THE 4TH INTERNATIONAL CONFERENCE ON ALUMINUM ALLOYS

Effect of Mg Addition on Precipitation Hardening of Al-Li-Cu-Zr Alloys
Containing Ag
Shoichi Hirosawa l , Tatsuo Satol, Akihiko Kamio l ,Kazunori Kobayashi2 and Toshimasa Sakamoto 2
1. Dept. of Metallurgical Engineering, Tokyo Institute of Technology,
2-12-1 O-okayama, Meguro-ku, Tokyo 152, Japan.
2. Alithium Ltd., 2-2-2 Ogoso-Higashi, Yokkaichi-shi, Mie 510, Japan.

Abstract-The AI-Li-Cu-Mg-Zr alloy containing Ag, known as Weldalite'lM049 alloy, exhibits ultrahigh mechanical strength caused by homogeneous precipitation of the Tl(AhCuLi) phase, andshows excellent weldability with marked natural aging response due to high Cu content. In
particular, small addition of Mg has a significant role to attain these characteristic properties.
The present work aims to investigate the effect of small amount of Mg addition on
precipitation hardening of AI-Li-Cu-Ag-Mg-Zr alloy at various aging temperatures. The
mierostructures observed with transmission electron microscopy and hardness changes indicate
the pronounced effects of Mg addition. In the temperature range from R.T. to 373K small
amount of Mg stimulates precipitation of GP(l) zone, resulting in the marked age-hardening.
At the more elevated temperatures, however, the Mg addition results in the change in the
predominant precipitates from the GP(2) zone and/or e' phase (Mg-free alloy) to the Tl phase:.\t1g-containing alloy). The mechanism for the effect of Mg addition is discussed based on~oth the nucleation of precipitates and the vacancy-solute atom interaction.
Introduction~he application of a low density and high modulus AI-Li alloy over a wide range of structuralcomponents requires both desirable strength-ductility combination and good weldabiIity. Fromthese alloy design concepts, Weldalite'lM049 alloy was developed by Martin Marietta
Corporation, having a high Cu/Mg(mass%) ratio and small addition of Ag with a normalcomposition of AI-( 4.5-6.3)Cu-l.3Li-0.4Ag-0.4Mg-O.14Zr (in mass%). The ultra-highstrength of the alloy (yield strength of ~ 690MPa) in either T8(stretch and artificial aging) or
T6(no stretch and artificial aging) temper is attributed to homogeneous precipitation of the Tl(AhCuLi) phase with a presence of 0' (AbLi), GP zone, e' (AI2Cu) and S' (AI2CuMg) phases.
In addition, this alloy exhibits excellent weldability due to both the marked aging response in
the T4(no stretch and natural aging) temper and a large amount of eutectic compounds to heal
hot tears formed during the welding operation. The high mechanical strength in the T4 temper(yield strength of ~ 440MPa) arises from a combination of refined GP zones and the 0'
precipitates, without any other phases recognized.

199

Small additions of Ag and Mg are utilized to stimulate fine distribution of the Tl phase in theartificial aging. Polmear et al.[l] reported that small amounts of Ag and Mg have marked
effects to increase the age-hardening response for AI-Cu alloys, resulting from fine dispersionof hexagonal-shaped thin plates of the Q phase on the {Ill} a planes. Furthermore, Taylor etal.[2] proposed that the nucleation of the Q phase takes place on the fine Mg3Ag particles
formed on the {Ill} a planes. This seems to be applicable to explain the nucleation of the Tl
phase in Weldalite lM 049 alloy.
The goal of this work is to examine the role of small amount of Mg addition on the
precipitation hardening of Weldalite™049 alloy in the temperature range from R.T.(roomtemperature) to 523K. The predominant phases effective to increase mechanical strength atvarious aging temperatures and time were identified using transmission electron microscope(TEM) technique. The measurement of precipitate size distributions was also performed todetermine precipitate populations and formation kinetics of each phase. From these results, themechanism for the effect of Mg addition is discussed based on both the nucleation of
precipitates and the vacancy-solute atom interaction.
Experimental procedure
The alloys were prepared from high-purity materials in argon gas atmosphere. The chemicalcomposition of the alloys is shown in Table 1. The ingots were homogenized at 723K for
36.4ks, then scalped, and hot and cold-rolled to 1.6mm thick sheets. All of the specimens were
solution-treated at 778K for 1.8ks in a salt bath, and followed by water quenching. The aging
treatment was carried out in an oil bath at temperatures from R.T. to 523K for various periods.
Hardness measurements were made on the as-polished specimens using a Micro Vickers
testing machine with a 500g load. The specimens for the tensile test were prepared from the
1.0mm thick sheets in the longitudinal direction with a gauge length of 25mm. Tensile tests
were performed with a Instron-type tensile test machine at a strain rate of 1.3xlO-4/s. The
specimens for transmission electron microscopy (TEM) analysis were prepared by the twin-jet
polishing technique in a solution of 25 vol. % nitric-acid and 75 vol. % methanol at about 250K.
TEM observation was carried out at an accelerating voltage of 200kV using a JEM200CX
transmission electron microscope.
Table 1 Chemical compositions of Weldalite and Mg-free alloys (mass%).
Weldalite
Mg-free

5.52
5.65

0.50
0.002

0.39
0.38

1.36
1.35

0.16
0.15

0.08

0.05

0.03

bal.
bal.

I.Aging behavior at R.T. to 373K
Fig.l shows the isothermal aging curves of hardness for Weldalite and Mg-free alloys aged at
R.T. to 373K. With increasing aging time Weldalite alloy displays a rapid and pronounced
hardening to attain a constant value at each aging temperature. On the other hand, although thehardncss increment is obviously retarded, Mg-free alloy also shows increased age-hardening
to attain the almost identical values of hardness after extended incubation periods (sec curvesof 373K shown in Fig.l).

200/

Fig.2(a) shows a bright field TEM image for Weldalitc alloy aged at R.T. for 864ks. The
corresponding diffraction pattern exhibits diffuse continuous strcaks and superlatticereflections, showing the presence of a combination of refined GP(l) zones and thc (5 I phase. It
is well known that the (5 I phase is rapidly formed from the supersaturated solid solution during
or just after water quenChing, whereas the growth of the (5 I phase occurs to it lesser extent at
low temperatures such as R.T.. Therefore, the marked natural aging response of
Weldalite™O49 alloy seems to be mainly due to the growth of GP(l) zones. In contrast, themicrostructure of Mg-free alloy aged at R.T. for 864ks reveals almost no phases except the (5 I
phase. Howevcr, increased aging time (~2.7xl()5ks) stimulates the formation of GP(l) zones
with the imagc contrast due to the strain field (Fig.2(b)). This implies that small addition of Mg
has a significant role to produce a rapid natural aging response for Weldalite™O49 alloy
associated with GP(J) zones. The similar results are also obtained for both Weldalite and Mgfree alloys aged at 373K. Thus, the addition of small amount Mg is found to accelerate theprecipitation of GP(\) zones resulting in the faster age-hardening. In this aging temperature
range the T1 phase is not observed.
1110
The tensile properties of two alloys arc'" •well correlative with the hardening
behavior. The stress-strain curves for two ii;"!(jLas-quenched alloys reveal serrated flow oIT'as frequently observed in the plastic ~'"' Ilf),(
deformation of AI-Li alloys. IIowever,
J(
I'
,I'
after natural aging for 81i.4ks the stress- c,lI){'
strain curve becomes smooth only for :¥Weldalite alloy, in contrast, that of Mgfree alloy still represents serrated flow.
Af)
10
l'l
The disappearance of the serrated flow in/I~Jlllq (Ull'!, r I k<;n1
the naturally aged Weldalite 049 alloy is
Fig.1
Isothermal aging curves of hardness foralso attributed to the effect of Mg
Weldalitc and Mg-frce alloys aged at R.T.
addition.
to 373K.

...
/
'J

I,

Fig.2

I,

Electron micrographs with diffraction pattern for (a) Weldalite

alloy aged at R.T. for 864ks and (b) Mg-frcc alloy aged at
R.T. for -2.7xl()5ks.

201

2.Aging behavior at 403 to 523K
Fig.3 illustrates the isothermal aging
2/0
curves of hardness for Weldalite and Mg20()
free alloys aged at 403 to 463K. The >I 18(J
results indicate that the small amount of "Mg reduces the aging time to reach peak ~I 150
hardness with a marked increase in peak :£"E 1'1()
hardness. This can be explained in terms of i'!
.fj 120the change in the microstructures produced ;;duringaging.
The
predominant b 100
strengthening precipitates identified by ;;. eo .
TEM for various aging conditions arc
1(J'1
p/
10'
10 J
listed in Table 2. After aging at 433K for
A.Q
IIJ
1\1]1119 tim!"!, 1/'5
36ks, for example, both Weldalite andMg-free alloys exhibit microstructures Fig. 3 Isothermal aging curves of hardness for
Weldalite and Mg-free alloys aged at 403
containing GP(l) zones and the 0' phases.
to 463K.
However, after aging for 346ks Welda lite
alloy alone produces the T1 phase with a Table 2 Precipitates observed iin Weldalite andMg-frce alloys aged at R.T to 523K.
presence of 0 ' phase and GP(2) zone,
Tcmp. Timclks
l\lg-frcc
Wcldalilc
whereas in Mg-free alloy the T1 phase isRT
8.6 X 10'
6'
6' Cil'( I)not detected until the time of overaging(6') GI'(I)
2.7X Hr'(1.9xl0 4 ks) (FigA(a),(b),(c)). The similar
3731(
6' GP(l)1.8X 10'
6' CiI'(I)
results are also demonstrated [or two alloys
0'GI'(I)
403K
1.8
X
10'
0'01'(1)aged at 463K. Thus, the addition of Mg
9.7X !O'0' OP(2) T,
the major
causesthechange inGI'B(2)
strengthening phase from the GP zones
6' Ul'( I)4331( 3.6X 10
6'Cil'(I)6' GI'(2)3.SX 10'
0'GI'(2) T,
and/or e' phase to the T1 phase, resulting(0') e' T,
1.9X W'in the marked acceleration of the T1 phase(6 ')GI'(2)
4631(
7.2
T,
formation.
Mg-containing precipitates
6 'OP(2) T,.l.6X 10T,such as GPB zone and the S' phase are not(0') 0' T,
1.9X Hr'T,
detected in Wcldalite alloy, except for5211(
T,
9.0X 10' T,
small amounts of GPB(2) zones observed
T,
1.8
T,
in the aging at 403K for 9.7x10 3ks.
~

Discussion/

1. Mechanism for Mg addition at RT. to 373K
In the temperature range from R.T. to 373K, an obvious difference in the aging respon£::between Weldalite and Mg-free alloys (Fig. I) can be mainly explained in terms of the chang::in the precipitation rate of GP(I) zones. It is well known that the addition of small amollnt of
Mg to AI-Cu alloys accelerates the formation of GP(l) zones resulting in the faster natural
age-hardening, whereas the additions of other clements such as Cd, In and Sn lower the natural
age-hardening rate by suppressing the formation of GP(l) zones. Entwistle et al.[3] proposedthat Cu/Mg/vacancy complexes are formed at the initial stage of aging due to a high Mgvacancy binding energy, and promote copper diffusion necessary to form GP(l) zones by aprocess of reshuffling of the associated vacancies. Furthermore, Wyss et a/.[ 4] suggested that

202

...

Fig.4

Electron micrographs with diffraction pattern for (a) Wcidalite alloy aged at 433K
for 346ks and Mg-free alloys aged at 433K for (b)346ks and (c) 1.9xl()1ks.

thc formation of Mg-vacancy clustcrs cnhances nucleation of GP(1) zones by increasing thenumber of nucleation sites or increasing the nucleation rate. It is likely that the similar
mechanisms for Mg addition arc applicable to the Li-containing Weidaiite™O49 alloy.
In general, the diffusion coefficient of Cu in AI is given as [5J
Deu = A exp(-EmlkTa) exp(-Ef/kTq),(1)

where Em and Ef are the activation energies for the migration and formation of vacancies, Ta
and Tq are the aging and solution treatment temperatures, k is Boltzmann's constant and A is aconstant. Howevcr, if the quench operation is carricd out quickly enough to retain the vacancy
concentration associated with Tq, the clustering process of Cu atoms should be mainly
governed by the first exponential term because an activation energy Ef for the formation ofvacancies becomes negligible. In fact, an activation energy for the formation of GP zones in
as-quenched AI-4%Cu alloy, O.54eV, agrees well with an activation energy Em for thc:-nigration of vacancies, O.51eV.
JOr--r--r--r--r--r--r--r---.
:n the present study, the activation energies
5 --- --.cor the formation of GP(1) zones in
Neldalite and Mg-free alloys can be:stimatcd by utilizing the isothermal agingo:urveS of hardness (Fig.l). That is, if the:lardness increment is proportional to the :2 -5
volume fraction of the formed GP(l) zones 5'
to a tta in i d en ti ca I v alu es of hard ness,
Arrhenius plots of the reciprocal time lit, -10
obtained from the cross-cut at 50')'0 agehardening from as-quenched hardness, ·15
would provide the activation energies forthe formation of GP(l) zoneS. In fact, from-20 ~~L...~L...~L-..~L...~L...~L......~L....~
0.003
0.004
o
0.002
0.001
the slope of each straight line through the
Ifr
Arrhenius plots in Fig.5, the activation
Fig.S Arrhenius plots of reciprocal time to 50% ageenergies are assessed to be O.38eY andhardening from as-quenched hardness against
reciprocal temperatures for Wcldalite and MgO.5geY for Weldalite and Mg-free alloys,
frec alloys.

203respectively. In contrast, the jump frequency
Table 3 The calculated enthalpy of solutionof vacancies or the nucleation frequency ofat infinite dilution, ~Ho in kJ/mol[8j.
GP(1) zones, given by an intercept on theIn(1/t) axis in Fig.5, is found to be
AI Li OJ Mg Ag Ql In Sn
approximately equal in two alloys. TheseAI-15 -34 -7 -17 12 24 14
results indicate that small amount of Mg in
Li -13Weidalite™049 alloy remarkably accelerate-19 -1 -55 -43 -41 -56
the growth rate of GP(l) zones, resulting in a
8 -3
OJ -28 -18
2-15 8
decreased activation energy for the formation,-42 -22 -14 -31
Mg -8 -1 -20
whereas the number of nucleation sites or the-8 -5 -10
Ag -18 -65 10 -40
nucleation rate is not affected. The similar
results arc also obtained for the cross-cut at
30% age-hardening, the activation energy for the formation of GP(1)zone in Weldalite alloy(0.40eV) is rather smaller than that in Mg-free alloy (0.56eV), which is almost equal to theactivation cnergy for the migration of vacancies (0.51eV).
These can be interpreted in terms of the Cu/Mg/vacancy complexes diffusion propose~ by
Entwistle et a/.[3]. Although a number of arguments have been done on a Mg-vacancy
interaction, the relatively high Mg-vacancy binding energy compared with Cu-vacancy
binding energy (values of 0.19cV[G] and 0.36 ± 0.04eV[7J arc mcasured) will contribute toform numcrous Mglvacancy pairs during the solution treatment. These pairs persist through thequench operation, and might acquire copper atoms because of the strong Mg-Cu interaction
predicted based on the calculated cnthalpy of solution at infinite dilution. Table 3 shows theextracted values from the lists reported by Niessen et a/.[8J. Therefore, the diffusivity ofcopper atoms in Weldalite'IM049 alloy increases due to the increased frequency of Cu atoms toencounter vacancies through Mg atoms. In contrast, the elements having high binding energies
with vacancics such as Cd(0.18eV), In(O.25eV) and Sn(O.22eV) do not have so strong
interactions between Cu atoms as shown in Table 3, thereby inhibit the formation of GP(l)zones by trapping vacancies available for Cu diffusion.
/
2. Mcchanism for Mg addition at 403 to 523K
On the other hand, in the temperature range from 403 to 523K, the precipitate size distributions
and precipitate densities of the strengthening phases such as the TI phase, GP(2) zone and the
IM 049 alloy. The
I phase were mcasured to determine the effect of Mg addition in Weldalite'
measurement was performed on the bright field TEM images in various aging conditions in
Weldalite and Mg-free alloys, assuming the thickness of the TEM foils to be - 10011111. The
average precipitate sizes were determined from the size distributions of edge-on precipitateson the {Ol1} a orientations, and the precipitate densities were assessed as the precipitate
numbers including other variants per lcm 3 .

eFig.G shows the size distribution changes of each strengthening phase produced during aging at
433K, in which the Mg addition exerts the pronounced influence on the aging behavior of
Weldalite™049 alloy. It is evident that the small amount of Mg causes the high density
precipitation of the TI phase at the peak-aging time as 34Gks (Fig.G(a». In contrast, althoughMg-free alloy has a rather high precipitate density of finc G P(2) zone, thc Tl phase is 110t

204

detected until the time of overaging (1.9xl04ks)(Fig.6(b),(c». This implies that the
10 r--',lL
small amount of Mg is effectively utilized for "'senhancing nucleation of the Tl phase, which ~ 8 f-results in the improved mechanical strength. ~ 7
Pickens et al.[9] suggested that local ordering
6
including Ag and Mg atoms could lead to the '",~5formation of the Tl nucleus, as has been '"n~ 4shown for the ~" phase in AI-Cu alloys [2].
~ 3However, the mechanism for the refined ..,'5. 2
distribution of the Tl phase is not fully 'uinterpreted in terms of the nucleation aids of P-.~ 1
such fine Mg3Ag particles because an Ag-free0(b)alloy containing identical composition with100
1M
Weldalite 049 alloy also exhibits the similar "'s 90
marked aging response caused by the ~ 80
homogeneous precipitation of the Tl phase.;:;o 70
The reduction in the peak hardness is only
60
about HVIO under the aging conditions of 1';- 50
both 433 and 463K[lO].
.~

mT

J

o GI'(2) zone

J

x

-

-

-I~ 10-I.

0(;P(2)zone

r--

x
~

r

f-- -

- -

-

r---

-

-

- - -

-

-

-

40

~ 30

--

As the Tl phase has a hexagonal structure !'l.- r-- 1- - - (a=0.497nm and c=O.93Snm) with a lattice ·.2l.
-u 20
1plane stacking sequence AOAOAO"', the10stacking faults are inevitably introduced ill the0(rlfcc structure of aluminum, AOCAOAOC"', for10
the nucleation and growth of the Tl phase. In "'s 9r-- r-- - II J
T ~- . CJ 0'
general, the stacking fault energy of aluminum,:;:' 8a relatively high value of ~ O.lSJm-2, is '0
Iknown to he lowered by trace additions of xf-- - - -t-- 1';elements such as zinc, magnesium, lithium, 'Uj. - -- - - - -silver
andgermanium
due
totheir ~
r-- - - -'-' - - --- - comparatively high solubilities at the usual a;r-r--- -. - r--aging temperatures[ll]. Thus, the lIucleation l:l- f-- - - - :and growth of the Tl phase is considered to be
2
tTIuch easier to take place in Mg-containing 0:'If] I• ., I.".D::l.
Weldalite alloy because of the reduced
0 0 0 0 000..,. '-?
., 0 0 '.D0 00..,. 0.,...,. 0ry 0 Nw~ N
stacking fault energy. However, it is also
'"
'"'0
0
a..,. "., '" 6'" <D 0..., ."..,. O'J ".,necessary to investigate in more details the'" n ;:: ,.,0 n N00 "'"Precipitate size / TIm
mechanisms of the enhanced nucleation andgrowth of the Tl phase by the small addition
Fig.6 Precipitate size distributions for (a) Wcldaliteof Mg from other points of views, i.e. the
~

J:

n

-~1

~

:E'

:[ IIJ1]

0')

if)

H

~,

'" '" '" '" '" '" '" '" '" '"0')metastable phase boundaries, the dislocation
loops amI heliees introdueed with the exeess
vacancies.

205

'"

0')

alloy aged at 433K for 346ks and Mg-freealloys aged at 433K for (b) 3~ks and (e) 1.9
x I O\s. Micrographs arc shown in PigA.

Conclusions
The addition of small amount Mg to AI-Li-Cu-Zr alloys contammg Ag, known as
Weldalite™049 alloy, significantly exerts an effect on the precipitation hardening over thetemperature range from R.T. to 523K. The marked increment of mechanical strength is foundto arise from the changes in the precipitation rate of GP(1) zones and the TI phases. In thetemperature range from R.T. to 373K the small amount of Mg accelerates the formation of
GP(1) zones resulting in the faster age-hardening. This can be explained in terms of a decreasein an activation energy for formation of GP(1) zone due to the Cu/Mg!vaeaney complexes
diffusion. On the other hand, at the more elevated temperatures the Mg addition causes theI phase to the TI phase,change in the major strengthening phase from GP zone and/or
resulting in the marked acceleration of the Tl phase formation. This is considered to be mucheasier in the nucleation and growth of the Tl phase due to not only the nucleation aids of the
Mg3Ag particles but also the reduced stacking fault energy caused by the small addition of Mg.

eReferences
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2. J.ATaylor, B.AParker and I.J.Polmear, Metal Science 12, (1978), 478.
3. K.M.Entwistle, J.H.FeIl and Kang II Koo, J. Inst. Metals, 91, (1962-63),84.
4. R.K.Wyss and R.E.Sanders, Jr., Metall. Trans. 19A, (1988), 2523.
5. AKeIly and R.B.Nicholson, Precipitation Hardening, Prog. Mater. Sci.lQ, (1961), 148.
6. AJ.Perry and E.M.Entwistle, 1. Inst. Metals, 96, (1968), 344.
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8. AK.Niessen, F.R.de Boer, R.Boom, P.F.de Chfttel and W.C.M.Mattens, CALPHAD, 1,(1983),51.
9. LR.Pickens, F.H.Heubaum T.J.Langan and L.S.Kramer, Proc. 5th Int. AI-Li Conf., VoJ.3,(1987), 1397.
10. AKamio, T.Sato, S.Hirosawa, K.Kobayashi and Y.Tsuji, Proc. 84th Conf. Japan Inst.
Light Metals, (1993), 389.
11. V.D.Scott, S.Kerry and R.L.Trumper, Mat. Sci. Tech.,~, (1987), 827.

/

206