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THE ROLE OF MAGNESIUM IN Al-Li-Cu-Mg-Zr ALLOYS
S. Harris, B. Noble, K. Dinsdale
To cite this version:
S. Harris, B. Noble, K. Dinsdale. THE ROLE OF MAGNESIUM IN Al-Li-Cu-Mg-Zr ALLOYS.
Journal de Physique Colloques, 1987, 48 (C3), pp.C3-415-C3-423. �10.1051/jphyscol:1987347�. �jpa-
00226578�
THE ROLE O F MAGNESIUM IN Al-Li-Cu-Mg-Zr ALLOYS
S.J. HARRIS, B. NOBLE and K. DINSDALE
Department of Metallurgy and Materials Science, University of Nottingham, University Park, GB-Nottingham NG7 2RD, Great-Britain
ABSTRACT
Within the U R , research has concentrated on alloys of composition A1--2.5%Li-1.2%Cu-0.7%Mg-O.l2%Zr (8090) which is a medium strength replacement for 2014. Ageing this lithium-containing alloy at. an elevated temperature produces a mixture of phases which are, in order of importance, S'(A13Li), S(A1,CuMg) and TI(A12CuLi). Changing the magnesium content of the alloy changes the relative proportions of the latter two phases, and also infl.uences the volume fraction of S' that is produced. A systematic study has been made of the effect of magnesium (in thc range 0-2%) on the relationship between the mi.crostructure and mechanical properties of these alloys.
INTRODUCTION
The addition of lithium to aluminium produces an alloy of low density and high elastic modulus and therefore the use of such alloys for various aircraft components results in substantial weight. savings (1). Two alloy systems AI-Li-Cu and A1-Li-Cu-Mg have received the most intensive study in respect of property-microstructure evaluation. Both types of alloy are capable of providing higher strength and tougher alloys than the binary A1-Li system. To both alloy systems an addition of 0.12% zirconium is made which retards recrystallisat.ion from taking place during conventional mechanical working and heat treatment procedures.
Alloys containing bet,ween 1.4 and 4 wt.% lithium are age hardenable due to the formation of a S'(A13Li) transition phase ( 2 , 3 ) . This phase has an Li2 ordered structure and forms as spherical parti.cles of a size 8--40 nm and volume fraction up to 304.
When copper is present in A1--Li alloys another phase, T1(A12CuLi), appears in the microstructure (4-7). The crystal structure of this phase is hexagonal and it grows as plates on (1.11) aluminium planes. If now small amounts of magnesium are added to the Al--Li-Cu alloy then the TI phase will still. precipitate; however, as the magnesium:copper ratio in(-reases to =2 the TI phase is replaced by S(A1,CuMg).
This latter phase is well characterised in the Al-Cu-Mg system where it has been shown to form as laths on {210} planes (8). Peel et a1 (9,lO) have demonstrated that in lithium-containing alloys the formation of S phase is very sluggish and it requires dislocat,ion sites for its nucleation. To promote S precipitation it is therefore necessary to mechanically deform the alloy before ageing (T8 treatment) in order to increase the number of nucleation sites. T1 precipitation is also encouraged by deformation prior to ageing.
Very little information is available on the relative importance of the TI and S phases on mechanical properties of A1-Li based alloys. The present work was instigated to provide such information. A base alloy of A1-2.5%Li-1.2%Cu--0.12%Zr has been used and to this a series of magnesium additions (0-2 wt.56) has been made.
The resulting compositions relative to a section of the AI-Cu-big ternary phase diagram are shown in fig. 1 (11).
Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987347
JOURNAL DE PHYSIQUE
Fig. 1. Phases present after long ageing times at 190'C in the A1-Cu-Mg system. X-composition of A1-2.5XLi- 1.2%Cu--Mg-0.12%Zr alloys investigated.
0 1 2 3
wt% Mg.
The compositions studied encompass those of commercial 8090 type alloys (A1--2.5%Li-1.2%Cu-0.7%Mg) and Al-Li-Cu alloys such as Al-3%Li-l%Cu; the work therefore provides information on the strengthening mechanisms that may operate in these commercial alloy systems.
EXPERIMENTAL
All alloys were prepared from high purity materials melted and cast to 32mm diameter ingots under argon. Alloy compositions are listed in Table I. The alloys were reheated to 520'C for lh and hammer forged to 15mm thick slabs. On reheating again to 520'C the slabs were hot rolled to 4mm strip. This strip was water quenched from 520'C and cold rolled to 2.5mm thick material. The final stage of forming was to reheat again to 520'C, water quench and cold roll to 1.6mm strip.
Table I Composition of Alloys
Composition (wt.%) Density
Alloy Code Li Cu Mg Zr (Mg/m3)
(Fe <0.035% and Si (0.02%)
Heat treatment of the 1.6mm strip involved solution treatment at 520'C for 20 minutes followed by cold water quenching. Where aapropriate, solution treated material was cold stretched 2.5% prior to ageing. Ageing was carried out for 16h at 190'C in oil baths with material in the unstretched (T6) or stretched (T8) conditions. For the purpose of comparing the properties of the alloys before ageing (ST condition) and after ageing, samples were refrigerated prior to carrying out the mechanical tests.
Tensile tests were carried out on specimens machined parallel to the rolling direction and at an initial plastic strain rate of 8 x 10-5/s. The average of at least two tests are reported for each alloy composition and heat treatment condition. The microstructures of the alloys were examined by transmission
Hardness tests
A survey of hardness (fig. 2) was carried out on all alloys heat treated in the ST, T6 and T8 conditions. In the ST condition the hardness of the alloy steadily, increased with magnesium content (range 0-2%) from 73 to 85 HV. By ageing to the TF condition the hardness of the magnesium-free alloy increased by 61 HV.
Magnesium additions up to 2% produced further increases in hardness of 20 HV to 153 HV. This represents an increase in hardness increment above that achieved by magnesium in the ST condition. The introduction of a 2.5% stretch prior to ageing (T8 condition) increased the hardness of the Al-Li-Cu alloy by 8 HV and this increment was maintained at a reasonably constant level, above that achieved by T6 treatment, as the magnesium content increased.
.
T8 peakT6 peak
I Fig. 2. The effect of magnesium on
120 - the hardness of A1-2.5%Li-1.2XCu-klg-
- 0.12%Zr alloys in the as-quenched, T6 and T8 peak aged conditions. T6 and
100
-
T8 aged at 19O'C.Fig. 3. The effect of magnesium on the tensile properties of A1--2.52Li- 1.2%Cu-Mg-0.12%Zr alloys in the as- quenched condition. L direction.
JOURNAL DE PHYSIQUE
Tensile tests
The tensile results obtained on the alloys in the ST condition are shown .in fig. 3. Both 0.2% proof stress and tensile strength increased with magnesium content, the majority of the data fitting a straight line relationship over the composition range examined. There is a tendency for the difference between proof and tensile strength to increase with increasing magnes-ium content, i.e. from 145 MPa to 170 MPa. Elongation to failure values obtained on the various alloys were averaged around 14% and appeared to be independent of magnesium concentration in the alloy.
In the T6 condition the proof and tensile strength results (fig. 4) do no1 follow a linear relationship with magnesium content. Relative to the ST data the proof stress values rise more steeply (up to 1% Mg) and thereafter more slowly t.o the 2% Mg level. The tensile strength generally follows the proof stress versus magnesium content plot, but the difference between the two -is halved t.o 275 MPa when compared with the ST data. Elongation values are generally in the range 4-64 in this TG condition end once again no significant change can be detrc:f.ecI with increasing magnesium content.
Fig. 4. The effect of magnesium on the Fig. 5. The effect of magnesium on tensile properties of Al-Z.S%Li-1.2%Cu- the tensile properties of A1-2.5XLi- Mg-O.l2%Zr alloys in t.he T6 peak aged 1.2%Cu--Mg-0.12XZr alloys in the T8 condition. L direction. (2.5% stretch) peak aged condition.
L direction.
The introduction of a 2.5% stretch between solution treatment and ageing increases the proof stress and tensile strength values of the alloys in the T8 condition (see P l g . 5). The proof stress for the magnesium-free alloy is approximately 30 MPa higher than that measured in the T6 condition. Once again relative to the ST data the proof stress rises more steeply up to 1% magnesium.
The tensile strength plot follows that for proof stress, but the difference between
compared with the T6 condition, see fig. 5. Increasing the n~agnesiunt content up to 0.75% restores the elongation to around 5% and this level of ductility is maintained up to 2% magnesium.
Transmission electron microscopy was carried out on thin foils of alloys treated in the T6 and T8 conditions. In all of these alloys spherical S'(A13Li) precipitates in the size range 5-20 nm were present. The alloys had not recrystallised to a significant extent and subgrains were therefore present. of a size =Zm. Since these features were present in all the all.oys they wil.1 not. be referred to in detail, instead attention will be given to the size and distribution o-F the TI and S phases, the nature of which changed significantly with magaesium concentrat.ion.
The magnesium--free alloy (Al--Li-Cu-Zr) in t.he T6 condition showed T1 precipitation on grain and sub-grain boundaries, see fig. 6a. The introducti.on of a 2.52 stretch prior to ageing encouraged a wider distribution of T1 on dislocation sites ifia. 6b).
Fig. 6. Transmission electron micrographs of A1-2.5%Li-l.2%CuO.l2%Zr.
a) T6 TI precitation on grain boundaries. Beam direction [110];
b) T8 T1 precipitation within grain interiors. Beam direction [110].
Fig. 7. Transmission electron micrographs of A1-2.5%Li-~1.2%Cu-0.41%Mg-O.l2%Zr.
a) & b) T8 TI and S precipitation on dislocations. Beam direction [llO].
C3-420 J O U R N A L D E PHYSIQUE
The addition of small amounts of magnesium into the alloys did not appear lo significantly affect the microstructure. However, when the magnesium level had increased to 0.4%, changes in precipitation type were found, particularly in the T8 condition. Fig. 7 (a and b) shows the presence of both S and TI precipitation.
At 0.7% magnesium the alloys showed increased amounts of S precipitation in both the T6 and T8 conditions. These precipitates were restricted to grain and sub-boundary sites in T6 and were more widely dispersed within the grain in the T8 condition, see fig. 8. Further increase of the magnesium concentration to 2% (fig.
8b and c) did not change the distribution of S phase precipitation in either condition. However, from observations made on numerous foils from alloys containing various magnesium concentrations in the range 0.7-2% it appeared that the volume fraction of S phase increased up to 1% magnesium but then remained essentially constant thereafter.
graphs of A1-2.5%Li-1.2%Cu-Mg-0.1 S precipitation within grains.
a) 0.68%Mg, Beam direction [110].
b) O.9S0%Xg, Beam direction [110].
c) 1.43%Mg, Beam direction [110].
DISCUSSION
The relationship between microstructure and proof stress will be dealt with for each of the heat-treated conditions (ST, T6 and T8) that have been studied.
Consideration will also be given to the effect of magnesium concentration on the tensile strength and ductility of the alloys.
ST condition
The variation of hardness and proof stress is linearly related to the magnesium content of the alloy (figs. 2 and 3), the increment being 18 MPa/wt.%Mg. This value is very similar to that observed for a A1-2.5%Li-Zr alloy to which various additions of magnesium had been made (fig. 9). By comparing the two plots it can be deduced that the 1.3% copper addition to the 8090 type alloy has produced a strengthening increment of =18 MPa when tested in the solution treated condition.
Fig. 9. Effect of magnesium on the proof stress of A1-Li alloys.
a A1-2.5%12i-Zr, T4.
b A1-2.5%Li-1.2%Cu-Zr, T4.
c A1--2. 5%Li-Zr, T6.
d Al-2.5%Li-1.2tCu-Zr, TG.
e A1--2.5%Li-I. Z%Cu-Zr, T8.
T6 condition
Ageing the alloys caused a large increase in proof stress at all levels of magnesium.
Without any magnesium addition to the alloy the increment in proof stress, relative to the ST condition, was 200 m a . The electron microscope observations indicated that this increment was caused by 8' and TI precipitation. Tests on an A1-2.5%Li-Zr alloy produced a (T6-ST) increment of 185 MPa; assuming that copper does not influence S' precipitation (14) it can therefore be deduced that T1 precipitation is producing a strengthening increment of =15 MPa. This is comparable with the value obtained by other workers for TI precipitation (13).
When magnesium additions are made to the Al-Li-Cu-Zr alloys the proof stress in the T6 condition is not linear with increasing magnesium concentration (fig. 4).
This behaviour is quite different from that observed with Al-Li-Zr alloys where the proof stress increases linearly with magnesium concentration when the alloys are tested in the T6 condition (fig. 9). The proof stress change in these copper-free alloys is approximately 40 MPa/wt..% Mg.
On the basis of microstructural observations, the variation in proof stress with magnesium concentration in Al-Li-Cu-Zr alloys can be divided into three regions 0-0.4%Mg, 0.4-l.O%Mg, and 1.0-2.O%Mg.
In the first region, only S 7 and Tl precipitation occurs and the increase in proof stress with magnesium concentration is small, approximately 20 MPa/wt.%Mg.
This proof stress increment with increasing magnesium content is probably due to solid solution strengthening, the magnitude of which is similar to that observed when the alloys were tested in the ST condition (i.e. 18 MPa/wt.%Mg).
In the second region, 0.4-l.O%Mg, a rapid change in proof stress takes place, the increase being ~ 7 5 MPa/wt.%Mg. This rapid increase corresponds with the
C3-422 J O U R N A L DE PHYSIQUE
appearance of S precipitation and the virtual disappearance of TI. The volume fraction of S continues to increase up to the 1% magnesium level. The rapid rise in the proof stress over this region will be due to several effects. Firstly, dispersion hardening from the S phase. Secondly, because the TI phase is being eliminated there will be more lithium available to form S ' . Thirdly, not all the magnesium will be used in S precipitation and therefore some solid solution strengthening from magnesium wi.11 occur. Fourthly, magnesium in solid solution reduces the amount of lithium that can be held in solution (14) and this will therefore increase the volume fract.i.on of S' that occ:urs during ageing.
In the region 1.0-2.O%Mg the change in proof stress with magnesium concentration reduces to a value around 12 MPa/wt.%Mg. Electron microscopy shows the volume fraction of S phase to be approximately constant over this region, so that the major contribution to the strengthening increment will be from magnesium entering solid solution. At the 2% magnesium level the alloy composition is situated in the a: + S + T phase field (where T is the A16CuMg4 phase); t,he S phase should therefore be starting to be replaced by T. However, in the present work no evidence was found for the presence of T. The presence of lithium may have changed the position of the phase boundaries or it may be suppressing nucleation of T as it does with many other phases in var-ious aluminium alloys.
T8 condition
Artificially ageing the stretched alloy t.o produce the T8 condition produces a microstructure that is different from that developed by a T6 heat treatment. In the magnesium-free alloy there is an increase in the volume fraction of TI; the precipitates are small.er and are nucleated on matrix dislocations introduced by the stretching operation. The increment. of proof stress in this alloy relative t.o the T6 condition was 35 MPa, and this will result from the increased dislocation density and dispel-sion strengthening from TI.
In alloys that contain magnesium the- (TX - T 6 ) strengthening increment increases slightly with increasing magnesium concentration so that
>
1. O%Mg the increment is around 45 MFn. The change in microstructure with increasing magnesium content jn the T8 condition is similar to that after T6 heat treatment, i.e. Tl precipitation occurs up to 0.4%?1g, thereafter S precipitation is dominant up to l.O%Mg, and from 1.0-2.O%Mg the volume fract-ion of S i.s approximately constant. However, the distribution of S is different in the T6 and T8 conditions; in T6 the S is nucleated on sub-boundaries, in the T8 the S is more uniformly distribul.ed on matrix dislocations introduced by the stretch. The explanation of the shape of the proof stress curve with magnesium concentration will be the same as that for the T6 condition but with the added strengthening increment from the increased dislocation density and increased volume fraction of S.Tensile strength and ductility
The tensile strengths of the alloys increase with magnesium concentration, the trend being similar to that for proof stress for the ST, T6 and T 8 conditions. The most noteworthy point is the increasing difference between proof stress and tensile strength as the magnesium concentration changes from 0-2%, and this occurs without an accompanying increase in elongation to failure. The reason for this increasing difference between tensile strength and proof stress
+
a more rapid rate of work hardening and loss necking deformation as the magnesium concentrat.ior1 rises.Similar effects have been noted in previous work (15) on Al-Mg and A1-Li--Mg alloys.
Elongations to failure after ST and T6 heat treat.ments arc unchanged as the magnesium concentrat ion increases. However, in the T8 condition the elongation in the magr~esium-free alloy is low, but ductility improves as the magnesium increases to 0.6%. This suggests that a dispersion of S phase is more effective than TI at dispersing the planar slip that takes place in these alloys.
A1-2.5%Li-1.3%Cu-Zt. a l l o y i n b o t h t h e TG and T8 c o n d i t i o n s , w i t h o u t l o s s o f d u c 1 : i l i t y .
2 . D F I I S ~ ? . ~ f a 1 . 1 ~ 1 i n e a r l y w i t h magnesium concent.r.at ion s o t h a t t h e s p e c i f j c s t r e n g t h s o f an a l l o y c o n t a i n i n g Z%Mg i n T6 and T8 c o n d i t i o n s a r e -10% above t h o s e found f o r t h e 8090 c o m p o s i t i o n (0.7XMg).
3. T e n s i l e s t r e n g t h i n c r e a s e s w i t h magnesium c o n t e n t a t a f a s t e r r a t . e t h a n d o e s p r o o f s t r e s s due t o an i n c r e a s e d work h a r d e n i n g r a t e and more u n i f o r m p l a s t i c d e f o r m a t i o n a t h i g h e r magnesium c o n c e n t r a t i o n s .
4 . The m a j o r s t r e n g t h e n i n g i n c r e m e n t o c c u r s between 0 . 4 and 1 . 0 % luagnesiun~ and t h i s c o r r e s p o n t l s t.o r e p l a c e m e n t o f t h e TI p h a s e by S p r e c i p i t a t i o n .
5 . E l o n g a t i o n t o f a i l u r e i s independrint. o f inagricsium c o n c e n t r a t i o n except. f o r a l . l o y s i n t h e T8 c o n d i t i o n where low c o n c e n t r a t i o n s o f magnesium p r o d u c e low v a l u e s of d u c t i l i t y . T h i s s u g g e s t s that d i s p e r s i o n s of TI p h a s e are l e s s e f f e c t i v e t h a n S p h a s e a t d i s p e r s i n g p l a n a r s l i p .
Tho a u t h o r s e x p r e s s t h e i r a p p r e c i a t i o n t o t h e Procurement. E x e c u t i v e o f R.A.E.
Fartiborough f o r f i n a n c i a l s u p p o r t , and t o C.J. P e e l and B . Evans f o r many h e l p f u l d i s t : u s s i o n s on t h e work.
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