Investigation and Nanomechanical Behavior of the Microconstituents of Al-Si-Cu alloy After Solution and Ageing Heat Treatments

The effects of heat treatment on the microstructural evolution and mechanical behavior of Al-Si-Cu alloy were investigated. Solution heat treatment was performed at 500oC/6h and 540oC/6h, followed by quenching in warm water at 60oC. Lastly, artificial ageing treatment was performed at 160oC/6h, 190oC/4h and 210oC/2h. The microstructural evolution of the samples before and after heat treatments was analyzed by optical microscopy, scanning electronic microscopy (SEM) and Energy Dispersive Spectroscopy (EDS). Analyses revealed the presence Al5FeSi, Al15(FeMn)3Si2, Al2Cu and Si particles. The morphologies of the present phases were dependent on the heat treatment. Solution treatment at 540oC/6h resulted in significant softening of the alloy compared with the treated samples at 500oC/6h. Vickers hardness analysis indicated that the hardness peak of the alloy after heat treatment was about 125 HV. Mechanical properties of different phases have been investigated by nanoindentation technique, which indicated that the highest hardness and modulus of elasticity were found for the Si particle, especially those distributed in the microstructures of the samples submitted to higher solution and ageing temperatures.


Introduction
The hypoeutectic Al-Si-Cu alloys present excellent mechanical properties, including a good strength versus weight ratio, which make these alloys excellent choices for automotive industries 1 . The Al-Si-Cu alloys in the as-cast condition present microstructures composed by primary aluminum matrix with predominantly dendritic grains and Al-Si eutectic region 2 . The presence of elements as Fe, Cu, Mg and Mn in their chemical composition may promote the formation of intermetallic compounds as Al 5 FeSi, Al 15 (FeMn) 3 Si 2 and Al 2 Cu on the microstructures of the aluminum alloys 3 . Some intermetallics may be responsible for decreasing the mechanical properties on the materials. Some of Fe-rich compounds present needle-shaped morphology, which may cause material embrittlement by stress concentration 4 . However, the presence of other microconstituents contributes to improve mechanical properties. The presence of Al 2 Cu and Al 15 (FeMn) 3 Si 2 on the microstructure, promotes a good combination of strength and ductility in the alloys 5 .
Another alternative to obtain improved mechanical properties to Al-Si alloys is the employment of solution heat treatment which promotes necking, fragmentation and spheroidization of eutectic Si particles, modifying the morphology from needle to globular-shapes, reducing the detrimental effects of the former phase 6,7 . Wang et al. 8 observed the coarsening process of the Si particles in eutectic Al-Si alloy after solution heat treatment at 560 o C. According to Samuel et al. 9 , ultimate tensile stress and ductility are greatly enhanced after solution treatments during different periods of time, followed by ageing treatment at 150 o C for 5h. Enhancement of the mechanical properties of the alloy studied in their work was attributed to the Mg 2 Si precipitation, dissolution of Si within the Al-matrix, and change in the Si particle morphology to spheroid shape. Thus, modification of the primary aluminum grains from dendrites to globular shapes, as well as changes of the Al-Si eutectic region are options found in the solution treatments.
Ageing treatments performed at higher temperatures and longer times promote the formation of Guinier-Preston (GP) zones that may act as precursors to nucleate and grow precipitates. On the other hand, ageing at low temperatures contributes to the delay of precipitates growth, resulting in the formation of a higher quantity of smaller particles, which enhances the strength of the alloy. Microhardness tests are utilized in Al-Si alloys after solution treatment and ageing to investigate the influence of the heat treatment parameters on the mechanical properties of the treated alloys. In previous literature 10, 11 , ageing treatment in Al-Si-Mg alloys lead to the formation of finer precipitates in the supersaturated primary *e-mail: silvanoleal@fatecsp.br Al matrix in Al-Si-Mg alloys, resulting in a good combination of strength and ductility. The hardening peak obtained in alloy under ageing treatment is related the formation of GP zones and precipitated phases in the microstructure. These phases play an important role in material strengthening due to the strong reduction of dislocation movements caused by semi-coherent and needle-shaped precipitates 12 .
The different phases formed in the microstructure of the Al-Si-Cu-Fe alloys due to the presence of alloying elements need to be controlled and studied. The individual phases, such as primary Al, Fe-rich intermetallics and Si particles present different hardness and elastic modulus. Thus, the individual mechanical properties of each microconstituent contributes directly to the mechanical properties of the alloy 13,14 . Fe-rich intermetallics, for example, may cause embrittlement of Al-Si-Cu alloys due their morphology 15 . Nanoindentation tests are, in this sense, a powerful tool to investigate the mechanical behavior of each individual phase. Nevertheless, there are few reports on hardness properties and elastic modulus of the intermetallic compounds and particles that constitute the microstructures of the Al-Si-Cu alloys 16,17 .
In spite of the relevant findings accumulated so far on the macrohardness behavior of Al-Si-Cu alloys, the concomitant effect of the solution temperature and ageing treatment on the hardness and elastic modulus response in each individual phase is rarely found in the literature.
In this sense, this work aims to investigate the mechanical behavior of the individual phases obtained from solution / ageing treatments. In order to achieve this goal, nanoindentation technique was employed to obtain hardness and elastic modulus of the following individual phases: Al matrix, Si particles and Al 15 (FeMn) 3 Si 2 . We also correlate the mechanical properties found from nanoindentation tests to the solution temperature and ageing treatment. Finally, the mechanical behavior of the alloy is investigated and correlated to the individual contribution of each microconstituent.

Material and preparation
The Al-Si alloy used in this work was the DIN EN AC 46000, widely utilized in automotive applications 18,19 . The alloy was melted in a Si-C crucible at 720ºC, using a resistance furnace. The melted alloy was poured into a rectangular cross section metallic mold (2 x 4 x 25 cm). Billets shape samples (2 x 2 x 2 cm) were obtained. The chemical composition of the alloy, shown in Table 1, was obtained by optical Emission Spectrometry (Spectrolab).
Samples were submitted to solution heat treatments, carried out at 500ºC and 540ºC, both for 6 hours, and quenched in warm water at 60 ºC. The samples were, afterwards submitted to ageing at 160ºC, 190ºC and 210ºC holding at 6, 4, and 2h, respectively. The aged samples were cooled in air to room temperature. The thermal cycle for each stage of the T6 heat treatment is represented in Figure 1.

Microstructural and microhardness analysis
Differential scanning calorimetry (DSC) measurements under argon atmosphere were performed to evaluate the phase transformation temperatures of the alloy in the as-cast condition under cooling using a Setaram Setsys Evolution 1200 equipment with a sample weight of approximately 40 mg. Temperature was raised from room temperature up to 700°C at 10°C·min −1 . After reaching this temperature, a cooling rate of 10°C·min −1 was applied to the molten alloy and the heat flux was measured during solidification back to room temperature.
Microstructure analyses were performed by optical microscopy (Zeiss, Axio) and scanning electron microscopy (Jeol, JCM 6000 Plus). For the microstructural analysis, samples were previously grounded with SiC abrasive paper up to 600 grit, polished with diamond paste up to 1 µm and etched with Keller's solution (0.5% HF in 50 ml H 2 O) for 10 s. X-ray energy dispersive spectroscopy (EDS) was employed for the analysis of chemical composition of the phases and intermetallic compounds. High Resolution X-ray diffraction (XRD) was also employed to identify the microconstituents in the alloy (Bruker AXS D8).
Microhardness measurements were carried out in a Vickers microhardness tester (Mitutoyo, HM 210) with an applied load of 1 kgf. Hardness values were obtained from the average of seven indentation readings taken in each sample.

Nanoindentation technique
Nanoindentations tests were performed to access the mechanical properties of individual phases within the Al-Si alloy and correlate the results with the macroscopic behavior of the materials. Nanoindentation tests were performed in an Anton Paar Nanoindenter, model NHT3, operated in a continuous stiffness measurement mode, under 20 and 50 mN loads, using a diamond Berkovich indenter. From load-displacement curves obtained from nanoindentation tests, we could obtain the nanohardness value, H, for each phase, according to the method of Oliver and Pharr 20,21 : where Max P is the maximum indentation load and C A is the projected area of the hardness impression. The elastic modulus of each phase may also be retrieved from nanoindentation tests, as follows: . r E is the reduced modulus, where S is the slope of the load/indentation curve during the unloading cycle. The reduced modulus is a combination of effective contact between the indenter and the material test 16 .

Thermal analysis
The DSC curve of Al-Si-Cu alloy in the as-cast condition is shown in Figure 2. The exothermic peak centered at ~505ºC (peak 1) is related to the formation of the copper rich phase (Al 2 Cu) 22 . The peak centered at ~552ºC (peak 2) may be attributed to the solidification of the eutectics (Si -α-Al) 23 and Fe-containing phases 24 . The peak centered at ~577ºC (peak 3) corresponds to the solidification of the primary α-Al phase. Finally, the peak centered at ~614 ºC (peak 4), indicates the presence of the primary Fe rich phase 25 .

Microstructural Characterization
The microstructural features of the as-cast sample are shown in Figure 3a. Si particles with needle-like morphology dispersed on α-Al matrix are observed. Besides, the presence of Al 5 FeSi intermetallic phase with acicular shape was also noted. Acicular phases shown in the as-cast alloy are quite detrimental to the mechanical properties of the material, because the tips of these particles may act as stress concentrators 26 . Similar microstructural features were also reported for as-cast Al-Si alloys 27 .
The microstructures of the samples submitted to solution treatments at 500ºC/6h and ageing at 160 o C/6h, 190 o C/4h and 210 o C/2h showed, sequentially, necking, fragmentation and spheroidization of eutectic Si particles. In this sense, the heat treatments caused a modification in Si particles, from acicular to globular shapes, when compared with the as-cast sample, as shown in the Figure 3b-d. Results showed that solution treatment was able to increase the distortion energy near curvature regions of the Si particles, causing quick fragmentation and spheroidization of eutectic Si particles 28 . Spheroidization, by the other hand, occurs due to the reduction in the total interfacial energy, promoted by the instability of interfaces between two different phases. In a higher temperature, diffusion in the interface regions between Si particles and primary Al matrix is promoted, which favors the increase in the curvature radii reducing the free energy of the Si particles surface 28,29 . Solution treatment at 500 o C, in all ageing conditions proved, thus, to be adequate to reduce the detrimental effect of the acicular Si particles.
Figure 3e-f shows the microstructures for the samples submitted to solution treatment at 540ºC. It is notable the occurrence of exaggerated growth of Si particles and Al 5 FeSi intermetallic compound. Al-Si alloys with silicon content near the eutectic composition usually exhibit long rod/acicular eutectic Si particles 30 . However, primary silicon particles with polyhedral shape were found in hypoeutectic Al-Si alloys after heat treatment at 540 ºC 31 . This morphology is predominantly present in the microstructure eutectic matrix of cast hypereutectic Al-Si alloy 32-34 . In Figure 3g, a significant expansion of the eutectic region (Al-Si) was also noted, mainly after ageing at 210 o C/2h. Coarsening of the Si particles occurs in detriment of the acicular Si particles present in the eutectic region. This is in conformity to the fact that diffusion transport of silicon atoms is enhanced as temperature rises 35 . It was also noted the growth of Fe-rich particles distributed in the eutectic region, for the solution treatment at 540 o C in all ageing temperatures, independently of the holding times. Menargues et al. 36 observed the growth of eutectic Si particles in Al-Si alloys, during a short T6 treatment. The observed growth of the Si particles may be attributed to Ostwald ripening phenomenon 37 , that occurs when the particles sizes are greater than the critical volume, growing at the expense of smaller particles. Si particles coarsening for solution treatments above 500 o C was already reported in the literature [38][39][40] .
SEM images shown in Figure 4a-c, show the microstructures for samples in the as-cast condition, as well as for samples submitted to solution heat treatment at 500 o C and 540 o C. EDS microchemical analysis exhibited the presence of Si particles and the eutectic Al 2 Cu + Si. Besides, iron-rich intermetallics  Al 5 FeSi and Al 15 (FeMn) 3 Si 2 were also identified in the samples. Microstructural evolution shows, as discussed previously, that the detrimental needle-like particles, present in the ascast sample, turns into globular shape particles after solution heat treatment, performed at 500ºC. The efficiency of surface diffusion on the fragmentation of eutectic Si particles depends strongly on the existence of interface defects between the Si particles and primary Al phase during solution treatment 28 . The change of Si particles morphology occurs when the eutectic structure is exposed to a thermal treatment, in which the temperature is enough to activate shape perturbations in the Si particle, eventually leading to the fragmentation into a series of nearly spherical particles. The overall mechanism indicated that the fragmentation of the needle-like particles was followed by spheroidization after solution treatment at 500ºC. Nevertheless, exaggerated growth occurred for solution treatment at 540ºC, resulting in the formation of the intermetallic Al 15 (FeMn) 3 Si 2 with plate shape, as illustrated in Figure 4c. This increase in the particles size is due to the higher diffusion coefficient, which resulted in the rise of the coarsening rate, in comparasion to the spherodization rate. This is easily perceived when we compare the mean size of the particles under heat treatments at 500 and 540ºC. We also recall that the presence of intermetallic particles such as Al 2 Cu and Al 5 FeSi in the microstructure of Al-Si-Cu alloys after solution treatments and ageing were already reported in the literature 41,42 .
XRD analysis was also performed ( Figure 5) and the observed peaks were in identified as Fe-rich intermetallics Al 5 FeSi and Al 15 (FeMn) 3 Si 2 , as well as Al 2 Cu, Si and Al phases.
The presence of intermetallic with Chinese script morphology on the microstructure of sample at 500ºC/6h and ageing at 160ºC/6h was also perceived and ilustrated in the Figure 6. EDS analysis indicated that Chinese Script is constituted by the intermetallic phase Al 15 (FeMn) 3 Si 2 . Manganese present in the Al-Si alloys may contribute for the change in the morphology of Fe-intermetallic phases from lamellar to massive polyhedron 43 . Figure 7 details the microstructure of the sample submitted to solution treatment at 540 o C/6h and ageing at 190 o C/4h. EDS spectra, once more, helped to reveal the chemical composition of each microconstituent present in the material. The spectrum of each selected region indicated the presence of the Si particle (spectrum 1), Fe-intermetallics, such as Al 15 (FeMn) 3 Si 2 (spectrum 2) and Al 5 FeSi (spectrum 3), as well as Al 2 Cu + Si (spectrum 4) distributed in the primary Al matrix of the Al-Si alloy. These kinds of complex intermetallic compounds were already reported in similar alloys 44,45 . The presence of Al 2 Cu precipitate in aluminum alloys after solution treatments and ageing was also observed for Al alloys 46,47 . In this sense, alloying elements can promote nucleation and growth of complex intermetallic compounds in the microstructure of the alloy.

Effect of heat treatment on microhardness of the Al-Si-Cu alloy
The average hardness of the as-cast alloy and the samples submitted to solution/ageing heat treatments were investigated by Vickers microhardness measurements. The results are shown in the Figure 8. Samples submitted to solution treatment at 540 o C/6h presented smaller hardness (100, 105 and 112 HV), when compared to samples at 500 o C/6h (108, 109 and 125 HV). This result could be attributed to the softening of the Al-Si alloy due to the growth and morphological change of the primary Al grains from dendritic to globular shape and modification of the eutectic region. Consequently, the temperature rise of solution treatment contributed for a reduced effect of the pinning force, due to the dissolution of precipitates in the grain boundaries, leading to the growth of the primary Al grains and reduction of the hardness of the alloy.
Ragab et al. 48 reported that solution temperatures above at 540 o C could cause higher homogeinization of the microstructure, with probable full dissolution of the Al 2 Cu precipitates in regions of grain boundaries, decreasing the mechanical strength of the treated alloy. The hardness peak was obtained in samples aged at 190 o C during 4h for solution treatments at 500 o C (125 HV) and 540 o C (112 HV). This gradual increase in hardness for ageing temperatures ranging from 160 to 190 o C may be associated to the formation of GP zones (Guinier-Preston), associated with the precipitattion of Al 2 Cu phase 45 . However, a drop in the microhardness value was observed for ageing temperatures above 190 o C for both solution temperatures. This can be attributed to the higher thermodynamic stability, that resulted from the growth of the Al 2 Cu precipitate. Sjölander and Seifeddine 49 reported simillar results in a Al-Si-Cu alloy, when submitted a heat treatment in the T5 condition. Barbosa et al. 50 reported the occurrence of the maximum hardness in an intermediate ageing time, very similar to the observations made in this work.

Nanoindentention tests
Differently form microhardness tests, nanaoindentation tests are able to access the mechanical behaviour of individual phases within the Al-Si alloys. In this sense, nanoindentation tests were used to measure nanohardness and determine elastic modulus of α-Al, Al 5 FeSi and Si particles microconstituents. Multiple tests were performed in order to obtain only indentation marks covering 100% of the particles. Figure 9a shows load-displacement curves for the microcontituents present in the as-cast Al alloy. A maximum penetration depth of ~ 945 nm occurred for Al matrix, indicating this is the softer phase among all analyzed microconstituents. The load-displacement curve of the Al 5 FeSi intermetallic registered a maximum penetration depth of ~ 799 nm; maximum penetration depth for Si particle was ~ 639 nm. The plastic deformation caused by the indenter tip in each phase revealed that Al matrix has a greater capacity of elastic and plastic deformation, which was expected since soft phases present better mobility for dislocations. Discontinuities were also observed in the loading curves. These events are known as "pop-ins" and are more pronounced for brittle intermetallic phases, such as Al 5 FeSi and Si particles. Pop-in events were indicated by red and blue arrows on the loading curves of the as-cast sample presented in Figure 10a. The appearance of pop-ins in the loading curves is influenced by the homogeneous dislocation nucleation, as reported elsewhere 46,47 .
For the sample submitted to solution treatment at 500 o C/6h and aged at 160 o C/6h, Figure 9b shows that the Si particle presented the smallest penetration depth and a more pronounced elastic recovery, when compared with Al 5 FeSi intermetallic. These features are compatible with the high hardness value obtained for the Si particle (H IT = 11.59 ± 0.04 GPa). In fact, Si particle also presented the smallest penetration depth (~362) nm among all samples treated at 540 o C/6h, regardless of the conditions used for ageing treatments, as seen in Figure 10c. Unloading curves of the Si particles in heat treated samples also showed "elbows", also known as "pop-outs" phenomena. Pop-outs may be attributed to plastic deformation caused by dislocation activity, as well as pressure-induced phase transformation 51,52 . Nevertheless, cracks in the Si phases were also observed after unloading ( Figure 10) and may also originate the observed pop-outs, also revealing the brittle behavior of this microconstituent.
The results of hardness obtained by nanoindetation measurements are summarized in Figure 11.   Nanohardness of α-Al phase presented small variations among the heat treated samples. However, a higher hardness (H IT = 1.30 ± 0.13 GPa) was observed for the sample submitted to the solution treatment at 500ºC and ageing at 190ºC/4h. This is also the sample that presented the smaller Al 2 Cu and α-Al particles, as a consequence of particle fragmentation. The hardness of α-Al phase presented a slight reduction in samples submitted to the solution treatment at 540ºC, as it was noted on the aged sample at 190ºC/4h (H IT = 1.14 ± 0.03 GPa). This reduction in the hardness could be attibuted to the coarsenig of α-Al grains and presence of Al 2 Cu eutectic regions. The smallest hardness for α-Al phase was observed for the as-cast sample (H IT = 1.08 ± 0.01 GPa). The values of hardness obtained for α-Al by nanoindentation tests in room temperature are compatible with values of hardness obtained elsewhere 46,47 .
The hardness peak obtained for the Al 5 FeSi intermetallic phase was (H IT = 9.31 ± 0.52 GPa) for the sample under solution treatment at 540ºC/6h and ageing at 190ºC/4h. The Si particle presented a higher hardness value (H IT = 19.53 ± 0.54 GPa) for the treated sample at 540ºC/6h and ageing at 210ºC/2h. The increase of the hardness of the Si particle and Al 5 FeSi intermetallic compound may be attributed to the intense diffusion process acting in the conversion of small particles to big platellets (Al 5 FeSi) and disks (Si) during the solution treatment at 540 o C/6h. The formation of big Si particles discs in the eutectic region indicates that coarsening occurred via Ostwald Ripening growth mechanism 27 .
Elastic modulus from each phase could also be retrieved from nanoindentaion tests. Results showed that elastic modulus of the Al, Al 5 FeSi and Si increased for almost all the samples submitted to solution / ageing heat treatments, as shown in Figure 12. Among the phases submitted to the indentation tests, α-Al presented the smallest elastic modulus value for the heat treated samples (E IT = 105 ± 3 GPa for the sample treated at 540ºC and ageing at 210ºC/4h). However, a small variation of the elastic modulus values of α-Al phase was observed among the treated samples. Al 5 FeSi intermetallic phase presented an intermediary value of elastic modulus (E IT = 182 ± 5 GPa, for the heat treated sample at 540ºC and ageing at 160ºC/6h); Si particles showed the highest elastic modulus among all microconstituents (E IT = 267 ± 15 GPa, for the treated sample at 540ºC with ageing at 210ºC/2h). The elastic modulus profiles of different phases revealed that the Si particle presented higher elastic modulus than Al 5 FeSi intermetallic in all heat treatment conditions.
It was also noted that the Al 5 FeSi intermetallic and Si particles within the samples submitted to solution treatment at 540ºC possess higher elastic modulus when compared with the samples heat treated at 500 o C. This clearly reveals that the elastic modulus of the microconstituents is sensitive to the morphological changes promoted by the solution treatment.

Conclusions
The effect of processing parameters employed during solution and ageing treatments on the microstructural evolution and hardness properties was investigated and the conclusions may be drawn: -As-cast Al-Si alloy revealed a microstructure composed of primary dendritic α-Al matrix grains, along with Al-Si eutectic region as well as Al 5 FeSi, Al 15 (FeMn) 3 Si 2 and Si particles distributed on the α-Al matrix. -Small coarsening of the dendritic grains was observed in samples submitted to solution treatment at 500ºC/6h and ageing at 160ºC/6h, 190ºC/4h and 210ºC/2h presented small. Also, modification of Si particles from acicular to globular morphologies was observed, probably due to the presence of fragmentation and coarsening mechanisms. Si particles exhibited the highest elastic modulus (E IT = 267 GPa) among all analysed microconstituents, independently of the heat treatment conditions. In summary, we could demonstrate how the conditions used for solution and ageing treatments affect each individual phase in the microstructure of the Al-Si alloy. Nanoindentation technique proved to be a powerful tool to access the role that each microconstituent play on the determination of the overall mechanical properties of the alloy. We remark that studies comprising the employment of nanoindentation techniques along with the traditional tests used to characterize the macromechanical properties of Al alloys are not common in the literature. Nevertheless, the knowledge that can be obtained from the detailed investigation of each microconstituent in the alloy may be used to improve the alloy manufacturing in order to achieve superior mechanical properties for the engineering of advanced materials.