) the solution-treatment time; and(b) the dimensionless diffusion time, using SDAS to calculate the radius of the diffusion field. The dotted line shows the calculated curve.

Fig. 9—Concentration increase in the center of the dendrite for theAl-Si-Cu[24] and Al-Si-Cu-Mg alloys as a function of the solution-treatment time.


Not all Mg-rich phases dissolve in the Al-Si-Cu-Mgalloy, and the Mg concentration increases rapidly to avalue ranging from 0.22 to 0.24 wt pct where it stabilizesfor longer solution-treatment times. To be able to modelthe increase of Mg in the Al-Si-Cu-Mg alloy, thesolubility of Mg in the a-Al phase at 768 K (495 �C)must be known. ThermoCalc predicts a solubility of0.32 wt pct Mg in the a-Al phase at 768 K (495 �C),[21]which is much higher than the measured concentration.According to ThermoCalc the Q phase, the Al7Cu2Fephase and the b-Fe phase are present at 768 K (495 �C).The existence of these phases was confirmed by micro-structural investigations. However, phase transforma-tions for all three phases take place in the temperatureregion around 768 K (495 �C), and this might decreasethe accuracy of the calculation. Changing the solubilitylimit at 768 K (495 �C) to 0.23 wt pct Mg and usingSDAS to calculate the radius of the diffusion field gave agood fit between the measured and the calculated Mgconcentrations for the Al-Si-Cu-Mg alloy, see Fig-ure 8(b). The change in solubility of Mg with temper-ature needs to be known to allow using the model fortemperatures other than 768 K (495 �C) for Al-Si-Cu-Mg alloys. The concentrations of Mg in the matrixafter the solution treatments for 40 h at 748 K, 758 K,768 K, and 778 K (475 �C, 485 �C, 495 �C, and505 �C) were measured for the coarsest microstructure.The Mg concentration increases with solution-treat-ment temperature, but is on a lower level comparedwith concentrations calculated using ThermoCalc,[21]

see Figure 11.Using the distance between the Al2Cu phases to

calculate the diffusion radius gave a much too slowincrease in Cu concentration for the Al-Si-Cu-Mgalloy, see Figure 7(b). This has also been observedpreviously by these authors for an Al-Si-Cu alloy.[24]

The deviation might be due to the very high diffusivityof Cu in Si crystals, resulting in a diffusivity of Cu inthe Al-Si eutectic being four to five times higher than inthe primary a-Al phase.[30] The microstructure consistsof about 65 vol pct Al-Si eutectic according to a mass

balance calculation assuming that 1.5 wt pct Si isdissolved in the primary a-Al phase, which is ameasured average in the dendrite center for samplesused in this investigation. The diffusion field wasdivided into different parts consisting of primary a-Alphase and Al-Si eutecticum, see Figure 10(b). Thedivision was made based on the appearance of themicrostructure, with about three secondary dendritearms between two Al2Cu phases and the primary a-Alphase being the closest to the Al2Cu phase. The Al-Sieutectic was given a diffusivity five times higher thanfor the primary a-Al phase. The result of the calcula-tion using the modified geometry is in better agreementwith measurements, see Figure 7(b). The measuredconcentration increase is, however, still faster thanthe modeled one. One reason for the deviation is thetoo low initial concentration after solidification calcu-lated using Scheil segregation.

VI. VALIDITY OF THE MODEL

The focus of this investigation has been the concen-trations of solute elements in the a-Al matrix aftersolution treatment, which are important to determinethe yield strength obtained after artificial aging. Itshould be remembered that the solution treatment alsoinfluences the spheroidization and coarsening of eutecticSi particles which influences the alloy ductility. Thetimes mentioned in this article are those needed toachieve complete dissolution and homogenization, andlonger times are probably needed to achieve an increasein ductility. Zhang et al.,[29] for example, report that atime of 10 minutes at 813 K (540 �C) is sufficient toachieve a high and homogenous concentration in an Al-7Si-0.3Mg alloy having SDAS around 50 lm, while atime of 30 minutes was needed to achieve an increase inductility.The experimental results are valid for Sr-modified

and grain-refined alloys cast with the gradient solidi-fication process. Sr modification is reported [14,16] toslow down the dissolution of the Al2Cu phase in Al-Si-

(a) (b)

Fig. 10—(a) Geometry used for the solution-treatment model. r0 isthe radius of the particle, and rs is the radius of the spherical diffu-sion field. (b) Modified geometry for diffusion of Cu.

Fig. 11—Solubility of Mg at various temperatures. Measured con-centrations and calculated concentrations using ThermoCalc.[21]


Cu-Mg alloys due to the segregation of Cu duringsolidification leading to the formation of the blockyAl2Cu phase which is more difficult to dissolve than theeutectic (Al-Al2Cu) phase. The distance between Cu/Mg-rich phases is also influenced by grain refinementand the solidification process. The model can, however,be used for all kinds of aluminum alloys when thedistance between the dissolving phases can be measuredor calculated.

VII. CONCLUSIONS

The time needed to reach complete dissolution andhomogenization can be predicted if the distancebetween regions which are rich in alloying elements isknown either from measurements or calculations. Itcould either be the distance between particles whichdissolve during solution treatment or the distancebetween regions having a high concentration of alloy-ing elements in solid solution after solidification. ForAl-Si-Mg alloys, the distance between p-Fe phasesshould be used. For the Al-Si-Cu-Mg alloy, thedistance between Al2Cu phases should be used todescribe the Cu diffusion and SDAS to describe the Mgdiffusion.

All Mg dissolves into the matrix for the Al-7Si-0.3Mg alloy when solution treated at 803 K (530 �C)because the p-Fe phase is unstable and transforms intoshort b-Fe plates which releases Mg. The situation isdifferent in the Al-8Si-3Cu-0.5Mg alloy where Mg isbound in the Q phase in the as-cast condition. The Qphase does not dissolve completely at 768 K (495 �C)for this alloy, and the Mg concentration in the dendritemainly increases because of diffusion within the a-Alphase. An Mg concentration ranging from 0.22 to0.25 wt pct in the matrix is reached. Most of the Cu inthe Al-8Si-3Cu-0.5Mg alloy is bound in the Al2Cuphase in the as-cast condition, which dissolves duringsolution treatment. Some Cu is bound in the Q phaseand in the Fe-rich phases which are stable at 768 K(495 �C).

ACKNOWLEDGMENTS

The financial support received from the EuropeanProject NADIA (New Automotive componentsDesigned for and manufactured by Intelligent process-ing of light Alloys) is gratefully acknowledged.

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optimization of solution treatment of cast al-7si-0.3mg and al-8si-3cu-0.5mg alloys

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