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Home / Review on Melt Treatment of Aluminium and Its Alloys

Review on Melt Treatment of Aluminium and Its Alloys

Open access peer-reviewed chapter

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

Submitted: December 8th, 2011 Reviewed: May 29th, 2012 Published: September 26th, 2012

DOI: 10.5772/50282

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1. Introduction

The trend of the automotive manufacture goes toward the construction of high-powered, comfortable, economical, ecological and safe vehicles. A few Al alloys containing Cu and a few containing Mg and Si are oestrus treatable in the cast condition due to the precipitation strengthening mechanisms. Ii of the major families of heat treatable aluminum alloys containing magnesium and silicon are the 6xxx series in wrought aluminum alloys, and 3xx series in casting aluminum alloys. Al-Si-Cu/Mg alloys are well studied and there exists a lot of publication virtually the effect of alloying elements and solidification rate on the microstructure formation [1-iii]. The influence of heat handling on the mechanical backdrop including hardness and tensile strengths is besides well studied, while the influence on plastic deformation beliefs and elongation to fracture is less studied.

Although the benefit of heat treatment is undisputed, there be several challenges for estrus treatment operators, including marketplace expectations of higher performance and reliability, lower production costs and energy use, every bit well as concern over ecology impacts. The heat treatment of age hardenable aluminum alloys involves solutionizing the alloys, quenching, and so either aging at room temperature (natural aging) or at an elevated temperature (bogus aging). The enhancement in mechanical properties subsequently thermal handling has largely been attributed to the germination of non-equilibrium precipitates inside primary dendrites during aging and the changes occurring in Si particles characteristics from the solution treatment. The historic period hardening response depends on the fraction size, distribution and coherency of precipitates formed. Al-Si-Cu-Mg alloys and Al-Si-Mg alloys generally have a high age hardening response, while Al-Si-Cu alloys take a slow and low age hardening response.

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2. Solidification process

During the solidification from a melt, chemic thermodynamics and kinetics are generally considered in terms of the enthalpy and Gibbs gratis energy changes, the solidification path, composition changes, and phase transformations etc.. Chemical thermodynamics describes the most stable phases at equilibrium weather condition (i.due east. temperature, pressure, compositions etc.) relating to only the initial and final states of a system. Accordingly, the solidification rate in a metallurgical system can be estimated past the enthalpy (H) and heat chapters (Cp), and how these thermodynamic properties reflect the system thermal state and estrus free energy requirements. Chemical equilibrium is controlled by the Gibbs free energy (G) of the system which is minimized for equilibrium conditions. In contrast, the dynamic system transformation between initial and final states controlled past chemic kinetics, indicates the path and phase changes of a chemical reaction in a arrangement when the limited atomic movement (i.e. in solids, depression temperatures, etc.) becomes ascendant in a short process time. Hence the solidification rate under a real fourth dimension status will be greatly influenced by the nucleation efficiency and the cantlet diffusion between phases [4].

The solidification charge per unit determines the coarseness of the microstructure including the fraction, size and distribution of intermetallic phases and the segregation profiles of solute in the α-Al phase. Big and brittle intermetallic phases grade during a slow solidification, which may initiate or link fracture, decreasing elongation to fracture. Additional, the defect size such as pore size, is also controlled to some extent by the solidification rate. The influence of defects on the elongation to fracture depends on their size, shape, distribution and fraction. Dendrite artillery with smaller radius may remelt into the molten into the molten liquid along with the decreasing full interfacial energy. The Ostwald-ripening effect on the germination of dendrite arm spacing (DAS) is determined by local solidification fourth dimension, allowing smaller particles to grow and merge into the larger ones due to the reduced total surface energy in the organisation. DSA, which is proportional to (average cooling rate)-n where n =ane/2 and 1/3 for the primary and secondary dendrites respectively, generally ranging from ten to 150 mm and which are controlled mainly past the solidification rate [5]. To gain an optimum holding of an alloy, the DAS therefore must be minimized and distributed homogeneously.

The major phases in as-bandage microstructure of Al-Si alloys are large size grains and primary α-Al, acicular eutectic Si, coarse principal Si, and also other harmful intermetallic phases such as needle like β-AlvFeSi, with uncontrolled and unevenly distributed porosities etc. [half dozen]. Table i summarizes the sequence of stage precipitation in hypoeutectic Al-Si alloys [vii]. Al in the eutectic has been reported to have mainly the same crystallographic features as the primary α-Al dendrites in unmodified alloys [8]. Figure 1.a indicates a basic construction of hypoeutectic Al-Si alloys consisting of grains (sizes at 1~ten mm in general), dendrites (typical DAS - 10~150 µm), and eutectic Si which can be in acicular shapes as long as 2 mm or circular particle as small every bit ane µm. The acicular Si might be chemically modified to a fibrous morphology past using effective modifiers. Heterogeneous nucleation should be the major arroyo to refine the grains, which nucleate on some of the foreign nuclei sites and grow slowly in the melt. Effective grain refiners, such as TiAl3 and TiBii, must match their lattice perfect coherently to the Al matrix with their lattice coherencies (Effigy 1.b). In contrast, particles with a poor lattice matching have little influence on increasing the nucleation of grains (Figure 1.c), resulting in an unrefined grain construction [9]. Typical examples of the microstructure of unmodified, Sr-modified, and Sb-modified alloys are shown in Effigy 2.

Temperature (C) Phases precipitated Suffix
650 Primary Al15(Mn, Iron)3Si2 (sludge) Pre-dendrite
600 Aluminum dendrites
and (Alfifteen(Mn, Fe)3Si2)
and /or AlfiveFeSi		 Dendritic
Postal service-dendritic
Pre-eutectic
550 Eutectic Al + Si
and Al5FeSi
MgtwoSi		 Eutectic
Co-eutectic
500 CuAl2 and more circuitous phases Post-eutectic

Tabular array 1.

Sequence of phase atmospheric precipitation in hypoeutectic Al-Si alloys [7]

Effigy 1.

Schematics of a) Three essential elements (grains, Al dendrites, DAS, and eutectic Si in a basic hypoeutectic Al-Si microstructure; b) Perfect grain refiner particles (squares) with one to one lattice matching to Al atoms (points); c) Poor lattice matching [9].

Figure 2.

Comparison of the silicon morphology in: (a) unmodified; (b) Sr-modified (300 ppm Sr); and (c) Sb-modified (2400 ppm Sb), hypoeutectic aluminum-silicon alloys [10].

Copper forms an intermetallic phase with Al that precipitates during solidification either as blocky CuAl2 or as alternate lamellae of α-Al + CuAl2 [11]. During solidification, in the presence of iron, other copper containing phases form, such as Cu2FeAl7 or Q-Al5Cu2Mg8Si6 [12]. The CuAltwo phase can be blocky shape or finely dispersed α-Al and CuAl2 particles inside the interdendritic regions, as shown in Figure 3. The presence of nucleation sites, such as FeSiAl5 platelets or loftier cooling rates during solidification can effect in fine CuAl2 particles [11]. The blocky CuAl2 phase particles are difficult to dissolve during solid solution heat handling, dissimilar the fine CuAlii stage particles that can dissolve within two hrs solid solution heat treatment [13]. Magnesium is present equally Mg2Si in Al-Si-Mg alloys if Mg is not in solution. Mg tin can besides form a true quaternary chemical compound Cu2Mg8Si6Al5 with other alloy elements in Al-319 alloy. In the absence of Cu, loftier Fe and Mg outcome in the appearance of π-FeMg3SiviAl8. The π phase is hard to dissolve during solid solution rut treatment [8].

Effigy three.

Cu-rich phases in as-cast 319 alloy: (a) Eutectic Al2Cu and (b) blocky AltwoCu [14].

A comparative written report of the mechanical properties of Al-Si-Cu-Mg alloys was carried out by Cáceres et al . [15] to investigate the effects of Si, Cu, Mg, Fe, and Mn, as well as solidification rate. The authors observed that increasing the Cu and Mg content by and large resulted in an increase in strength and a subtract in ductility, whereas an increased Atomic number 26 content (at an Fe/Mn ratio of 0.5) dramatically lowered the ductility and forcefulness of low-Si alloys. They also reported that the Cu + Mg content of the alloys determines the atmospheric precipitation strengthening and the volume fraction of the Cu-rich and Mg-rich intermetallics obtained.

Yi [16] adopt the enhanced solid diffusion coefficient of Cu in his model. The diffusion coefficient of Cu in α-Al phase is increased past 4-fold. The presence of Si-phase also has great influence on the diffusion of Cu in the matrix. It is assumed that the diffusion coefficient of Cu increases by xx-fold due to the presence of Si-phase. The distribution of Mg and Si across the dendrite arm spacing also changes due to the increase of Cu diffusion in the matrix. This is attributed to the alter of solidification path.

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3. Heat treatment of cast al alloys

Heat-treatment is of major importance since it is normally used to modify the mechanical properties of cast aluminum alloys. Heat-treatment improves the forcefulness of aluminum alloys through a procedure known as precipitation-hardening which occurs during the heating and cooling of an aluminum alloy and in which precipitates are formed in the aluminum matrix. The improvement in the mechanical properties of Al alloys equally a result of estrus handling depends upon the change in solubility of the alloying constituents with temperature. Effigy 5 shows the major steps of the rut handling which are normally used to improve the mechanical properties of aluminum. The alloy should commencement be solution treated at a temperature but below the eutectic temperature for long enough to allow solutionizing of the second phase. So it should be quenched to room temperature. Finally it should exist heated to a lower temperature to let precipitation. Tabular array 2 details a few of the more commonly applied heat treatments.

The T6 oestrus treatment is illustrated in Effigy 4 for an Al-Si-Cu alloy as an example. The evolution of the microstructure is shown; from (1) atoms in solid solution at the solution handling temperature, through (2) a supersaturated solid solution at room temperature after quench, to (3) precipitates formed at the artificial ageing temperature.

Figure four.

Diagram showing the three steps for precipitation hardening.

Handling Solution Quench Aging
T4 Yes Yes Room Temperature simply
T5 No No Elevated Temperatures
T6 Yes Yes Elevated (to yield increased strength)
T7 Yes Yes Elevated (to yield dimensional stability)

Tabular array 2.

Common aluminum heat treatment designations

Figure five.

The T6 heat treatment process [17].

three.1. Solution oestrus treatment

Solution heat handling must be applied for a sufficient length of time to obtain a homogeneous supersaturated construction, followed by the application of quenching with the aim of maintaining the supersaturated construction at ambient temperature. In Al-Si-Cu-Mg alloys, The solution treatment fulfils three roles: [18,19]

  1. Homogenization of equally-cast structure.

  2. Dissolution of certain intermetallic phases such as AliiCu and MgtwoSi.

  3. Modify of the morphology of eutectic silicon.

The segregation of solute elements resulting from dendritic solidification may have an adverse effect on mechanical properties. The time required for homogenization is determined by the solution temperature and past the dendrite arm spacing. Hardening alloying elements such as Cu and Mg brandish pregnant solid solubility in heat-treatable aluminum alloys at the solidus temperature; this solubility decreases noticeably as the temperature decreases.

The changes in the size and morphology of the silicon phase take a significant influence on the mechanical properties of the alloy. Information technology has been proposed that the granulation or spheroidization process of silicon particles through heat treatment takes place in two stages: (i) fragmentation or dissolution of the eutectic silicon branches and (ii) spheroidization of the separated branches [20]. During solution treatment, the particles undergo changes in size and in shape. In the initial stages, the unmodified silicon particles undergo necking and split into segments, which retain their original morphology. Every bit a result of the separation, the average particle size decreases and the fragmented segments are somewhen spheroidized. The spheroidization and the coarsening of eutectic Si tin occur meantime during the second stage.

The solution handling process needs to be optimized because also brusk a solution treatment fourth dimension means that not all alloying elements added will exist dissolved and made available for precipitation hardening, while likewise long a solution handling means using more energy than is necessary. The solution heat handling may be carried out in either a unmarried step or in multiple steps. Single-stride solution handling is normally express to about 495˚C, in view of the fact that higher temperatures lead to higher thermal stresses induced during quenching and the risk of the incipient melting of the Cu-rich phases [21-23]. This incipient melting tends to lower the mechanical properties of the casting. Solution treatment at temperatures of 495˚C or less, even so, is non capable of maximizing the dissolution of the copper-rich phases, nor is it able to modify the silicon particle morphology sufficiently. In Al-Si-Cu-Mg alloys having a depression magnesium content (0.5 wt.%), Ouellet et al . [24] used a solution temperature of 500oC because, at 505oC, fusion of depression melting betoken phases tin occur; Wang et al . [25], on the other hand, reported that, for a like blend with a solution temperature of 520ºC, mechanical properties increase without any observable localized melting.

Based on conventional solution treatment rules, the solution temperature of Al-Si-Cu-Mg alloys is restricted to 495oC, in lodge to avoid incipient melting of the copper-rich phase [26,27]. The time at the nominal solution handling temperature must be long enough to homogenize the alloy and to ensure a satisfactory degree of precipitate solution. In alloys containing high levels of copper, complete dissolution of the AliiCu phase is non usually possible. The solution time must and then be chosen carefully to allow for the maximum dissolution of this intermetallic phase, bearing in mind nevertheless, that solutions treating the alloy for long times are expensive and may not be necessary to obtain the required alloy force. Moreover, the coarsening of the microstructural constituents and the possible formation of secondary porosity which event afterward prolonged annealing at such temperatures can take a deleterious effect on the mechanical properties [28].

Studies by Gauthier et al . [19] on the solution heat handling of 319 blend over a temperature range of 480oC to 540oC, for solution times of upward to 24 hours, showed that the best combination of tensile force and ductility was obtained when the as-cast material was solution heat-treated at 515oC for 8 to 16 hours, followed by quenching in warm water at lxoC. A college solution temperature was seen to event in the partial melting of the copper phase, the formation of a structureless form of the phase and related porosity upon 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 rising 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; as well the rate of dissolution increases with Sr concentration.

A ii-step solution treatment, namely, conventional solution treatment followed by a high-temperature solution handling, equally 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 as well to improvements in the mechanical properties. The belongings time for the first stage and the solution temperature of the 2d stage are both meaning parameters. Sokolowski et al . [32] studied the improvement in 319 aluminum blend casting immovability past means of high temperature solution handling. Their results showed that a two-step solution handling of 495ºC/2h followed past 515ºC/4h produced the optimum combination of strength and ductility compared to the traditional unmarried-step solution treatment of 495ºC/8h.

Dissolve the micro-segregation of Mg and Si elements to class a supersaturated solid solution in the primary Al matrix in club 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 other alloying elements in the primary Al matrix.

3.2. Quenching

Following solution heat treatment, quenching is the next important pace in the estrus-treatment 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 within the atomic lattice [33,34].

The quench rate is especially critical in the temperature range between 450 °C and 200 °C for most Al-Si casting alloys where precipitates class chop-chop due to a loftier level of supersaturation and a high diffusion rate. At higher temperatures the supersaturation is also low and at lower temperatures the improvidence charge per unit is too low for atmospheric precipitation to be critical. 4°C/s is a limiting quench rate above which the yield strength increases slowly with farther increase in quench rate [35-37].

Faster rates of quenching retain a higher vacancy concentration enabling college 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 role distortion after quenching. A dull rate of quenching would reduce rest stresses and distortion in the components, notwithstanding, it causes detrimental effects such as precipitation during quenching, localized over-ageing, reduction in grain boundaries, increment tendencies for corrosion and result in a reduced response to ageing handling [38,39].

The best combination of strength and ductility is achieved from a rapid quenching. Cooling rates should be selected to obtain the desired microstructure and to reduce the elapsing time over certain disquisitional temperature ranges during quenching in the regions where diffusion of smaller atoms can lead to the atmospheric precipitation of potential defects [xl]. The effectiveness of the quench is dependent upon the quench media (which controls the quench charge per unit) and the quench interval. The media used for quenching aluminum alloys include water, brine solution and polymer solution [41-43]. Water used to exist the dominant quenchant for aluminum alloys, but water quenching most often causes distortion, cracking, and residual stress issues [44,45]. It has been reported that the water temperature affects the properties of the cast aluminum blend A356 subjected to T6 heat handling once the water exceeds 60-70oC, with UTS and YS being significantly more than sensitive than ductility. Detailed TEM investigations on A356 alloy, reported elsewhere [46], revealed that, at the tiptop-aged condition and with a water quench at 25°C, the α-Al matrix consists of a large number of needle-shaped and coherent β″-Mg2Si precipitates. The size of the precipitates is approximately 3 to 4 nm in diameter and 10 to xx in length. With a h2o quench at 60°C, they observed how the density of the precipitates decreases and the size of the precipitates increases slightly; at the same time a significant number of fine Si precipitates resulting from precipitation of backlog Si could exist observed in the α-Al matrix.

With a irksome quenching in air, very different atmospheric precipitation features are normally evidenced. By air quenching, the fabric remains at high temperatures for a longer period, which enhances the diffusion of silicon and magnesium. As well a high density of fine β″-Mg2Si precipitates, the α-Al matrix besides contained a large number of areas with fibroid rods β′-Mg2Si grouped parallel to each other [46]. While the first precipitates have an average size approximately ii to iii nm in diameter and around forty nm in length, the latter show an boilerplate size ~fifteen nm in diameter and 300 nm in length.

3.three. Aging

Age-hardening has been recognized as ane of the nigh important methods for strengthening aluminum alloys, which involves strengthening the alloys by coherent precipitates which are capable of beingness sheared by dislocations [47]. By decision-making the aging fourth dimension and temperature, a wide diversity of mechanical backdrop may be obtained; tensile strengths can be increased, residual stresses can be reduced, and the microstructure tin can be stabilized. The atmospheric precipitation process can occur at room temperature or may exist accelerated past bogus aging at temperatures ranging from 90 to 260oC.

After solution treatment and quench the matrix has a loftier 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 go along to abound 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.

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 atmospheric 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 furnishings of underaging, peak-aging, and overaging on hardness [48-fifty], 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 AliiCu-based precipitates. The Al2Cu precipitation sequence is generally described every bit follows: [53-55]

GP Zones → → → (Al2Cu)

The sequence begins with the decomposition of the solid solution and the clustering of Cu atoms; the clustering then leads to the germination of coherent, disk-shaped GP zones. At room temperature aging conditions, GP zones arise homogeneously; these zones manifest equally ii-dimensional, copper-rich disks with diameters of approximately 3-5 nm. Every bit 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 to a higher place 100oC, the GP zones dissolve and are replaced past the precipitate. This precipitate is a 3-dimensional deejay-shaped plate having an ordered tetragonal system of Al and Cu atoms; also appears to nucleate uniformly in the matrix, and is coherent with the matrix in binary Al-Cu alloys. The loftier degree of coherency causes extensive coherency-strain fields to arise [56], giving height strength to the cloth at this time.

Every bit aging proceeds, the starts to dissolve, and begins to form past nucleating on dislocations and/or prison cell walls [54,55]; as well has a plate-similar shape and is equanimous of Al and Cu atoms in an ordered tetragonal structure; loses coherency with the matrix, notwithstanding, as information technology grows. Thus, since the long-range coherency-strain fields exercise not arise, a decrease in strength properties may be observed, while connected crumbling causes the equilibrium (Al2Cu) precipitate to occur. Tetragonal in shape, the 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 institute 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 precipitation-hardened alloys which hinder the movement of dislocations may be either the strain field effectually the GP zones resulting from their coherency with the matrix, or the zones and precipitates themselves, or both. The dislocations are and then forced to cutting through them or go around them forming loops. The preceding thus implies clearly that there are 3 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 handling, 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. Past increasing the aging temperatures, however, expansion is produced equally a upshot of the transformation of the metastable phases into equilibrium phases.

Shivkumar et al . [59] take studied the parameters which control the tensile properties of A356 alloy in the T6 atmosphere. 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 five, can exist described as follows:

  1. Precipitation of GP zones, (needles nearly 10 nm long);

  2. Intermediate stage β′′-MgtwoSi, (homogeneous precipitation);

  3. Intermetallic stage β′-Mg2Si, (heterogeneous precipitation);

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

The maximum alloy strength (tiptop-crumbling) is achieved just before the precipitation of the incoherent β-platelets. Apelian et al . [60] studied the crumbling behaviour of Al-Si-Mg alloys and observed that the atmospheric precipitation of very fine β′-Mg2Si during aging leads to a pronounced improvement in forcefulness properties. Both aging fourth dimension and temperature make up one's mind the final backdrop, encounter Figures vi and 7. Their study also established that increasing the aging temperature by 10oC is equivalent to increasing the crumbling time by a factor of two. The event of natural ageing on atmospheric precipitation can be justified equally follows. A high concentration of quenched in vacancies enhances the rate of solute clustering in the early on 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 upshot. Patently, this is non 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 increment the solute concentration, while larger clusters that are stable remove solute by growing into GP zones that get 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, equally the Q'' phase and the θ' stage may too form. Cu can increment the fraction of the β'' phase formed, simply information technology can also form the Q'' phase, which has a lower forcefulness contribution compared to the β'' stage. The β'' phase is therefore preferred, rather than the Q'' phase. Information technology is even so not clearly stated when the Q'' stage forms at the expense of the β'' phase in bandage alloys. For wrought alloys it has been shown that the fraction of the Q'' phase increases with natural ageing and bogus 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 equally Mg and Si atoms are attracted kickoff to themselves (cluster) so to each other to form precipitates GP(I), sometimes too called initial-β''. GP(I) zones either further evolve directly to a stage β'' and then to a number of other metastable phases labelled β', B', U1, U2 (some other one, U3, has been postulated theoretically), or get-go grade an intermediate phase called pre-β''.

Figure 7.

TEM images of Al-Si-Mg alloy subjected to 2 different estrus 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 atmospheric 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 stage limerick. 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.

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four. Modelling of the heat handling process

Designing an alloy and a estrus treatment process for a fabric that meets specified requirements for a certain component can be facilitated by the use of models. Evolution of models tin can likewise aid in the search for new alloys equally knowledge is gained most the influence of a specific part of the microstructure on the alloy backdrop. The first model where the yield forcefulness is coupled to the evolution of the microstructure during artificial ageing was developed past Shercliff and Ashby in 1990 [69]. They defined their model as a mathematical relation between the procedure 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 (due east.g. thermodynamics, kinetics of precipitation, strengthening mechanisms etc.).

More refined models have been developed since then for prediction of yield forcefulness [70,71] and elongation to fracture [72] after artificial ageing. To be able to model the tensile strength after heat handling, the evolution of the microstructure has to be modelled from casting to artificial ageing. Empirical equations equally 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 draw the plastic deformation behavior. The KM strain hardening theory has already been successfully used to couple the plastic deformation behavior to the microstructure for oestrus treatable wrought alloys and Al-Si-Mg casting alloys.

The Scheil equation is a simple model giving off-white results for segregation profiles and fraction of particles formed during solidification for aluminum alloys. The Scheil equation assumes no improvidence in the solid and complete improvidence in the liquid [75]. The correctness of the predictions of the Scheil segregation model depends on the diffusivity of the alloying elements in the α-Al phase.

From the as-cast microstructure the time needed for dissolution and homogenization tin can exist modelled. The model developed past Rometsch et al. [76], which handles solution treatment of Al-Si-Mg alloys, is an example of a simple, merely efficient model. The evolution of the microstructure during artificial ageing involves nucleation, growth and coarsening. Two main approaches are used; precipitates having an average radius or precipitates having a size distribution. For the instance of precipitates which a size distribution, coupled nucleation, growth and coarsening tin exist calculated, while for an average radius growth is sequentially followed by coarsening.

The strength of an alloy derives from the power of obstacles, such equally precipitates and atoms in solid solution, to hinder the motion of mobile dislocations. The strength contributions from atoms in solid solution and from shearable and non-shearable precipitates modify during ageing, while contributions from lattice, dislocations and grain boundaries are constant. Small and not too hard precipitates are usually sheared by moving dislocations, see Figure 8.a. When the precipitates are larger and harder the moving dislocations pass the precipitates by bowing, leaving a dislocation ring around the precipitate, run across Figure 8.b. The strength of the precipitates increases with size as long equally information technology is sheared by dislocations. When dislocations pass the precipitates by looping, the blend strength decreases with increasing radius of the precipitates. Figure 8.c shows the different strength contributions to the total yield forcefulness for different ageing times.

Figure nine.

Dislocations passing a precipitate by a) shearing and b) looping (Orowan machinery) [77] c) Illustrates the unlike strength contributions to the total yield force

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five. Conculsions

Although many previous investigations into the thermal processing of Al-Si-Cu/Mg casting alloys have been carried out, most focus on a single aspect of the overall process and a comprehensive experimental study because all oestrus treatment stages is still required. This review shows that it is of vital importance to take the whole heat treatment process into consideration in guild to achieve the optimal mechanical backdrop of an alloy. It is non sufficient to consider only the solution treatment and the artificial ageing parameters. Furthermore, the development of process models for the prediction of microstructure and mechanical property changes in aluminum alloys has focused on wrought alloys, while casting alloys that contain more complex microstructures have been disregarded and the evolution of the solution treated microstructure and its influence on subsequent ageing behaviour has not been incorporated into the models.

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Acknowledgement

The authors would like to give thanks Professor Mariam Al-Maadeed, Head of Materials Technology Unit, MTU, at Qatar Academy for her help and support.

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