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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
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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
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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
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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.
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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)
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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
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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
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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
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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)
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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)
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-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.
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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)
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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)
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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)
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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)
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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.
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