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THE PHYSICAL METALLURGY OF ALUMINIUM-LITHIUM-COPPER-MAGNESIUM-ZIRCONIUM

ALLOYS-8090 AND 8091W. Miller, J. White, D. Lloyd

To cite this version:W. Miller, J. White, D. Lloyd. THE PHYSICAL METALLURGY OF ALUMINIUM-LITHIUM-COPPER-MAGNESIUM-ZIRCONIUM ALLOYS-8090 AND 8091. Journal de Physique Colloques,1987, 48 (C3), pp.C3-139-C3-149. �10.1051/jphyscol:1987317�. �jpa-00226547�

JOURNAL DE PHYSIQUE Colloque C3, suppl6ment au n09, Tome 48, septembre 1987

THE PHYSICAL METALLURGY OF ALUMINIUM-LITHIUM-COPPER-MAGNESIUM- ZIRCONIUM ALLOYS-8090 AND 8091

W.S. MILLER, J. WHITE and D.J. LLOYD

Alcan International Ltd., Southam road, Banbury, GB-Oxon OX16 7SP, Great-Britain

ABSTRACT

Aluminium alloys containing up to 3wt% lithium and cast using an ingot metallurgy route are currently being developed for aerospace applications. The objective of this paper is to review the current status of the metallurgical understanding of these alloys. Particular emphasis is placed on the alloy system Al-Li-Cu-Mg-Zr ,which includes the alloys 8090 and 8090 developed in the United Kingdom by the Royal Aircraft Establishment and Alcan. Comparison is made with other alloy systems, where appropriate, e.g. Al-Li,Al-Li-Mg,Al-Cu-Li.

INTRODUCTION

Aluminium alloys containing up to 3wt% lithium and cast using an ingot metallurgy route are currently being developed for aerospace applications. The objective of this paper is to review the current status of the metallurgical understanding of these alloys. Particular emphasis is placed on the alloy system Al-Li-Cu-Mg-Zr and comparison is made with other alloy systems where appropriate.

The Al-Li-Cu-Mg-Zr system was developed in the United Kingdom by the Royal Aircraft Establishment and Alcan (1). Two alloys are currently produced under the generic name "LITAL" ; 8090 (aimed for damage tolerant and medium strength applications) and 8091 (aimed for high strength applications). The compositional ranges (wt%) for the main constituents of the alloys are:-

MICROSTRUCTURAL ASPECTS

The main microstructural features of these alloys can be divided into three groups; grain structure, matrix precipitation and grain boundary precipitation.

Grain Structure

Chromium,manganese and zirconium have all been utilized in attempts to control grain structure in AL-Li based alloys. Of these zirconium has received the greatest attention and is the most powerful recrystallization inhibitor in these alloys. It is consequently used for grain structure control in all commercial A1-Li base alloys. It is,well established that the cubic (LIZ) A4Zr particles retard recrystallization . This most likely arises because gram boundary migration is retarded during recrystallization by the requirement for AllZr particles to transfohn from coherent to incoherent.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987317

JOURNAL DE PHYSIQUE

Matrix Precipitation

Within the 8090 and 8091 composition ranges the levels of copper and magnesium do not significantly affect the nature of the precipitation of 6' and 6 phases (2). This behaviour contrasts with Al-Cu-Li alloys, such as 2090, where the 61 distribution is altered by the formation of the lithium containing T, phase (AkCuLi) which precipitates as plates on (111) planes. The 6- phase also nucleates on 8' (AhCu)(3). Both these factors reduce the volume fraction of 6' compared with a copper free alloy. A further reduction in the 6' volume fraction occurs when a stretching treatment is used prior to ageing and widespread nucleation of T, takes place on dislocations (4).

In 8090 and 8091 the presence of copper and magnesium results in the additional precipitation of S1(AkCuMg) phase during artificial ageing. S' precipitates as laths along <001> with an orthorhombic structure (5). Figure 1 shows the microstructure of a peak aged 8090 alloy where St co-precipitates with T, and b' . However,the balance of St relative to T, is critically dependent on the relative concentrations of all three main alloying additions. High copper and lithium contents relative to magnesium result in a dominance of TI whereas high copper and magnesium relative to lithium produce a dominance of S'. Magnesium content > 0.5wt % are necessary to ensure a dominance of S' in alloys containing up to 2.7% Lithium and 1.4% copper. As with TI, S 1 nucleates heterogeneously and hence deformation prior to ageing is required to ensure widespread nucleation within the matrix. In contrast to TI, however, heterogeneities in the S' distribution do not manifest themselves in the 6' distribution as the phases precipitate independently.

Recent work has shown that the natural ageing behaviour of 8091 alloy can influence its artificial ageing behaviour (6). This has been attributed to the strong interaction between lithium and vacancies and the liberation of vacancies during 6' growth. Figure 2 shows the natural ageing behaviour of 8090 and 8091 in the unstretched (T4) and stretched (T3) conditions. Comparison is made with an A1-2.5%Li-0.12%Zr and an Al-1.2%Cu-0.7ZMg-O.l2%Zr alloy. The Al-Li-Zr alloy shows no significant natural ageing in either condition whereas the 8090 and 8091 alloys show a strong natural ageing response. This together with the increased natural ageing response of 8091 compared with 8090 suggests that the response is not due to the growth of 6' but rather due to a precursor to the S' phase. Further evidence for this hypothesis is that the ageing response of the lithium free alloy is similar to the 8090 alloy of similar copper and magnesium content. The magnitude of the natural ageing response is ,however,reduced in 8090, compared with the lithium free alloy. This is because of the lower excess vacancy concentration in 8090, probably due to the higher lithium-vacancy binding energy. Stretching prior to natural ageing reduces the natural ageing characteristics of 8090 and 8091 but increases that of the lithium free Al-Cu-Mg-Zr alloy. Further work is required in this area, including small angle neutron experiments,to understand the mechanisms and explain the differences between the natural ageing response of T3 and T4 tempers.

Grain Boundary Precipitation

In alloys where a second matrix precipitate is present heterogeneous nucleation of Q', T or Sr occurs on low angle grain boundaries. The presence of the lithium containing TI phase leads to the formation of a 6' precipitate free zone (PFZ)(see figure 3). The formatiqn of a PFZ in this manner is dependent on the ratio of copper to lithium(3). In alloys where S' is the dominant second phase (8090 and 8091) no PFZ is formed at a low angle grain boundary (Figure 4) except in overaged tempers where the precipitation of equilibrium phases causes localized lithium denudation.

The equilibrium 6 phase has been reported as the principal high angle grain boundary phase in all A1-Li base alloys(7). The presence of this phase leads to the formation of a 6' PFZ at these grain boundaries. 6 phasc also precipitates on sub-grain boundaries and within the matrix in overaged tempers (Figure 5). In this case incoherent AIJZr particles have been shown to be the main nucleating site for 6(8). No systematic study of copper containing phases on high angle grain boundaries has been carried out. However it is clear that 6 phase is not the only phase formed and that other phase may dominate at

the grain boundaries in 8090 and 8091 alloys in specific ageing treatments. One phase identified is the icosohedral I-phase (AkCu(Li,Mg (9).This phase can form during the quench and grows as 'snake-like' precipi ? ates along grain boundaries (Figure 6). The structure of the I-phase is the subject of extensive research as evidence by the numerous papers at this conference. Other complex phases can form at grain boundaries and within the matrix of overaged alloys e.g. the C-phase (10). The presence of these phases influences the PFZ width and are considered detrimental to fracture toughness.

The effect of the various grain boundary precipitates on the engineering properties is not well established. Work by the authors has shown that large 'snake-like' precipitates of I-phase, formed in slowly quenched material, are detrimental to fracture toughness, in underaged tempers (11). In the peak and overaged tempers the area fraction of grain boundary precipitates is large and fracture toughness is low irrespective of the type of precipitate formed (Figure 7). In a private communication from Professor E.A.Starke (Univ. of Virginia) it has been suggested that the formation of I-phase during ageing is suppressed at low ageing temperatures (e.g. 150 C). Alloy chemistry appears to play a significant role in the formation of grain boundary precipitates. Increased I-phase formation ,especially during the quench, occurring as the copper content increases. Consequently alloys with high copper (e.g. 2090) are much more quench sensitive than low copper alloys such as 8090. Within the AL-Li-Cu-Mg alloy types 8091 is more susceptible to I-phase formation than 8090.

In view of the importance of the grain boundary structure in controlling engineering properties further detailed studies of grain boundary microstructure are clearly essential to the understanding of the microstructure /property relationship in these alloys.

THE INFLUENCE OF MICROSTRUCTURE ON ENGINEERING PROPERTIES

Modulus and Density

Since the main driving force for the development of A1-Li based alloys is their reduced density and increased modulus compared to conventional alloys it is pertinent to show how these properties are affected by composition.

Peel et a1 (12) developed the following empirical formula to calculate density:-

where the atomic symbols represent the concentration of that element in weight percent. Using the above formula the density ranges for the Alcan alloys are 2.52 to 2.56 g/cc (8090) and 2.55 to 2.58g/cc (8091). This is confirmed by hydrostatic weighing. These values are approximately 10% and 8% lighter respectively than conventional aluminium alloys (typical density 2.80 g/cc).

The influence of alloying additions, in particular lithium, on the modulus of elasticity (E) is well established (13). Recent work has shown that during ageing E maximizes just before peak strength is achieved and thereafter decreases (14). This is attributed to the higher modulus of the metastable ' precipitates compared with the equilibrium phases (e.g. AlLi,I-phase) that form in overaged tempers. The 8090 and 8091 alloys have a modulus 10% greater than conventional alloys.

Strength

Several different mechanisms contribute to the strength of A1-Li base alloys(l5):

7 s solid solution hardening - due to Li,Cu,Mg in solution in the matrix 7 B substructure hardening - due to grain boundaries,substructure and

dislocations p precipitation hardening - due to bq,S' and T, precipitated on ageing

The flow stress of an aged alloy is:-

JOURNAL DE PHYSIQUE

The particle strengthening term is itself composed of components due to the individual precipitates 6', s' and TI . Other precipitates are coarse and incoherent (e.g.6,I-phase) and do not contribute significantly to strength. As considered in recent reviews (15,16), the individual precipitate components not should be linearly added, but treated as the sum of squares.

For the case of an A1-2.5%Li alloy the solution treated strength r 9 is 30MPa. Ageing to peak strength raises the strength to 75MPa , so that

761 45MPa. Addition of 0.12IZr further increases the peak strength to 100MPa. Slnce the volume fraction of AlyZr is too small to contribute much to strength, the increase in strength is due to grain and sub-grain boundary strengthening, thus T B =25MPa. In the case of 8090 6' and S f are present after ageing. Ageing a lithium-free version of the alloy produces an increase of 27MPa, which is due to S t . Using the strengthening for b' from the binary alloy together with this S' strengthening value in equation 2 we have:

7 p -. (45)a+ (27)* = 52.5MPa

The experimentally observedTp of 55MPa , is in good agreement with this analysis. It must be appreciated ,however, that these are minium values for the individual strengthening components because we have ignored the decrease in solution hardening as ageing proceeds. Nevertheless, the analysis serves to emphasize that several different components are contributing to the strengthening of these alloys. Within the 8090 composition limits 6: is the main strengthening phase with St and sub-boundary strengthening approximately equivalent. As the copper level is increased the strengthening due to S f increases and in 8091 is larger than that attributable to 6' In stretched 2090 experiments indicate that T, is the main strengthening phase ,.-.. 111).

The importance of sub-structure is clearly shown in this analysis. It plays an important role not only in providing intrinsic strengthening but also provides a high density of nuclei for the precipitating phases S f and T . Recrystallized structures therefore will not achieve the strength levels of unrecrystallised material except in the case of very fine grained material, which cannot be obtained in most product forms.

Deformation and Fracture

In any alloy strengthened by coherent particles there is a tendency for strain to localize into bands,due to the fact that a dislocation pile up can shear through the particle more easily than can individual dislocations. In alloys with ordered precipitates (e.g. 6') this tendency for strain localization is enhanced because dislocations move as pairs ; the first dislocation creating disorder, the second dislocation restoring order. The result is intense shear bands (Figure 8). The local strains in these shear bands have been estimated as 3. These intense shear bands lead to high stress concentrations at grain boundaries and the low ductility of A1-Li base alloys has been attributed to this effect(l8). Attempts have been made to disperse slip. Introducing dispersoids or increasing the particle matrix misfit are possible ways of dispersing slip but they have not proved effective in A1-Li base alloys(19,20). Instead in alloys such as 8090 and 8091 the presence df S r promotes cross slip , as shown by the irregular slip lines produced in peak aged 8090 compared to the binary (Figure 9).

The fracrure mode in 8090 and 8091 is intergranular with extensive grain boundary splitting along grain boundaries parallel to the tensile axis and linkage along sub-boundaries and shear bands (Figure 10). With this high susceptibility to grain boundary failure the short transverse ductl~lty and fracture toughness can be particularly low. These factors have been attributed variously to strain localisation in slip bands(l8),strain localisation in PFZ1s (21) ,grain boundary embrittlement due to the precipitation of equilibrium phases (22)and impurity segregation effects(23).

Strain localisation and grain boundary precipitation certainly reduce ductility and toughness by promoting grain boundary failure but they are not essential for the splitting type of fracture. Solution treated or lightly aged material when tested at -193 C exhibits the same form of fracture as peak aged material deformed at 25 C (Figure 11)(24). In the solution treated condition there is no PFZ,no grain boundary precipitation and strain localization is no more extensive than in conventional high strength aluminium alloys.

Table 1, details the effect of test temperature on the tensile properties of a 50mm 8090 plate in the unaged and peak aged condition, shows that at strength increases in the longitudinal and short-transverse directions as the test temperature decreases. Ductility increases with decreasing test temperature for the longitudinal direction, but falls in th short- transverse direction. The fracture stress in solution treated material tested at -193 C (542Mpa for the longitudinal direction approaches that observed in peak aged material tested at room temperature (552 MPa). This is due to an increase in the thermal component of the flow stress and an increase in the work hardening rate. The failure mode observed (Figure 11) indicates that the grain boundary strength is lower than conventional alloys.While sodium and/or potassium segregation to grain boundaries (23) and the occurrence of a form of liquid metal embrittlement (LME)(25)have been suggested as a reason for the failure mode extensive high resolution scanning Auger microscopy of in-situ fracture surfaces of commercially produced 8090 have failed to detect any impurity segregation (24).(It should be noted that if the alloys are cast under a sodium and/or potassium containing flux isolated flux particles can be incorporated into the alloy during casting and may occur as segregated regions on the fracture surface).

The observation that the tensile ductility increases at low temperature is common in FCC materials provided that strain localization mechanisms are involved in the fracture process. This is because the work hardening rate and the strain rate sensitivity increase with decreasing temperature and these inhibit strain localisation. In those cases where the brittle fracture component is totally dominating (e.g the ST direction),the failure ductility will decrease with decreasing temperature since an improvement in ductility from an increase in plastic flow resistance would not be expected .The data in Table 1 shows that there is a small increase in longitudinal ductility with decreasing temperature but little change in the short transverse direction. Thus there is little evidence to the suggestion (25) that the increase in ductility at low temperatures supports a LME mechanism.

The present understanding of the fracture process suggests that the tendency for strain localisation and brittle equilibrium phases at grain boundaries contribute to the low ductility observed in these alloys In addition,the observations of the fracture mode in unaged material at -193 C indicates that the grain boundary strength of these alloys is lower than other high strength A1 alloys. This may be due to the segregation of lithium itself, either as solute strains or as clusters or precursors to the grain boundary equilibrium precipitates.

As-cast grain size control and metal cleanliness is critical in these-alloys ,because of the lack of recrystallization in the fabrication cycle with the resulting highly fibrous grain structure and the ease of cavitation at inclusions(26). Attention by commercial manufactures to these aspects is reflected in the steady improvement in ductility and fracture toughness and this is expected to continue.

Within the confines of this paper it is not the intention to discuss ways in which, by control of strain localisation and grain boundary precipitation, ductility and fracture toughness can be improved. 'These aspects are dealt with by the authors elsewhere in this conference (11,27).

Stress corrosion and corrosion

In complex A1-Li base alloys the presence of copper-rich precipitates largely controls their corrosion and stress corrosion (SCC) behaviour. In general the higher the copper content the more susceptible are the alloys. This is attributed to the copper-rich phases promoting anodic dissolution at locally preferential sites i.e. grain boundaries.

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Overageing 8090 and 8091 results in equilibrium phases precipitating on sub-structure and within the matrix thereby decreasing the potential difference between the grain boundary and the matrix. This reduces the driving force for preferential d'issolution of the copper-rich phases. These aspects are covered in more detail in the paper by A. Gray at this conference (28).

SUMMARY

This paper has reviewed the essential features of the microstructure of 8090 and 8091 and has demonstrated how these microstructural features influence the engineering properties (strength ,toughness and SCC). However it is clear that there are certain areas in which the current knowledge of these alloys needs extending:-

1. An understanding of the natural ageing response. 2. The structure at the grain boundary , in particular grain

boundary precipitate type, and more importantly the subsequent influence on material properties.

3. A clearer understanding of the actual fracture mode,in particular at low temperatures in the unaged condition.

Even without this complete understanding it is apparent that the metallurgy of Al-Li-Cu-Mg-Zr alloys has become sufficiently advanced to enable the development of an extremely attractive balance of properties (29).

The authors are grateful for the permission of Alcan International to publish this paper.

REFERENCES

1. C.J.Pee1,B.Evan.s and D.McDarmaid. Proc. 3rd. Int. Aluminium Lithium Conference . Oxford,U.K. Ed. C.Baker, P.J.Gregson,S.J.Harris and C.J.Pee1 Pub. Inst. of Metals London (1986). p26.

2. G.E.Thompson and B.Noble J.I.Met. 101 (1973) p1793. 3. M.H.Tosten,A.K.Vasudevan and P.R.Howel1 in ref 1. p483. 4. R.J.Rioja,P.E.Bretz,R.R.Sawtell,W.H.Hunt and E.A.Ludwiczak In Aluminium

Alloys - Their Physical and Mechanical Properties . Proc Int. Conf. at Charlottsville VA. Ed. E.A.Starke and T.H.Sanders Pub. EMAS. Vol 3 (1986) p1781.

5. P.J.Gregson and H.M.Flower Acta. Met. 33 (1985) p527. 6. H,M.Flower,P.J.Gregson,C.N.J.Tite and A.Makhopadhyay in Ref 4 Vol 2 p743. 7. J.C.Huang and A.J.Ardel1 in ref 1 p455. 8. W.S.Miller,J.White and D.J.Lloyd in ref 4 Vol 3 p1799. 9. W.A.Cassada,G.J.Shiflet and E.A.Starke in ref 4 Vol 2 p695. 10. A.Loiseau and G.Lapasset J.de Physique, 47 (1986) pC3-331. 11. J.White and W.S.Miller this conference. 12. C.J.Peel,B.Evans,C.Baker,D.A.Bennett,P.J.Gregson and H.M.Flower. Proc.

2nd. Int. Aluminium Conference Monterey CA. Ed. T.H.Sanders and E.A.Starke TMS-AIME (1984) p363.

13, In Aluminium Properties and Physical Metallurgy Ed J.E.Hatch Pub ASM (1984) 14. F.Broussaud and M.Thomas in ref 1 p442. 15. D.J.Lloyd ICSMA 7. Montreal 1986 Ed. H.J.McQueen,J.P.Bailon,J.I.Dickson,

J.J.Jonas and M.G.Akben Pub. Pergamon Press Vol 3 (1986) p1745. 16. A.J.Ardel1 Met.Trans 17A (1986) pl. 17. J.C.Huang and A.J.Ardel1 Proc Int Con£ "Aluminium Technology 1986"

Ed. T.H.Sheppard Pub. Inst of Metals (1986) p434. 18. T.H.Sanders and E.A.Starke Acta. Met. 30 (1982) p927. 19. S.F.Bauman and D.B.Wi1liams in ref 12 p17. 20. K.K.Sankaran,S.M.Shastry and J.E.OtNeal Met.Trans 14a (1983) p2174. 21. F.S.Lin,S.B.Chakrabortty and E.A.Starke Met. Trans. 13A (1982) p401. 22. S.Suresh and A.K.Vasudevan in ref 1 p595. 23. A.K.Vasudevan,A.C.Miller and M.M.Kersker in ref 12 p181. 24. W.S.Miller,M.P.Thomas,D..J.Lloyd and D.Creber Mat.Sci.Tech, 2 (1986)

~1210.

25. D,Webster in ref 1 p602. 26. N.J.Owen,D.J.Field and E.P.Butler in ref 1 p576. 27. W.S.Miller,J.White,M.A.Reynolds,D.S.McDarmaid and G.M.Starr This Conference. 28. A.Gray This Conference. 29. R.Grimes This conference.

TABLE 1

Influence of Test Temperature on the Mechanical Properties of 50 mm Plate from an A1-2.5%Li-1.2%Cu-0.7XMg-0.12%Zr Alloy

Figure 1. Dark and bright field transmission electron micrographs showing the main matrix precipitate types found in Al-Li-Cu-Mg-Zr alloys.

Test Temperature OK

295 233 183 153 7 7

T351 Temper T651 Temper

L-Direction

0.2%PS (MPa)

197 205 195 217 232

S-T Direction L-Direction

El (%)

15.6 15.9 14.4 17.6 22.5

0.2%PS (MPa)

162 161 174 179 182

O.Z%PS (MPa)

417 440 417 433 445

S-T Direction

El (%)

16.5 16.2 17.3 17.0 10.0

El (%)

5.7 6.3 6.3 6.8 8.6

0.2%PS (MPa)

335 337 330 342 356

El (%)

2.3 2.25 2.25 1.40 1.50

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0 100 250 500 1000 1500 NATURAL AGEING TIME (h)

Figure 2. Natural ageing response of 8090,8091,Al-2.5%Li-0.12%Zr and A1-1.2%Cu -0.7XMg-0.12XZr alloys in the unstretched ( T 4 ) and stretched (T3) conditions.

Figure 3. Sub-grain boundary PFZ Figure 4. Sub-grain boundary observed in aged Al-Cu-Li precipitation of Sf-phase alloys. in aged Al-Li-Cu-Mg

alloys.

Figure 5. Equilibrium precipitation of AlLi on sub-grain boundarigs and within the matrix of an overaged Al-Li-Cu-Mg-Zr alloy.

1r:r;ure 6. I-phase formed duriuy slov qurnching of an Al-Li-Cu-Mg-Zr alloy.

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Figure 7. Grain boundary precipitation present in two samples of 8090 alloy after ageing for 16h at 190 C with similar toughness ; (a) cold water quenched prior to. ageing .- the AlLi phase predominates and (b) air quenched prior to ageing - I-phase predominates.

Figure 8. Intense shear bands in a peak aged Al-2.5%Li alloy.

Figure 9. Irregular slip lines found in peak aged 8090 (a) in comparison with those observed in an A1-2.5%Li-0.12%Zr alloy (b) .

Figure 10. Fracture mode observed in Figure 11. Fracture mode observed in peak aged 8090 plate when solution treated 8090 tested at room temperature. plate tested at -193 C.