**3.3. Aging**

62 Heat Treatment – Conventional and Novel Applications

single-step solution treatment of 495ºC/8h.

other alloying elements in the primary Al matrix.

concentration.

**3.2. Quenching** 

within the atomic lattice [33,34].

increase in quench rate [35-37].

quenching, with a consequent deterioration of the tensile properties. A two-stage solution heat treatment suggested by Sokolowski *et al*. [29] is reported to reduce the amount of the copper-rich phase in the 319 alloys significantly, giving rise to better homogenization prior to aging and improving mechanical properties. Also, Crowell *et al*. [30] stated that the blocky Cu phase in Al-Si-Cu alloys dissolves with increasing solution time at the recommended solution temperature of 495oC; also the rate of dissolution increases with Sr

A two-step solution treatment, namely, conventional solution treatment followed by a hightemperature solution treatment, as suggested by Sokolowski *et al*., [31, 32] is reported to reduce the amount of the copper-rich phase in 319 alloys significantly, thereby giving rise to better homogenization prior to aging and thus also to improvements in the mechanical properties. The holding time for the first stage and the solution temperature of the second stage are both significant parameters. Sokolowski *et al*. [32] studied the improvement in 319 aluminum alloy casting durability by means of high temperature solution treatment. Their results showed that a two-step solution treatment of 495ºC/2h followed by 515ºC/4h produced the optimum combination of strength and ductility compared to the traditional

Dissolve the micro-segregation of Mg and Si elements to form a supersaturated solid solution in the primary Al matrix in order to enable the formation of a large number of strengthening precipitates during subsequent natural and artificial ageing processes. Homogenize the casting, and attain a globular morphology of the eutectic Si phase to impart improved ductility and fracture toughness to the component. Reduce micro-segregation of

Following solution heat treatment, quenching is the next important step in the heattreatment cycle. The objectives of quenching are to suppress precipitation during quenching; to retain the maximum amount of the precipitation hardening elements in solution to form a supersaturated solid solution at low temperatures; and to trap as many vacancies as possible

The quench rate is especially critical in the temperature range between 450 °C and 200 °C for most Al-Si casting alloys where precipitates form rapidly due to a high level of supersaturation and a high diffusion rate. At higher temperatures the supersaturation is too low and at lower temperatures the diffusion rate is too low for precipitation to be critical. 4°C/s is a limiting quench rate above which the yield strength increases slowly with further

Faster rates of quenching retain a higher vacancy concentration enabling higher mobility of the elements in the primary Al phase during ageing. An optimum rate of quenching is necessary to maximize retained vacancy concentration and minimize part distortion after quenching. A slow rate of quenching would reduce residual stresses and distortion in the Age-hardening has been recognized as one of the most important methods for strengthening aluminum alloys, which involves strengthening the alloys by coherent precipitates which are capable of being sheared by dislocations [47]. By controlling the aging time and temperature, a wide variety of mechanical properties may be obtained; tensile strengths can be increased, residual stresses can be reduced, and the microstructure can be stabilized. The precipitation process can occur at room temperature or may be accelerated by artificial aging at temperatures ranging from 90 to 260oC.

After solution treatment and quench the matrix has a high supersaturation of solute atoms and vacancies. Clusters of atoms form rapidly from the supersaturated matrix and evolve into GP zones. Metastable coherent or semi-coherent precipitates form either from the GP zones or from the supersaturated matrix when the GP zones have dissolved. The precipitates grow by diffusion of atoms from the supersaturated solid solution to the precipitates. The precipitates continue to grow in accordance with Ostwald ripening when the supersaturation is lost. The length of each step in the sequence depends on the thermal history, the alloy composition and the artificial ageing temperature.

A Review on the Heat Treatment of Al-Si-Cu/Mg Casting Alloys 65

precipitates themselves, or both. The dislocations are then forced to cut through them or go around them forming loops. The preceding thus implies clearly that there are three sources for age hardening: strain field hardening, chemical hardening and dispersion hardening. Gloria *et al*. [58] investigated the dimensional changes occurring during the heat treatment of an automotive 319 alloy by means of T6 and T7 tempers involving solution treatment, quenching and artificial aging. They observed that increasing the solution temperature has the greatest influence in the dimensional change of samples due to dissolution of the Al-Cu(θ) eutectic phase. By increasing the aging temperatures, however, expansion is produced as a result of the transformation of the metastable phases

Shivkumar *et al*. [59] have studied the parameters which control the tensile properties of A356 alloy in the T6 temper. The improvement in the alloy strength has been attributed to the precipitation of negligible phases from a supersaturated matrix. The sequence of

The maximum alloy strength (peak-aging) is achieved just before the precipitation of the incoherent β-platelets. Apelian *et al*. [60] studied the aging behaviour of Al-Si-Mg alloys and observed that the precipitation of very fine β′-Mg2Si during aging leads to a pronounced improvement in strength properties. Both aging time and temperature determine the final properties, see Figures 6 and 7. Their study also established that increasing the aging temperature by 10oC is equivalent to increasing the aging time by a factor of two. The effect of natural ageing on precipitation can be justified as follows. A high concentration of quenched in vacancies enhances the rate of solute clustering in the early stages of natural ageing, and this clustering of solute leads to a reduced supersaturation of solute in the matrix. The solute clusters have a fine distribution within the matrix, and if they were to act as successful nuclei for the formation of β'' during subsequent artificial ageing, a fine precipitate distribution would result. Evidently, this is not the case, a reason being that many of the clusters are below the critical size for stability at artificial ageing temperatures. Furthermore, a lower solute supersaturation is expected to reduce the kinetics of precipitation. Thus, during artificial ageing, the dissolution of unstable clusters increase the solute concentration, while larger clusters that are stable remove solute by growing into GP zones that become nucleation sites for β''. Therefore, the solute supersaturation is maintained at a relatively low level during artificial; ageing and the density of the b'' is much lower than that occurring in alloys

The precipitation sequence for Al-Si-Cu-Mg alloys is similar, but more complex, as the Q'' phase and the θ' phase may also form. Cu can increase the fraction of the β'' phase formed, but it can also form the Q'' phase, which has a lower strength contribution compared to the

precipitation in Al-Si-Mg alloys, see Figure 5, can be described as follows:

iv. Equilibrium phase β-Mg2Si, FCC structure (a= 0.639), rod or plate-shaped.

i. Precipitation of GP zones, (needles about 10 nm long); ii. Intermediate phase β′′-Mg2Si, (homogeneous precipitation); iii. Intermetallic phase β′-Mg2Si, (heterogeneous precipitation);

into equilibrium phases.

without natural ageing.

The phenomenon of precipitation was originally discovered by Ardel in 1906 [48]. He found that the hardness of aluminum alloys which contained magnesium, copper, and other trace elements increased with time at room temperature, which was later explained by precipitation hardening. Over the years, much research was carried out to understand the aging kinetics of T4 and T6 heat treatments and to study the effects of underaging, peak-aging, and overaging on hardness [48-50], ultimate tensile strength, crack propagation behavior [51], and the cyclic stress-strain response of cast aluminum-silicon alloys [52].

The precipitation sequence for an Al-Si-Cu alloy, such as 319, is based upon the formation of Al2Cu-based precipitates. The Al2Cu precipitation sequence is generally described as follows: [53-55]

$$\alpha\_{ss} \to \text{GP Zones} \to \theta' \to \theta' \to \theta \text{ (Al:Cu)}$$

The sequence begins with the decomposition of the solid solution and the clustering of Cu atoms; the clustering then leads to the formation of coherent, disk-shaped GP zones. At room temperature aging conditions, GP zones arise homogeneously; these zones manifest as two-dimensional, copper-rich disks with diameters of approximately 3-5 nm. As time increases, these GP zones increase in number while remaining approximately constant in size. With regard to the Al-Cu alloys, as the aging temperature is increased above 100oC, the GP zones dissolve and are replaced by the θ′′ precipitate. This precipitate is a threedimensional disk-shaped plate having an ordered tetragonal arrangement of Al and Cu atoms; θ′′ also appears to nucleate uniformly in the matrix, and is coherent with the matrix in binary Al-Cu alloys. The high degree of coherency causes extensive coherency-strain fields to arise [56], giving peak strength to the material at this time.

As aging proceeds, the θ′′ starts to dissolve, and θ′ begins to form by nucleating on dislocations and/or cell walls [54,55]; θ′ also has a plate-like shape and is composed of Al and Cu atoms in an ordered tetragonal structure; θ′ loses coherency with the matrix, however, as it grows. Thus, since the long-range coherency-strain fields do not arise, a decrease in strength properties may be observed, while continued aging causes the equilibrium θ (Al2Cu) precipitate to occur. Tetragonal in shape, the θ phase is completely incoherent with the matrix; this fact, combined with its relatively large size and coarse distribution, reduces the strength properties significantly [56].

Increases in Cu were found mainly to reduce ductility and change the morphology of the Cu-containing phases [57]. The strength of an age-hardenable alloy is governed by the interaction of moving dislocations and precipitates. The obstacles in precipitationhardened alloys which hinder the motion of dislocations may be either the strain field around the GP zones resulting from their coherency with the matrix, or the zones and precipitates themselves, or both. The dislocations are then forced to cut through them or go around them forming loops. The preceding thus implies clearly that there are three sources for age hardening: strain field hardening, chemical hardening and dispersion hardening. Gloria *et al*. [58] investigated the dimensional changes occurring during the heat treatment of an automotive 319 alloy by means of T6 and T7 tempers involving solution treatment, quenching and artificial aging. They observed that increasing the solution temperature has the greatest influence in the dimensional change of samples due to dissolution of the Al-Cu(θ) eutectic phase. By increasing the aging temperatures, however, expansion is produced as a result of the transformation of the metastable phases into equilibrium phases.

Shivkumar *et al*. [59] have studied the parameters which control the tensile properties of A356 alloy in the T6 temper. The improvement in the alloy strength has been attributed to the precipitation of negligible phases from a supersaturated matrix. The sequence of precipitation in Al-Si-Mg alloys, see Figure 5, can be described as follows:

i. Precipitation of GP zones, (needles about 10 nm long);

64 Heat Treatment – Conventional and Novel Applications

alloys [52].

follows: [53-55]

atoms;

θ

equilibrium

As aging proceeds, the

θ

dislocations and/or cell walls [54,55];

precipitates grow by diffusion of atoms from the supersaturated solid solution to the precipitates. The precipitates continue to grow in accordance with Ostwald ripening when the supersaturation is lost. The length of each step in the sequence depends on the thermal

The phenomenon of precipitation was originally discovered by Ardel in 1906 [48]. He found that the hardness of aluminum alloys which contained magnesium, copper, and other trace elements increased with time at room temperature, which was later explained by precipitation hardening. Over the years, much research was carried out to understand the aging kinetics of T4 and T6 heat treatments and to study the effects of underaging, peak-aging, and overaging on hardness [48-50], ultimate tensile strength, crack propagation behavior [51], and the cyclic stress-strain response of cast aluminum-silicon

The precipitation sequence for an Al-Si-Cu alloy, such as 319, is based upon the formation of Al2Cu-based precipitates. The Al2Cu precipitation sequence is generally described as

> θ′′ → θ′ → θ

The sequence begins with the decomposition of the solid solution and the clustering of Cu atoms; the clustering then leads to the formation of coherent, disk-shaped GP zones. At room temperature aging conditions, GP zones arise homogeneously; these zones manifest as two-dimensional, copper-rich disks with diameters of approximately 3-5 nm. As time increases, these GP zones increase in number while remaining approximately constant in size. With regard to the Al-Cu alloys, as the aging temperature is increased above 100oC, the

θ

′′ also appears to nucleate uniformly in the matrix, and is coherent with the matrix

θ

θ

dimensional disk-shaped plate having an ordered tetragonal arrangement of Al and Cu

in binary Al-Cu alloys. The high degree of coherency causes extensive coherency-strain

however, as it grows. Thus, since the long-range coherency-strain fields do not arise, a decrease in strength properties may be observed, while continued aging causes the

incoherent with the matrix; this fact, combined with its relatively large size and coarse

Increases in Cu were found mainly to reduce ductility and change the morphology of the Cu-containing phases [57]. The strength of an age-hardenable alloy is governed by the interaction of moving dislocations and precipitates. The obstacles in precipitationhardened alloys which hinder the motion of dislocations may be either the strain field around the GP zones resulting from their coherency with the matrix, or the zones and

(Al2Cu)

′′ precipitate. This precipitate is a three-

′ begins to form by nucleating on

′ loses coherency with the matrix,

phase is completely

θ

′ also has a plate-like shape and is composed of Al

history, the alloy composition and the artificial ageing temperature.

α

GP zones dissolve and are replaced by the

θ

and Cu atoms in an ordered tetragonal structure;

distribution, reduces the strength properties significantly [56].

*ss* → GP Zones →

fields to arise [56], giving peak strength to the material at this time.

′′ starts to dissolve, and

(Al2Cu) precipitate to occur. Tetragonal in shape, the

θ


The maximum alloy strength (peak-aging) is achieved just before the precipitation of the incoherent β-platelets. Apelian *et al*. [60] studied the aging behaviour of Al-Si-Mg alloys and observed that the precipitation of very fine β′-Mg2Si during aging leads to a pronounced improvement in strength properties. Both aging time and temperature determine the final properties, see Figures 6 and 7. Their study also established that increasing the aging temperature by 10oC is equivalent to increasing the aging time by a factor of two. The effect of natural ageing on precipitation can be justified as follows. A high concentration of quenched in vacancies enhances the rate of solute clustering in the early stages of natural ageing, and this clustering of solute leads to a reduced supersaturation of solute in the matrix. The solute clusters have a fine distribution within the matrix, and if they were to act as successful nuclei for the formation of β'' during subsequent artificial ageing, a fine precipitate distribution would result. Evidently, this is not the case, a reason being that many of the clusters are below the critical size for stability at artificial ageing temperatures. Furthermore, a lower solute supersaturation is expected to reduce the kinetics of precipitation. Thus, during artificial ageing, the dissolution of unstable clusters increase the solute concentration, while larger clusters that are stable remove solute by growing into GP zones that become nucleation sites for β''. Therefore, the solute supersaturation is maintained at a relatively low level during artificial; ageing and the density of the b'' is much lower than that occurring in alloys without natural ageing.

The precipitation sequence for Al-Si-Cu-Mg alloys is similar, but more complex, as the Q'' phase and the θ' phase may also form. Cu can increase the fraction of the β'' phase formed, but it can also form the Q'' phase, which has a lower strength contribution compared to the β'' phase. The β'' phase is therefore preferred, rather than the Q'' phase. It is however not clearly stated when the Q'' phase forms at the expense of the β'' phase in cast alloys. For wrought alloys it has been shown that the fraction of the Q'' phase increases with natural ageing and artificial ageing time and temperature [64-66].

**Figure 6.** Sequence of phases found during age hardening of Al-Mg-Si alloys [60-62]. Supersaturated solid solution (SSSS) decomposes as Mg and Si atoms are attracted first to themselves (cluster) then to each other to form precipitates GP(I), sometimes also called initial-β''. GP(I) zones either further evolve directly to a phase β'' and then to a number of other metastable phases labelled β', B', U1, U2 (another one, U3, has been postulated theoretically), or first form an intermediate phase called preβ''.

A Review on the Heat Treatment of Al-Si-Cu/Mg Casting Alloys 67

**Figure 8.** TEM-BF micrographs show precipitated phases in association with T6 over-ageing period; (a) 100 hours, (c) 300 hours. (b) and (d) are corresponding SADPs from (Mg2Si) particles indicated by the

Designing an alloy and a heat treatment process for a material that meets specified requirements for a certain component can be facilitated by the use of models. Development of models can also help in the search for new alloys as knowledge is gained about the influence of a specific part of the microstructure on the alloy properties. The first model where the yield strength is coupled to the evolution of the microstructure during artificial ageing was developed by Shercliff and Ashby in 1990 [69]. They defined their model as a mathematical relation between the process variables (e.g. alloy composition, heat treatment temperature and time), and the mechanical response of the alloy (e.g. yield strength, hardness), based on physical principles (e.g. thermodynamics, kinetics of precipitation,

More refined models have been developed since then for prediction of yield strength [70,71] and elongation to fracture [72] after artificial ageing. To be able to model the tensile strength after heat treatment, the evolution of the microstructure has to be modelled from casting to artificial ageing. Empirical equations as the Hollomon's [73] and the Ludwigson's [74] and equations where the parameters are coupled to the microstructure as in the KM strain hardening theory can be used to describe the plastic deformation behavior. The KM strain hardening theory has already been successfully used to couple the plastic deformation behavior to the microstructure for heat treatable wrought alloys and Al-Si-Mg casting alloys.

arrows in (a) and (c), respectively.

strengthening mechanisms etc.).

**4. Modelling of the heat treatment process** 

**Figure 7.** TEM images of Al-Si-Mg alloy subjected to 2 different heat treatments. (a) solutionising and quenching, immediate aging at 180°C for 540 min, (b) solutionising and quenching, natural ageing for 10,000 min at 20°C, aging at 180°C for 540 min [63].

The precipitation of metastable Mg-rich phases depends on the Mg-to-Si ratio. The excess of Si in solid solution can significantly alter the kinetics of precipitation and the phase composition. In other words, equilibrium phases are enriched in Mg and metastable phases are enriched in Si. Silicon precipitates are observed if stable phases are formed [67,68].

**Figure 8.** TEM-BF micrographs show precipitated phases in association with T6 over-ageing period; (a) 100 hours, (c) 300 hours. (b) and (d) are corresponding SADPs from (Mg2Si) particles indicated by the arrows in (a) and (c), respectively.
