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Effect of retrogression and reaging on the precipitates in an 8090 Al–Li alloy
Materials Science and Engineering A242 (1998) 39 – 49
Effect of retrogression and reaging on the precipitates in an 8090 Al–Li alloy V. Komisarov a, M. Talianker a,*, B Cina b a
Department of Materials Engineering, Ben Gurion Uni6ersity of the Nege6, Beer-She6a, Israel b Metallurgy Group, Israel Aircraft Industries Limited, Lod-70100, Israel Received 6 May 1997; received in revised form 8 July 1997
Abstract The microstructure of the precipitates in an 8090 Al – Li alloy has been studied qualitatively and quantitatively by techniques of transmission electron microscopy through the stages of retrogression and reaging (RR). The retrogression was performed for various times at temperatures of 240 and 260°C. The reaging treatment was carried out at 170°C. The behavior of the precipitates was specific for each temperature of retrogression: while for retrogression at 240°C there was only partial dissolution of d% particles, total dissolution of d% was observed after 30 min of treatment at 260°C. Reaging following retrogression at 240 and 260°C resulted respectively in the full and only partial restoration of the volume fraction of the d% phase. Change in the d% precipitate structure was the dominant factor which was responsible for the changes in hardness of the Al – Li 8090 alloy in the course of RR treatment for a retrogression temperature of (RR-240). On the other hand, changes in the hardness of the alloy subjected to RR-260 treatment could not be explained simply in terms of changes in the d% particle structure. The precipitation of an additional phase, T1, concomitant with the dissolution of d% phase at 260°C, also contributes to the hardness on this retrogression stage of the treatment. © 1998 Elsevier Science S.A. Keywords: Retrogression; Reaging; Precipitates
1. Introduction Aluminium–lithium alloys are of considerable potential interest to the aerospace industry because of their significantly higher values of specific strength and modulus compared with conventional aluminium alloys. Although such Al – Li alloys have been developed intensively in the past 10 years or so, their utilization has been much less than anticipated. This is because of problems of reduced ductility and fracture toughness, marked anisotropy of mechanical properties especially in non-recrystallised structures and the inability, as yet, to obtain resistance to stress corrosion akin to that obtainable from the high strength 7075-aluminium alloy in a T73 temper [1 – 4]. Some years ago a non-conventional two-stage heat treatment process known as retrogression and reaging (RR) was developed by one of the present authors [5,6]. * Corresponding author. Tel.: + 972 7 461475; fax: + 972 7 472946. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 0 2 - 9
This process overcame the problem of susceptibility to stress corrosion for aluminium alloys of the 7000 type. In contrast to material in the T73 temper which results in overaging of alloys, material in the RR temper combines the mechanical properties of the T6 temper (of maximum tensile strength) with the full resistance to stress corrosion of the T73 temper. RR has been shown to be applicable to a broad range of aluminium alloys of the 7000 type [7]. Since Al–Li alloys constitute an age-hardening system wherein the alloys are hardened by the precipitation of d%, a coherent ordered phase of composition Al3Li, theoretically this system, when retrogressed and subsequently subjected to reaging, should tend to re store its hardness. Some preliminary results on the effect of RR heat treatments on the susceptibility to stress corrosion of a commercially available 8090 Al–Li alloy were reported by Hu et al. [8]. Recently the present authors undertook a more systematic investigation of this alloy [9] with the aim of determining to what extent retrogression and reaging could be applied
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in order to achieve an optimum combination of strength and resistance to stress corrosion. It was found that retrogression and reaging heat treatment of an 8090 Al–Li alloy, initially in the T8771 temper, can substantially improve its resistance to stress corrosion. The present paper presents the results of the second part of this work which covers the transmission electron-microscopic investigation of the microstructural evolution and the change in hardness of an 8090 Al–Li alloy subjected to RR treatment.
2. Experimental procedures The material investigated was a plate of an 8090 Al–Li alloy, 50.8 mm thick, manufactured by Alcan, UK. The material was supplied in the T8771 temper, i.e. solution treated, stretched 6% and aged for 32 h at 170°C. The chemical composition of the alloy in weight percent is given in Table 1. All RR treatments were performed on specimens, 22× 22 by 50.8 mm long, which were taken from the S-T direction of the plate, having their length perpendicular to the rolling direction. The retrogression of the alloy was performed in a molten salt bath (50% sodium nitrate/50% sodium nitrite) at two temperatures, 240 and 260°C, for time intervals ranging from 5 to 40 min. The bath size was large enough (1× 2 ×1 m) to provide very rapid heating of specimens to the desired temperature. The reaging treatment was the same for all retrogressed specimens and was performed in an air furnace at 170°C for 32 h. One series of specimens was retrogressed and then quenched in water. Another series of specimens was retrogressed, quenched in water and subsequently reaged. All heat treatments were monitored by the measurement of the Rockwell ‘B’ hardness of the specimens. The microstructure of the specimens was studied in a transmission electron microscope operating at 200 Kv. Samples for TEM analysis were cut in the form of slices, 0.2 mm thick, using an electric spark erosion machine, and final thinning was achieved by means of a twin-jet electropolishing unit (Tenupol), using a solution of 30% nitric acid and 70% methanol at − 18°C. The diameters of the spherical d% precipitates were measured from the dark-field electron micrographs using the d% superlattice reflection. A statistical treatment of the size and density of the d% particles per unit area Table 1 Chemical composition of the 8090 Al–Li alloy (wt%) Li
Cu
Mg
Fe
Si
Zr
Na
Al
2.35
1.15
0.74
0.05
0.03
0.11
0.0002
Balance
of electron micrograph was performed using Optilab image processing and an analysis software program [10].
3. Results
3.1. Hardness measurements Fig. 1(a) shows the change of hardness versus time of retrogression for two temperatures, 240 and 260°C. The variation of the hardness of retrogressed and reaged material as a function of retrogression time is shown in Fig. 1(b). It can be seen (Fig. 1(a)) that the general shape of the retrogression curves obtained in the present investigation for the 8090 Al–Li alloy is very similar to those obtained for a 7000 type aluminium alloy [6]. The hardness curves in Fig. 1(a) decrease very rapidly during retrogression until a minimum occurs, after which some increase in hardness is observed and thereafter there is no further change in hardness for the times investigated. After retrogression at 240°C and subsequent reaging (Fig. 1(b)), the original hardness level obtained prior to RR treatment was restored and this did not change significantly with retrogression time. For retrogression performed at 260°C, followed by reaging, the original hardness was initially restored but this began to decrease after only 10 min of retrogression when followed by reaging.
3.2. The beha6ior of the d% precipitates Dark field micrographs showing d% particles were taken from regions having the same foil thickness (of the order of 100 nm) and for the same diffraction conditions (zone area [001]), using the same superlattice d% reflection (100). Typical examples of dark-field images of the d% precipitates for the alloy subjected to different retrogression treatments are shown in Fig. 2(a)–(e). The d% precipitate structure of the alloy in the as-received condition is shown in Fig. 2(a). It is seen that the d% particles are fairly uniformly distributed in the matrix, are small in size (0.0015–0.02 mm) and that the number of particles per unit area of the image is large. Fig. 2(b) and (c) show the d% precipitates after retrogression at 260°C for 5 and 30 min respectively. Fig. 2(d) and (e) show the d% precipitates in a specimen retrogressed at 240°C for 20 and 40 min, respectively. Comparison with the as-received condition indicates that a relatively short time (5 min) of retrogression at 260°C causes a significant decrease in the density of the d% precipitates. Almost all the finer d% precipitates and many of the larger ones have now dissolved. Longer periods of retrogression led to further decrease in the density of the d% particles, and after 30 min at 260°C d% precipitates disappeared almost entirely (Fig. 2(c)). On
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Fig. 1. Variation of hardness of an 8090 Al–Li alloy as a function of (a) time of retrogression and of (b) time of retrogression followed by reaging, for retrogression at 240 and 260°C .
the other hand, the structure of the specimens retrogressed at 240°C still contained d% precipitates, albeit mainly the larger ones, even after prolonged periods of retrogression (Fig. 2(e)). The dark field images of the d% precipitates obtained after different times of retrogression followed by reaging are shown in Fig. 3(a) – (d). Fig. 3(a) and 3(b) show the d% precipitates in the samples retrogressed at 260°C for 5 and 30 min, respectively, and then subsequently reaged. Fig. 3(c) and (d) show d% precipitates in samples subjected to retrogression at 240°C for 20 and 40 min respectively, and then reaged. It can be easily seen that both retrogression and reaging treatments result in substantial changes in precipitate size distribution. A clear quantitative understanding of all these changes in size and number of d% precipitates is ob-
tained from the histograms in Figs. 4 and 5. These show the overall size distribution of d% particles, as measured from the dark field images, in terms of the number of the particles of different size per unit area of the electron micrograph. Fig. 4(a) represents material in the as-received condition, Fig. 4(b)–(d) are for retrogression periods at 260°C and Fig. 4(e) and (g) relate to material retrogressed at 260°C as above and then reaged. The histograms in Fig. 5(a)–(c) are for periods of retrogression at 240°C. Fig. 5(d)–(f) are for retrogression at 240°C as above and then reaged. It can be seen from the examination of these histograms that retrogression up to 30 min at 260°C results in the progressive dissolution of the d% particles until they have almost entirely disappeared. On subsequent reaging of a specimen retrogressed for 5 min at 260°C there
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Fig. 2. Dark-field TEM micrographs showing d% particles in an 8090 Al – Li alloy. (a) As-received condition, (b) retrogressed at 260°C for 5 min, (c) retrogressed at 260°C for 30 min, (d) retrogressed at 240°C for 20 min, and (e) retrogressed at 240°C for 40 min.
is formed a broad range of size of d% particles, as can be seen in Fig. 4(e). This probably represents mainly simple coarsening of pre-existing d% particles, but also the nucleation and growth of new fine d% particles. For specimens retrogressed for 20 and 30 min at 260°C, subsequent reaging results primarily in the precipitation of new fine d% particles as seen in Fig. 4(f) and (g). These have formed from a retrogressed matrix largely denuded of d% particles. Fig. 5(a)–(c) show that retrogression for up to 40 min at 240°C results in the gradual dissolution of d% particles, however, even after 40 min retrogression, there is still a substantial amount of undissolved d% particles over the whole initial range of size. Concomitant with the above dissolution there occurs some coarsening of d% particles. Fig. 5(d)–(f) show that on reaging after retrogression at 240°C there is a tendency for the d% particles to
coarsen at the expense of the total number of finer particles. In addition to the data represented in the histograms, the observed area fraction covered by the d% particles on the electron micrographs was estimated. Then the volume fraction of the d% phase in the alloy subjected to various retrogression and reaging treatments was calculated using the expression given by Underwood [11] for spherical particles of diameter D fd% = − 2 ln(1− A)D/(D +3t) where A is the area fraction of the d% particles as measured from the dark-field TEM images and t is the thickness of the foil. The diameter D of the particles was replaced by the average particle diameter D obtained from: D= % Ni Di
,
% Ni
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Fig. 3. Dark-field TEM micrographs showing d% particles in an 8090 Al – Li alloy which was retrogressed and subsequently reaged. (a) Retrogressed at 260°C for 5 min, (b) retrogressed at 260°C for 30 min, (c) retrogressed at 240°C for 20 min, and (d) retrogressed at 240°C for 40 min.
where Ni is the number of observed particles with the size Di. Graphical representation of the results of the calculation of fd% is given in Fig. 6. It can be seen from Fig. 6(a) that for retrogression at 240°C the curve R-240 displays a minimum which occurs after approximately 4 min of heating, similarly to the behavior of the hardness curve R-240 in Fig. 1(a). The volume fraction of d% particles in the reaged material (curve RR-240) is practically restored. In contrast to this, for retrogression at 260°C, the volume fraction of the d% precipitates gradually decreases with the time, and the reaging treatment only partially restores the amount of the d% phase (Fig. 6(b)).
3.3. d Phase precipitates The d phase particles observed are generally without specific morphology although they tend to be elongated. In as-received material, and in that after relatively short times of retrogression, the d particles appear mainly at grain and sub-grain boundaries. During further retrogression the number of grain boundary d precipitates increases and this is accompanied by their appreciable coarsening. The evolution of the d phase particles with time of retrogression can be followed in Fig. 7(a)–(d). The insert in Fig. 7(a) shows a typical
selected area diffraction pattern taken from a d phase particle in as-received material. For long periods of retrogression at 260°C, the d particles grow also in the interior of the grains (Fig. 7(b)), however pronounced growth of the d phase in the body of the grains occurs in the course of reaging as shown in Fig. 7(c). Fig. 7(d) shows the d particles which have formed after 40 min of retrogression at 240°C followed by reaging. The d particles within the grains were identified in TEM by analysis of selected area diffraction patterns while the increase in their amount was confirmed by X-ray diffraction as will be seen in Fig. 9(a). The coarsening of the d particles results in depletion of lithium solute atoms in the surrounding matrix. Therefore such zones in the matrix are free of d% precipitates, as seen in Fig. 3(b) and (d). The dark areas in these dark field electron micrographs are associated with the presence of the d particles. The concentration of the d particles after RR-240 treatment is significantly higher for specimens retrogressed at 260°C than it is at 240°C.
3.4. The T1 and S phase precipitates The T1 and S precipitates were observed by TEM in the as-received alloy and at all stages of retrogression and reaging treatment. In the 100-oriented foils the
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Fig. 4. Histograms of the distribution of the size (radius) of the d% particles in an 8090 Al – Li alloy. (a) As received condition, (b)–(d) for retrogression of 5, 20 and 30 min at 260°C respectively, (e) – (g) as (b) – (d) but after reaging respectively.
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Fig. 5. Histograms of the distribution of the size (radius) of the d% particles in an 8090 Al – Li alloy. (a) – (c) For retrogression of 5, 20 and 40 min at 240°C respectively, (d)–(f) as (a)–(c) but after reaging respectively.
S-particles appeared as fine laths with a 100A1 growth direction, while the orientation of traces of T1 plates was 110A1. This was in keeping with the fact that T1 precipitates have a {111} habit plane [12] and the S phase grows on {210}A1 planes in the 001A1 direction [13]. For example, Fig. 8 shows the S and T1 precipitates in the as-received alloy. Because of the overlapping of the S and T1 precipitates it was difficult to estimate from the TEM micrographs how retrogression and reaging treatments influence their amount. This information was obtained,
however, for the T1 phase only, from the X-ray diffractograms taken at different stages of the treatment. It follows from the diffractograms in Fig. 9(a) that retrogression at 260°C favors growth of the T1 phase, as manifested by the appearance of the (10.3)T1 peak after 2 min of retrogression. This peak becomes stronger after 5 min of retrogression, and on retrogression for 40 min both the (10.0)T1 and (20.0)T 1 peaks grow. The disappearance of the (10.3)T1 peak after 40 min of retrogression can be ascribed to reorientation of the aluminium grains of the matrix during the prolonged
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Fig. 6. The volume fraction of the d% phase in an 8090 Al – Li alloy as calculated from the electron micrographs. (a) For retrogression at 240°C and (b) for retrogression at 260°C . The curves R relate to retrogression only, and the curves RR relate to material which was retrogressed and reaged.
time of heat treatment. It should be noted that there is no significant change in intensity of the T1 peaks after reaging. This indicates that the T1 phase grows only on the retrogression stage. As to the S phase, the latter was not detected in the diffractograms. In contrast to retrogression at 260°C, retrogression at 240°C does not result in growth of the T1 phase. Again, the peaks of the S phase were not detected. This is clearly seen from the diffractograms in Fig. 9(b) where only the changes of d and d% peaks are visible.
4. Discussion The first and probably main process occurring during retrogression heat treatment of the Al – Li 8090 alloy at 240 and 260°C is the dissolution of d% particles. This
dissolution should result in an increase of Li-content in the matrix. The decrease in the number of d% precipitate particles at 240° is, however, accompanied by the coarsening of larger d% particles at the expense of smaller ones as can be inferred both from the electron micrographs and from the histograms in Fig. 5(a)–(c). For 260°C, such coarsening is observed only for very short times of retrogression (see histogram, Fig. 4(b)). Despite the relatively small difference in temperature between 240 and 260°C, there is a substantial difference in the dissolution kinetics of d% particles at these two temperatures. Whereas, after 40 min of retrogression at 260°C, d% precipitates have largely disappeared, after 40 min of retrogression at 240°C a considerable number of d% precipitates of a broad range of size is still present. This observation would be in keeping with a solvus temperature of approximately 250°C for the d% phase as
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Fig. 7. Bright-field TEM micrographs showing d particles in an 8090 Al – Li alloy. (a) As-received condition, (b) retrogressed at 260°C for 40 min, (c) retrogressed at 260°C for 40 min and subsequently reaged, (d) retrogressed at 240°C for 40 min and subsequently reaged.
determined by Noble and Thomson [14] for an Al– 2%Li alloy. Thus for retrogression temperatures above 250°C the d% phase should disappear and for retrogression at 240°C d% precipitates should always remain, as was in fact observed. As follows from Fig. 6(b), the total volume fraction of the d% particles in the alloy substantially decreases during retrogression treatment at 260°C and after 30 min d% phase is completely dissolved. This suggests that the Li-solute, which dissolves in the matrix during retrogression, could then be available to other Li-containing phases, such as d and T1, thus allowing their growth in the matrix. This is precisely what was ob-
Fig. 8. Bright-field TEM micrograph showing S and T1 precipitates in an 8090 Al – Li alloy as received.
served in the X-ray diffractograms in Fig. 9(a): the higher peaks for both d and T1 show that these phases have grown in the matrix in the course of retrogression at 260°C. For a retrogression temperature of 240°C, however, growth of only the equilibrium d phase was observed. This is probably due to insufficient increase in Li-concentration in the matrix because of the incomplete solution of the d% phase and the concomitant coarsening of some of the d% particles. On the reaging following retrogression at 240°C, the Al-matrix, still enriched by the lithium solute, supplies Li-atoms to large undissolved d% precipitates and to the new nuclei of the d% phase. The lithium-containing equilibrium d phase also grows, however, as can be deduced from the X-ray diffractograms, no T1 phase is formed. As follows from the analysis of the dark field micrographs in Fig. 3, there is a significant difference in the precipitation behavior of the d% phase on reaging after retrogression at 240 and 260°C . Whereas precipitation following retrogression at 240°C is largely by the growth of existing d% particles, the characteristics of precipitation following retrogression at 260°C are dependent on the period of retrogression. Thus for a short period of retrogression, 5 min at 260°C, there is both a coarsening of existing d% precipitates and also the formation of new ones on reaging. For an extended period of retrogression at 260°C, more than 30 min, when the d% particles were completely dissolved, precipitation on reaging is entirely of new d% particles.
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Fig. 9. X-ray diffractograms taken at different stages of heat treatment of an 8090 Al – Li alloy. (a) Patterns R-260/2, R-260/5 and R-260/40 relate to retrogression for 2, 5 and 40 min at 260°C respectively. (b) Patterns R-240/20 and R-240/40 relate to retrogression for 20 and 40 min at 240°C respectively. RR denotes that the material was retrogressed and reaged.
The retrogression curves in Fig. 1 can now be explained in terms of the microstructural changes observed in the alloy. Turning to the changes in the volume fraction of d% precipitates after retrogression at 240°C, as shown in Fig. 6(a), a correlation will be seen between these two figures. The reduction in hardness in R-240 specimens and the reduction in volume fraction of d% precipitates both show a minimum after about 3–4 min of retrogression. Thereafter both parameters increase slightly in value and subsequently
become stabilized for the periods of retrogression of up to 40 min. For RR-240 specimens both the volume fraction of the d% precipitates and the hardness values virtually all return to their values prior to RR treatment for all the periods of retrogression investigated. This suggests that the change in the d% precipitate structure is the dominant factor which is responsible for the changes in strength of the Al–Li 8090 alloy in the course of RR treatment for a retrogression temperature 240°C.
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For the RR treatment which employed a retrogression temperature of 260°C, the relationship between hardness and changes in the volume fracture of d% particles can not, however, be explained simply in terms of the amount of the d% phase. The hardness curve in Fig. 1(a), for retrogression at 260°C, shows a minimum after about 2 min, thereafter the hardness increases slightly and after 5 min it more or less stabilizes, although the volume fraction of d% particles is still progressively decreasing. This implies that an additional hardening mechanism must operate continuously after 2 min of retrogression at 260°C . First of all, it might be that the increasing amount of Li atoms in solution increases the contribution of solution hardening to the strength of the material. Secondly, the precipitation of the T1 phase, as was observed by X-ray diffraction, could also be responsible for the cessation of softening on retrogression for periods longer than 2 min. The presence of the T1 phase may, probably, explain the hardness level of about 60RB (Fig. 1(a)) obtained after retrogression for 40 min at 260°C, at which point the volume fraction of d% particles is almost zero (Fig. 6(b)). This conclusion is consistent with the results of Huang and Ardell [15,16] who indicated that the T1 precipitates in an Al – Li – Cu alloy can be more potent strengtheners than d%, this even though their volume fraction is small compared with that of the d% precipitates. With regard to the possible effect of the S phase on changes in hardness on RR treatment, this could only have been relatively small since the S phase is not present in sufficient amount to be detected by X-ray diffraction. On reaging after retrogression at 260°C, although the volume fraction of d% precipitates increases, it is always substantially less than that of d% prior to RR treatment (Fig. 6(b)). A comparison of the histogram for material prior to RR heat treatment (Fig. 4(a)) with that for material retrogressed for 30 min at 260°C and then reaged (Fig. 4(g)), shows that the latter has almost the same total number of d% particles as the former, but the hardness of the RR treated specimen is considerably less than that prior to RR treatment. Since the vast bulk of the d% particles in the RR-260°C (30 min) specimen is very fine, the implication is that very fine d% precipitates are less effective in their contribution to hardness than coarser particles, for sizes in the range 1 –14 nm. This conclusion is in keeping with the observation by Watkinson and Martin [17] that mini d% particles formed at 80°C would have little effect on the yield strength of an 8090 Al – Li alloy.
5. Summary and conclusions (1) The microstructure of the precipitates in a commercial . Al–Li alloy has been studied through the stages
49
of retrogression and reaging. For the two temperatures of retrogression employed, 240 and 260°C, the behavior of the precipitates was specific for each temperature. At 240°C there was partial dissolution of d% particles, but this was accompanied by a simultaneous coarsening of others. At 260°C total dissolution of d% was observed after 30 min of retrogression. (2) Reaging following retrogression at 240 and 260°C resulted respectively in the full and only partial restoration of the volume fraction of the d% phase. (3) Change in the d% precipitate structure is the dominant factor which is responsible for the change in hardness of the Al–Li 8090 alloy in the course of RR treatment for a retrogression temperature of 240°C. (4) For retrogression treatment at 260°C, changes in the hardness of the alloy could not be explained simply in terms of changes in the d% particle structure. The precipitation of an additional phase, T1, concomitant with the dissolution of d% phase, possibly contributes to the hardness on this retrogression stage of the treatment. On the reaging step, following retrogression at 260°C, reprecipitation of the d% particles causes only a partial restoration of the hardness.
References [1] J.B. Lumsden, A.T. Allen, Corrosion Sci. 44 (1988) 527. [2] M.A. Reynolds, A. Gray, E. Creed, R.M. Jordan, A.P. Titchener, in: C. Baker et al. (Eds.), Aluminium – Lithium III, The Institute of Metals, London, 1986, pp. 57 – 65. [3] A. Gray, J. de Physique, Colloque 3 (48) (1987) 891–904. [4] N.J.H. Holroyd, A. Gray, G.M. Scamans, R. Hermann, in: C. Baker et al., (Eds.), Alminium – Lithium III, The Institute of Metals, London, 1986, pp. 310 – 320. [5] B. Cina, US Patent No. 3856584, December 24, 1974. [6] B. Cina, B. Ranish, Paper No. XXV, Aluminum Industrial Products, American Society for Metals, Pittsburgh, PA, October, 1974. [7] B. Cina, F. Zeidess, Materials Science Forum, Trans. Tech. Publications, vol. 102/104, Switzerland, 1992, pp. 99–108. [8] Z.Q. Hu, Y. Zhang, Y.L. Liu, Z.Y. Zhu, Corrosion 49 (1993) 491. [9] V. Komisarov, M. Talianker, B. Cina, Mater. Sci. Eng. A221 (1996) 113. [10] Optilab (v 2.0) — Image Processing and Analysis Software, Copyrighted 227 by Graftek France, 1992. [11] E.E. Underwood, in: M. Cohen (Ed.), Quantitative Stereology, Addison – Wesley, Reading, MA, 1970, p. 178. [12] J.C. Huang, A.J. Ardell, in: C. Baker, P.J. Gregson, S.J. Harris, C.J. Peel (Eds.), Aluminium – Lithium Alloys III, The Institute of Metals, London, 1986, pp. 455 – 70. [13] R.N. Wilson, P.G. Partridge, Acta Metall. 13 (1965) 1321. [14] B. Noble, G.E. Thompson, Metal Sci. J. 5 (1971) 114. [15] J.C. Huang, A.J. Ardell, in: G. Champier, B. Dubost, D. Minnay, L. Sabetay, (Eds.), Proceedings of the Fourth Aluminium– Lithium Conference, vol. 48, Les Edit. de Physique, Paris, J. Phys., 1987, pp. 373 – 383. [16] J.C. Huang, A.J. Ardell, Acta Metall. 36 (1988) 2995. [17] P.C. Watkinson, J.W. Martin, Materials Characterization 33 (1994) 11.
Effect of retrogression and reaging on the precipitates in an 8090 Al–Li alloy V. Komisarov a, M. Talianker a,*, B Cina b a
Department of Materials Engineering, Ben Gurion Uni6ersity of the Nege6, Beer-She6a, Israel b Metallurgy Group, Israel Aircraft Industries Limited, Lod-70100, Israel Received 6 May 1997; received in revised form 8 July 1997
Abstract The microstructure of the precipitates in an 8090 Al – Li alloy has been studied qualitatively and quantitatively by techniques of transmission electron microscopy through the stages of retrogression and reaging (RR). The retrogression was performed for various times at temperatures of 240 and 260°C. The reaging treatment was carried out at 170°C. The behavior of the precipitates was specific for each temperature of retrogression: while for retrogression at 240°C there was only partial dissolution of d% particles, total dissolution of d% was observed after 30 min of treatment at 260°C. Reaging following retrogression at 240 and 260°C resulted respectively in the full and only partial restoration of the volume fraction of the d% phase. Change in the d% precipitate structure was the dominant factor which was responsible for the changes in hardness of the Al – Li 8090 alloy in the course of RR treatment for a retrogression temperature of (RR-240). On the other hand, changes in the hardness of the alloy subjected to RR-260 treatment could not be explained simply in terms of changes in the d% particle structure. The precipitation of an additional phase, T1, concomitant with the dissolution of d% phase at 260°C, also contributes to the hardness on this retrogression stage of the treatment. © 1998 Elsevier Science S.A. Keywords: Retrogression; Reaging; Precipitates
1. Introduction Aluminium–lithium alloys are of considerable potential interest to the aerospace industry because of their significantly higher values of specific strength and modulus compared with conventional aluminium alloys. Although such Al – Li alloys have been developed intensively in the past 10 years or so, their utilization has been much less than anticipated. This is because of problems of reduced ductility and fracture toughness, marked anisotropy of mechanical properties especially in non-recrystallised structures and the inability, as yet, to obtain resistance to stress corrosion akin to that obtainable from the high strength 7075-aluminium alloy in a T73 temper [1 – 4]. Some years ago a non-conventional two-stage heat treatment process known as retrogression and reaging (RR) was developed by one of the present authors [5,6]. * Corresponding author. Tel.: + 972 7 461475; fax: + 972 7 472946. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 0 2 - 9
This process overcame the problem of susceptibility to stress corrosion for aluminium alloys of the 7000 type. In contrast to material in the T73 temper which results in overaging of alloys, material in the RR temper combines the mechanical properties of the T6 temper (of maximum tensile strength) with the full resistance to stress corrosion of the T73 temper. RR has been shown to be applicable to a broad range of aluminium alloys of the 7000 type [7]. Since Al–Li alloys constitute an age-hardening system wherein the alloys are hardened by the precipitation of d%, a coherent ordered phase of composition Al3Li, theoretically this system, when retrogressed and subsequently subjected to reaging, should tend to re store its hardness. Some preliminary results on the effect of RR heat treatments on the susceptibility to stress corrosion of a commercially available 8090 Al–Li alloy were reported by Hu et al. [8]. Recently the present authors undertook a more systematic investigation of this alloy [9] with the aim of determining to what extent retrogression and reaging could be applied
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in order to achieve an optimum combination of strength and resistance to stress corrosion. It was found that retrogression and reaging heat treatment of an 8090 Al–Li alloy, initially in the T8771 temper, can substantially improve its resistance to stress corrosion. The present paper presents the results of the second part of this work which covers the transmission electron-microscopic investigation of the microstructural evolution and the change in hardness of an 8090 Al–Li alloy subjected to RR treatment.
2. Experimental procedures The material investigated was a plate of an 8090 Al–Li alloy, 50.8 mm thick, manufactured by Alcan, UK. The material was supplied in the T8771 temper, i.e. solution treated, stretched 6% and aged for 32 h at 170°C. The chemical composition of the alloy in weight percent is given in Table 1. All RR treatments were performed on specimens, 22× 22 by 50.8 mm long, which were taken from the S-T direction of the plate, having their length perpendicular to the rolling direction. The retrogression of the alloy was performed in a molten salt bath (50% sodium nitrate/50% sodium nitrite) at two temperatures, 240 and 260°C, for time intervals ranging from 5 to 40 min. The bath size was large enough (1× 2 ×1 m) to provide very rapid heating of specimens to the desired temperature. The reaging treatment was the same for all retrogressed specimens and was performed in an air furnace at 170°C for 32 h. One series of specimens was retrogressed and then quenched in water. Another series of specimens was retrogressed, quenched in water and subsequently reaged. All heat treatments were monitored by the measurement of the Rockwell ‘B’ hardness of the specimens. The microstructure of the specimens was studied in a transmission electron microscope operating at 200 Kv. Samples for TEM analysis were cut in the form of slices, 0.2 mm thick, using an electric spark erosion machine, and final thinning was achieved by means of a twin-jet electropolishing unit (Tenupol), using a solution of 30% nitric acid and 70% methanol at − 18°C. The diameters of the spherical d% precipitates were measured from the dark-field electron micrographs using the d% superlattice reflection. A statistical treatment of the size and density of the d% particles per unit area Table 1 Chemical composition of the 8090 Al–Li alloy (wt%) Li
Cu
Mg
Fe
Si
Zr
Na
Al
2.35
1.15
0.74
0.05
0.03
0.11
0.0002
Balance
of electron micrograph was performed using Optilab image processing and an analysis software program [10].
3. Results
3.1. Hardness measurements Fig. 1(a) shows the change of hardness versus time of retrogression for two temperatures, 240 and 260°C. The variation of the hardness of retrogressed and reaged material as a function of retrogression time is shown in Fig. 1(b). It can be seen (Fig. 1(a)) that the general shape of the retrogression curves obtained in the present investigation for the 8090 Al–Li alloy is very similar to those obtained for a 7000 type aluminium alloy [6]. The hardness curves in Fig. 1(a) decrease very rapidly during retrogression until a minimum occurs, after which some increase in hardness is observed and thereafter there is no further change in hardness for the times investigated. After retrogression at 240°C and subsequent reaging (Fig. 1(b)), the original hardness level obtained prior to RR treatment was restored and this did not change significantly with retrogression time. For retrogression performed at 260°C, followed by reaging, the original hardness was initially restored but this began to decrease after only 10 min of retrogression when followed by reaging.
3.2. The beha6ior of the d% precipitates Dark field micrographs showing d% particles were taken from regions having the same foil thickness (of the order of 100 nm) and for the same diffraction conditions (zone area [001]), using the same superlattice d% reflection (100). Typical examples of dark-field images of the d% precipitates for the alloy subjected to different retrogression treatments are shown in Fig. 2(a)–(e). The d% precipitate structure of the alloy in the as-received condition is shown in Fig. 2(a). It is seen that the d% particles are fairly uniformly distributed in the matrix, are small in size (0.0015–0.02 mm) and that the number of particles per unit area of the image is large. Fig. 2(b) and (c) show the d% precipitates after retrogression at 260°C for 5 and 30 min respectively. Fig. 2(d) and (e) show the d% precipitates in a specimen retrogressed at 240°C for 20 and 40 min, respectively. Comparison with the as-received condition indicates that a relatively short time (5 min) of retrogression at 260°C causes a significant decrease in the density of the d% precipitates. Almost all the finer d% precipitates and many of the larger ones have now dissolved. Longer periods of retrogression led to further decrease in the density of the d% particles, and after 30 min at 260°C d% precipitates disappeared almost entirely (Fig. 2(c)). On
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Fig. 1. Variation of hardness of an 8090 Al–Li alloy as a function of (a) time of retrogression and of (b) time of retrogression followed by reaging, for retrogression at 240 and 260°C .
the other hand, the structure of the specimens retrogressed at 240°C still contained d% precipitates, albeit mainly the larger ones, even after prolonged periods of retrogression (Fig. 2(e)). The dark field images of the d% precipitates obtained after different times of retrogression followed by reaging are shown in Fig. 3(a) – (d). Fig. 3(a) and 3(b) show the d% precipitates in the samples retrogressed at 260°C for 5 and 30 min, respectively, and then subsequently reaged. Fig. 3(c) and (d) show d% precipitates in samples subjected to retrogression at 240°C for 20 and 40 min respectively, and then reaged. It can be easily seen that both retrogression and reaging treatments result in substantial changes in precipitate size distribution. A clear quantitative understanding of all these changes in size and number of d% precipitates is ob-
tained from the histograms in Figs. 4 and 5. These show the overall size distribution of d% particles, as measured from the dark field images, in terms of the number of the particles of different size per unit area of the electron micrograph. Fig. 4(a) represents material in the as-received condition, Fig. 4(b)–(d) are for retrogression periods at 260°C and Fig. 4(e) and (g) relate to material retrogressed at 260°C as above and then reaged. The histograms in Fig. 5(a)–(c) are for periods of retrogression at 240°C. Fig. 5(d)–(f) are for retrogression at 240°C as above and then reaged. It can be seen from the examination of these histograms that retrogression up to 30 min at 260°C results in the progressive dissolution of the d% particles until they have almost entirely disappeared. On subsequent reaging of a specimen retrogressed for 5 min at 260°C there
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Fig. 2. Dark-field TEM micrographs showing d% particles in an 8090 Al – Li alloy. (a) As-received condition, (b) retrogressed at 260°C for 5 min, (c) retrogressed at 260°C for 30 min, (d) retrogressed at 240°C for 20 min, and (e) retrogressed at 240°C for 40 min.
is formed a broad range of size of d% particles, as can be seen in Fig. 4(e). This probably represents mainly simple coarsening of pre-existing d% particles, but also the nucleation and growth of new fine d% particles. For specimens retrogressed for 20 and 30 min at 260°C, subsequent reaging results primarily in the precipitation of new fine d% particles as seen in Fig. 4(f) and (g). These have formed from a retrogressed matrix largely denuded of d% particles. Fig. 5(a)–(c) show that retrogression for up to 40 min at 240°C results in the gradual dissolution of d% particles, however, even after 40 min retrogression, there is still a substantial amount of undissolved d% particles over the whole initial range of size. Concomitant with the above dissolution there occurs some coarsening of d% particles. Fig. 5(d)–(f) show that on reaging after retrogression at 240°C there is a tendency for the d% particles to
coarsen at the expense of the total number of finer particles. In addition to the data represented in the histograms, the observed area fraction covered by the d% particles on the electron micrographs was estimated. Then the volume fraction of the d% phase in the alloy subjected to various retrogression and reaging treatments was calculated using the expression given by Underwood [11] for spherical particles of diameter D fd% = − 2 ln(1− A)D/(D +3t) where A is the area fraction of the d% particles as measured from the dark-field TEM images and t is the thickness of the foil. The diameter D of the particles was replaced by the average particle diameter D obtained from: D= % Ni Di
,
% Ni
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Fig. 3. Dark-field TEM micrographs showing d% particles in an 8090 Al – Li alloy which was retrogressed and subsequently reaged. (a) Retrogressed at 260°C for 5 min, (b) retrogressed at 260°C for 30 min, (c) retrogressed at 240°C for 20 min, and (d) retrogressed at 240°C for 40 min.
where Ni is the number of observed particles with the size Di. Graphical representation of the results of the calculation of fd% is given in Fig. 6. It can be seen from Fig. 6(a) that for retrogression at 240°C the curve R-240 displays a minimum which occurs after approximately 4 min of heating, similarly to the behavior of the hardness curve R-240 in Fig. 1(a). The volume fraction of d% particles in the reaged material (curve RR-240) is practically restored. In contrast to this, for retrogression at 260°C, the volume fraction of the d% precipitates gradually decreases with the time, and the reaging treatment only partially restores the amount of the d% phase (Fig. 6(b)).
3.3. d Phase precipitates The d phase particles observed are generally without specific morphology although they tend to be elongated. In as-received material, and in that after relatively short times of retrogression, the d particles appear mainly at grain and sub-grain boundaries. During further retrogression the number of grain boundary d precipitates increases and this is accompanied by their appreciable coarsening. The evolution of the d phase particles with time of retrogression can be followed in Fig. 7(a)–(d). The insert in Fig. 7(a) shows a typical
selected area diffraction pattern taken from a d phase particle in as-received material. For long periods of retrogression at 260°C, the d particles grow also in the interior of the grains (Fig. 7(b)), however pronounced growth of the d phase in the body of the grains occurs in the course of reaging as shown in Fig. 7(c). Fig. 7(d) shows the d particles which have formed after 40 min of retrogression at 240°C followed by reaging. The d particles within the grains were identified in TEM by analysis of selected area diffraction patterns while the increase in their amount was confirmed by X-ray diffraction as will be seen in Fig. 9(a). The coarsening of the d particles results in depletion of lithium solute atoms in the surrounding matrix. Therefore such zones in the matrix are free of d% precipitates, as seen in Fig. 3(b) and (d). The dark areas in these dark field electron micrographs are associated with the presence of the d particles. The concentration of the d particles after RR-240 treatment is significantly higher for specimens retrogressed at 260°C than it is at 240°C.
3.4. The T1 and S phase precipitates The T1 and S precipitates were observed by TEM in the as-received alloy and at all stages of retrogression and reaging treatment. In the 100-oriented foils the
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Fig. 4. Histograms of the distribution of the size (radius) of the d% particles in an 8090 Al – Li alloy. (a) As received condition, (b)–(d) for retrogression of 5, 20 and 30 min at 260°C respectively, (e) – (g) as (b) – (d) but after reaging respectively.
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Fig. 5. Histograms of the distribution of the size (radius) of the d% particles in an 8090 Al – Li alloy. (a) – (c) For retrogression of 5, 20 and 40 min at 240°C respectively, (d)–(f) as (a)–(c) but after reaging respectively.
S-particles appeared as fine laths with a 100A1 growth direction, while the orientation of traces of T1 plates was 110A1. This was in keeping with the fact that T1 precipitates have a {111} habit plane [12] and the S phase grows on {210}A1 planes in the 001A1 direction [13]. For example, Fig. 8 shows the S and T1 precipitates in the as-received alloy. Because of the overlapping of the S and T1 precipitates it was difficult to estimate from the TEM micrographs how retrogression and reaging treatments influence their amount. This information was obtained,
however, for the T1 phase only, from the X-ray diffractograms taken at different stages of the treatment. It follows from the diffractograms in Fig. 9(a) that retrogression at 260°C favors growth of the T1 phase, as manifested by the appearance of the (10.3)T1 peak after 2 min of retrogression. This peak becomes stronger after 5 min of retrogression, and on retrogression for 40 min both the (10.0)T1 and (20.0)T 1 peaks grow. The disappearance of the (10.3)T1 peak after 40 min of retrogression can be ascribed to reorientation of the aluminium grains of the matrix during the prolonged
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Fig. 6. The volume fraction of the d% phase in an 8090 Al – Li alloy as calculated from the electron micrographs. (a) For retrogression at 240°C and (b) for retrogression at 260°C . The curves R relate to retrogression only, and the curves RR relate to material which was retrogressed and reaged.
time of heat treatment. It should be noted that there is no significant change in intensity of the T1 peaks after reaging. This indicates that the T1 phase grows only on the retrogression stage. As to the S phase, the latter was not detected in the diffractograms. In contrast to retrogression at 260°C, retrogression at 240°C does not result in growth of the T1 phase. Again, the peaks of the S phase were not detected. This is clearly seen from the diffractograms in Fig. 9(b) where only the changes of d and d% peaks are visible.
4. Discussion The first and probably main process occurring during retrogression heat treatment of the Al – Li 8090 alloy at 240 and 260°C is the dissolution of d% particles. This
dissolution should result in an increase of Li-content in the matrix. The decrease in the number of d% precipitate particles at 240° is, however, accompanied by the coarsening of larger d% particles at the expense of smaller ones as can be inferred both from the electron micrographs and from the histograms in Fig. 5(a)–(c). For 260°C, such coarsening is observed only for very short times of retrogression (see histogram, Fig. 4(b)). Despite the relatively small difference in temperature between 240 and 260°C, there is a substantial difference in the dissolution kinetics of d% particles at these two temperatures. Whereas, after 40 min of retrogression at 260°C, d% precipitates have largely disappeared, after 40 min of retrogression at 240°C a considerable number of d% precipitates of a broad range of size is still present. This observation would be in keeping with a solvus temperature of approximately 250°C for the d% phase as
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Fig. 7. Bright-field TEM micrographs showing d particles in an 8090 Al – Li alloy. (a) As-received condition, (b) retrogressed at 260°C for 40 min, (c) retrogressed at 260°C for 40 min and subsequently reaged, (d) retrogressed at 240°C for 40 min and subsequently reaged.
determined by Noble and Thomson [14] for an Al– 2%Li alloy. Thus for retrogression temperatures above 250°C the d% phase should disappear and for retrogression at 240°C d% precipitates should always remain, as was in fact observed. As follows from Fig. 6(b), the total volume fraction of the d% particles in the alloy substantially decreases during retrogression treatment at 260°C and after 30 min d% phase is completely dissolved. This suggests that the Li-solute, which dissolves in the matrix during retrogression, could then be available to other Li-containing phases, such as d and T1, thus allowing their growth in the matrix. This is precisely what was ob-
Fig. 8. Bright-field TEM micrograph showing S and T1 precipitates in an 8090 Al – Li alloy as received.
served in the X-ray diffractograms in Fig. 9(a): the higher peaks for both d and T1 show that these phases have grown in the matrix in the course of retrogression at 260°C. For a retrogression temperature of 240°C, however, growth of only the equilibrium d phase was observed. This is probably due to insufficient increase in Li-concentration in the matrix because of the incomplete solution of the d% phase and the concomitant coarsening of some of the d% particles. On the reaging following retrogression at 240°C, the Al-matrix, still enriched by the lithium solute, supplies Li-atoms to large undissolved d% precipitates and to the new nuclei of the d% phase. The lithium-containing equilibrium d phase also grows, however, as can be deduced from the X-ray diffractograms, no T1 phase is formed. As follows from the analysis of the dark field micrographs in Fig. 3, there is a significant difference in the precipitation behavior of the d% phase on reaging after retrogression at 240 and 260°C . Whereas precipitation following retrogression at 240°C is largely by the growth of existing d% particles, the characteristics of precipitation following retrogression at 260°C are dependent on the period of retrogression. Thus for a short period of retrogression, 5 min at 260°C, there is both a coarsening of existing d% precipitates and also the formation of new ones on reaging. For an extended period of retrogression at 260°C, more than 30 min, when the d% particles were completely dissolved, precipitation on reaging is entirely of new d% particles.
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Fig. 9. X-ray diffractograms taken at different stages of heat treatment of an 8090 Al – Li alloy. (a) Patterns R-260/2, R-260/5 and R-260/40 relate to retrogression for 2, 5 and 40 min at 260°C respectively. (b) Patterns R-240/20 and R-240/40 relate to retrogression for 20 and 40 min at 240°C respectively. RR denotes that the material was retrogressed and reaged.
The retrogression curves in Fig. 1 can now be explained in terms of the microstructural changes observed in the alloy. Turning to the changes in the volume fraction of d% precipitates after retrogression at 240°C, as shown in Fig. 6(a), a correlation will be seen between these two figures. The reduction in hardness in R-240 specimens and the reduction in volume fraction of d% precipitates both show a minimum after about 3–4 min of retrogression. Thereafter both parameters increase slightly in value and subsequently
become stabilized for the periods of retrogression of up to 40 min. For RR-240 specimens both the volume fraction of the d% precipitates and the hardness values virtually all return to their values prior to RR treatment for all the periods of retrogression investigated. This suggests that the change in the d% precipitate structure is the dominant factor which is responsible for the changes in strength of the Al–Li 8090 alloy in the course of RR treatment for a retrogression temperature 240°C.
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For the RR treatment which employed a retrogression temperature of 260°C, the relationship between hardness and changes in the volume fracture of d% particles can not, however, be explained simply in terms of the amount of the d% phase. The hardness curve in Fig. 1(a), for retrogression at 260°C, shows a minimum after about 2 min, thereafter the hardness increases slightly and after 5 min it more or less stabilizes, although the volume fraction of d% particles is still progressively decreasing. This implies that an additional hardening mechanism must operate continuously after 2 min of retrogression at 260°C . First of all, it might be that the increasing amount of Li atoms in solution increases the contribution of solution hardening to the strength of the material. Secondly, the precipitation of the T1 phase, as was observed by X-ray diffraction, could also be responsible for the cessation of softening on retrogression for periods longer than 2 min. The presence of the T1 phase may, probably, explain the hardness level of about 60RB (Fig. 1(a)) obtained after retrogression for 40 min at 260°C, at which point the volume fraction of d% particles is almost zero (Fig. 6(b)). This conclusion is consistent with the results of Huang and Ardell [15,16] who indicated that the T1 precipitates in an Al – Li – Cu alloy can be more potent strengtheners than d%, this even though their volume fraction is small compared with that of the d% precipitates. With regard to the possible effect of the S phase on changes in hardness on RR treatment, this could only have been relatively small since the S phase is not present in sufficient amount to be detected by X-ray diffraction. On reaging after retrogression at 260°C, although the volume fraction of d% precipitates increases, it is always substantially less than that of d% prior to RR treatment (Fig. 6(b)). A comparison of the histogram for material prior to RR heat treatment (Fig. 4(a)) with that for material retrogressed for 30 min at 260°C and then reaged (Fig. 4(g)), shows that the latter has almost the same total number of d% particles as the former, but the hardness of the RR treated specimen is considerably less than that prior to RR treatment. Since the vast bulk of the d% particles in the RR-260°C (30 min) specimen is very fine, the implication is that very fine d% precipitates are less effective in their contribution to hardness than coarser particles, for sizes in the range 1 –14 nm. This conclusion is in keeping with the observation by Watkinson and Martin [17] that mini d% particles formed at 80°C would have little effect on the yield strength of an 8090 Al – Li alloy.
5. Summary and conclusions (1) The microstructure of the precipitates in a commercial . Al–Li alloy has been studied through the stages
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of retrogression and reaging. For the two temperatures of retrogression employed, 240 and 260°C, the behavior of the precipitates was specific for each temperature. At 240°C there was partial dissolution of d% particles, but this was accompanied by a simultaneous coarsening of others. At 260°C total dissolution of d% was observed after 30 min of retrogression. (2) Reaging following retrogression at 240 and 260°C resulted respectively in the full and only partial restoration of the volume fraction of the d% phase. (3) Change in the d% precipitate structure is the dominant factor which is responsible for the change in hardness of the Al–Li 8090 alloy in the course of RR treatment for a retrogression temperature of 240°C. (4) For retrogression treatment at 260°C, changes in the hardness of the alloy could not be explained simply in terms of changes in the d% particle structure. The precipitation of an additional phase, T1, concomitant with the dissolution of d% phase, possibly contributes to the hardness on this retrogression stage of the treatment. On the reaging step, following retrogression at 260°C, reprecipitation of the d% particles causes only a partial restoration of the hardness.
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