Effect of Zn additions on precipitation during aging of alloy 8090

Scripta METALLURGICA et MATERIALIA

Vol. 25, pp. 243-248, 1991 Printed in the U.S.A.

Pergamon Press plc All rights reserved

EFFECT OF Zn ADDITIONS ON PRECIPITATION DURING AGING OF ALLOY 8090

R.J. Kilmer, G.E. Stoner Center for Electrochemical Sciences and Engineering Department of Materials Science University of Virginia Charlottesville, VA 22903-2442 (Received August 3, 1990) (Revised November 7, 1990)

Introduction AI-Li base alloys have received considerable attention as potential lightweight replacements for conventional AI-base alloys in aerospace applications. The low density and high stiffness of alloy 8090 (AIU-Cu-Mg) earmarks it as an especially attractive candidate. Alloy 8090 derives its strength from the coprecipitat|on of AI3Li (~'), AI2CuMg (S') and AI2CuLi 0°1) during artificial aging. Stretching prior to aging enhances the heterogeneous precipitation of S' and T 1 greatly improving the fracture toughness properties in AI-U-Cu-(Mg) alloys (1). In applications where a pre-aging stretch is not feasible, additions of Zn to the AI-Li-Cu-Mg system have been investigated as a means to improve toughness. (2) The present paper reports the precipitation events of three stretched alloys whose compositions fall within the 8090 composition window and contain varying amounts of Zn additions up to 1.07 wt %. The alloys were examined under two different aging conditions using transmission electron microscopy (TEM) andthe results were compared with those of alloy 8090 employed as a baseline. Differential scanning calorimetry (DSC) was employed to provide further insight into the effects that varying the Zn content has on the precipitation events. ExDerimental Procedure The compositions (in wt %) of the wrought alloys used in the present investigation are shown in Table 1. The alloys were solution heat treated for 1 hour at 543C followed by a cold water quench. The alloys were specialty cast by Alcoa, scalped, forged and rolled to 2.5mm sheet. The four alloys were stretched to obtain a T3 condition. The aging hardening characteristics were traced by Vickers pyramid hardness measurements utilizing a 500 gm load with the stress axis normal to the longitudinal direction of the sheet. Microstructural studies were performed by TEM on the samples in both the as-received (T3) condition and after aging for 100 hours at 160C. This aging time and temperature represents the peak aged cond t on based on hardness tests. Specimens for TEM were prepared using standard methods and examined with a Phillips 400T electron microscope operating at 120 kV. Thermal analysis of the alloys was undertaken in a Perkin-Elmer DSC using standard procedures immediately after quenching from solution heat treatment utilizing a constant heating rate of 10C/min. TABLE I.

Chemica] compositions of alloys (wt-Z)

Alloy code A B

{}090 Baseline 6090 + 0.21 Zn

hi

Cu

Mg

Zn

2.44

J.06 0.63

2.27

1 , 0 7 0 . 6 O 0.21

'' -

Zr 0.I0 0.10

C

8090 + 0.58 Zn

1.91

1.07

0.6~

0.58

0.10

D

8090 4- 1.07 Zn

2109

l.O0

0.59

1.07

0.10

243 0036-9748/91 $3.00 + .00 Copyright (c) 1991 Pergamon Press plc

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Results and Discussion Age hardening curves, as shown in Figure 1, illustrate the aging response of the alloys at 160C. The additions of Zn, in general appear to delay the onset of peak hardnesses in the alloys. The baseline 8090 which d d not contain Zn, reached a peak hardness of -146 VHN after 48 hours thereafter dropping to -128 VHN after 101.5 hours. The Zn containing alloys attained peak hardness after approximately 100 hours at 160C. The T3 tempers resulted in hardness values which ranged between 82-89 VHN for all the alloys and the peak hardness values ranged from 145-155 VHN. From the Vickers aging profile peak aging was determined to be 100 hours at 160C for the alloys with Zn additions. Additions of Zn to alloy 8090 had a significant influence on the precipitation process. The baseline 8090 (alloy A) developed ~'-free zones at both high and low angle grain boundaries after 100 hours at 160C. Precipitate free zones have previously been attributed to the existence of vacancy gradients near the boundaries (which act as vacancy sinks) (3)and 6'-free zones have additionally been associated with the growth of I_i-containing precipitates on the boundaries which consume the U solute in the metastable ~' [4]. Both T 1 and S' phases may reside on low angle (subgrain)boundaries and on dislocations introduced during the pre-age stretch. It was determined from Kikuchi patterns that the sheet alloys contained a relatively large number of high angle grain boundaries (> -10-12°). Figure 2 illustrates a typical set of misorientations between grains in alloy C. Alloy A developed 6'-free zones along the majority of subgrain boundaries after aging for 100 hours at 160C. The zones were nearly continuous and flaired out around precipitates on the boundary as shown in Figure 3a. In the instances where no 6'-free zone was apparent along the boundaries, the boundaries themselves did not contain a significant degree of precipitation. A number of these boundaries were analyzed via Kikuchi patterns to determine their degree of misorientation and it was found that they were very low angle boundaries, being approximately 1° or less in magnitude. In alloys B and C the 6'-free zones were notably absent along the subgrain boundaries when aged 100 hours at 160C. A typical boundary is shown in Figure 3b in alloy C. In the alloy of highest Zn content alloy D, 6 -free zones were again observed along subgrain boundaries. Furthermore, these boundaries contained a relatively high density of coarse precipitates which were found via EDS to be rich in Cu and Zn. A typical subgrain boundary precipitate is shown in Figure 4. It was also observed that subgrain boundary precipitation of S' and T 1 was for the most part absent. Apparently Zn additions can have a marked influence on the precipitation processes in these alloys, most notably upon subgrain boundary precipitation. Baumann, et. al. (4) have previously demonstrated that Zn additions to the AI-Li system result in an increase in volume fraction of 6' for a given Li content. The Zn additions lower the solubility of Li in AI and were found to partition to the 6' resulting in an increased ,~/6 misfit strain. Given this information it became evident that to obtain further insight into the role Zn plays in the precipitation process, DSC experiments should be performed on all four alloys. DSC results (Figure 5) indicated that the exothermic peaks associated with the precipitation of 6' at position E shifted tohigher temperatures as the concentration of Zn increases in the four alloys. The exothermic peak for 6' is typically seen at - 170C (5). As the Zn content is increased to 1.07 wt. % the exothermic peak is shifted to approximately 215C. This change is attributed to, and is consistent with Zn incorporation into the 6' increas=ng the =/6' misfit. Evaluation of the endothermic peaks requires some caution. The positions of the peaks at F did not appear to be significantly altered in the alloys. These endotherms are associated with the dissolution of 6' and S' (5,6). Normalization of these energies to a per gram basis reveals a five fold increase in the endothermic energies as Zn content is increased to .58 wt. % as illustrated in Figure 6. The normalized endothermic energy subsequently drops in alloy D. The magnitude and trends of these energy changes are consistent with and likely a combination of the following: changes in the volume fraction of the ~ andS' and/or an alteration in the energies associated with dissolution of these precipitates. It seems unlikely that a five fold increase may be solely attributed to a volume fraction change in 6' and S', although it is known that Zn additions increase the volume fraction of 6' and that the S' precipitation also changes with Zn additions. Figure 7 indicates the S' distribution in alloys A and C after 100 hours at 160C. Alloy A formed a

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widespread distribution of relatively fine S' laths, nucleating on dislocations at the grain interiors. Alloy C also formed a widespread distribution of S' but the precipitates were more coarse in morphology. This effect was also seen in the other Zn containing alloys B and D. The energy changes therefore are likely the reflections of both volume fraction changes in 6' along with changes in the entha p c makeup of the 6. The drop in energy in alloy D may be influenced by the formation of the coarse Zn containing precipitates which form on the boundaries and in the grain interiors, resulting in a decrease in Zn incorporation into the 6'. If trace additions of Zn can alter subgrain boundary precipitation then the possibility exists to design alloys with improved properties such as SCC resistance and toughness. The possibility of improving the relatively poor SCC behavior is suggested in these alloys, and is currently under investigation.

Relatively small additions of Zn can significantly alter the precipitation events in the four alloys studied. DSC results indicated that Zn was incorporatedinto the 6' which may stabilize it as evidenced by the lack of a 6'-free zone in alloys B and C after 100 hours at 160C. Both the 6' and S' precipitates were influenced by Zn additions and that coarse Zn containing precipitates can form on the boundaries and within the interiors of the grains when the Zn content is 1.07 wt. %. Acknowledgements This work was supported by a grant from NASA Langley Research Center under grant NAG-745-1-4, B.A. Lisagor and D.L Dicus project monitors. It was co-sponsored by Alcoa, S.C. Byrne, E.L. Colvin and J.J. Witters technical advisors. The authors would also like to extend their thanks to G.J. Shiflet and A.K. Mukhopadhyay for their assistance. References 1.

2. 3. 4. 5. 6.

R.F. Ashton, D.S. Thompson, E.A. Starke Jr. and F.S. Un, AI-Li Alloys III, (ed. C.Baker et al), Institute of Metals, p.66. (1985) P.J. Gregson, K. Dinsdale, S.J. Harris, B. Noble, Mater. Sci. and Tech.,3,7,1987 P.N.T. Unwin, G.W. Lorimer and R.B. Nicholson, Acta Met, 17,1363,1969 S.F. Baumann and D.B. Williams, AI-Li Alloys II, (ed. T.H. Sanders and E.A. Starke), 17;1984, Warrendale,Pa, The metallurgical Society of AIME A.K. Mukhopadhyay, C.N.J. Tite, H.M. Flower, P.J. Gregson and F. Sale, Journal De Physique, Colloque C3, 48,1987 S. Abis, E. Evangelista, P. Mengucci and G. Riotino, AI-Li V, (ed. T.H. Sanders and E.A. Starke), 681;1989,

150 -

o o.oow/o,-,,[,,,,,yA a

O#.X ~./o h

14o -

o

o.~ ,,/o ,.., [ " " T c]

-

,~ F-

[J,lloy seJ

~,,.,.'~m"V--'~

..~'-~

u

I

\

120

110 eO

,,,7¥~ , 10 -s

, ,,,,,,I

,

, ,,,,,,I

10 °

10 [

.,~i~ Time (hours) Figure 1. Age hardening response of alloys A-D at 160C.

....

,,,, 10a'

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FIG. 2. TEM micrograph illustrating typical misorientations between grains in alloy C.

0

FIG. 3. TEM centered dark field micrograph of a) alloy A and b) alloy C after aging after 100 hours at 160C. [100] matrix beam direction.

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FIG. 4. TEM centered darkfield of 6' and coarse subgrain boundary precipitate and S' in neighboring grain. [100] matrix beam direction.

t

10.0 9.0-

Alloy A

F

Alloy B LO-

"6

~

7.0Alloy C a 0 . S.04.0LO-

F |.8-

tO~O Temperature

FIG. 5. DSC thermograms of alloys A-D.

(°C)

j

Alloy D

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DSC thermograms of alloys A-D. 3OOO

2OOO

1000

I Alloyk

I J~rB

I AlloyC

I AlloyD

Sample I.D.

Figure 6. Normalized energies associated with endotherms at position F.

Figure 7. TEM micrograph of S' matrix distribution in a)alloy A and b)alloy C after aging 100 h at 160C near the [112] AI orientation.