Controlled Diffusion Solidification (CDS) of Al-Zn-Mg-Cu (7050):Microstructure, heat treatment and mechanical properties

Reza Ghiaasiaan, Xiaochun Zeng, Sumanth Shankar n

Light Metal Casting Research Center (LMCRC), Department of Mechanical Engineering, 1280 Main Street West, JHE 316, McMaster University, Hamilton, ON,Canada L8S 4L7

a r t i c l e i n f o

Article history:Received 20 September 2013Accepted 26 November 2013Available online 4 December 2013

Keywords:Control diffusion solidificationAluminum wrought alloys7050MicrostructureHeat treatmentMechanical properties

a b s t r a c t

The tilt pour gravity casting process coupled with the Controlled Diffusion Solidification (CDS) processtechnology was employed to demonstrate the ability to cast AA7050 wrought alloy into high integritynear net shaped components with high strength and ductility. The CDS technology involves mixing twoprecursor alloys at different thermal mass and subsequently casting the resultant mixture into near netshaped cast components. The process enables casting of the high performance Al wrought alloys intonear net shaped components by circumventing the problem of hot tearing by obtaining a non-dendriticmorphology of the primary Al phase. This study presents the process and alloy parameters necessary forthe casting of 7050 Al wrought alloy (Al-Zn-Mg-Cu) using the CDS process technology. The uniaxialtensile properties after various heat treatment conditions such as as-cast, solutionizing and annealingmandatory to the development of the artificial ageing were investigated and presented along with in-depth and quantitative microstructural analyses.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Among all of the wrought and cast aluminum alloys, the 7xxxseries alloys (Al-Zn-Mg-Cu) and 2xxx series alloys (Al-Cu-Mg) arethe most applicable for structural automotive and aerospacecastings; in part because of their high strength to weight ratio[1] coupled with good ductility in operation. Typically, the 7xxxseries Al wrought alloys in their as-cast condition do not offertheir best mechanical properties and performance due to inherentmicrostructural deficiencies such as the coarse intermetallicphases, coarse grains, significant elemental micro-segregation,solid solubility limitations and, above all, macro-segregation andhot cracking (due to the inability of the solidifying liquid metal toaccommodate the strain field imposed by contracting solidfraction) [2–4]. Among these inherent casting defects, the hottearing or hot cracking is predominantly responsible for theinability to manufacture near net shaped cast components fromthese alloys. In theory, the hot tearing defect could be overcomeduring solidification by two methods; significant refinement of theSecondary Dendrite Arm Spacing (SDAS) of the primary Al phaseand by altering the morphology of this phase to a non-dendriticone. The former objective is achieved by one or many of precisecontrolled process parameters such as pouring temperature(superheat), cooling rates (design of cast part) and grain

refinement by alloying elements; all of which reduce the compli-cation associated with the solidification network of dendrites ofthe primary phase, which in turn would allow the solidifyingliquid phase to better accommodate the strain field gradientsbrought about by the shrinking solid fraction [5,6]. The morpho-logical modification of the primary Al phase in the solidifiedcasting could be achieved by employing one of the several Semi-Solid Metal (SSM) processes such as thixoforming, thixocastingand rheocasting; all of which are carried out by using external ornatural forces to physically or thermally alter the morphology ofthe solidifying Al phase during the intermediate stages of solidi-fication [7–9]. However, these SSM processes are not commerciallyfavored by industry due to the complexity and cost of theseprocesses coupled with the lack of repeatability and reproduci-bility of the commercial processes [10,11]. Further, some research-ers [12,13] have argued that the difficulties associated with castingof these alloys are probably related to their fluidity and viscositylimitations, as well. Thus, the Al shaped casting industry remainsunable to fully benefit from the superior properties of Al7xxxwrought alloys. All the alloy compositions in this publication willbe in weight percentage of the respective elements except other-wise mentioned.

2. Background

Controlled Diffusion Solidification (CDS) is a novel and innova-tive process that enables Al wrought alloys to be cast into a near

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n Corresponding author. Tel.: þ1 905 525 9140x26473; fax: þ1 905 572 7944.E-mail address: shankar@mcmaster.ca (S. Shankar).

Materials Science & Engineering A 594 (2014) 260–277

net shaped product [14] by circumventing the typical castingproblems of hot tearing during solidification by favorably alteringthe morphology of the primary Al phase to non-dendritic. Fig. 1(a) and (b) illustrates one of common casting defects of aluminumwrought alloys (2xxx and 7xxx series) known as “hot tearing” inboth its macro- and micro-scale appearances, respectively [14,15].

The CDS process involves mixing two precursor liquid alloyshaving different thermal mass (temperature and solute concentra-tion) and subsequently casting the alloy mixture into a shapedmold [15]. Fig. 2 shows a schematic illustrating the CDS process;the Alloy3 is the desired final alloy composition which is obtainedby mixing the precursor Alloy1 at a specific temperature into theAlloy2 at a different and lower temperature and subsequently castinto the product with high integrity.

2.1. CDS prior art: novelty and mechanism

In the later part of the last century, Apelian et al. [16] employedthe concept of diffusion solidification for rapid-cycle casting ofsteel. Recently, Saha et al. [17] developed a variation of thediffusion solidification in the CDS and cast the Al–4.4Cu with anon-dendritic primary phase morphology by mixing the pure Al(high thermal mass precursor alloy) at a temperature of 938 K (5 Ksuperheat) and eutectic Al–33Cu alloy (low thermal mass pre-cursor alloy) at a temperature of 823 K. Subsequently, the CDSprocess was successfully employed with the 2014, 4145, 5056 and7050 Al wrought alloys, and 222 and 319 aluminum cast alloys toobtain a nearly non-dendritic morphology of the primary Al phasein the solidified conditions. As the superheat of the Alloy1

Nomenclature

Alloy1 Pre-cursor alloy with higher thermal mass (highertemperature and higher mass)

Alloy2 Pre-cursor alloy with lower thermal massAlloy3 Resultant mixed alloy

TL1, TL2 and TL3 Liquidus temperature of Alloy1, Alloy2 andAlloy3, respectively

T1, T2 and T3 Melt temperature of Alloy1, Alloy2, and Alloy3,respectively

m1 and m2 Mass of Alloy1 and Alloy2 respectively

Fig. 1. Typical images of hot tearing in the as cast condition in both macroscopic and microscopic scales for (a) Al-Mg-Cu (AA2014) [14] and (b) Al-Zn-Mg-Cu (AA7050)showing the hot tearing, respectively.

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increases the primary aluminum phase becomes more dendritic[17,18]. Apelian et al. set a theoretical framework for the CDSprocess by specifying three considerations: thermodynamics,liquid mixing and solidification [19]; Symeonidis [15], subse-quently, investigated the CDS process with Al–4.5Cu and postu-lated a three-stage mechanism for the process: mixing, nucleation,and growth. Symeonidis' three-stage mechanism was later revisedby Khalaf et al. as shown in Fig. 3 [14]: (I) mixing stage, (II)re-distribution of thermal and solute fields, and nucleation, and(III) growth of stable nuclei.

Khalaf et al. [14] succeeded in obtaining non-dendritic Almorphology of various aluminum wrought alloys series such as2024, 6082, 7005 and 7075 using the CDS process in laboratoryscale experiments. Moreover, he proposed a set of conditions for asuccessful CDS process [14,20–22]: (a) the difference between theliquidus temperature of the precursor alloys before mixing shouldbe more than 50–80 K, (b) the maximum temperature attainedduring mixing of the two alloys should preferably be more thanthe liquidus temperature of the resultant alloy to facilitate com-plete filling of the casting mold, (c) the mass ratio between thetwo precursor alloys should be at least 3. The three-step

mechanism proposed by Khalaf et al. [14,20] could be brieflydescribed as below and with reference to Fig. 3(a).

2.1.1. Segment AB (stage I)At this stage of “mechanical mixing”, the Alloy1 continuously

entering the Alloy2 will break down into small volume liquidpockets (at temperature T1) in the resultant mixture (at a tem-perature less than T1) to form the resultant Alloy3. The exposure ofhot liquid of Alloy1 to the low temperature liquid mixture in thisstage of mixing results in the nucleation of the broken liquidpockets of Alloy1 shown by the morphology of the “AREA 1” inFig. 3(b). The “AREA 1” nuclei forming in the AB segment stage ofFig. 3 have stable growing interface (as opposed to perturbedinterface which typically leads to a dendritic structure) and that isbecause the solute back diffuses toward the solid liquid interface,which is opposite to what occurs in a conventional solidificationprocess wherein the solute piles up in front of the growingsolidification front; as shown comparatively in Fig. 4 [14,18].

In Fig. 4, the two thermal graphs shows the negligibleconstitutional supercooling for the CDS process when compared

Fig. 2. Illustration of the Controlled Diffusion Solidification (CDS) process technology; (a) alloy phase diagram, (b) schematic of process and (c) typical product in ABS brakehousing [13].

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to the typically significant one in the conventional solidificationprocess; caused by the reversal of the diffusion paths of the solutein the former and leading to a different solute redistribution field.

2.1.2. Segment BCD (stage II: (a) and (b))At this stage of “re-distribution of the thermal and solute fields in

the resultant mixture”, convective cells (which are akin to Bernardcells [23]) form in the mixture where the nucleation of the grainsshown as “AREA 2” in Fig. 3(b) takes place. These hexagonal cellsresemble a plate-like colloidal phase of Alloy1 (impoverished insolute) enclosed by the liquid Alloy2 (enriched in solute). TheseAREA2 nucleation sites act similarly to those created in stage AB(AREA1), in that the solute back diffuses from Alloy2 in the walls ofthe cells toward the central part which is predominantly Alloy1; ina direction opposite to the heat extraction. Nucleation in thesecells occurs when the actual temperature of the respective solutecomposition region is undercooled below the liquidus tempera-ture [14,18].

2.1.3. Point D nucleation (stage III)At this stage the final nucleation events in CDS process occur.

At this stage, as the temperature field and solute concentrationfield of the residual liquid have reached a nearly homogenizedstate, where they can nucleate and grow as “AREA3”, which issimilar to conventional solidification with constitutional under-cooling resulting in growing interface instabilities; however, thespatial constraints existing during solidification does not allow forgrowth of large dendritic networks but rather a mildly perturbedset of primary phase grains similar to those observed in the centralequiaxed region of a Direct Chill (DC) cast ingot [24,14,18].

2.2. 7xxx series Al wrought alloys

High strength 7xxx aluminum alloys (Al-Zn-Mg-Cu) are com-mercially used in the form of plates, bars and rods in wroughtcondition. These products undergo a number of processing stagesnamely ingot casting, homogenization, preheating, hot rolling,solutionizing and ageing. Precipitation and dissolution of secondphase precipitation occurring during each step impact the micro-structural evolution in the ensuing stages. In order to design anoptimal and pragmatic process, it is vital to understand the originof microstructural evolution appearing in the final product [25].

Fig. 3. The three stages in the mechanism of the CDS technology. (a) Schematic ofthe typical thermal profile observed in the CDS process with three distinct stagesstarting with mixing of alloys at point A, and (b) the three morphologies of theprimary Al phase which are associated with the three stages shown in (a) for theAl–4.5Cu alloy [13].

Fig. 4. Comparative schematic solute and temperature redistribution regimes ahead of the solidifying solid–liquid interface are presented for (a) the CDS process and(b) conventional solidification [19].

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2.3. Microstructural evolution during post homogenization coolingin 7xxx Al series alloys

Precipitation during intermediate stages of the processing of7xxx aluminum series wrought alloys usually occurs during thenon-isothermal condition; this precipitation in the high strength7xxx alloys during cooling from the solution treatment tempera-ture has been studied by many researchers [26]. These non-isothermal precipitations are industrially important because theyremove solute from the supersaturated solid solution matrix andform coarse precipitates which have a detrimental effect on thesubsequent age-hardening response of the material; these pre-cipitates nucleate on heterogeneous sites such as grain boundaries,sub-grain boundaries, dispersoids and dislocations [25].

In the complex alloy system of Al-Zn-Mg-Cu, several precipita-tion phases could evolve, namely, the η phase (MgZn2), M/Sigmaphase Mg(Zn,Cu,Al)2, T phase {(Al2Mg3Zn3) or Al32(Mg,Zn)49},S phase (Al2CuMg), θ phase (Al2Cu), irregularly shaped ironcontaining intermetallic of Al7Cu2Fe (bulk-shape), needle/plateshape iron containing intermetallic of Al3Fe (or Al13Fe4) andMg2Si; some of which could evolve below the solidus temperature[10,27–29]. The hexagonal structured phase of η or M are mostlyobserved in the as-cast microstructures whereas the orthorhombicstructured phases of S and T phase are more common in solidsolution states with the extended composition ranges containingall four elements, i.e., Al, Zn, Mg, and Cu [10,25,30].

2.4. Strengthening mechanism in different ageing conditions for 7xxxAl series alloys

Al-Zn-Mg-Cu series alloys are typically age-hardening alloysstrengthened by precipitation reactions; the ageing treatment is acritical process to achieve desired microstructures and resultantproperties. The peak hardness and tensile strength could be obtainedwith a T6 type of temper; however, this increases the susceptibility forstress corrosion cracking (SCC) in the component. In the one-stepageing processes for pre-stretched thick plate1 with two differentageing temperatures, the peak ultimate strengths of 595 and 575MPacould be attained after 22 and 6 h at 393 K and 408 K, respectively.The over-aged temper condition (T7) increases corrosion resistance bymodifying the microstructures which are gained at the price ofsacrificing 10–15% of the tensile strength compared to the peak-agedcondition (T6 temper) [31]. The AA 7075 alloy is one of the mostpopular 7xxx Al wrought series alloys which has been the subject ofmany studies during recent years because of (a) its quench sensitivity(i.e., the reduction in age hardening capacity due to low quenchingrate), (b) its high susceptibility to SCC in T6 temper, and (c) its strengthdeterioration through the wall thickness of the plate. With slightchemistry modifications in the AA7075, the resulting AA7050 over-comes some of these limitations [7]. The mechanical properties ofAl-Zn-Mg-Cu alloys make them a good choice for the force bearingcomponents in structures. Their highest mechanical properties couldbe achieved by forging coupled with the T6 temper condition [32]. Forsemi-solid metal (SSM) products of Al-Zn-Mg-Cu aluminum alloys,Chiarmett reported themechanical properties in the T6 condition havea tensile strength of 405MPa and elongation of 6.6%. Lu et al. reportedthat the T6 treatment mechanical property of SSM 7075 Al alloy has atensile strength of 474MPa. Comparing the reported properties of SSM7xxx Al alloy in T6 condition with conventionally forged wroughtcondition that has a UTS of 570MPa and elongation of 11%, it isnoticeable that the mechanical properties of SSM products, especially

in performance (ductility), are very poor. The main difference betweenSSM and conventional forging products of 7xxx series of Al alloys isthe difference in elemental content and phase composition within theSSM billet; Table 1 tabulated some of mechanical properties of AA7050in different commercial process conditions [31–35].

The main objective of this study was to investigate thefeasibility of manufacturing near net shaped components of Al7050 wrought alloy using the CDS process with industrial facilitysuch as the tilt pour gravity casting machine and further evaluatethe tensile mechanical properties of the as-cast, T4 and T6 heattreated samples. Furthermore, an intensive microscopic study wasconducted in order to characterize the precipitation formationsduring solutionizing heat treatment and non-isothermal annealingprocedures on the CDS cast products of the AA7050 alloy.

3. Materials and experimental procedure

Table 2 presents the nominal composition of the Al 7050 usedin this study; Si and Mn are impurity elements with maximumallowable limits and Cr is typically added to enhance the grainboundary strengthening effects that are possible in a solid-statetransformation processes. Therefore, Si, Mn and Cr were notspecifically added in the alloy used for the CDS process.

All the precursor alloys used in this study to cast componentsof the AA7050 resultant alloy were freshly prepared usingAl master alloy ingots and several raw material such as Al–50Mg,Al–28Cu, Al–33Cu, Al–36Si, Al–25Mn, Al–50Fe, TiBor and Pure Zincingots. The Ti was added to the Alloy1 to as to effect grainrefinement of the primary Al phase to further improve themechanical properties of the resultant castings. In casting withthe CDS technology, the final resultant alloy is prepared by mixingtwo precursor alloys with different elemental composition. Table 3presents the elemental composition of the two precursor alloysand the final desired AA7050 alloy in this study; it is notable thatthe range of elemental composition presented in this table denotesthe range obtained during several experiment and it was ensuredthat the composition of the final AA7050 alloy in the castcomponents did not deviate from the nominal shown in Table 2.

This CDS experiment was initially carried out in a laboratoryscale to optimize the process parameters such as superheatstemperatures of the initial precursor alloys and mass ratiobetween them. The experiments was then replicated on anindustrial scale in an actual foundry2 using a tilt pour gravitycasting machine (TPGC) with cavity mold (containing the ASTMstandard tensile bars), making use of the process parametersoptimized in the laboratory experiments.

Fig. 5 shows the TPGC machine used to cast the CDS samplesfor this study along with a list of the different operationalcomponents of the machine and the metal mold used in themachine to cast the two tensile and one fatigue test bar specimenwhich were designed according to the ATM standard numbersASTM B557 and ATM E466-96, respectively.

The freshly prepared pre-cursor alloys were firstly degassed intwo respective electric furnaces separately with ultrahigh purityArgon gas purged into the molten alloys at a flow rate between6 and 8 L min�1 using a rotary degasser at 120 RPM for 30 min.The alloy melts were then held at about 10 K superheat tempera-tures above their respective liquidus temperatures. The initialprecursor alloys were mixed at their specific pouring temperaturesinto the preheated metal mold maintained at 648 K, using twofixed volume ladles. The fixed volume ladles were designed to

1 Traditionally ingot cast slabs were, homogenized, scalped, hot-rolled to platesof 40 mm in thickness, solid solution treated at 743 K, water quenched (a roller-type spray quenching equipment was used) to room temperature, andpre-stretched (residual stress relieving) [30].

2 Orlick Industries Limited, 411 Parkdale Avenue North Hamilton, Ontario,Canada L8H5Y4, www.orlick.on.ca.

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maintain a constant mass ratio between the initial precursoralloys. For instance, in our case of AA7050 the mass ratio, mr,between Alloy1 to Alloy2 was three. The gross weight of finalcasting shot in the TPGC machine is about 900 g.

As a benchmark, the AA7050 alloy was cast directly using theTPGC equipment without the CDS process and will be referred toas conventional casting throughout this publication. The alloy meltof AA7050 was maintained at 75 K above the liquidus temperatureof 988 K in an electric resistance furnace and the melt wasdegassed with Ar gas through a rotary degasser prior to casting.

The castings from the TPGCmachinewere subjected to the uniaxialtensile test provided that they can pass through a mandatory inspec-tion stage, including a visual inspection for macro-scale hot tearing onthe cast sample surface andmicrostructural observations to ensure thenon-dendritic morphology of the primary Al phase. Notable, all thecastings in this study were deemed successful without any hot tearingphenomena and a non-dendritic primary Al phase in the microstruc-ture. The uniaxial tensile test machine used was the Instron 8800 withan MTS Frame equipped with a 250 kN Model 312 MTS Load cellcoupled with an on-line Extensometer of 50mm gauge and connectedto an on-line data acquisition system; all uniaxial tensile tests werecarried out a load speed of 1 mm/min.

Microstructure evaluation of the sample sections were carried outusing Nikon light optical microscope and Scanning Electron Micro-scope Model JEOL 7000. The samples for microstructural evaluationwere sectioned from the gauge length of the tensile test bar castings.Various heat treatment procedures were carried out on these castings:solution treatment (T4), annealing (O) and age-hardening (T6), anddetails of the heat treatment procedures are presented in Table 4.

The heat treatment procedures were carried out in an electricresistance convection furnace wherein, an internal fan maintainedthe constant thermal condition to within 2 K of the set temperatures.The temperature was set at 750 K which is about 6 K below the meansolidus temperature which measured separately by thermal analysisduring the solidification of the re-melted and typical resultantcasting of the AA7050 alloy cooled at an average rate of 0.1 K s�1.Samples from each casting were sectioned from the in-gate andriser areas to be re-melted to obtain thermal curves during solidifica-tion at 0.1 K s�1; enabling a separate method to verify the finalalloy composition, as well. Additionally, the Differential Scanning

Calorimeter (DSC) experiments were conducted at a heating rate of0.3 K s�1 on two representative samples from the castings for thepurpose of validating the solidus temperature prior to heat treatment;the solidus temperature was obtained as 755.7 K as shown in thetypical graph in Fig. 6 from the DSC experiments and the result wasfurther validated with data from the literature [25,36].

The SEM was used for both the topographical studies on thefracture surface and phase characterization along with the quan-titative EDS measurements.

4. Results and discussion

Several isopleth of the multi-component phase diagram of theAl-Zn-Mg-Cu system were investigated in order to select the mostappropriate compositions and initial temperatures of Alloy1 andAlloy2, respectively, such that the temperature difference betweenthe liquidus of initial alloys, TL1 and TL2, is greater than 328 K [14].Mixing of Alloy1 into Alloy2 yields the desired composition of Alloy3at a temperature T3 near the value of the liquidus temperature, TL3.The initial temperatures of Alloy1 and Alloy2 were typically about 5–10 K above the respective liquidus temperatures of TL1 and TL2.

In order to obtain the nominal composition of AA7050, twoprecursor alloys, Alloy1 and Alloy2, were developed using thermo-dynamic phase diagrams. A mass ratio of 3 between Alloy1 andAlloy2 was selected for the CDS casting parameters in thelaboratory experiment. The required weight of Alloy3 (AA7050)was 290–340 g adjusted by the laboratory furnace sizes, thecrucible sizes and the density of the AA7050 alloy. The favorablephase diagram isopleths of the Al-Zn-Mg-Cu (AA7050) simulatedby Pandat database3 is presented in Fig. 7, and the experimentalthermal curves obtained during solidification at 0.1 K s�1 for thetwo precursor alloys and the resultant alloy after the CDS processare presented in Fig. 8. The thermal data during solidification forthe AA7050 alloy was obtained by remelting the castings from theCDS casting trials and casting again with the TPGC process. The

Table 1Typical tensile mechanical properties of Al 7050 wrought alloys by various processing routes.

Alloy/heat treat Process/product UTS MPa (ksi) YS MPa (ksi) El% Ref.

7050/T7651 12.7 mm (0.5 in.) thick plate 524 (76) 455 (66) 8 [33]7050/T7451 150 mm hot-rolled thick plate 496 (71) 437 (63) 12.4 [34]7050/T6 Forging 570 (82) 469 (68) 11 [35]7050/T6 (393 K–22 h) Pre-stretch hot-rolled thick plate 595 (86) – – [31]7050/T6 (408 K) 575 (83) – –

7050/T6 Semi-solid thixoforming 405 (59) – 6.6 [32]

Table 2Nominal composition (wt%) of Al 7050 alloy.

Alloy/element Zn Mg Cu Fe Ti Si Mn Cr Al

7050 5.7–6.7 1.9–2.9 2–2.6 o0.15 o0.06 o0.12 o0.1 o0.04 Balance

Table 3Elemental composition as evaluated by GDOES for the precursor and the resultant alloy.

Alloy type Zn Mg Cu Fe Si Ti Al

Alloy1 – 2.65-2.97 – – – 0.13–0.15 BalanceAlloy2 23.03–24.06 – 8.23–8.82 – – – BalanceAlloy3 (AA7050) 5.66–6.09 2.11–2.32 2.13–2.49 0.10–0.14 0.18–0.2 0.03 Balance

3 PanAluminum Database; Computherm LLC., Madison, WI, USA; http://www.computherm.com/.

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liquidus temperature and corresponding choice of pouring tem-perature (TL1 and T1) for this alloy with a mass ratio of three arepresented in Table 5.

The resultant castings from the TPGC machine were consideredsuccessful and sound if the following conditions were met:

� No visual hot-cracking or hot-tearing on the cast component.� No visual shrinkage or defect feature on the surface of the cast

test bars.

� A non-dendritic morphology of the primary Al phase is presentin the microstructure obtained from the cross-section of thegauge in the tensile test bar.

� Reasonably sound tensile properties of the as-cast samples.� Reasonably compact features in the optical low magnification

micrograph of the fracture surface of the tensile bar.

The typical composition of Alloy1, Alloy2 and Alloy3 for the CDScasting trials as measured with the Glow Discharge Optical

Fig. 5. Industrial casting equipment for the study, (a) Tilt Pour Gravity Casting Equipment (TPGC) and (b) metal mold cavity of tilt pour gravity casting (TPGC) machineconsisting of two tensile test bars and one fatigue test bar according to ASTM standard numbers ASTM B557 and ASTM E466-96, respectively [21].

Table 4Typical heat treatment temper conditions used on the cast samples of tensile test bars.

Temper Solutionizing Natural ageing Artificial ageing

F (as-cast) None 496 h NoneO (anneal) Soaked @ 413 1C for 2 h ,then cool

down to 150 1C @ the rate of 20 1C/hNone None

T4 496 h @ RT NoneT6 Soaked @ 477 1C for 24 h and

quenched in water at 25 1CNone A combination of soaking @

120 1C for 6 h and @ 180 1C for12 h.

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Emission Spectroscopy (GDOES) are presented in Table 5, alongwith the liquidus temperatures of the respective alloys as measuredby thermal analysis during the solidification of these individualalloys as in Fig. 8 and the corresponding temperatures obtainedfrom the thermodynamic Pandat database presented in Fig. 7.

Fig. 9 shows the photographs of typical castings of AA7050 alloyobtained from the CDS and conventional casting processes, respec-tively; wherein, the castings from the CDS process were sound withthe absence of any discernible casting defect such as hot tears asshown in Fig. 9(a), while the castings from the conventional processwere consistently defective, specifically, due to the presence ofseveral hot cracks in the castings as illustrated in Fig. 9(b).

Optical micrographs in low and high magnifications arepresented for the CDS cast samples and the conventional cast

samples in Figs. 10 and 11, respectively; notably, both the castingshave identical chemical composition and both were cast using thesame TPGC process. The casting defect of hot cracking, in micro-scopic scale, is evident in Fig. 11 for the conventionally cast samplewhile absent in Fig. 10 for the samples cast by the CDS technology.Additionally, the non-dendritic morphology of the primary Alphase is notable in Fig. 10 while the morphology in Fig. 11 ispredominantly dendritic. In Fig. 10, the hot cracking or hot tearingdefect had been mitigated in the cast sample microstructure dueto the non-dendritic morphology of the primary Al phase whichsignificantly improves the feedability of the liquid during

Fig. 6. Typical results from the Differential Scanning Calorimeter (DSC) experi-ments of solidification of the AA7050 casting at a heating rate of 0.3 K s�1.

Fig. 7. Typical isopleths of multi-components phase diagrams representing the two initial precursor alloys and the final desired alloy as simulated using the Pandat softwarewith the PanAl8 elemental database. (a) Al–Mg system showing Alloy1 (Al–2.8Mg), (b) Al–Zn–Cu systems showing the Alloy2 (Al–23.8Zn–8Cu) and (c) Al-Cu-Mg–Zn systemshowing the resultant Alloy3 (Al–6Zn–2.18Cu–2.1Mg) or AA7050.

Fig. 8. Typical thermal data obtained during solidification of the two precursoralloys and the final desired AA7050 alloy re-melted from the cast material.

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277 267

solidification and alleviates the high strain fields in the micro-structure that typically causes hot tearing during solidification.

Figs. 12 and 13 presents the typical microstructure of the CDS castsamples obtained from the center of the gage section in a tensile testbar in as-cast and T4 condition, respectively; wherein the predomi-nantly non-dendritic morphology of the primary Al phase is observed;the light grey area is the non-dendritic primary Al matrix and the darkgrey coupled with the nearly black phases are the eutectic phases.

In Fig. 13, the morphology of the primary Al phase is non-dendriticafter heat treatment to the T4 temper; the heat treatment resulted insignificant solute redistribution resulting in a nature and morphologyof the solidified phases shown in Fig. 12 and the details of thesechanges to the phases will be further analyzed and discussed in thesubsequent sections of this manuscript. It is noteworthy that the lightgrey areas in high magnification images of Fig. 12(b) are soluteelements micro-segregation, also known as coring, within the grainarea of the as-cast microstructure which would have been leveled outafter solutionizing treatment shown in Fig. 13; moreover, the darkblack areas in Fig. 12 are eutectic phases which will be furtheranalyzed and discussed in more details by electron microscopy results(SEM) in the subsequent sections of this manuscript.

4.1. Mechanical property of AA7050 CDS castings

Table 6 and Fig. 14 present the average uniaxial tensile proper-ties along with the respective standard deviations for the AA7050alloy cast using the CDS technology; in F (as-cast), T4 and T6conditions. Fig. 14 shows a typical graph obtained during thetensile curves of one of the cast samples.

The uniaxial tensile properties for the CDS process shown inTable 6 are in the similar regime of magnitude as that of otherproducts manufactured by this alloy as shown in Table 1. However, theelongation obtained in the samples in Table 6, after the T6 temper islower than those obtained for products with AA7050 alloy manufac-tured by solid state transformation processes such as forging andextrusion. The T6 temper conditions used in this study were derived

from the background literature for this wrought alloy; the samples inthis study are near net shaped castings and hence, there could be asignificant variation in both the solute redistribution and texturing ofthe primary Al grains in the microstructure of these samples in the ascast and heat treated conditions when compared to their counterpartproducts manufactured by conventional forging or extrusion processeswhich would necessitate a redesign of the heat treatment process tooptimize the process parameters in order to maximize the tensileproperties, specifically the elongation at failure. This optimization ofthe heat treatment for the products cast using the CDS technology iscurrently underway. It is notable that the tensile properties presentedin Table 6 are not optimized and merely presents a first-off snap shotof the properties for these castings manufactured by the CDStechnology. Optimization of the heat treatment process for thecastings in this study would require a critical understanding of themicrostructure evolution, nature of solute redistribution and precipita-tion of strengthening phases during solidification and subsequentstages of the heat treatment process.

4.1.1. Annealing treatment on AA7050 CDS castingsThe annealing heat treatment (O temper) was carried out to

provide a better understanding of the microstructural changesfrom the as-cast condition to the heat-treated one. Carrying outthe annealing treatment on the AA7050 castings from the CDSprocess had a significant influence on the mechanical properties aspresented in Table 7 and graphically shown in Figs. 15 and 16,where typical graphical results of the tensile testing response forthe various heat-treated conditions are presented. Furthermore,Fig. 17 shows the typical optical micrographs of the different heat-treatment conditions before and after the annealing heat treat-ment, and illustrates the microstructural evolution during eachheat treatment for the purpose of interpreting the correlationbetween microstructure and mechanical properties.

The uniaxial tensile properties for the CDS process of AA7050 invarious heat treatment conditions are tabulated in Table 7 and alsographically illustrated in Figs. 15 and 16; both show the softeningeffect of annealing process on the solution treatment (T4þO).Having shown the modification of the microstructure inducedduring annealing process in Fig. 17, the respective effects on thetoughness–strength compromise can be classically attributed tothe following microstructural features:

(a) Inter-granular grain boundary precipitation occurring duringeither the quenching or annealing process controls thestrength and ductility of the grain boundaries and has detri-mental effect on the toughness of the material. The presence ofthese coarse precipitates in the grain boundaries decreasestheir coherency, making them vulnerable to inter-granularfracture, resulting in a loss in strength [37].

(b) Trans-granular coarse precipitates nucleated on the dispersoids4

also appear to have a strong influence on the toughness of the

Fig. 9. Typical results of the visual inspection of the AA7050 casting; (a) the CDSprocess and (b) conventional process.

Table 5Process parameters for the TPGC casting of Al 7050 alloy using the CDS technology.

Alloy# Nominal chemical composition Tmix (1C) TL (1C) experiment TL (1C) phase diagram (Pandat data base) mr

Alloy1 Al–2.8Mg–0.15Ti 660 650.2 652.7 3Alloy2 Al–23.8Zn–8.68Cu 600 588.6 591.6Alloy3 Al–5.95Zn–2.18Cu–2.1Mg–0.11Ti – 632 632.5

4 These dispersoids (which are most likely of impurity elements with anequilibrium partition coefficient k41 therefore they tend to concentrate at thecenter of grains or dendrites, such as chromium, titanium, vanadium, andzirconium [24,26]) precipitates in the form of well-defined bands towards theedge of original cast dendrites within the grain areas [34].

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277268

material, as they favor the nucleation and growth of voidsinside the grains and therefore decrease the intrinsic tough-ness of the grain interiors. This depletion of solute content ofsolid solution in the immediate surroundings of these coarseprecipitates (that nucleate and grow on the aforementioneddispersoids) increase the difference in yield stress between thegrain interiors and the precipitate free zones (PFZs) – forming

both around the grain boundaries and in the vicinity of thesecoarse precipitates [37].

(c) Fine-scale hardening precipitates forming homogeneously inthe grain interiors5 may also play a crucial role in the

Fig. 11. Typical microstructure of AA7050 alloy casting from the conventional casting process using the TPGC process as obtained by a light optical microscope. The hotcracking defect was prevalent in the entire sample microstructure due to the complex dendritic network of the primary Al phase.

Fig. 12. Typical light optical micrographs of as-cast samples for AA7050 alloy cast using the TPGC process with CDS technology.

Fig. 10. Typical microstructure of AA7050 alloy castings from the CDS technology in the TPGC process as obtained by a light optical microscope wherein the morphologyof the primary Al phase is non-dendritic and no discernible evidence of hot cracking defect was observed in the entire sample microstructure.

5 GP zones-η′-metastable precipitates-η-stable precipitates.

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277 269

toughness of the material via the homogeneous depletion ofsolid solution, resulting in an increase in the yield stress of thematerial due to a decrease in strain hardening rate [37].

(d) The intermetallic phases consisting mainly of Al7Cu2Fe, as wellas undissolved Mg2Si, may also have a detrimental effect onthe fracture toughness of the material, although their effect isless than the aforementioned factors given that the volume

Fig. 13. Typical light optical micrographs of solution heat treated (T4 temper) of the AA7050 castings from Fig. 12.

Table 6Uniaxial tensile properties of Al 7050 alloy cast with the CDS process.

Heat treatment UTS MPa (ksi) YS (0.2%) MPa (ksi) El. (%)

F (as-cast) 230.0 (33.4) (725) 213.9 (30.9) (70.7) 0.3 (70.03)T4 447.5 (65) (714.3) 315.9 (46) (79.8) 7.3 (71.7)T6 551 (80) (710.1) 540 (78) (77.3) 1.2 (70.2)

Fig. 14. Typical stress–strain curve for Al 7050 wrought alloy using the CDS processfor both T4 and T6 heat treatment conditions (refer to Table 6 for average data).

Table 7Uniaxial tensile test results of CDS samples of AA7050 in different heat treatmentcombinations between ageing, solutionizing and annealing.

Heat treatment UTS (MPa) YS (0.2%) (MPa) Elongation (%)

F (as-cast) 230.0 (725) 213.9 (70.7) 0.3 (70.03)FþO 174.1 (75.04) 108.7 (72.4) 2.03 (70.2)FþOþT4 390.6 (741.3) 261.7 (730.7) 9.1 (72.1)FþT4 447.5 (714.3) 315.9 (79.8) 7.3 (71.7)FþT4þO 209.0 (714.2) 71.0 (717.2) 10.2 (72.6)FþT4þOþT4 462.9 (714.9) 303.5 (713.9) 8.0 (71.2)

Fig. 15. Comparative bar graph of the tensile properties of the AA7050 CDS castsamples under various heat treatment conditions from Table 7 showing thesoftening effect of annealing treatment as an intermediate process.

Fig. 16. Graphical uniaxial tensile testing results for typical samples of the AA7050alloy cast with the CDS technology, (a) T4, T4þO and T6 tempers, and (b) T4, T4þOand T4þOþT4 tempers. The various heat treatment processes are described in Table 4.

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277270

fraction of these intermetallic particles is reported to beapproximately less than 1% in AA7050 [37].

(e) Texturing of the primary Al grains during the annealing(O) treatment coupled with the evolution of the strengtheningand non-strengthening intermetallic phases in the matrix andgrain boundary regions. The FþO treatment did not show thehigh elongation as that by the FþT4þO treatment because ofthe significant differences in the evolution of the intermetallic

phases after the homogenization of solute elements in thematrix.

4.2. Microstructure analysis

The in-depth analysis of the sample microstructure in thevarious heat treatment conditions presented in Table 4 are pre-sented in this section of this manuscript along with detailed

Fig. 17. Typical light optical micrographs of as-cast samples for AA7050 alloy cast using the TPGC process with CDS technology in high- and low-magnitude: (a) and (c) forannealing heat treatment condition (T4þO), and (b) and (d) for solutionizing heat treatment condition (T4), respectively, showing the precipitation within the grain interiorsof the samples. Notably, all these microstructures were obtained using the same level of etching with the Keller's agent.

Fig. 18. Projection sections (isopleths) of multi-component phase diagram of Al-Zn-Mg-Cu-Fe for (a) maximum and (b) minimum set element compositions evaluated in theAA7050 castings from the CDS process.

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277 271

discussion of the correlation between the microstructure devel-opment with heat treatment to the uniaxial tensile propertiespresented in the previous section.

4.2.1. Microstructure of F condition of AA7050 CDS castingsThe AA7050 has a range of elemental composition as shown in

Table 2, and Fig. 18 presents the simulated isopleths from the multi-

Fig. 19. Typical SEM microstructures of as-cast (F) samples of AA7050 alloy from the CDS process. (a) Low magnification image showing two typical eutectic regions inRegions A and B, (b) eutectic structure from Region A along with qualitative elemental maps and (c) eutectic structure from Region B along with qualitative elemental maps.

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277272

component phase diagrams of the Al-Zn-Mg-Cu–Fe systemwith thetwo extreme composition ranges of elements that were measuredfrom the cast samples in this study: the maximum and minimumamounts of the element addition sets. There are two main regionsin the microstructure of the as-cast tensile test bar samples fromthe CDS process and these are shown in the low magnificationimage in Fig. 19(a) by two boxes on the image: Fig. 19(b) is amagnified section of the box region A along with the results of theelemental mapping carried out by the EDS system in the SEM forthis magnified region and Fig. 19(c) is that of the box region B.Fig. 19(b) shows a system of eutectic phases in AlþAl2Cu (θ)þMg(Zn,Cu,Al)2 (Sigma) along with the Mg2Si phase evolved due to theSi impurity in the alloy as shown in Table 3; the respectivequantitative elemental compositions of these phases are shown inTable 8 along with the respective phase identification numbers.Fig. 19(c) shows another system of the eutectic phases solidifiedfrom the last solidifying pockets of liquid wherein the AlþCuMgAl2(S)þMg (Zn,Cu,Al)2 (Sigma) forms during the eutectic reactionalong with the Cu2FeAl7 phase evolved due to the Fe impurity inthe alloy as shown in Table 3; the respective quantitative elementalcompositions of these phases are shown in Table 8 along with therespective phase identification numbers.

4.2.2. Microstructure of T4 condition of AA7050 CDS castingsFig. 20(a) and (b) shows typical low magnification SEM micro-

structure images of the T4 heat treated cast samples of the AA7050alloy from the tensile bars cast with the CDS process along withthree demarcated regions of interest, Regions C, D and E; Fig. 20(c) and (d) show magnified sections of the Regions C and E,respectively showing the nature of the phases that exist, Fig. 20(e) and (f) show the elemental maps of the Regions C and D inFig. 20(a), respectively. The typical results represented by Fig. 20present the following salient observations and analyses of thetransformation in the microstructure during heat treatment:

� The Al2Cu (θ) phase that predominantly existed in the typicaleutectic region shown by Region A in Fig. 19 does not feature inthe T4 heat treated microstructure shown in Fig. 20. The Fe andMg atoms have redistributed by homogenization during thesolutionizing treatment and further reacted with the Al2Cu (θ)phase to form the Cu2FeAl7 (rod-like) and CuMgAl2 (S) (bulky)phases, respectively as shown in Fig. 20 and the quantitativeelemental composition of these phases are shown in Table 9.

� The Mg(Zn,Cu,Al)2 (Sigma) phase which occurred as the divorcedeutectic phase morphology in the F condition shown in Fig. 19seem to have disappeared during the T4 heat treatment andreappears entrapped inside the CuMgAl2 (S) (bulky) phases duringthe cooling of the sample after T4 treatment as shown in Fig. 20(d).Robson [25] evaluated (simulation) the volume fraction (distribu-tion) of the precipitation phases in the microstructure as a functionof the sample temperature during the solutionizing treatment;Fig. 21 presents the results of this study wherein the S-CuMgAl2phase is predominant at temperatures over 673 K while the Sigma-Mg(Zn,Cu,Al)2 phase is predominant at lower temperatures in thesample microstructure. Accordingly, in agreement with Fig. 21,Fig. 20 shows that the predominant phase in the sample micro-structure of the AA7050 alloy sample after the T4 solutionizingtreatment and subsequent quenching is the S phase with evidenceof entrapped Sigma phase distributed within the S phase. Table 9presents the result of the quantitative EDS analyses of these phases.

� The volume fraction of the impurity intermetallic phase ofMg2Si is not more than 1%. Since the phase remains unchangedduring the solutionizing heat treatment from the non-equilibrium as-cast structure, they do not play an importantrole in the mechanical properties of the materials [37].

4.2.3. Microstructure of T4þO condition of AA7050 CDS castingsThe annealing treatment (O) was carried out to throw more

light on the sequence of precipitation reactions after the solutio-nizing treatment (T4). The microstructure obtained after theannealing process (T4þO) of the AA7050 castings using TPGCprocess are shown in Fig. 22 along with the EDS elementalanalyses on specified areas presented in Table 10.

The SEM images presented in Fig. 22 are consistent with theliterature pertaining to the phases detected in the post-coolingprocesses after solutionizing (T4þO) on AA7050 alloy [25,30]. Themajor microstructural change observed after the T4þO treatmentas compared to the T4 treatment alone is the copious precipitationof the Sigma-Mg(Zn,Cu,Zl)2 phase within all the primary Al grainswith a precipitate depleted zone at the grain boundaries of about5 μm thick; there is no discernable change to the morphology anddistribution of the large S, Sigma and Fe bearing intermetallicphases in the microstructure.

The low magnification image presented in Fig. 22(a) shows asignificant fraction of coarse intermetallic phases decorating thegrain boundaries after the annealing process; and these phases

Table 8EDS elemental analyses of the areas referenced in Fig. 19.

Area# Elem. Possible phase

Mg Al Fe or Si Cu Zn

wt% at% wt% at% wt% at% wt% at% wt% at%

1 0.45(70.12)

0.73(70.21)

45.45(70.39)

65.96(70.25)

– – 52.46(70.67)

32.33(70.27)

1.63(71.15)

0.98(70.69)

Al2Cu (θ)

2 18.59(70.85)

31.75(71.30)

17.58(70.69)

27.06(71.07)

– – 34.23(71.71)

22.38(71.16)

29.60(71.65)

18.81(71.09)

Mg(Zn,Cu,Al)2(Sigma)

3 13.44(71.68)

20.71(72.99)

36.21(74.26)

50.05(75.03)

– – 25.36(71.33)

14.94(71.05)

24.99(71.65)

14.30(71.21)

Mg(Zn,Cu,Al)2(Sigma)

4 49.43(76.43)

55.51(74.87)

14.75(79.78)

13.76(79.46)

Si:31.28(76.71)

Si:30.59(76.40)

0.31 0.134 – – Mg2Si

5 2.59(71.41)

3.95(72.09)

46.86(70.75)

64.95(72.05)

Fe:13.44(74.70)

Fe:8.96(72.99)

29.04(77.03)

17.15(74.36)

6.11(71.50)

3.50(70.88)

Cu2FeAl7

6 1.14(70.77)

1.86(71.26)

43.21(71.63)

63.47(71.87)

– – 51.24(72.85)

31.98(71.90)

4.33(72.76)

2.63(71.70)

CuMgAl2 (S)

7 17.05(71.03)

29.28(72.31)

19.14(74.75)

29.33(76.15)

– – 29.82(72.59)

19.56(71.62)

34.00(75.52)

21.83(74.29)

Mg(Zn,Cu,Al)2(Sigma)

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277 273

Fig. 20. Typical SEM microstructures of the AA7050 alloy castings from the CDS process after the T4 solution heat treatment and quenching. (a) and (b) Low magnificationimages showing three typical eutectic regions in Regions C, D and E, (b) and (c) microstructure images from Regions C and E, respectively, and (e) and (f) elemental mapsfrom Regions C and D, respectively.

Table 9EDS elemental analyses of the numbered phases in Fig. 20.

Area# Elem. Possible phase

Mg Al Fe/Si Cu Zn

wt% at% wt% at% wt% at% wt% at% wt% at%

1 18.59 (70.39) 31.70 (70.89) 17.82 (71.20) 27.37 (71.66) – – 31.5(70.86)

20.56(70.425)

32.10(71.66)

20.37(71.19)

Mg (Zn, Cu,Al)2(entrapped Sigma)

2 48.50 (77.73) 51.452 (76.37) 13.55 (710.48) 13.50 (710.16) Si:32.58(75.41)

31.84(75.10)

0.32 0.132 – – MgSi2

3 14.93 (70.68) 22.56 (70.78) 36.14 (71.40) 49.19 (71.23) – – 45.11(72.20)

26.11(71.73)

3.81(70.60)

2.14(70.33)

CuMgAl2(S)

4 0.32 (70.58) 0.48 (70.87) 50.21 (76.61) 69.19 (75.18) Fe:12.61(72.15)

8.52(71.76)

34.73(76.36)

20.65(74.63)

2.11(71.58)

1.18(70.85)

Cu2FeAl7(rod-like)

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277274

have been retained from after the T4 solution treatment processprior to the O treatment. The phases defining the grain boundarieshave been identified as either insoluble iron-containing phases(Cu2FeAl7) and/or undissolved S phases (CuMgAl2). The Fe-bearingintermetallic phase appears in two distinct morphologies as can beseen in Fig. 22(c) and (d): irregular bulky and acicular (needle).

In addition to the large intermetallic phases at the grainboundary, Fig. 22(a) and (b) shows significant volume fraction ofintermetallic phases precipitated within the primary Al grainsduring the annealing treatment (O); the morphological pattern ofthese precipitates resemble the “widmanstatten pattern” [38] andhave been termed as pseudo-widmanstatten precipitates in thispublication. These precipitates form during the slow coolingtemporal regime in the annealing treatment wherein adequatethermodynamic and kinetic environments are presented for com-plete precipitation of favorable phases from the super saturatedsolid solution in the primary Al grains. An observable feature ofthese pseudo-widmanstatten precipitates is that they evolve ingroups of preferred orientation with the grain as typically demar-cated in Fig. 22(b) by thick broken black line boundaries. The EDSelemental analyses (detailed in Table 10) of these pseudo-widmanstatten precipitates coupled with the information gath-ered from the background literature [25] suggest that they are allthe Sigma phase (Mg(Zn,Al,Cu)2) which are rich in Zn.

Given the findings of this microstructural study, the noticeable lossin toughness–strength response of the annealed (T4þO) samples ofAA7050 CDS castings, illustrated in Figs. 15 and 16, can be attributedto two microstructural features: (a) the intergranular grain boundaryprecipitations including the insoluble Cu2FeAl7 phase and undissolvedcoarse S-phase (CuMgAl2), both of which decrease the coherency ofthe grain boundaries; and (b) the pseudo-widmanstatten precipitatesof the Sigma (Mg(Zn,Al,Cu)2) phase which are typically in over-agedcondition and completely incoherent with the primary Al grainmatrix. The formation of these precipitations causes the creation ofprecipitation free zones (PFZ) both around the precipitates within thegrain and around the grain boundaries. The existence of the PFZ in themicrostructure suggests that the entire Al matrix is nearly pure afterthe annealing treatment (O).

5. Summary

The microstructure in the as-cast condition (F) suggest that theAl matrix has a reasonable high solubility of the alloying elements Ta

ble

10ED

Selem

entalan

alysis

referred

tothesp

otsin

theSE

Mim

ages

specified

inFig.

22.

Area#

Elem

.Po

ssible

phase

Mg

Al

FeCu

Zn

wt%

at%

wt%

at%

wt%

at%

wt%

at%

wt%

at%

17.57

(72)

10.72(7

3.45

)56

.2(7

9.58

)70

.11(7

8.57

)–

–7.4(7

0.82

)3.97

(70.58

)28

.84(7

7.17

)15

.19(7

4.69

)Mg(Zn,C

u,A

l)2(Sigmaphase)

21.04

(70.19

)1.19

(70.22

)93

.5(7

0.73

)96

.48(7

0.42

)–

––

–5.47

(70.64

)2.33

(70.28

)Dep

letedsolid

solution

PFZarou

ndGB

32.57

(70.97

)3.95

(71.5)

55.2

(713

.45)

71.78(7

10.71)

9.78

(73.97

)7.01

(72.36

)32

.91(7

10.06)

19.06(7

7.09

)3.5(7

0.20

)1.73

(70.10

)Cu2Fe

Al 7

(bulkysh

ape)

416

.05(7

1.64

)27

.35(7

3.52

)21

.94(7

7.41

)32

.98(7

9.39

)–

–13

.81(7

6.91

)9.19

(75.11

)48

.2(7

4.09

)30

.49(7

3.39

)Mg(Zn,C

u,A

l)2(Sigma-phaseon

Sphaseat

GB)

515

.65(7

0.45

)23

.48(7

0.60

)36

.14(7

1.12

)48

.84(7

1.10

)–

–48

.22(7

1.22

)27

.69(7

0.98

)–

–CuMgA

l 2(S-phaseat

GB)

62.8(7

0.4)

4.1(7

0.4)

64.1

(78.9)

79.1

(77.2)

8.6(7

2.1)

5.2(7

1.6)

24.8

(77.1)

13.3

(74.5)

3.5(7

0.2)

1.7(7

0.1)

Cu2Fe

Al 7

(rod

-like)

Fig. 21. Simulated equilibrium phase fraction for Direct Chilled cast 7050 ingot; Hcorresponds to the homogenization temperature (753 K). The calculation is limitedto equilibrium conditions and does not present the effect of metastable precipita-tion reactions. “H” denotes the recommended temperature for solutionizingtreatment [24].

R. Ghiaasiaan et al. / Materials Science & Engineering A 594 (2014) 260–277 275

such as Zn, Cu and Mg in it while some of these elements combineto evolve as phases during solidification and occur as bulkymorphologies on grain boundary regions; these phases aretypically θ (Al2Cu), Sigma (Mg(Zn,Al,Cu)2), S (CuMgAl2), Cu2FeAl7and Mg2Si (as in Fig. 19). The uni-axial tensile properties of the as-cast alloy present a weak (YS-214 MPa) component with lowelongation (1%) (as in Fig. 15) that is typical of the inhomogeneousand random distribution of solute elements and phases with fairlyfaceted interface with the Al matrix in the microstructure. Duringthe T4 solutionizing heat treatment process the significant re-distribution of the alloying elements in the microstructure at thehigh temperature results in a super saturates solid solution (SSSS)matrix of primary Al phase along with refinement of the bulkygrain boundary phases to result in only the bulky S (CuMgAl2)phase with pockets of Sigma (Mg(Zn,Al,Cu)2) phase within it andcoupled with some minor volume fraction of the Fe-bearingCu2FeAl7 phase decoration the grain boundaries of the Al matrix(as in Fig. 20). The tensile mechanical properties of the alloy in theT4 heat treated condition present a significant increase in thestrength (YS-315 MPa) along with a nominal increase in elonga-tion (7.3%) (as in Fig. 15) caused by the complete homogenizationof the Al matrix to SSSS and fewer number and distribution of thebulky grain boundary phases with significantly rounded interfacewith the Al matrix. The annealing treatment (O) results indenuding the SSSS of the primary Al matrix off of the soluteelements by copious precipitation of the Sigma (Mg(Zn,Al,Cu)2)phase with incoherent interface with the Al matrix coupled withthe retention of the incoherent and well rounded bulky S(CuMgAl2) phase with pockets of Sigma (Mg(Zn,Al,Cu)2) phasewithin it; the primary Al matrix throughout the microstructureafter the O treatment in nearly pure as indicated by the PFZ (as inFig. 22). During the uni-axial tensile tests of these alloy samplesafter the O treatment, the load is simply applied to a pure Almatrix populated with incoherent Sigma and S phases in various

morphologies which results in a very low strength (YS-71 MPa)and large elongation (410%) (as in Fig. 15). Carrying out the T4solutionizing treatment after the annealing (O) treatment resultsin rendering the pseudo-widmanstatten Sigma phase precipitatesback into the Al matrix resulting in a SSSS and consequentlyreturning the strength and elongation to higher and nominalvalues, respectively (as in Figs. 15 and 16).

This study has amply demonstrated the feasibility of castingAA7050 wrought alloy into near-net shaped casting componentsusing the novel technology of control diffusion solidification (CDS)in conjunction with the tilt pour gravity casting process. Further,the high integrity and properties of these castings under variousdemonstrative heat treatment procedures present viable optionsfor application of this technology in manufacturing Al componentsfor structural applications. This study only serves as a proof ofconcept to further investigate, understand and hone the variousindependent parameters in both the casting and heat treatmentprocesses such as alloy compositions, thermal cycles and temporalcycles to attain significant favorable refinement in the microstruc-ture leading to improved mechanical properties and performanceof such near net shaped alloy castings.

Acknowledgement

The authors express their gratitude to the Natural Sciences andEngineering Research Council (NSERC) of Canada for their financialsupport through the Discovery Grant program.

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Fig. 22. SEM micrographs from the grain boundary and central areas of annealed and furnace cooled solutionized cast parts of CDS 7050 with elemental map by EDSdetector.

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