Effect of chemical composition and T6 heat treatment on the mechanical properties and fracture behaviour of Al-Si alloys for IC engine components

Microstructural examinations of Al-Si alloys intended to manufacture IC engine components revealed a complex phase composition in all samples. Polyhedral crystals of primary silicon were detected in Al-12.5Si alloy, besides the α-Al phase, eutectic silicon, and several intermetallic phases, identified in the cast samples of both alloys. Better tensile properties were found for samples of Al-11Si. Predominantly intercrystalline fracture with features of ductile failure was observed in both alloys. In as-cast specimens of Al-11Si alloy, the cracks were formed by the decohesion mechanism between the particles of the intermetallic phase AlCuFeNi and α-Al phase.The microcracks initiated on the interface were spread along the branches of α-Al 15 (Fe,Mn,Cu) 3 Si 2 particles. After T6 treatment of Al-11Si alloy, almost half the intermetallics quantity presented Al 3 Ni phase, while the iron-based phases were observed in a small amount. Spheroidized eutectic silicon, less portion of Al 5 Cu 2 Mg 8 Si 6 , and a more considerable quantity of Al 3 (Fe,Mn,Cu,Ni,Co) were detected for T6 specimens of Al-12.5Si alloy. The rounded crystals of eutectic silicon contributed to the improvement of their tensile properties. Larger and deeper dimples of mostly polygonal shapes were observed in the samples of Al-11Si alloy after T6 treatment. The microcracks have occurred at the boundary of the intermetallic phase/α-Al solid solution.


Introduction
Aluminium-silicon alloys have been recognised for a long period of time as the essential material in automotive, transportation and several other industries [1]. Aside from their excellent castability, good corrosion resistance and low specific gravity [2], those alloys exhibit a high strength to weight ratio as well as an excellent thermal conductivity [3]. Nowadays, aluminium-silicon castings widely substitute the cast iron for manufacturing the engine components [3]. Designing of Al-Si alloys with a good combination of castability and plasticity is based on the silicon concentration defined in the range from 4 % to 22 % [4]. Those materials are generally classified according to their silicon content into three major groups: hypoeutectic (below 12 %), eutectic (12 % -13 %) and hypereutectic (14 % -25 %) [5]. The alloys containing about 50 vol. % to 90 vol. % of the eutectic phase are mostly used in the automobile industry [6,7]. The primary purpose of the hypoeutectic alloys is the manufacturing of the engine blocks, cylinder heads, and wheels [8]. The hypereutectic Al-Si alloys find application in cylinders, cylinder liners, cylinder heads, air compressor cylinders, engine pistons, pumps, and piston rings [9][10][11][12]. Hard primary silicon crystals provide superior wear resistance of hypereutectic alloys but also reduce tool life in machining process [9]. In order to achieve favourable wear performance and simultaneously minimise issues related to hypoeutectic with 10.93 % of silicon and hypereutectic containing 12.50 % of silicon. The weight ratio of Cu/Mg was planned as 1.1 in the hypoeutectic and 1.6 in the hypereutectic alloys. A similar content of nickel, iron, cobalt and manganese was chosen, while the concentration of molybdenum was quite different. The microstructure, mechanical properties and fracture behaviour of multicomponent Al-Si alloys have been studied in as-cast and T6 heat-treated samples.

Experimental procedure
Two alloys Al-11Si and Al-12.5Si used in this study have been prepared by melting AlSi10Mg master alloy, pre-alloys (Al-33% Cu, Al-60% Mn, and Al-75% Fe), technically pure silicon, nickel and molybdenum in a 20-kW electric resistance furnace using a graphite crucible. The melting temperature was kept at 780 o C ± 5 o C. The TAL -2 consisting of potassium chloride, sodium chloride and cryolite reacting at 690 o C was added for the fluxing melts in the amount of about 2 % of the quantity of each melt. The degassing of molten Al-Si alloys was performed by Hexachloroethane tablets in the amount approximately equal to 0.25% of their quantity. The Al-10% Sr master alloy was added to the melts for the modification of eutectic silicon. The alloys were poured at 740 o C ± 5 o C into the permanent grey cast iron moulds which had not been previously preheated. After solidification, the castings were removed from the moulds and further cooled down to room temperature. The chemical composition of the produced alloys is listed in Table 1. The T6 heat treatment was performed to achieve high strength, necessary for application conditions of engine pistons [27,40] although an increase in tensile strength may be accompanied by a decrease in elongation [4]. Alloys were solution treated at 520 o C ± 5 o C for 6 hours followed by quenching in water (at 70 o C) as shown in Figure 1. The subsequent ageing treatment was performed at 205 o C ± 5 o C for 7 hours while specimens were finally cooled in ambient air.  The tensile test specimens with 4.0 mm gauge diameter and 45.0 mm gauge length were prepared from the as-cast and T6-treated alloys. The machine 1195 Instron was used for the determination of the tensile strength and relative elongation of the samples at room temperature, while the testing machine produced by Karl Frank Gmbh, type 38532 was applied for measuring their hardness as Brinell hardness HB 5/250/30. The microstructures of as-cast and T6-treated specimens were revealed by an optical microscope and scanning electron microscope JEOL JSM-6460 LV, equipped with energy dispersive X-ray spectroscopy. The fractured samples of as-cast and T6-treated alloys were observed using JEOL JSM-6460 LV in secondary electron mode in order to study the fracture behaviour of alloys.

Microstructural characterisation
The complex phase composition was revealed during microstructural examinations of the as-cast and T6-treated specimens of Al-Si alloys, as shown in Table 2. The optical micrographs of examined alloys are shown in Figures 2, 4, 6 and 7, while corresponding SEM-EDS spectra from the α-Al phase, silicon particles and intermetallic phases are presented in Figures 3 and 5. Almost 71 vol. % of α-Al phase, 21 vol. % of eutectic silicon and around 8 vol. % of several intermetallic phases were found in as-cast samples of Al-11Si alloy ( Table 2). Coarse dendrites of α-Al phase were unevenly distributed within the crosssection of the as-cast samples of hypoeutectic Al-11Si alloy. While the directional growth of α-Al dendrites was observed in the surface of the specimen (Figure 2a), in its central part, several small areas without primary α-Al crystals around the porosity could be recognised (Figure 2b), indicating higher cooling rate during crystallisation. Modified eutectic silicon with significantly diverse morphology formed a fine mixture with Al5Cu2Mg8Si6 phase (Figures 2d-2f). Two soluble Mg-bearing phases Mg2Si and Al5Cu2Mg8Si6 may be expected in Al-11Si alloy in accordance with the magnesium content. Since Mg2Si was not revealed in the microstructure of the cast specimens of Al-11Si alloy, it could be stated that magnesium which was not dissolved in the α-Al phase has formed Al5Cu2Mg8Si6 phase with copper. The high magnesium content in the alloy (1.1 %) has led to precipitation of an Al5Cu2Mg8Si6 in the amount, which presents almost 32 % of all intermetallic phases. Although the particles of Al5Cu2Mg8Si6 phase are hard to dissolve, they contribute to the strengthening of the alloy. A small quantity of Al2Cu phase was detected in the fine mixture with Al5Cu2Mg8Si6, α-Al phase and eutectic silicon (Figure 2e), which is probably the result of the ternary eutectic reaction (α-Al + Al2Cu +Si) or quaternary eutectic reaction (α-Al + Al2Cu + Al5Cu2Mg8Si6+Si). Furthermore, copper created compact or rod-shaped particles of AlCuFeNi phase (Figures 2e and 2f). Hence, the appearance of iron-based phase which bonded copper decreased its strengthening effect. Most of the nickel was found in the form of Al3Ni phase, but tiny needle-shaped particles of Al9FeNi were also recognised. The Al3Ni phase was found to be 3.04 vol. %, accounting for 38.14 % of the total amount of intermetallic phases. Two iron-bearing phases were also revealed as needle-shaped β-Al5FeSi(Mn) phase (Figures 2d and 2f) and Chinese script shaped α-Al15(Fe,Mn,Cu)3Si2 phase (Figures 2c, 2e and 2f). In the surface area of the as-cast specimens of Al-11Si alloy, α-Al15(Fe,Mn)3Si2 phase which may dissolve copper was discovered as an almost triple quantity of β-Al5FeSi(Mn). The amount of manganese equal to one half of iron content, generally desirable from the standpoint of the transformation of β-Al5FeSi to α-Al15(Fe,Mn)3Si2 [41], was not achieved for the examined alloys. However, the appearance of α-Al15(Fe,Mn,Cu)3Si2 phase may be the result of the influence of other elements which were added for correction of the harmful effect of the iron. Molybdenum and cobalt moved equilibrium towards the formation of α-Al15(Fe,Mn,Cu)3Si2 phase, but because of the conditions of crystallization that favour the appearance of β-Al5FeSi(Mn) phase, all iron was not bounded in the phase of more favourable form.

Figure 2. The microstructure of as-cast Al-11Si alloy at the surface of the specimen (a), central part of the specimen's cross-section (b), as well as other positions of examination (c) -(f)
Hypereutectic microstructure of as-cast Al-12.5Si alloy is shown in Figure 4. Around 3.35 vol. % of primary silicon, 22.62 vol. % of eutectic silicon and 65.65 vol. of α-Al phase were detected in as-cast samples of Al-12.5Si alloy ( Table 2). Polyhedral crystals of the primary silicon are dominantly formed in the central part of the cross-section of the as-cast specimen close to the porosity, while narrow zones with dendrites of α-Al solid solution within eutectics could be recognised. Total amount of several intermetallic phases was 8.38 vol. %. Modified eutectic crystals coarser (Figures 4a, 4b and 4c) than in Al-11Si alloy are mostly present in the very fine mixture with Al5Cu2Mg8Si6 phase (Figures 4d and 4e). Lower quantity of Al5Cu2Mg8Si6 phase in Al-12.5Si alloy than in Al-11Si is caused by magnesium content about 0.8 %. The phase AlCuFeNi was not revealed in the microstructure of as-cast specimens of Al-12.5Si alloy. Compared to Al-11Si, a higher quantity of copper-bearing phases is expected in Al-12.5Si alloy since it contains more copper (1.3 %). However, the formation of a small quantity of Al2Cu was detected ( Figure 4e). Hence, the excess of copper should be discovered within some other phase. Definitely, it created large, star-like crystals or regular polygonal crystals of Al3(Fe,Mn,Cu,Ni,Co) phase in an amount of 2.09 % which were observed as aggregates (Figures 4d and 4e) and often accompanied by the primary silicon crystals ( Figure  4c). The high content of molybdenum and cobalt decreased the deleterious effect of iron. However, iron formed needle-like β-Al5FeSi(Mn) phase ( Figure 4d) as well as Al9FeNi which is most often seen with a massive compact iron-bearing phase (Figures 4c and 4d) and fine rounded particles of Al3Ni phase (Figures 4d and 4e). Particles of α-Al15(Fe,Mn,Cu)3Si2 phase could be recognised in small quantity.

Figure 3. SEM image of phases (a) and corresponding EDS spectra from phases in as-cast samples of Al-11Si alloy (be)
After performing the T6 treatment, a small quantity of α-Al15(Fe,Mn,Cu)3Si2 phase was observed in the specimens of Al-11Si alloy. Rounded eutectic silicon is interdendritic, distributed as could be seen in Figures 6a and 6b. It suggests that purely eutectic zones are not observed in as-cast specimens of Al-11Si alloy. Hence, the amount of eutectic silicon in T6treated specimens is lower compared to as-cast specimens of Al-11Si alloy. While AlCuFeNi phase was not found, small rounded particles of copper-bearing phases Al5Cu2Mg8Si6 and Al2Cu were revealed in T6-treated specimens of Al-11Si alloy (Figures 6c and 6d) in quantity lower by 0.97 vol. % compared to as-cast samples. Platelets of β-Al5FeSi(Mn) phase and massive coagulated particles of Al3Ni in aggregates with eutectic silicon and Al5Cu2Mg8Si6 were also observed (Figures 6c and 6d). Al3Ni phase was revealed in the amount of 3.02 vol. % which is more than half of the total quantity of detected intermetallic phases. In the microstructure of T6-treated samples of Al-12.5Si alloy, large, compact and rounded primary silicon crystals were rarely seen, but predominantly in surface area ( Figures  7a and 7c), whereas eutectic silicon was rounded and mostly spheroidal (Figures 7c and 7d). Massive, compact star-like crystals of Al3(Fe,Mn,Cu,Ni,Co) phase (Figures 7a, 7c and 7d) are revealed in 1.5 times greater quantity than in as-cast specimens. Short needles of Al9FeNi phase (Figure 7d), fine rounded particles of Al3Ni (Fig. 7d) and acicular β-Al5FeSi(Mn) phase ( Figure 7c) are also found in heat-treated specimens of the Al-12.5Si alloy. The copperbearing phases are dissolved, and hence the quantity of Al5Cu2Mg8Si6 phase was decreased compared to cast specimens of Al-12.5Si alloy. Al2Cu phase was dissolved during solution treatment, while β-Al5FeSi(Mn) phase was almost unaffected by T6 heat treatment in both alloys. Taylor [43] was found that α and β phases in Al-Si alloys undergo to transformation, spheroidisation and Oswald ripening with the increase in the solution treatment time, while only little signs of phase transformations were observed at the typical solution treatment temperature of about 540 o C.  Table 3 lists the mechanical properties of as-cast and T6-treated specimens of the examined alloys which are the result of the intermetallic phase composition, the presence of α-Al dendrites, morphology of eutectic and primary silicon as well as the central porosity appearance. While similar values of hardness were identified for both alloys, higher values of tensile strength and elongation were exhibited by samples of Al-11Si compared to Al-12.5Si alloy. Magnesium and copper have formed Al5Cu2Mg8Si6 phase which has been hard to dissolve. The particles of Al5Cu2Mg8Si6 and Al2Cu phases promoted the tensile properties of both alloys. The Al3Ni phase was detected at least as one-third of the total amount of the intermetallic phases in examined specimens. The presence of larger particles of eutectic silicon in Al-12.5Si than in Al-11Si alloy may be considered as one of the causes of its lower tensile properties due to the creation of high levels of stress concentration [23]. It is worth to mention that hypereutectic Al-12.5Si alloy contained polyhedral crystals of primary silicon in as-cast specimens which were rarely seen after performing the T6 heat treatment. Those crystals certainly influenced the level of mechanical properties of Al-12.5Si alloy. In addition, intermetallic phases based on iron formed in all samples (Figures 2 and 4) have had deleterious effects on their mechanical properties, but nickel and cobalt mitigated that influence to some extent. The iron-bearing phase α-Al15(Fe,Mn,Cu)3Si2 was revealed in higher quantity than needle-like β-Al5FeSi(Mn) in as-cast specimens of alloy Al-11Si which contained 0.52 % of iron, 0.67 % of nickel, 0.80 % of cobalt and 0.20 % of molybdenum. On the other hand, a small quantity of α-Al15(Fe,Mn,Cu)3Si2 and β-Al5FeSi(Mn) phases was detected in as-cast Al-12.5Si alloy containing 0.43 % of iron, 0.58 % of nickel, 0.85 % of cobalt and 0.50 % of molybdenum. Besides, needle-like β-Al5FeSi(Mn) and Al9FeNi phase were also revealed in all samples. The air and gasses entrainment during the filling phase probably have caused the porosity observed in as-cast samples of Al-11Si and Al-12.5Si alloys. Dinnis et al. argued that the iron content contributes to the increase in the porosity level, but depending on silicon and copper concentration in alloy [42]. In order to avoid a decrease in mechanical properties, an increase in the porosity level and loss of ductility, the concentration of iron (in wt. %) in Al-Si alloys should be kept below the critical iron content that can be calculated [43] as follows:

Mechanical properties
≈ 0.075 • [ . % ] − 0.05 (1) As regards our alloys, Fecrit was determined as 0.77 % and 0.89 % for Al-11Si and Al-12.5 Si alloys, respectively. Since the content of iron was below the critical level, it could not be recognised as the major cause of the central porosity of the examined alloys. At the same time, the content of copper of about 3 %, which was defined as critical for the creation of microporosity, was not reached in our alloys. Therefore, it could be concluded that the pouring procedure contributed to the porosity appearance in the examined samples.
The as-cast specimens of Al-11Si and Al-12.5Si alloys exhibited a tensile strength of 228.8 MPa and 192.5 MPa and a low elongation of 2.4 % and 0.8 %, respectively. Compared to our alloys, Shivaprasad et al. have reported better values of tensile stress of 237.0 MPa as well as a significantly higher elongation of 9.1% for Al-12Si alloy due to the refinement of grains and modification of the eutectic silicon [10]. The values of tensile strength and hardness of our alloys are higher than Al-12Si-xCu-1Mg-1.78Ni (x = 0, 1, 2, 3 and 4) alloys studied by Pratheesh et al. [37] as well as Al-Si-Cu-Mg-Ni alloys with 10.5 %, 12 %, 18 % and 24 % of silicon examined by Zeren [44]. Nevertheless, Al-Si-Cu-Mg-Ni alloys with 10.5 %, 12 %, 18 % and 24 % of silicon have exhibited an elongation below 2.4 %, while the lowest value (about 0.5 %) was recorded for the alloy with the highest content of silicon (24 %) [44].
The tensile testing of T6-treated samples has revealed an increase in tensile strength and elongation as well as a decrease in the hardness of both alloys. While a negligible increase was recognised for Al-11Si alloy, the strength of Al-12.5Si alloy was improved by 33.5 MPa.
That enhancing of strength was caused in a great extent to the distribution of rounded and spheroidal eutectic silicon. After performing T6 heat treatment, the elongation of Al-12.5Si alloy was improved from 0.8 % to 2.3 %, whereas the samples of Al-11Si alloy have exhibited a slight increase in elongation. The results of the hardness measurement indicated a decrease after performing T6 heat treatment which may be caused by wide dispersion of eutectic silicon particles. According to Iskah et al. [45], the area of widely dispersed silicon particles that is available for distribution of the forces applied by indenter during hardness measuring is less compared to the area of the soft α-Al phase, resulting in lower values of sample hardness.
After the T6-treatment, our alloys had higher values of tensile strength and similar hardness compared to Al-12Si-xCu-1Mg-1.78Ni alloys without copper or containing 1 % and 4 % of copper [37]. However, for Al-12Si-xCu-1Mg-1.78Ni alloy with 3 % of copper, the tensile strength above 250 MPa and hardness above 120 BHN [37] were found, which was better than the values exhibited by our alloys. The values of the hardness of the T6-treated samples of Al-Si-Cu-Mg-Ni alloys with 10.5 % and 12 % of silicon [44] and our Al-11Si and Al-12.5 Si alloys were similar. Same alloys have shown tensile strength of above 250 MPa [44] which was better compared to our specimens. However, the increase in silicon content in Al-Si-Cu-Mg-Ni alloys up to 18.1 % and 24.0 % contributed to a higher hardness (above 104 BHN), but a lower tensile strength (below 200 MPa) compared to our alloys [44].

Fracture analysis
Fracture surfaces of the as-cast and T6-treated Al-11Si and Al-12.5Si alloys are shown in Figures 8 to 13. The plane surface was observed indicating a brittle fracture (Figures 8, 10a,  10b, 11a, 11b and 13a). However, examination under higher magnification revealed predominantly intercrystalline fracture with features of ductile failure in some areas.  (Figures 9c, 9d, 9e and 9f). The cracks were formed by the decohesion mechanism between the particles of the intermetallic phase/α-Al solid solution rather than by the initiation of microcracks within the particles. When we consider elongated crystals, it can be concluded that microcracks are formed on the boundary surface of the particles and α-Al solid solution and spread through the long axis of the particle ( Figure  9c) followed by the sideways spread and finally to fracture of the particles (Figure 9e). In the case of phase α-Al15(Fe,Mn,Cu)3Si2, the microcracks initiated on the boundary surface ( Figure  9d) and spread along the branches of particles with Chinese script shape. Around the brittle particles of the intermetallic phase, shallow or deep dimples appeared, indicating ductile fracture of α-Al matrix as well as the increased resistance of eutectic silicon crystals to fracture. The dimple areas are larger after the T6 heat treatment (Figure 10). Due to the rounding and coarsening of eutectic silicon crystals, the dimples are broader and deeper and mostly polygonal in shape. Microcracks occurred at the surface boundary of the intermetallic phase/α-Al solid solution, but also due to the fracture of particles of the intermetallic phase without expanding into the surrounding α-Al solid solution. The appearance of the primary silicon crystals and, to a lesser extent, modified eutectic silicon crystals in the as-cast specimen of Al-12.5Si alloy (Figures 11a and 11b), contributed to a greater brittleness compared to the cast Al-11Si alloy. The microcracks were formed by a fracture of the primary silicon crystals (Figures 11d and 11e) and massive particles of star-like and polyhedron forms of Al3(Fe,Mn,Cu,Ni,Co) phase (Figures 11f, 12a, 12b and 12c). However, it should be noted that Al3(Fe,Mn,Cu,Ni,Co) phase has exhibited high resistance to fracture. Areas of ductile fracture were observed with tiny polygonal dimples (Figures 11c, 11f and 11a), while in other places, elongated dimples with a planar, shallow bottom covering the width of the eutectic silicon lamellas were discovered (Figures 11e and 12c). When we consider the tensile testing of the alloys it should be kept in mind that the cracking initiation is affected by a very pronounced porosity, which is observed macroscopically for all specimens. The pores in the central part of the sample's cross-section are deep and irregular in shape (Figures 8a, 10a, 10b  and 10c). However, there are also planar, more rounded pores (Figures 11a, 13a and 13c) where dendrites of the α-Al solid solution can be clearly seen (Figure 10b).
A higher ductility of the Al-12.5Si alloy is produced by the T6 heat treating (Figure 13c), which is manifested in a larger area with polygonal dimples of different sizes (Figures 13b  and 13d). Finer dimples were observed on the fracture surface of the T6 heat-treated samples of Al-12.5Si compared to Al-11Si alloy. This is in agreement with the size of rounded crystals of eutectic silicon in the structure of these two heat-treated alloys. However, the increase in the fraction of ductile fracture has also been caused by the appearance of few primary silicon crystals, which are easily broken (Figure 13e). Although massive compact particles with iron recognised in the structure contribute to the brittle intercrystalline fracture, microcracks initiated in them are less visible (Figures 13d and 13f). On the other side, compared to as-cast samples, the tensile strength of the T6-treated specimens of Al-12.5Si alloy was increased by 33.5 MPa, whereas the elongation was raised from 0.8 % to 2.3 %. The tensile properties were enhanced due to the distribution of rounded and spheroidal eutectic silicon. The decrease in the hardness values after the T6 heat treatment of both alloys, probably was caused by the wide dispersion of eutectic silicon particles. 4. The pronounced porosity has affected the cracking initiation. However, the iron content in the examined alloys was not the primary cause of the central porosity. The massive particles of intermetallic phases have indicated the intercrystalline fracture in Al-11Si alloy. On the other hand, silicon mainly in the form of primary crystals and in a lesser extent as eutectic crystals contributed to the brittle features in Al-12.5Si alloy. The harmful effect of iron has been decreased due to the addition of manganese, nickel, molybdenum and cobalt. The α-Al15(Fe,Mn)3Si2 phase was discovered as an almost triple quantity of needle-shaped β-Al5FeSi(Mn) phase. 5. Predominantly intercrystalline fracture with features of ductile failure was recognised for both alloys. In as-cast specimens of Al-11Si alloy, the cracks are formed by the decohesion mechanism between the particles of intermetallic phase AlCuFeNi and α-Al solid solution. Shallow or deep dimples appeared around the brittle particles of the intermetallic phase indicated a ductile fracture of α-Al matrix as well as the increased resistance of eutectic silicon crystals to fracture. 6. The primary silicon crystals mainly cause a high brittleness of as-cast Al-12.5Si alloy. The microcracks have arisen from the fracture of the primary silicon crystals and massive particles of star-like and polyhedron forms of Al3(Fe,Mn,Cu,Ni,Co) phase. Areas of ductile fracture are observed in some places with tiny polygonal dimples. Besides, elongated dimples with a planar, shallow bottom covering the width of the eutectic silicon lamellas are revealed. 7. Large dimples areas have appeared after the T6 heat treatment of specimens of Al-11Si alloy. Microcracks have occurred at the surface boundary of the intermetallic phase/α-Al solid solution. Also, without expanding into the surrounding α-Al solid solution, the fracture of particles of the intermetallic phase has influenced the fracture behaviour. 8. The T6 heat treating improves the ductility of the Al-12.5Si alloy. The increase in the fraction of ductile fracture has been caused by the appearance of few, easily broken primary silicon crystals. Despite the contribution of the massive compact particles with iron to the brittle intercrystalline fracture, microcracks initiated in them were less visible.  Figure 1. The temperature-time scheme of the T6 heat treatment procedure