ScriptaMatuialia, Vol. 37, No. 6, pp. 851-859,1997 Ekevia Science Ltd Copyright0 1997 Acta MctallurgicpInc. Printedin the USA. All rights -cd 1359-6462/97 s17.00 + .oo Pergamon PII S1359-6462(97)00186-3 STUDY OF THE 6’ REVERSION PROCESS IN 8090 ALLOYS O.A. Larnbri #, J.I. PCrez-Landazhbal, M.L. N6’ and J. San Juan Dpto. .FIsica de la Materia Condensada, Fat. Ciencias, Universidad de1 Pais Vasco, Apdo. 644,48080 Bilbao, Spain #Permanent address Institute de Fisica Rosario (CONICET-UNR), Avda. 27 de febrero 2 lObis, (2000) Rosario, Argentina ‘Dpto. Fisica Aplicada II, Fat. Ciencias Universidad de1 PaisVasco Apdo. 644,48080 Bilbao, Spain (Received November 29, 1996) (Accepted April 9,1997) Introduction Al-Li based alloys have been widely studied for their potential applications in the aerospace industry, owing to the fact that they exhibit a higher Young’s modulus and lower density than the conventional Al-based alloys. The 6’ phase, ALLI, is a strengthening phase in the 8090 Al-Li alloys [l, 21 and consequently, its zstability during the aging treatments is critical for optimizing processes of technological interest. In fact, the reversion process of the 6’ phase could present an interesting possibility for forging and cold working tasks owing to the decrease in hardness of the sample [3]. The reversion processes have been largely studied in Weldalite alloys [3-61 in order to optimize the microstructural control for strengthening the alloy from the T4 temper through reversion to the T6 temper. However the 6’ reversion processes have not been extensively studied in 8090 alloys. In particular, a previous work [7] has examined the influence of the reversion process (storage) on the toughness of this alloy. Nevertheless, in this work [7], only the reversion on a previously high temperature (170°C) treated alloy was evaluated. In this work, the microstructure evolution between room temperature (RT) aging of the as-quenched sample, and the microstructure associated with high temperature treatments in a 8090 alloy has been studied. Thermoelectric-power measurements (TEP), X-rays diffraction (XRD) and Vickers micro hardness test (VMH) were employed in this work in order to know simultaneously the microstructural state and the mechanical behaviour. Due to the precision of the TEP measurements, the present study shows clearly the kinetics of the 6’ reversion process. Therefore, the best reversion treatment in order to reduce the hardness of the as-quenched sample after a solid solution treatment can be accurately determined. Experimental Procedure Samples were obtained from a commercial 8090 alloy (Al-2.4wt?hLi- 1.2 1wl??Ku-0.77wt%Mg-0.05 wt%Zr). Two binaries alloys Al-2.5wt?hLi-O.lwt%Zr and Al-O.Swt?/oLi were also employed. The zyxwvuts 851 852 STUDYOF THE 6’ REVERSIONPROCESS Vol. 37, No. 6 material was supplied by Pechiney as a rolled plate with the T3 treatment. They were solution treated at 530°C for 30 min. in a salt bath and quenched into iced water. The high temperature aging were performed in a salt bath at 12O’C and 2OO’C. Subsequently, the samples were quenched in water at R.T. and they were examined at RT employing the different techniques. Then, the samples were introduced again into the salt bath to follow the aging treatment. The TEP measurement unit, TECHMETAL, is equipped with pure Aluminum blocks. The resolution of the equipment is H.002 pV/K. In order to follow the early evolution of the microstructure at RT immediately after quenching, we carried out RT aging “in situ” in the TEP unit with the low temperature block at T = (15.0 f zyxwvutsrqponmlkjihgfedcbaZYXWVUTSRQPONMLKJIHGFEDCBA 0.1)oC and with a difference between the low and high temperature block of AT = (10.0 f O.l)oC [8,9]. About two minutes were needed to start the measurement after quenching from solid solution because the sample surface needed to be polished. However the quenched sample was held in the iced water environment, even during mechanical polishing. The %P values measured refer to the absolute value of the Aluminum reference (-1.4 pV/K). Samples for TEP measurements were cut from the plate as 60 mm length strips. The maximum section of samples was less than 1.5 mm’ to avoid temperature gradients between it and the thermoelectric unit block. The micro-hardness measurements were performed employing an automatic LECO-AMH equipment. The samples employed in this test were plates 20 mm long, 5 mm wide and 1.2 mm thick. A layer of 0.15 mm was removed from the samples surface, by means of mechanical polishing, for avoiding the effect of lithium loss from the surface during the solid solution treatment [lo]. X-ray measurements were carried out in a STOE powder diffractometer working in transmission mode at RT and employing the sample rotation device to eliminate texture effects. The measurements conditions were 40 kV and 20 mA using the K, of Cu as incident radiation. The employed polycrystalline samples were polished to reduce the thickness below 0.3 mm after the solid solution treatment, in order to obtain a good diffraction spectra in transmission mode. zyxwvutsrqponmlkjihgfedcbaZYXWVUTSR 103 104 103 106 107 time(s) Figure 1. (a) TEP evolution at room temperaturefor both, an Al-O.S%wtLi, curve (l), and Al 2.25% Li, curve (2). (b) TEP evolution duringaging at room temperature,curve (l), and Vickqs micro-hardnesscurve (2), for a 8090 alloy. 853 zyxwvut STUDYOF THE 6’ REVERSIONPROCESS Vol. 37, No. 6 Results In order to follow the precipitation kinetics of Al-Li alloys during RT aging, the stability of the TEP unit was first verified using an Al-O.S%Li alloy which has a Li concentration lower than the solubility limit. Consequently, this sample should not show any precipitation processes after quenching. In fact, the curve (1) in Fig. la does not show any TEP evolution as a function of time during an in situ RT aging, as it is expected. Therefore, the high stability of the used TEP unit has been corroborated. In contrast, Al 2..5% Li (+Zr) binary alloys with the same thermal treatment (quenching plus in situ RT aging) show a IEP evolution which presents two stages, curve (2) in Figure la. The fmt stage below 5 x lo3 s is related to the nucleation of 6’ precipitates by a spinodal decomposition and the second one that starts at about 3 x lo4 s is related to the 6’ precipitates growth [ 1I]. On the other hand, the curve 1 in Fig. lb shows the TEP evolution at RT as a function of aging time in an as-quenched 8090 sample, under the same experimental conditions employed for measuring the two binary alloys. The TEP values reduce progressively following only one stage that shows the characteristic shape of a precipitation process in a logarithmic scale. This behaviour is different from the evolution in the Al-2.5 Li (+Zr) binary alloy where two stages can be observed (curve (2) in Fig. la). Curve 2 of Figure lb shows the microhardness evolution of the 8090 alloy during the RT aging. Despite the dispersion found in the microhardness data, it can be observed that as long as the TEP values decrease, the hardening of the sample increases. In order to follow the microstructural changes, X-ray powder diffraction measurements were performed in a 80190sample in the as-quenched condition, Fig. 2a, and after aging during lo6 seconds at RT, Fig. 2b. The 6’ precipitated state was characterized by the (l,O,O) diffraction peak at 28= 21.9“ which is the more intense diffraction peak of the 6’ superlattice not superimposed to the Al matrix diffraction zyxwvutsrqponmlkjihgfedcbaZYXWVUTSRQPONMLKJIHGFEDCBA peaks[10,121. In the as quenched sample the (l,O,O) reflection can not be seen, but after RT aging this reflection develops and the peak intensity increases as a function of time, although it is even a wide peak. zyxwvutsrqponmlkjihgfedcbaZYXWVUTSRQPONMLKJIHGFEDCBA 11:should be pointed out that, transmission electron microscopy (TEM) does not allow to follow the 6’ evolution as clearly as X-ray diffraction since TEM images are quickly saturated and consequently these are not suitable to quantify the evolution of the 6’ precipitated mass fraction. On a) 2 Quenched + 106 s at R.T. 10 15 20 25 30 35 40 45 50 28 Figure 2. X-ray lspectra in 8090 alloy, corresponding to the as quenched sample (upper plot) and after aging at room temperature during lo6 s (lower plot). 854 STUDY OF THE 6’ REVERSION PROCESS 100 101 102 103 104 10s Vol. 37, No. 6 106 zyxwvutsrqponmlkjihgfedcbaZYXW 107 time ( s ) Figure 3. TEP evolution in 8090 alloy duringaging at 2OO”C,curve (a) and at 120°C, curve (b). the contrary, X-ray diffraction spectrum refined by the Rietveld method allows to quantify the evolution of the mass fraction of 6’ precipitates [ 121. Figure 3 shows the time evolution of the TEP values in a 8090 sample at 2OO”C,curve (a), and at 12O”C, curve (b). The 200°C aging curve shows an initial slight increase of the TEP values until 10 s and then a two-stage decrease of the TEP values. The first one develops at about 20 s and last until about lo3 s and for longer times than 10’ s, the second one takes place. In the 120°C aging curve the TEP values decrease along a clear stage until lo5 s which was the maximum measured aging time. This stage is slowly shifted towards longer times comparing with the first stage of the 200°C curve. The correspondiig microstructural evolution of the sample at 200°C after aging at 6 x lo* s (a), 2 x 10’ s (b), 1.1 x 10’ s (c) and 1 x 106s (d) is shown in the X ray diffraction patterns of Figure 4. For increasing the aging time in the range of the first stage of the TEP curve (Fig. 3a), the 6’ (1 ,O,O)reflection becomes the more and more intense and sharp. Finally for aging times which correspond to the second stage in TEP curve, a new peak arise at 28 = 24.17” related to 6 stable phase, and the 6’ peak disappears. It is important to point out that there is not evidence of S’ precipitation in 8090 alloys in the time range studied. Nevertheless, the presence of S’ phase can be well determined by X-ray powder difhaction in Al-Cu-Mg and Al-Li-Cu-Mg alloys [ 131. This behaviour clearly indicates that the first stage is related to the growth of the 6’ precipitates and the second one is related to the re-dissolution of the 6 precipitates and the precipitation of the stable 6 phase according to the previous results reported in the literature [14,15]. During this process, the stable 6 phase pick-up lithium atoms from the matrix and reduces the concentration of Li in solid solution. As a consequence, the metastable 6’ phase dissolves in order to maintain the equilibrium with the surrounding solid solution. Nevertheless, the 6 phase capacity of pumping Li atoms from solid solution is higher than the ability of 6’ in releasing Li atoms, and the global observed effect is a slow decrease of TEP values (from 1 to 24 days) during the second stage in curve a of Fig. 3. It should be mentioned again that, the above presented results on the 6’ precipitation kinetics will be used in order to check accurately the reversion process. The TEP evolution during the aging treatment at 200% and 120°C performed on 8090 samples previously RT aged for lo6 s, i. e. at the end of curve 1 in Fig. lb, has been shown in Figure 5. The sample aged at 200°C, curve a, shows a relatively fast increase of the TEP values, where its maximum value is achieved for a time about 20 s, and subsequently the TEP decrease continuously. The aged sample at 120°C, curve b, shows only a slight increase of TEP value up to a maximum located around 5 x 10’ s, later on it decreases again. The TEP curves of Fig. 3, without previous RT aging, are also shown in Figure 5 for 855 STUDY OF THE 6’ REVERSION PROCESS Vol. 37, No. 6 b) d) o c 15 20 25 30 2e Figure 4. X-ray patternsin 8090 alloy aged at 200°C during6 x ld s (a), 2.1 x 10’ s (b), 1.1 x 10’s (c), and 1 x 10”s (d). The 6’ and S characteristicpeaks are identified and their evolution should be comparedwith the TEP evolution curve of Fig. 3. both temperatures (curves a’ and b’) for comparison. Moreover the TEP evolution for longer aging times (~10~ s) carried out in samples aged at RT (curves a and b) overlap the TEP values of samples in the as-quenched condition (curves a’ and b’). The VMH levolution during the aging treatment for the samples in Figure 5 has been plotted in Figure 6. The minimum values correspond approximately to the same times for which the maximum TEP values are achieved in Figure 5. The decrease in hardness in samples which were previously aged at R.T. is larger than in the quenched samples. Moreover, it can also be observed that the hardness of sample aged at 120°C practically does not decrease in agreement with the TEP values for the asquenched sample aged at 120°C (Figure 5b). Discussion As was alreaidy mentioned, the formation of 13’in Al-Li binary alloys after quenching from higher temperatures develops through a spinodal decomposition process, associated with a reduction in the TEP values as was shown in curve 2 of Figure la [ 111, but this stage did not appear in the quaternary 856 STUDYOF THE6'REVERSIONPROCESS Vol. 37,No. 6 -1 .o -1.l (b) A A ~~ time ( s ) Figure 5. TEP evolution of samples aged at 200°C (a) and 120% (b) after quenching and previously aged at room temperatwe during lo6 s. The TEP evolution of samples aged at 2OO’C(a’) and 12O’C(b’) after quenching, are also plotted for comparison. 8090alloy. The non-appearance of the spinodal stage could be explained by the presence of Cu and Mg which eliminates the spinodal decomposition since Cu and Mg atoms impede the formation of ordered clusters of minimal energy. However the observed variation in the TEP values as a function of aging time is connected with the precipitation process. The X-ray spectra shown in Figure 2 confirm this situation. In fact, Figure 2a shows the spectrum corresponding to the as-quenched state on a 8090 sample, but due to the delay during the slewing of the angle in the measurement, the (1 ,O,O)peak located at 28 = 21.9” is measured around 2 hours at RT after quenching. Consequently, this spectrum corresponds to the early stage of the TEP curve and qualitatively indicates that there exist little quantity of small precipitates in the matrix. After 1O6s at RT the X-ray spectrum shows a clear (1 ,O,O)6’ phase peak. Although this peak is smaller than 1% of the most intense aluminium peaks, the Rietveld method allows to quantify the mass fraction of 6’ [ 121, which in this case corresponds to 9 wt%. The difference between both spectra is patent. The large width of the peaks shows that the 6’ phase precipitates are very small in size as it could be expected. The formation of 6 precipitates at RT should finish when the metastable equilibrium is achieved. The RT hardening is associated to the increase in 150 zyxwvutsrqponmlkjihgfedcbaZYXWVUTSRQPONMLKJIHGFEDCBA 140 g 130 G $j 120 110 100 90 80 70 Figure 6. VMH evolution of samples aged at 200% (a) and 120°C (b) after quenchingand previously aged at room tempemtwe during lo6 s and VMH evolution of samples aged at 200°C (a’) and 12oOC(b’) after quencKmg. Vol. 37, No. 6 857 zyxwvutsrq SlWDY OF THE 6’ REVERSION PROCESS 0.7 0.6 (a): A.Q.+106 s at R.T. (b): (a)+10 s at 2Oo’C. (c): (a)+3 hours at 2Oo’C 0.1 0 16 19 20 21 22 23 24 25 28 Figure 7. X-ray powder diffraction spectra showing the evolution of the (l,O,O) reflections during the reversion process. Experimentaldata BF~plotted as black dots. the mass fraction of 6’ precipitates present in the alloy. This kind of behaviour can be predicted applying the hardening theory to the case of 6’ precipitates finely dispersed through the matrix where the straight dislocations move in pairs [ 16,171. GPB are not expected in low Cu and low Mg 8090 alloys [ 181 since the Lithium atoms trap vacancies during quenching due to the high Li-Vacancy binding energy (0.25 zyxwvutsrqponmlkjihgfedcbaZYXWVUTSRQPONMLKJIHGFEDCBA eV)l [ 19,201. On the other hand, from the high temperature aging which are shown in Fig. 3, it should be pointed out that the 200°C aging precipitation is faster than the 12O”C, which is well known. Moreover, the TEP values at the end of the 6’ precipitation process are lower when the aging is performed at 120°C than at 200°C, which indicates that the mass traction precipitated is larger for the 120°C than for the 200°C as it could be expected by the lever rule. The decreasing in hardness, curves (a’) and (b’) of Figure 6, could be indicating that some reversion process appears which is in agreement with the behaviour reported for Weldalite alloy [2-61. Moreover the decreasing in hardness in the curves (a) and (b), which correspond to samples which were previously RT aged, show a larger degree of softening. This fact is in agreement with a reversion process when the high temperature aging begins owing to the fact that the RT aging promotes the formation of higher quantities of 6’ precipitates, in comparison with the as-quenched state. Furthermore, the TEP evolution shown in Figure 5 for samples (a) and (b) corroborates that a reversion process exists, since the increase in TEP values is larger in the samples which were previously RT aged (samples a and b) than in the as-quenched ones (samples a’ and b’). In fact, when a previously RT aged sample is subjected to a high temperature aging, some 6’ precipitates dissolve increasing the Li concentration in solid solution due to two causes principally. Those are: (a) to reduce the interface energy and (b) for reducing the quantity of precipitated 6’ phase according to the lever rule. Therefore, a reversion process takes place at high temperatures. Particularly at 2OO”C,curve a in Fig. 5, where the driving force is higher, the quantity of 6’ phase in meta-equilibrium at 2OO’C is smaller than at 120°C and at RT. In effect, at 12O’C (curve b’ in Fig. 5) the reversion process is smaller than at 2OO“Cand the maximum TEP value shifts to longer times. The maximum of the reversion process takes place at 20 s and lo4 s in samples aged at 200°C and 120°C respectively. For longer aging times at high temperatures, the influence of the previous natural aging can be neglected in agreement with previous results [7]. The involved reversion time in 8090 alloys is shorter than the one observed in Weldalite alloys [3-51 owing to the Li concentration in Weldalite being smaller than the Li concentration in 8090 alloys, decreasing 858 STUDYOF THE 8’ REVERSIONPROCESS Vol. 37, No. 6 consequently, the reversion kinetics. Furthermore, this reversion in 8090 is observed directly in the evolution of (l,O,O) reflection associated to the 6’ phase. A sample aged at RT during lo6 s shows a small and wide (l,O,O) reflection, Fig. 7(a), which after Rietveld analysis corresponds to around 9O/owt of 6’. The same sample aged at 200°C during 10 seconds shows a clear reduction in the integrated intensity of the (1 ,O,O)peak, Fig. 7(b), which corresponds to a quantity of the 6’ phase precipitated, about 6wt%. The observed reversion process in this work is more important than that reported by Pitcher et al. [7], in a similar 8090 alloy, because in the present work the sample was aged at a lower temperature (RT), than in the work of Pitcher et al. [7] (170°C) before a reversion treatment. It should be also pointed out that the reversion process in 8090 alloy does not produce the complete redissolution of 6’ precipitates like in some Weldalite alloys [6] due to the higher concentration of liihium in the 8090 alloy. Finally after lo4 s aging, curve c of Figure 7, a mass fraction of 10% has been obtained which is similar to the RT aged sample. Besides this, when the sample is aged during a longer time the peak height increases and the width reduces, which is in agreement with the TEP and VMH interpretations given above. Moreover, it is also in agreement with the results reported in Refs.[4,7]. It should be pointed out that according to the study of Gregson et al. [ 181, on a low Cu and Mg concentration alloy similar to that employed in the present work, a very high temperature of solid solution treatment (-58O’C) is required in order to obtain the critical vacancy concentration needed to promote homogeneous precipitation of the S’ phase. Consequently, for the relatively low solid solution temperature used (-53O’C) in the present work, no homogeneous coprecipitation of S’ phase is expected below 10’ s. For longer aging times at 200°C the 6 phase appear as was shown in Figure 3 and 4. Nevertheless, some small quantity of other phases such as S’ and T, could be formed simultaneously to the 6 precipitation process (second stage on curve a of Figure 3) but they should be irrelevant because they are not detected by X-ray Diffraction (Fig. 4). It should be highlighted that, the microscopic configuration of 6’ precipitates is very different in both states even if the precipitated mass fraction is almost the same. The microstructural state of the RT aged sample is composed by a very fine distribution of 6’ precipitates as shown by the large width of the (l,O,O) peak (curve a in Figure 7). In contrast, after reversion and aging at 200°C the precipitates size increases associated with a reduction in their number [ 10,211. This coarsening is evidenced by the sharpening of the (1 ,O,O)reflection in Figure 7c. This microstructure evolution allows to explain the hardness evolution along the aging process, including the decrease of hardness during the reversion process. In fact this hardness decreasing is related to the reduction of the precipitated mass fraction but holding the very small size of precipitates, as can be expected from curve b of Figure 7. Conclusions The transition between the RT microstructures and the high temperature microstructure has been followed by TEP, X-ray and VMH techniques through the reversion of the 6’ formed during natural aging. The kinetics of the 6’ precipitation at RT and the kinetics of the 8 reversion at 120°C and 200°C of the unstable microstructure obtained after RT aging have been carefully determined. Acknowledmuents This work has been carried out with the financial support of the Spanish Ciencia y Tecnologla” (CICYT) in the framework of the “Plan National ber MAT 89-0554~CO2-02). O.A.L. wishes to express his appreciation fellow and to UPV/AECI for allowing his stay during research. Besides Pechiney Company for supplying the alloy. “Comision Interministerial de de Materiales” (Project numto CONICET by an external the authors wish also to thank Vol. 37, No. 6 STUDY OF THE 6’ REVERSION PROCESS 859 References 1. Aluminium-Litmum alloys, Aluminium and Aluminium alloys. ASM Speciality handbook. ASM International. The Materials Information Society. Ed. by J.R. Davis, Davis & Associates, 1993. 2. P.J. Gregson, K. Dinsdale, J. Harris and B. Noble, J. Mater. Sci. Technol., 3,7 (1987). 3. J.R. Pickens, F.H. 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