artigo carlos

Upload: jcpolicarpi

Post on 28-Feb-2018

216 views

Category:

Documents


0 download

TRANSCRIPT

  • 7/25/2019 Artigo Carlos

    1/18

    Effects of Additives on the Microstructures and Tensile Properties ofa New Al-Cu Based Alloy Intended for Automotive Castings

    E.M. Elgallad, A.M. SamuelUniversit du Qubec Chicoutimi, Chicoutimi (QC), Canada

    F.H. Samuel

    Universit du Qubec Chicoutimi, Chicoutimi (QC), CanadaCenter of Excellence for Research in Engineering Materials, King Saud University, Riyadh, Saudi Arabia

    H.W. DotyGeneral Motors Powertrain Group, Metal Casting Technology, Inc., Milford, New Hampshire

    Copyright 2010 American Foundry Society

    ABSTRACT

    This paper discusses the effects of melt treatment and

    addition of alloying elements on the tensile properties of anew Al-2.0%Cu-1.0%Si-0.4%Mg cast alloy in the as-cast

    and heat treated conditions. The additives involved

    include Sr, TiB2, Zr, Ag, Fe, Mn, Sn and Bi. The resultsshow that the role of Sr in refining the morphology of the

    -Fe Chinese script phase causes a slight improvement in

    ductility. The addition of Zr produces a significant

    improvement in the tensile properties as a result of itsgrain refining action. Excess amounts of Fe increase the

    precipitation of Chinese script -Fe particles and thereby

    decrease the tensile properties. The addition of silver does

    not induce considerable increase of strength. This may beascribed to the presence of Siwhich hinders the vital role

    of silver in precipitation-hardening. The softening effect

    of Sn and the replacement of Si with Sn in the Mg-

    hardening phases, as well as the formation of porosity

    arising from the melting of Sn during solution heattreatment were all found to decrease the strength

    properties of Sn-containing alloys. The addition of Bi

    reduces the strength properties in heat-treated conditions

    as a result of the Bi-Mg interaction which suppresses theprecipitation of the Mg-hardening phases.

    INTRODUCTION

    Aluminum-copper based alloys containing Si and Mg areused for the manufacturing of vehicle and airplane parts

    because of their superior mechanical properties,

    castability, weldability and machinability. As in most

    aluminum alloys, the mechanical properties of Al-Cu-Mg-Si alloys can be improved through the use of various

    metallurgical parameters including melt treatment,

    alloying element additions and heat treatment. The

    machinability of such alloys can be metallurgically

    improved so that the chips would flow freely from theircast specimens during machining operations. The present

    work was undertaken to study the effects of additives on

    the microstructures and tensile properties of a new Al-Cubased alloy intended for free-machining automotive

    castings. The additives in question include Sr, TiB2, Zr,

    Fe, Mn and Ag, as well as Sn and Bi as free-cutting

    elements, used to improve the machining behavior of the

    alloy under investigation.

    Melt treatments, such as eutectic silicon modification and

    grain refinement, improve both the casting and the

    mechanical properties of cast Al-Si alloys. Chemicalmodification, using trace additions of strontium, is the

    most common method of modification as a result of which

    the morphology of the silicon particles is changed from

    coarse, acicular plates to finer interconnected fibrous

    ones.1 This change in morphology reduces the stress-

    raising capacity of the silicon particles and significantly

    improves the mechanical properties, particularly

    ductility.2, 3 The addition of grain refiners creates largenumbers of nuclei in the melt thereby inducing the

    formation of small equiaxed grains of -Al. Grain refining

    leads to the even distribution of second phase constituentsand microporosity in the cast structure which in turn

    improves mechanical properties and machinability.4, 5

    Generally speaking, Al-Ti, Al-B, and Al-Ti-B master

    alloys are efficient grain refiners for cast aluminumalloys.6, 7

    Only a few scattered studies are available, to date, on the

    subject of the effects of Zr on cast aluminum alloys.

    Zirconium is used as a grain refiner to reduce the as castgrain size and consequently to improve strength and

    ductility.8 It was also reported that a minor addition of

    0.15 wt% Zr can significantly improve the hardness ofA319 aluminum alloys in both as solutionized and age

    hardened conditions because of the precipitation of the

    coherent coarsening-resistant Al3Zr dispersoids duringsolution heat treatment.

    9, 10 Yin et al.

    11 found that the

    simultaneous addition of 0.1% Zr and 0.2% Sc to Al-

    5%Mg increases strength values by 150 MPa whereas the

    ductility remains at a high level. These authors attributed

    the increments in strength mainly to grain-refinementstrengthening, to Al3(Zr, Sc) dispersive strengthening, and

    to substructure strengthening.

    Iron is one of the most common impurities to be found in

    aluminum alloys and which frequently appears as

    Paper 10-042.pdf, Page 1 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    39

  • 7/25/2019 Artigo Carlos

    2/18

    intermetallic second phases in combination withaluminum and other elements. The more outstanding and

    commonly observed Fe-rich intermetallic phases are -

    Al15(Fe,Mn)3Si2 and -Al5FeSi.12 The brittle -Fe

    intermetallic phase platelets act as stress raisers during

    service and adversely affect mechanical properties andmachinability.13, 14 Neutralization of Fe through the

    promotion of the less harmful -Fe script phase at the

    expense of the brittle platelet-like -Fe phase is soughtwith the goal of improving strength, ductility and other

    properties. Small amounts of manganese (usually

    wt%Mn/wt%Fe 0.5) play a positive role in combining

    with iron to form the less harmful -Fe script phaseinstead of the brittle-Fe phase.

    14, 15

    The addition of silver to Al-Cu-Mg alloys has been

    known to promote the formation of a hexagonal-shaped

    -strengthening phase replacing the precipitationsequence of Al-Cu based systems.16-18 The phase,

    believed to be a variant of the equilibrium (Al2Cu)

    phase, is most commonly found in Al-Cu-Mg-(Ag) alloys

    and substantially improves high-temperature strengthvalues because of its considerable thermal stability. Zhu et

    al.19 stated that, in Al-Cu-Mg-Ag alloys, Ag has an

    overwhelming tendency to form co-clusters with Mg and

    this leads to Ag-Mg-Cu co-clusters, which then act as

    precursors for precipitates. In Al-Mg-Si alloys the

    addition of Ag was found to increase peak hardness and to

    reduce the width of precipitate-free zones (PFZ).20, 21

    It has been reported that small quantities of Sn, of the

    order of 0.05 wt%, have a definite influence on the courseof the precipitation of copper in an Al-4%Cu-0.15%Ti

    alloy. The natural aging of the alloy then becomes

    depressed, while both the response to artificial aging and

    the absolute strength tend to increase.22Tin is one of the

    microalloying elements which is most effective in

    facilitating the nucleation of '.23,24Silcock et al.23found

    that the hardening of the Sn-containing Al-Cu alloy

    proceeds through a single stage at aging temperatures of

    130C (266F) and 190C (374F) as a result of thenucleation of the phase at the expense of Guinier-

    Preston (GP) zones and . Ringer et al.24observed that -Sn particles which precipitated in an Al-4%Cu-0.05%Sn

    alloy after quenching acted as heterogeneous nucleation

    sites for fine and uniformly dispersed phase

    precipitates. On the other hand, Grebenkin et al.25

    foundthat Sn and Pb are the electronic analogs of silicon and

    have been observed to replace it in magnesiumcompounds thereby impeding the formation of the Mg2Si-

    and AlxMg5Si4Cu4-hardening phases in Al-Cu-Si-Mg

    alloys.

    Only limited research has been carried out to date on the

    effects of Bi on the mechanical properties of Al castingalloys. It was demonstrated by a number of researchers

    that Bi could serve as an effective eutectic modifier in Al-

    Si casting alloys.26, 27It was also reported, however, thatincreasing the amounts of added Bi can counteract the

    modifying effects of Sr in A356.2 and A319 alloysbecause of the formation of Bi-Sr compounds which

    reduce the amount of free Sr available for Si

    modification.28, 29With regard to the effects of Bi on the

    aging of Al-Cu alloys, Hardy22found that this element has

    no influence on the aging behavior of the Al-4%Cu-0.15%Ti alloy. It was also suggested that the presence of

    undissolved Bi particles mechanically reduces the

    strength properties and elongation of the alloy studied.

    Heat treatment is one of the major techniques used to

    enhance the mechanical properties of aluminum casting

    alloys. The T6 and T7 tempers are the most commonly

    used tempers for the improvement of the mechanicalproperties of Al-Cu-Si-Mg casting alloys. The T6-temper,

    conducted at aging temperatures ranging from 150 to

    180C (302 to 356F), is applied to obtain the best

    compromise between strength and ductility.30, 31Whereas,

    the stabilizing T7-temper is conducted at higher aging

    temperatures of 200 to 240C (392 to 464F), causing

    overaging and thereby reducing hardness. This temper is

    usually carried out to improve some special characteristicsuch as corrosion resistance and to increase stability and

    performance at elevated temperatures.8 The precipitation-

    hardening characteristics of Al-Cu-Si-Mg alloys often

    appear to be relatively complex. This complexity can be

    attributed to the formation of several hardening phases

    including ' (Al2Cu), '' (Mg2Si), S' (Al2CuMg) and the

    quaternary phase which is designated Q (Al5Mg8Si6Cu2)or (Al5Mg8Si5Cu2).

    32-34Thus, it can be expected that the

    best combination of mechanical properties would be

    obtained when all these precipitates are present.

    This paper will investigate the effects of additives on themicrostructures and tensile properties of a new Al-

    2.0%Cu-1.0%Si-0.4%Mg cast alloy. Several alloys wereprepared from the base alloy with the intention ofstudying the effects of:

    1. melt treatment, namely modification and grainrefining using Sr, Ti and Zr additives,

    2. iron intermetallics by increasing the Fe and Mncontent of the base alloy,

    3. silver as a hardening element and4. free-cutting elements through the addition of Sn and

    Bi.

    The mechanical properties were studied in the as cast and

    in two different heat-treated conditions, namely T6 andT7 tempered conditions. The machining behavior of a

    number of these alloys will be investigated in asubsequent study.

    EXPERIMENTAL PROCEDURES

    ALLOYS AND MATERIALSThe nominal level of the alloying elements added to the

    base alloy and the codes of the resulting alloys togetherwith their classification are shown in Table 1. The actual

    composition of each of these alloys, as obtained from

    chemical analysis, is listed in Table 2. The alloys were

    Paper 10-042.pdf, Page 2 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    40

  • 7/25/2019 Artigo Carlos

    3/18

    Table 1. Nominal Composition and Codes for the Alloys Prepared in the Present Study

    Group I Group II Group III

    AlloyCode

    CompositionAlloyCode

    CompositionAlloyCode

    Composition

    A Base alloy A3 A + Sr + 0.10%Ti A4 A + 0.10%Ti + 0.20%Zr

    A1 A + Sr A31 A3 + 0.20%Fe A41 A4 + 0.15%Sn

    A2 A + 0.10%Ti A32 A3 + 0.20%Fe + 0.2%Mn A42 A4 + 0.50%Bi

    A3 A + Sr + 0.10%Ti A33 A3 + 0.50%Ag A43 A4 + 0.15%Sn + 0.50%Bi

    A4 A + 0.10%Ti + 0.20%Zr

    Note: Sr level = 100-150 ppm

    Table 2. Actual Chemical Composition of the Alloys Prepared for the Present Study

    Chemical Composition (% wt)AlloyCode Cu Si Mg Fe Mn Sr Ti Zr Ag Sn Bi Al

    A 2.09 1.32 0.42 0.58 0.59 0.000 0.07 0.00 0.00 0.00 0.00 bal.

    A1 2.13 1.28 0.42 0.58 0.60 0.013 0.08 0.00 0.00 0.00 0.00 bal.

    A2 2.18 1.23 0.40 0.61 0.61 0.000 0.15 0.00 0.00 0.00 0.00 bal.

    A3 2.11 1.23 0.40 0.52 0.60 0.011 0.16 0.00 0.00 0.00 0.00 bal.

    A4 2.24 1.28 0.41 0.61 0.58 0.000 0.15 0.20 0.00 0.00 0.00 bal.

    A31 2.17 1.22 0.40 0.84 0.59 0.014 0.16 0.00 0.00 0.00 0.00 bal.A32 2.09 1.17 0.39 0.82 0.79 0.010 0.18 0.00 0.00 0.00 0.00 bal.

    A33 2.09 1.21 0.39 0.57 0.60 0.010 0.16 0.00 0.50 0.00 0.00 bal.

    A41 2.31 1.33 0.43 0.63 0.59 0.000 0.16 0.20 0.00 0.22 0.00 bal.

    A42 2.31 1.26 0.45 0.52 0.61 0.000 0.18 0.20 0.00 0.00 0.51 bal.

    A43 2.24 1.24 0.47 0.45 0.61 0.000 0.17 0.20 0.00 0.24 0.55 bal.

    subdivided into three groups according to the alloyingadditions involved, namely Groups I, II and III.

    Group I will examine the effects of melt treatmentthrough the addition of Sr, Ti, Sr + Ti and Ti + Zr tothe Al-Cu base A alloy (A1, A2, A3 and A4 alloys,

    respectively). Group II will examine the effects of Fe, Fe + Mn and

    Ag, as a hardening alloying element,by adding them

    to the A3 alloy (A31, A32 and A33 alloys,

    respectively).

    Group III will examine the effects of free-cuttingelements through the addition of Sn, Bi and Sn + Bi

    to the A4 alloy (A41, A42 and A43 alloys,respectively).

    MELTING AND CASTING PROCEDURESThe base alloy A used in this study was supplied in the

    form of 12.5-kg ingots which were subsequently cut,

    dried and then melted in a SiC crucible of 40-kg capacityusing an electrical resistance furnace. The melting

    temperature was maintained at 750 5C (1382 41F)during which time the melt was grain-refined and

    modified with Al-5%Ti-1%B and Al-10%Sr master

    alloys, respectively. The elements Fe, Mn, Ag, Zr, and Bi

    were added in the form of Al-25%Fe, Al-25%Mn, Al-

    50%Ag, Al-15%Zr, and Al-50%Bi master alloys,respectively, whereas Sn was introduced in the form of

    the pure metal. The melt was degassed using pure dry

    argon for 15 min, injected into the melt by means of a

    graphite impeller rotating at 150 rpm. The surface oxides

    and/or inclusions were skimmed thoroughly prior to

    pouring. The melt was poured at ~735C (1355F) into an

    ASTM B-108 mold, which had been preheated to 450C(842F), so as to obtain castings for tensile test bars. Each

    casting provided two test bars. For each alloycomposition, fifty castings or one hundred tensile test bars

    were prepared. Samplings for metallographic observationand spectrochemical analysis were also taken for each

    alloy melt composition.

    HEAT TREATMENTThe one hundred tensile test bars obtained for each alloy

    composition were divided into twenty batches

    corresponding to the following alloy conditions (5 bars /

    condition):

    as cast condition;

    solution heat-treated condition carried out at 495C

    (923F) for 8 h; nine T6 heat-treated conditions corresponding to nineaging times and

    nine T7 heat-treated conditions corresponding to nineaging times.

    In both T6 and T7 tempers, the samples were solutionheat treated at 495C (923F) for 8 h, quenched in warm

    water at 65C (149F) and then artificially aged. Artificial

    aging of the samples was carried out at 180C (356F) and

    220C (428F) for T6 and T7 tempers, respectively, for

    Paper 10-042.pdf, Page 3 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    41

  • 7/25/2019 Artigo Carlos

    4/18

    aging times of 2, 4, 6, 8, 12, 16, 20, 24 and 48 h. All heattreatments were conducted in a programmable

    temperature controlled electric furnace.

    METALLOGRAPHYFor metallographic observations, 25 x 25 mm sampleswere cut from the castings prepared for this purpose and

    mounted in bakelite. The samples were ground and

    polished to the desired fine finish on 9, 6, 3 and 1 m

    diamond lap wheels. The microstructures were examined

    by means of an electron probe micro-analyzer (EPMA)

    and an optical microscope. The grain-size measurements

    were carried out using a Clemex image analyzer inconjunction with the optical microscope. The grain size

    was obtained from the average of 200 measurements

    taken over 20 fields (10 measurements per field) at 100xmagnification for each alloy sample. Volume fraction of

    the intermetallic phases was quantified using the electron

    probe micro-analyzer with built-in software for such

    measurements, based on phase brightness. The

    quantification process is based on the elimination

    technique which calculates the volume fraction of eachphase by subtracting the volume fraction of the brighter

    phases from the total volume fraction of the other phasesthat are present within the matrix. For each case, 15 fields

    were measured at 100X magnification.

    TENSILE TESTINGThe tensile test bars were pulled to fracture at room

    temperature at a strain rate of 4 x 10-4/s, using a

    Servohydraulic MTS Mechanical Testing machine. An

    extensometer with a 50.8 mm (2 in) gage length wasattached to the test bar to measure percentage elongation

    as the load was applied. The tensile properties, namely

    yield stress (YS) at a 0.2% offset strain, ultimate tensile

    strength (UTS) and fracture elongation (%El), were

    derived from the data-acquisition and data-treatmentsystems of the tensile testing machine used. The tensile

    properties of each alloy/heat-treatment condition were

    represented by the average %El, YS and UTS values

    which were calculated from the values obtained from thefive tensile test bars assigned to that specific alloy/heat

    treatment condition.

    RESULTS AND DISCUSSION

    MICROSTRUCTURESMicro-Constituents of the Base AlloyThe backscattered image of the as cast base A alloy,

    shown in Fig. 1, reveals the presence of Al2Cu,

    Al5Mg8SixCu2 and the Chinese script-like

    -Al15(Fe,Mn)3Si2 phases in the alloy microstructure(phases were identified using Wavelength Dispersive

    Spectroscopy [WDS] analysis). It seems that the low Si-

    content of the base A alloy was consumed in the

    formation of Al-Fe-Si and Al-Cu-Mg-Si intermetallic

    phases. The platelet-like -Al5FeSi phase was not inevidence because of the higher Mn/Fe ratio of the alloy

    (~1) which promotes the formation of the -Fe phase atthe expense of -Fe phase.

    Effects of Melt TreatmentThe effect of Sr addition on the microstructural

    characteristics can be understood by comparing themicrograph obtained from the base A alloy,(Fig. 2a), to

    the one obtained from the Sr-containing A1 alloy (Fig.

    2b). It would appear that the addition of Sr refines themorphology of the -Fe script phase to a certain extent in

    the A1 alloy, resulting in the even distribution of the

    particles of this phase within the matrix of the

    microstructure. Similar observations were also reported

    by Shabestari et al.35

    AlCuMgSi

    Al2Cu

    -Fe

    Fig. 1. Backscattered image obtained from the as-cast base A alloy.

    Paper 10-042.pdf, Page 4 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    42

  • 7/25/2019 Artigo Carlos

    5/18

    Fig. 2. Micrographs obtained from: (a) the unmodified base A alloy; and (b) the Sr-modified A1 alloyin as cast condition.

    Fig. 3. Micrographs obtained from: (a) A alloy; and (b) A4 alloy in as cast condition.

    The ZrTi particles formed in the Zr-containing A4 alloy

    act as nucleation sites for small equiaxed grains of -Al.

    Grain size measurements reveal that the combinedaddition of Ti and Zr causes a decrease in the grain size

    from 500m in the non-grain-refined base A alloy to

    160 m in the grain-refined A4 alloy. This difference in

    grain size is clearly evident upon comparing themicrographs of both the alloys,as shown in Fig. 3a and

    3b, respectively

    Effects of the Addition of Fe and Mn

    Increasing the Fe content to 0.8% in the A31 alloy wasfound to increase the precipitation of the -Fe script

    phase, as evidenced from volume fraction measurements.

    A typical micrograph is shown in Fig. 4a. The platelet-like -Fe phase, however, did not form since the Mn/Fe

    ratio of the alloy was still high enough (~0.7) to promote

    the formation of the -Fe phase rather than that of the -

    Fe phase. The further addition of 0.2% Mn to the A31alloy, namely the A32 alloy, did not lead to the

    precipitation of undesirable sludge particles which may

    form at higher Mn levels, as may be noted by their

    absence in the micrograph shown in Fig. 4b.

    Effects of the Addition of Sn and BiFigure 5a shows a high magnification backscattered

    image obtained from the as cast Sn-containing A41 alloy

    where the precipitation of Sn in the form of -Sn particlesmay be observed as the white phase. These particles

    appear as small non-uniformly distributed clusters usually

    solidified within the Al2Cu phase network. The presenceof Bi in the form of undissolved particles in the A42 alloy

    may clearly be observed in the high magnificationmicrograph presented in Fig. 5b. The presence of TiB2

    and ZrTi particles, which induce the grain refining effect,

    are also observed in this micrograph.

    The higher magnification micrographs, obtained from the

    T6-treated A41 alloy and illustrated in Fig. 6a and 6b,

    show, respectively, the morphology of the Mg2Sn phaseprecipitated in the alloy and a resoldified -Sn particle

    that had undergone incipient melting during the solution

    heat treatment.

    (a)

    Al2CuAlCuMgSi

    -Fe

    Al2Cu

    -Fe

    Al2Cu

    (b)

    (b)(a)

    Paper 10-042.pdf, Page 5 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    43

  • 7/25/2019 Artigo Carlos

    6/18

    Fig. 4. Micrographs obtained from: (a) A31alloy; and (b) A32 alloy in as cast condition.

    Fig. 5. High magnification backscattered images obtained from: (a) Sn-containing A41 alloy;and (b) Bi-containing A42 alloy in as cast condition.

    Fig. 6. Higher magnification backscattered images obtained from T6-treated A41 alloy showing:(a) morphology of Mg2Sn; and (b) resolidified -Sn particle which had undergone incipient melting.

    Fig. 7. Backscattered image obtained from the solutionized base A alloy.

    -Fe

    -Sn

    Al2Cu

    (a) Bi

    ZrTi

    TiB2

    Bi

    Bi

    (b)

    (a) (b)

    Al2Cu

    Al2Cu

    -Fe

    -Fe

    (b)

    (b)(a)

    Al2Cu

    Al2Cu

    -Fe

    -Fe

    Paper 10-042.pdf, Page 6 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    44

  • 7/25/2019 Artigo Carlos

    7/18

    Effects of Solution Heat TreatmentFigure 7 shows a backscattered micrograph obtained from

    the solutionized base A alloy. The scarcity of Cu-rich

    intermetallic phases implies that the solution heat

    treatment caused these phases to become almost

    completely dissolved into the solid solution. On the otherhand, the -Fe script phase particles, clearly seen in the

    microstructure, are not usually affected by such treatment.

    Figure 8 shows the effects of solution heat treatment on

    the volume fraction of the iron and copper intermetallics

    contained in the alloys studied. It will be observed that the

    solution heat treatment reduces the volume fraction of

    these intermetallics by approximately 35% corresponding

    to the dissolution of the Cu-rich intermetallic phases.Thus, the volume fractions of the -Fe script phase, which

    remained unaffected by the solution heat treatment, are

    those values plotted in the solution heat treated condition.It is also observed that the alloys have almost the same

    volume fraction value with regard to this phase, except for

    the higher values observed for both A31 and A32 alloys,

    because of the higher Fe content in the former and thehigher Fe + Mn content in the latter.

    TENSILE PROPERTIESEffects of Melt Treatment (Alloying Group I)The effects of melt treatment on the tensile properties of

    Alloying Group I are shown in Fig. 9a for the as cast

    condition, and in Figs. 10 and 11 for the 180C (356F) and220C (428F) aged conditions, respectively. These aging

    temperatures refer to the T6 and T7 tempers. (Here, it ought

    to be noted that, with respect to Figs. 10 and 11 (as well asFigs. 12-15), the Y-axis scales have been plotted according

    to the maximum/minimum values noted in each case to

    facilitate separation of the curves). It can be observed that

    the Sr-containing A1 alloy did not exhibit any noticeable

    change in the strength properties in the as cast and heattreated conditions, compared to the base A alloy. The

    improvement in the ductility of A1 alloy, especially in the

    as cast condition, can probably be ascribed to the role of Sr

    in refining the morphology of the -Fe script phaseappearing in the alloy microstructure.

    In the as cast condition, the grain-refined A2 alloy showsimprovement in the %El value compared to the base A

    Alloy. While in both T6 and T7 heat treated conditions,the UTS and %El were observed to be higher compared to

    the A alloy for most of the aging times studied.

    In spite of the Sr-modified A1 alloy and the grain-refined

    A2 alloy displaying improvements in the tensile

    properties, the modified grain-refined A3 alloy did notproduce tensile properties which were any better than

    those of the previously mentioned alloys whether in the as

    cast or heat treated conditions. This absence ofimprovement in tensile properties can be explained in

    terms of the interaction between Sr and B and/or Sr and Ti

    as reported by Liao et al.36, 37 These interactions have

    been known to cause mutual poisoning of the elements

    involved and, consequently, to suppress their modification

    and/or grain refining effects.

    The Zr-containing A4 alloy possesses the highest values

    for tensile properties among the alloys of Group I. This

    alloy displayed significant increases in the YS and UTSalong with a higher level of ductility in the as cast and

    heat treated conditions. There is a distinct possibility that

    the higher strength increment produced in this alloy,

    particularly in the as cast condition, may be attributed tothe strengthening mechanism stimulated by the grain-

    refining effect of Zr, as previously indicated by Mahmudi

    et al.9and Yin et al.

    11

    Effects of Iron Intermetallics and Silver (AlloyingGroup II)The effects of the addition of Fe, Fe + Mn and Ag on the

    tensile properties of Alloying Group II are shown in Fig.

    9b for the as cast condition; and in Figs. 12 and 13 for the

    180C (356F) and 220C (428F) aged conditions,respectively. Increasing the Fe content to 0.8% in the A31

    alloy causes a decrease in tensile properties, particularly

    ductility in the as cast and heat treated conditions. This

    decrease was predictable based on the increase in the

    volume fraction of iron intermetallic phases, mainly the

    Fig. 8. Effect of solution heat treatment on the volume fraction (%) of copper and iron intermetallic phases.

    0

    1

    2

    3

    4

    5

    6

    A A1 A2 A3 A4 A31 A32 A33 A41 A42 A43

    Alloy

    VolumeFraction

    %

    As-Cast SHT

    Paper 10-042.pdf, Page 7 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    45

  • 7/25/2019 Artigo Carlos

    8/18

    Fig. 9. Tensile properties in as cast condition: (a) Alloying Group I, (b) Alloying Group II and (c) Alloying Group III

    50

    75

    100

    125

    150

    175

    200

    225

    250

    A A1 A2 A3 A4

    Alloy

    YS,U

    TS(MPa)

    2.00

    2.50

    3.00

    3.50

    4.00

    4.50

    %El

    YS UTS %El(a)

    A = base alloyA1 = A + SrA2 = A + 0.10%TiA3 = A + Sr + 0.10%TiA4 = A + 0.10%Ti + 0.20%Zr

    50

    75

    100

    125

    150

    175

    200

    225

    250

    A3 A31 A32 A33

    Alloy

    YS,

    UTS(MPa)

    2.00

    2.50

    3.00

    3.50

    4.00

    4.50

    %El

    YS UTS %EL(b)

    A3 = A + Sr + 0.10%Ti

    A31 = A3 + 0.20%FeA32 = A3 + 0.20%Fe + 0.20%MnA33 = A3 + 0.50%Ag

    50

    75

    100

    125

    150

    175

    200

    225

    250

    A4 A41 A42 A43

    Alloy

    YS,UTS(MPa)

    2.00

    2.50

    3.00

    3.50

    4.00

    4.50

    %El

    YS UTS %El(c)

    A4 = A + 0.10%Ti + 0.20%ZrA41 = A4 + 0.15%SnA42 = A4 + 0.50%BiA43 = A4 + 0.15%Sn + 0.50%Bi

    Paper 10-042.pdf, Page 8 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    46

  • 7/25/2019 Artigo Carlos

    9/18

    Fig. 10. Variations in tensile properties of Alloying Group I after aging at 180C (356F): (a) YS, (b) UTS and (c) %El

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A = base alloy

    A1 = A + Sr

    A2 = A + 0 .10% Ti

    A3 = A + S r + 0.10% Ti

    A4 = A + 0 .10% Ti + 0 .20% Zr

    (a)

    240

    260

    280

    300

    320

    340

    360

    380

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A = base alloy

    A1= A + Sr

    A2 = A + 0.10% Ti

    A3 = A + Sr + 0.10% Ti

    A4 = A + 0.10% Ti + 0 .20% Zr

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    %E

    l

    A = base alloy

    A1 = A + Sr

    A2 = A + 0.10%Ti

    A3 = A + Sr + 0.10%Ti

    A4 = A + 0.10%Ti + 0.20%Zr

    (c)

    Paper 10-042.pdf, Page 9 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    47

  • 7/25/2019 Artigo Carlos

    10/18

    Fig. 11 Variations in tensile properties of Alloying Group I after aging at 220C (428F): (a) YS, (b) UTS and (c) %El

    100

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    320

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A = base alloy

    A1 = A + Sr

    A2 = A + 0.10%Ti

    A3 = A + Sr + 0.10%Ti

    A4 = A + 0.10%Ti + 0.20%Zr

    (a)

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A = base alloy

    A1 = A + S r

    A2 = A + 0 .10% Ti

    A3 = A + S r + 0.10% Ti

    A4 = A + 0 .10% Ti + 0.20% Zr

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    %E

    l

    A = base alloy

    A1 = A + Sr

    A2 = A + 0.10% Ti

    A3 = A + Sr + 0.10% Ti

    A4 = A + 0.10% Ti + 0.20% Zr

    (c)

    Paper 10-042.pdf, Page 10 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    48

  • 7/25/2019 Artigo Carlos

    11/18

    -Fe Chinese script phase, caused by the addition of iron,as mentioned in the discussions on solution heat treatment

    and microstructure.

    The subsequent addition of 0.2% Mn to the A31 alloy

    (producing the A32 alloy) slightly increases the YS andUTS values, which were previously reduced by increasing

    the iron content in the A31 alloy. The %El was not,

    however, affected in any marked way. The positive role ofMn, in promoting precipitation of the less harmful -Fe

    Chinese script phase particles instead of the brittle

    platelet-like -Fe phase particles (the latter were not

    detected in the microstructure), may explain the marginal

    improvement caused in the YS and UTS values of the

    A32 alloy.

    The addition of silver to the A3 alloy which produces the

    A33 alloy, did not, in fact, change the tensile properties ofthe as cast condition. In heat treated conditions, however,

    the addition of silver did increase the UTS but not as

    expected. This situation can be attributed to the presence

    of Si in the base alloy, which favors the formation of theMg-Si phases during the early stages of aging, in turnexhausting the supply of magnesium and reducing the

    number of Mg-Ag co-clusters known to act as nucleation

    sites for hardening precipitates. It has been reported thatthe precipitation of the phase may be hindered by the

    presence of small concentrations of Si in Al-Cu-Mg-(Ag)

    alloys.38-40Matsuda et al.21 found that the addition of Ag

    to Al-Mg-Si alloy containing Si content in excess of that

    required for Mg2Si precipitates did not produce any

    substantial improvement in age-hardening characteristics.The presence of Si prevents the formation of Mg-Ag

    clusters which provide a lot of nucleation sites for fine

    and more dispersed '' phase precipitates.20 Elevatedtemperature tensile testing may be recommended,

    nevertheless, so as to evaluate the tensile properties of

    A33 alloy, in view of the fact that the phase, which

    favors precipitation at the expense of the (Al2Cu) phase

    in the presence of Ag, was reported to improve the

    mechanical properties at higher temperatures.41

    Effects of Free-Cutting Elements (Alloying Group III)The effects of the addition of free-cutting elements,namely Sn and Bi, as well as a combination of both, on

    the tensile properties of Alloying Group III,are shown in

    Fig. 9c for the as cast condition and in Figs. 14 and 15 forthe T6 and T7 heat treated conditions, respectively.

    The addition of 0.15%Sn to the A4 alloy, namely the A41

    alloy, causes a decrease in the YS and UTS, but increases

    the %El in the as cast condition as a result of the softening

    effect of the soft Sn-bearing phases, dispersed within thealloy microstructure. In heat treated conditions, the

    noticeable reduction occurring in the YS and UTS of the

    Sn-containing A41 alloy can be explained in terms of thefollowing effects, which were previously confirmed by

    the examination of the microstructure: (1) the softening

    effect of the soft Sn-rich phases; (2) the replacement of Si

    by Sn in Mg compounds (formation of Mg2Sn) which inturn diminishes the precipitation of Mg2Si and/or

    Al5Mg8SixCu2 hardening phases and (3) the increase in

    the percentage porosity arising from the melting of the

    low melting point Sn-phases during solution heat

    treatment. These effects were also confirmed by the workof Mohamed et al.,42 who found that increasing the Sn

    content in B319.2 and A356.2 alloys decreases their

    mechanical properties in the heat-treated condition. Thisis due to the increase in the percentage porosity resulting

    from the melting of -Sn particles in the B319.2 alloy

    during solution heat treatment and to the formation of

    Mg2Sn in the A356.2 alloy, which lessens the amount of

    Mg required for the formation of Mg hardening phases.

    The increase in ductility, resulting from the softening

    effect of the soft Sn-bearing phases, may balance out the

    reduction caused by the increase in the percentage

    porosity, thus explaining why the ductility of the Sn-containing A41 alloy was not significantly affected by the

    addition of Sn in both T6- and T7-tempers over all the

    aging times applied.

    The bismuth-containing A42 alloy exhibits considerable

    deterioration of its tensile properties in the as cast andheat treated conditions. The presence of Bi particles

    within the alloy microstructure reduces the tensile

    properties as reported by Hardy.22 It is important tomention that the effectiveness of Bi addition, which

    improves the machinability of 6262 Al-Mg-Si alloys, was

    found to be reduced by the loss of Bi in the formation of

    Bi2Mg3 particles.43 The Bi-Mg-Sr interaction was also

    confirmed in research carried out by Elhadad et al.29Based on these observations, the reduction caused in the

    strength properties of A42 alloy in heat treated conditions

    can be explainedin terms of the Bi-Mg interaction, whichconsumes the Mg available for the formation of Mg-

    hardening precipitates.

    The deterioration, occurring in the tensile properties of the

    A43 alloy containing Sn and Bi in the as cast and heattreated conditions, was expected in light of the

    detrimental effects of the individual additions of Sn and

    Bi on the tensile properties of the A41 and A42 alloys,respectively. It can be concluded that the reductions

    caused in the strength values, whether YS or UTS, as a

    result of the combined addition of Sn and Bi are

    approximately the sum of those reductions caused by the

    individual addition of the A41 and A42 alloys,respectively.Moreover, the %El values of the A43 alloy

    appear too close to those of the Bi-containing A42 alloy.

    This observation indicates that the %El of the A43 alloywas only reduced by Bi, whereas Sn did not significantly

    affect the ductility for the same reason as is applicable to

    the %El of the A41 alloy. The slight increase in the %El

    of the A43 alloy, observed after T7 treatment, can be

    attributed to the harmful porosity effect, arising from the

    melting of Sn, being overridden by the beneficialsoftening effect of soft Sn-rich phases.

    Paper 10-042.pdf, Page 11 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    49

  • 7/25/2019 Artigo Carlos

    12/18

    Fig. 12. Variations in tensile properties of Alloying Group II after aging at 180C (356F): (a) YS, (b) UTS, and (c) %El

    100

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A3 = A + Sr + 0.10% Ti

    A31 = A 3 + 0.20% Fe

    A32 = A 3 + 0.20% Fe + 0.20% Mn

    A33 = A 3 + 0.50% Ag

    (a)

    220

    240

    260

    280

    300

    320

    340

    360

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A3 = A + Sr + 0.10% Ti

    A31 = A 3 + 0.20% Fe

    A32 = A 3 + 0.20% Fe + 0.20% Mn

    A33 = A 3 + 0.50% Ag

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    7.5

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    %E

    l

    A3 = A + Sr + 0.10% Ti

    A31 = A 3 + 0.20% Fe

    A32 = A 3 + 0.20% Fe + 0.20% Mn

    A33 = A 3 + 0.50% Ag

    (c)

    Paper 10-042.pdf, Page 12 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    50

  • 7/25/2019 Artigo Carlos

    13/18

    Fig. 13. Variations in tensile properties of Alloying Group II after aging at 220C (428F): (a) YS, (b) UTS, and (c) %El

    100

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A3 = A + Sr + 0.10% Ti

    A31 = A 3 + 0.20% Fe

    A32 = A 3 + 0.20% Fe + 0.20% Mn

    A33 = A 3 + 0.50% Ag

    (a)

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A3 = A + Sr + 0.10%Ti

    A31 = A3 + 0.20%Fe

    A32 = A3 + 0.2%Fe + 0.20%Mn

    A33 = A3 + 0.50%Ag

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    7.5

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    %E

    l

    A3 = A + Sr + 0.10% Ti

    A31 = A 3 + 0.20% Fe

    A32 = A 3 + 0.20% Fe + 0.20% Mn

    A33 = A 3 + 0.50% Ag

    (c)

    Paper 10-042.pdf, Page 13 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    51

  • 7/25/2019 Artigo Carlos

    14/18

    Fig. 14. Variations in tensile properties of Alloying Group III after aging at 180C (356F: (a) YS, (b) UTS, and (c) %El

    100

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A4 = A + 0. 10%Ti + 0. 20%Zr

    A41 = A 4 + 0.15% Sn

    A42 = A 4 + 0.50% Bi

    A43 = A 4 + 0.15% Sn + 0.50% Bi

    (a)

    200

    220

    240

    260

    280

    300

    320

    340

    360

    380

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A4 = A + 0.10% Ti + 0.20% Zr

    A41 = A 4 + 0.15% Sn

    A42 = A 4 + 0.50% Bi

    A43 = A 4 + 0.15% Sn + 0.50% Bi

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    0 2 4 6 8 12 16 20 24 48Aging Time (hrs)

    %E

    l

    A4 = A + 0.10%Ti + 0.20%Zr

    A41 = A4 + 0.15%S n

    A42 = A4 + 0.50%B i

    A43 = A4 + 0.15%S n + 0.50%B i

    (c)

    Paper 10-042.pdf, Page 14 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    52

  • 7/25/2019 Artigo Carlos

    15/18

    Fig. 15. Variations in tensile properties of Alloying Group III after aging at 220C (428F): (a) YS, (b) UTS, and (c) %El.

    80

    100

    120

    140

    160

    180

    200

    220

    240

    260

    280

    300

    320

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    YS(M

    Pa)

    A4 = A + 0.10%Ti + 0.20%Zr

    A41 = A4 + 0.15%S n

    A42 = A4 + 0.50%B i

    A43 = A4 + 0.15%S n + 0.50%B i

    (a)

    200

    220

    240

    260

    280

    300

    320

    340

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    UTS(MPa)

    A4 = A + 0.10%Ti + 0.20%Zr

    A41 = A4 + 0.15%S n

    A42 = A4 + 0.50%B i

    A43 = A4 + 0.15%S n + 0.50%B i

    (b)

    0.5

    1.5

    2.5

    3.5

    4.5

    5.5

    6.5

    0 2 4 6 8 12 16 20 24 48

    Aging Time (hrs)

    %E

    l

    A4 = A + 0. 10%Ti + 0.20% Zr

    A41 = A 4 + 0.15% Sn

    A42 = A 4 + 0.50% Bi

    A43 = A 4 + 0.15% Sn + 0.50% Bi

    (c)

    Paper 10-042.pdf, Page 15 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    53

  • 7/25/2019 Artigo Carlos

    16/18

    Effects of Age HardeningUpon comparing the two sets of tensile property curves,

    shown in Figs. 10, 12 and 14 with those in Figs. 11, 13

    and 15 corresponding to T6- and T7-tempers,

    respectively, it can be observed that the T6 treatment (i.e.

    aging at 180C [356F]) results in alloy hardening withaging time almost up to 20 h of aging whereas the T7

    treatment (i.e. aging at 220C [428F]) causes overaging

    and alloy softening after 2 h of aging. Therefore, YS andUTS decrease while %El increases. These results suggest

    that the T6-temper may be recommended for the new Al-

    Cu based alloys under investigation. It can also be

    observed that the aging time of up to 20 h does not

    significantly affect the tensile properties in the T6-

    tempered condition. This observation can be interpreted interms of the presence of several hardening phases in Al-

    Cu alloys, containing Si and Mg and including (Al2Cu),

    (Mg2Si) and Q(Al5Mg8Si6Cu2), which contribute to the

    precipitation hardening of these alloys.32-34

    CONCLUSIONS

    The effects of additives on the microstructures and tensile

    properties of an Al-Cu based alloy, having a low Si

    content, were investigated in the as cast and heat treatedconditions. From an analysis of the results obtained, the

    following conclusions may be drawn.

    1. The addition of Sr refines the morphology of the -FeChinese script phase which in turn contributes to a

    slight improvement in ductility.

    2. The addition of zirconium improves the tensileproperties in the as cast and heat treated conditionsconsiderably because of the strengthening induced by

    its grain refining effect.

    3. Increasing Fe content by 0.2% increases theprecipitation of -Fe Chinese script particles thereby

    reducing the tensile properties, particularly ductility.The subsequent addition of Mn marginally increases

    the YS and UTS without any observable change in

    the %El.4. The addition of silver does not produce any

    considerable increase of strength in heat-treated

    conditions. This result may be ascribed to the

    presence of Si which suppresses the vital role of

    silver in precipitation hardening.5. The addition of Sn lowers the YS and UTS but raises

    the %El in the as cast condition as a result of the

    softening effect of soft Sn-bearing phases. In the heattreated conditions, the reduction caused in the

    strength properties is attributed mainly to the

    formation of porosity associated with the melting of

    Sn during solution heat treatment and the

    replacement of Si by Sn in Mg compounds.This in

    turn hinders the precipitation of Mg-hardeningphases.

    6. The Bi-Mg interaction, which consumes the amountof Mg required to form the Mg-hardening phases, isresponsible for the reduction caused in the strength

    properties of Bi-containing alloys in the heat-treatedconditions.

    7. Applying a T6-temper at 180C (356F) produces asatisfactory compromise between strength and

    ductility. As a result of this treatment, the alloysshow hardening after up to 20 hours of aging time

    because of the presence of several hardening phases

    in the Al-Cu-Si-Mg alloy system. Applying a T7-

    temper at 220C (428F) causes overaging and alloysoftening after 2 hours of aging time.

    ACKNOWLEDGMENTS

    Financial assistance received from the Natural Sciences

    and Engineering Research Council of Canada (NSERC)

    and General Motors Powertrain Group (U.S.A.) is

    gratefully acknowledged.

    REFERENCES

    1. Gruzleski, J., Closset, B., The Treatment of Liquid

    Aluminum-Silicon Alloys, pp. 25-55, AmericanFoundrymens Society Inc., Des Plaines, IL (1990)

    2. Hafiz, M. and Kobayashi, T., Mechanical Propertiesof Modified and Non-modified Eutectic Al-Si

    Alloys, Journal of Japan Institute of Light Metals,vol. 44, no.1, pp. 28-34 (1994)

    3. Fat-Halla, N., Structural Modification of Al-SiEutectic Alloy by Sr and its Effect on Tensile andFracture Characteristics, Journal of Materials

    Science, vol. 27, pp. 2488-2490 (1989)

    4. Cibula, A., The Grain Refinement of Al AlloyCastings by Addition of Ti and B, Journal of the

    Institute of Metals, vol. 90, pp. 1-16 (1951-52)

    5. McCartney, D.G., Grain Refining of Aluminum andits Alloys Using Inoculants, International Materials

    Reviews, vol. 34, no. 5, pp. 247-260 (1989)

    6. Guzowski, M.M., Sigworth, G.K., Sentner, D.A.,The Role of Boron in the Grain Refinement of

    Aluminum with Titanium, Metallurgical &

    Materials Transactions A, vol. 18A, pp. 603-620

    (1987)

    7. Mohanty, P.S., Samuel, F.H., Gruzleski, J.E.,Studies on Addition of Inclusions to Molten

    Aluminum Using a Novel Technique, Metallurgical

    & Materials Transactions B, vol. 26, no. 1, pp. 103-109 (1995)

    8. Hatch, J.E. (Ed.),Aluminum Properties and Physical

    Metallurgy, 1st ed., American Society for Metals,Metals Park, Ohio (1988)

    9. Mahmudi, R., Sepehrband, P., Ghasemi, H.M.,Improved Properties of A319 Aluminum Casting

    Alloy Modified with Zr, Materials Letters, vol. 60,

    pp. 2606-2610 (2006)

    10. Sepehrband, P., Mahmudi, R., Khomamizadeh, F.,Effect of Zr Addition on the Ageing Behavior of

    A319 Aluminum Cast Alloy, Scripta Materialia,

    vol. 52, no. 4, pp. 253-257 (2005)

    Paper 10-042.pdf, Page 16 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    54

  • 7/25/2019 Artigo Carlos

    17/18

    11. Yin, Z., Pan, Q., Zhang, Y., Jiang, F., Effect ofMinor Sc and Zr on the Microstructure and

    Mechanical Properties of Al-Mg Based Alloys,

    Materials Science and Engineering, A280, pp. 151-

    155 (2000)

    12. Crepeau, P.N., Effect of Iron in Al-Si CastingAlloys: A Critical Review, AFS Transactions, vol.

    103, pp. 361-366 (1995)

    13. Couture, A., Iron in Aluminum Casting Alloys,AFS International Cast Metals Journal, vol. 6, no. 6,

    pp. 9-17 (1984)

    14. Bonsack, W., Discussion on the Effect of MinorAlloying Elements on Aluminum Casting Alloys,

    ASTM Bulletin, pp. 45-51(1942)

    15.ASM Handbook, Properties and Selection:Nonferrous Alloys and Special Purpose Materials,

    vol.. 2, ASM International, Materials Park, OH

    (1990)16. Polmear, I.J., The Effects of Small Additions of

    Silver on the Ageing of Some Aluminium Alloys,

    Trans. Met. Soc.,A230, p 1331 (1964)

    17. Vietz, J.T., Polmear, I.J., The Influence of SmallAdditions of Silver on the Ageing of AluminumAlloys. Observations on Al-Cu-Mg Alloys, Inst

    Metals J, vol. 94, no. 12, pp. 410-419 (1966)

    18. Garg, A., Chang, Y.C., Howe, J.M., Precipitation ofthe Omega Phase in an Al-4.0Cu-0.5Mg Alloy, Scr.

    Metall. Mater.,vol. 24, no. 4, pp. 677-680 (1990)

    19. Zhu, A., Gable, B.M., Shiflet, G.J., Strake Jr., E.A.,Trace Element Effect on Precipitation in Al-Cu-Mg-

    (Ag, Si) Alloys: a Computational Analysis, ActaMaterialia, vol. 52, pp. 3671-3679 (2004)

    20. Zou, Y., Matsuda, K., Kawabata, T., Himuro, Y.,Ikeno, S., Effects of Ag on Age-Hardening Behavior

    of Al-Mg-Si Alloys, Materials Forum - Institute ofMaterials Engineering Australasia Ltd, vol.. 28, pp.

    539-544 (2004)

    21. Matsuda, K., Fukaya, K., Young, Z., Kawabata, T.,Uetani, Y., Ikeno, S., Effect of Copper, Silver andGold on Tensile Behavior in Al-Mg-Si Alloy,

    Materials Forum - Institute of Materials Engineering

    Australasia Ltd, vol. 28, pp. 424-428 (2004)

    22. Hardy, H.K, The Effect of Small Quantities of Cd,In, Sn, Sb, Ti, Pb, or Bi on the Aging Characteristics

    of Cast and Heat Treated Al-4%Cu-0.15%Ti Alloy,

    Journal of the Institute of Metals, vol. 78, pp.. 169-

    194 (1950)23. Silcock, J.M., Heal, T.J., Hardy, H.K., The

    Structural Aging Characteristics of Ternary Al-CuAlloys with Cd, In, or Sn, Journal of the Institute of

    Metals, vol. 84, no. 1, pp. 23-31 (1955)

    24. Ringer, S.P., Hono, K., Sakurai, T., The Effect ofTrace Additions of Sn on Precipitation in Al-Cu

    Alloys: An Atom Probe Field Ion Microscopy

    Study,Metallurgical and Materials TransactionsA,vol. 26A, pp. 2207-2217 (1995)

    25. Grebenkin, V.S., Silchenko, T.V., Gorshkov, A.A.,Dzykovich, I.Y., Effect of Magnesium on the

    Distribution of Tin and Lead in Al-Si Alloys,

    Metallovedenie i Termicheskaya Obrabotka

    Metallov. (Metals Science & Heat Treatment), vol. 3,

    pp. 50-54 (1972)26. Salnikov, V.P., Zaigraikin, A.G., Effect of Bismuth

    Addition on Properties of Aluminum-Silicon Alloys,

    The Bulletin of the Bismuth Institute, no. 19 (1978)27. Pillai, N.P., Anatharaman, T.R., Elements of V

    Group as Modifiers of Aluminum-Silicon Alloys,

    Transactions of the Metallurgical Society of AIME,

    vol. 242, pp. 2025-2027 (1968)

    28. Cho, J.I., Loper, C.R., Jr., Limitation of BismuthResidual in A356.2 Al, American FoundrymensSociety, vol. 108, no. 64, pp. 359-367 (2000)

    29. Elhadad, S., Effect of Trace elements on theMicrostructure and Porosity Formation in 319 TypeAl-Si-Cu Alloys, M.Sc. Thesis, Universit du

    Qubec Chicoutimi, Canada (2003)

    30. Sigworth, G.K., Controlling Tensile Strength in

    Aluminum Castings, Private Communication,(2007)31. Jorstad , J.L., Aluminum Casting Technology, 2nd

    ed., American Foundrymens Society, Des Plaines,

    IL, USA (1993)32. Reif, W., Yu, S., Dutkiewicz, J., Ciach, R., Krol, J.,

    Pre-Ageing of AlSiCuMg Alloys in Relation to

    Structure and Mechanical Properties, Materials and

    Design, vol. 18, no. 4, pp. 253-256 (1997)

    33. Mishra, R.K., Smith, G.W., Baxter, W.J., Sachdev,A.K., Franetovic, V., The Sequence of Precipitationin 339 Aluminum Castings, Journal of Materials

    Science, vol. 36, no. 2, pp. 461-468 (2001)

    34. Li, R.X., Li, R.D., Zhao, Y.H., He, L.Z., Li, C.X.,Guan, H.R., Hu, Z.Q., Age-Hardening Behavior of

    Cast Al-Si Base Alloy, Materials Letters, vol.. 58,

    pp. 2096-2101 (2004)

    35. Shabestari, S.G., Gruzleski, J.E., Modification ofIron Containing Precipitates in AlSi12 Alloys withStrontium, Giesserei-Praxis (Germany), vol. 17, pp.

    385-394 (1997)

    36. Liao, H., Sun, G., Mutual Poisoning Effect BetweenSr and B in Al-Si Casting Alloys, Scripta

    Materialia, vol. 48, pp. 1035-1039 (2003)

    37. Liao, H., Sun, Y., Sun, G., Effect of Al-5Ti-1B onthe Microstructure of Near-Eutectic Al-13.0%Si

    Alloys Modified with Sr,Materials Science, vol. 37,

    pp. 3489-3495 (2002)38. Abis, S., Mengucci, P., Riontino, G., Influence of Si

    Additions on the Ageing Process of an Al-Cu-Mg-Ag

    Alloys, Philosophical Magazine A, vol. 70, no. 5,pp. 851-868 (1994)

    39. Gable, B.M., Shiflet, G.J., Strake Jr., E.A., TheEffect of Si Additions on Precipitation in Al-Cu-Mg-(Ag) Alloys, Scripta Materialia, vol. 50, pp.

    149-153 (2004)

    Paper 10-042.pdf, Page 17 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA

    55

  • 7/25/2019 Artigo Carlos

    18/18

    40. Muddle, B.C., Ringer, S.P., Polmear, I.J., HighStrength Microalloyed Aluminum Alloys,Advanced

    Materials93 VI/Frontiers in Materials Science

    Engineering, pp. 999-1023 (1994)

    41. Polmear, I.J., Pons, G., Octor, H., Sanchez, C.,Morton, A., Borbidge, W., Rogers, S., AfterConcorde: Evaluation of an Al-Cu-Mg-Ag alloy for

    Use in the Proposed European SST, Materials

    Science Forum, vol. 217/222, no. 3, pp. 1759-1764(1996)

    42. Mohamed, A.M.A., Samuel, F.H., Samuel, A.M.,Doty, H.W., Valtierra, S., Influence of Tin Addition

    on the Microstructure and Mechanical Properties ofAl-Si-Cu-Mg and Al-Si-Mg Casting Alloys,

    Metallurgical and Materials Transactions A, vol.

    39A, pp. 490-501 (2008)43. Couper, M.J., 6XXX Series Aluminum Alloy, U.S.

    Patent No. 6,364,969B1 (2002)

    Paper 10-042.pdf, Page 18 of 18AFS Transactions 2010 American Foundry Society, Schaumburg, IL USA