- Email: [email protected]
Low temperature superplasticity in 8090 AlLi alloys
Scripta METALLURGICA et MATERIALIA
Vol. 28, pp. 1125-1130, 1993 Printed in the U.S.A.
Pergamon Press Ltd. All rights reserved
LOW TEMPERATURE SUPERPLASTICITY IN 8090 A1-Li ALLOYS H-P Pu and J. C. H u a n g Institute of Materials Science and Engineering National Sun Yat-Sen University Kaohsiung, Taiwan, R.O.C. (Received January ii, 1993) (Revised February 23, 1993)
Introduction There have been a number of aluminum alloys which , after being subjected to appropriate thermomechanical treatments (TMTs), exhibit high strain-rate sensitive superplasticity while deformed at a temperature near 0.9T m [1]. Examples include the AI-Cu series (e.g. Supral 100 and 150), the AI-Mg series (e.g. 5083 A1), the A1-Zn-Mg series (e.g. 7075 and 7475 A1), and the A1-Li series (e.g. 2090 and 8090 A1). A m o n g these alloys, superplastic aluminum-lithium alloys are still in the early developing stage. Aluminum-lithium alloys have been recognized as attractive candidates for aircraft structure materials. The outstanding superplasticity behavior makes them even more competitive. Current research [2] has been directed toward exploring the superplastic forming (SPF) and diffusion bonding (DB) capability, so as to develop a complete SPF/DB process for engineering applications. The current optimum superplastic temperature for A1-Li alloys usually needs to be within the range of 510-530 oC [3,4]. Since the grain size in A1-Li base alloys is mainly stabilized by the small A13Zr particles, which usually do not act as highly effective grain growth inhibitors, at least not as effective as is the case in the the two-phase (x/[~ Ti base alloys, the grain growth problem cannot be prevented at such high temperatures. Meanwhile, a Li-depletion zone (and also Mg-depletion in the A1-Li-Cu-Mg base alloys)has been observed [5] near the surface of the SPF thin sheet and was considered to be detrimental to post-form properties. Thus, the development of lower-temperature superplastic A1-Li base alloys becomes interesting and meaningful. Research along this line has been reported in a number of papers by McNelley's group [6-10] for A1-Mg and A1-Mg-Li alloys. Although the result m a y not be of practical impact due to the much higher stress involved, it might be of scientific interest in understanding the superplasticity behavior. The current study was intended to develop a series of TMT processes which will allow the 8090 A1-Li base alloy to have superplastic capability (a) at temperatures below 450 oC so as to suppress to a certain degree of the extensive grain growth and Li-loss; (b) at temperatures above 300 °C so that the flow stress will not be too high to induce intensive cavitation; and(c) in excess of 400% elongation, thus reaching the level for certain commercial applications, if possible. Experimental The Alcan 8090 aluminum-lithium alloys in the form of thick plates (50 mm) were selected for TMT processes in this study. The as-received plates were in the T73 condition and did not possess any superplasticity. The composition was A1-2.4wt%Li-1.15%Cu-0.67%Mg-0.11%Zr-0.05%Fe. The trial TMT processes for each route varied in a fewparameters, such as the cold work amount before aging, overaging temperature and time, rolling reduction for each pass, reheat time between rolling passes, and the apphcation of cross-rolling or not. Rolling was performed mainly in the direction parallel to the rolling direction of the as-received plate. The as-received 8090-T73 thick plates were tirst hot rolled at 520°C to -15 m m thickness, followed by solution treatment and 0, 5 or 10% cold rolling. The overaging treatment was done at 350-400 °C for 4-20 h. Warm rolling was performed at 260-300 °C to 85-90% reduction. The controlling factor appeared to be the stored energy before tensile loading. No static recrystallization heat treatment was carried out before loading. Table -
- -
1125 0956-716X/93 $6.00 + .00 Copyright (c) 1993 Pergamon Press Ltd.
"
1
1126
LOW T SUPERPLASTICITY
IN AI-Li
Vol.
28, No.
summarizes the TMT routes listed in decreasing sequence of the stored energy. For comparison, a commercial 8090 SPF sheet (1.6 mm), purchased from Superform company and possessing normal high temperature superplasticity, was also evaluated. TABLE 1 Summary of the TMT Routes Tried Route Reduction per pass Reheat time Cross rollin[~
A high short no
B high medium no
C low medium no
D high medium yes
E high long no
F high long yes
Tensile samples were m a c h i n e d from the as-rolled sheets with tensile axes parallel to the rolling direction; gauge length was set to be 10 m m throughout the study. Tensile testing was conducted mainly at constant cross-head speed (except for m-value measurement). Initial strain rates in the range 10-4 to 10-3 s "1 were used. Test temperatures ranged from 300 to 525 °C. To d e t e r m i n e the strain rate sensitivity, m, sti'ain rate jump tests were c o n d u c t e d . Optical metallography and transmission electron microscopy (TEM) were used to back up the results. Thin foil examinatxon was performed on a JEOL 200CX STEM. Results and Discussion The strain rate sensitivity, m-value, was first evaluated over the strain rates from 8x10 -5 to 8x10 -j s -~ and temperatures from 350 to 450 oC. The results are plotted in Fig. 1. The m-value was determined to be within 0.3-0.4 over the strain rates of 2x10 -4 to 2x10 -3 s -1, and it reached 0.5 or 0.2 at lower or hiRher rates. It is evident that the m-value tends to increase with decreasin~ strain rate, suggesting the o p t i m u m rate to be 8x10-5 s-1 or less. This result is not surprising since tile lower test temperature w i H b e compensated b~ allower strain rate (thus correspondfng to a lower flow stress). ~mce strain rates lower than 2x10- s- appear to be impractical, the majority of current tests were thereby conducted at 8x10 -4 s -1. The uniaxial tensile results at 350, 400 and 450 °C and an initial strain rate of 8x10 -4 s -1 are resented in Table 2. No hydrostatic back pressure was imposed during loading. The flow stress ig. 2) was seen to be significantly higher than that of conventional superplastic a l u m i n u m alloys, which will be unfavorable for SPF practices. In comparison, the commercial 8090 SPF sheet (1.6 mm) purchased from Superform c o m p a n y was also tested u n d e r the same condition. This purchased material, which exhibited superplastic elongation over 500% at temperatures a r o u n d 525 °C, did not show similar superplastic behavior at 350 oC. TABLE 2 The Uniaxial Tensile Test Results Performed at an Initial Strain Rate of 8x 104 s1. The Test Results Typically Showed a Scatter Band of +50% Route
A
300 oC 350 oc 400 oc 450 oc 510 - 525 oc
. . 270% 350% 710% 50%
B .
. . 280% 410% 365% 120%
C . . . 270% 360% 505% 100%
D
E
230% 290% 320% 120%
130% 645% 552% 550% 90%
F
Purchased SPF 8090 AI . . . . 170% 120% 190% -200% -110% 550%
It can be seen from Table 2 that the elongation data at 350 °C are in the neighborhood of 250650%. Cross rolling d i d not seem to improve the superplastic properties, b u t actually degraded slightly the elongation while using the same total rolling reductxon ratio. As the test temperature
9
Vol.
28, No.
9
LOW T S U P E R P L A S T I C I T Y
IN AI-Li
1127
increased to 400 oC, flow stresses became lower and the elongation was within the range of 350550%. At 450 oC, the stress level d r o p p e d d o w n to 15 MPa o r l e s s and the superplastic elongation changed from 310% of path A to 710% of Path E. Figure 3 shows the tested specimens processed via path A a n d E. It can be seen that the deformation at 350-450 oC was sufficiently uniform. The variation of elongation as a function of test temperature can be seen in Fig. 4. Based on the data in Table 2, it seems a longer reheat time will render the rolled structure to better-defined subgrains, which m a y be transferred into highly-misoriented grains u p o n loading throu~gh recovery and partial recrystallization. This effect was seen to be especially a p p a r e n t for samples tested at 350 oC, where the test temperature itself was not high enough to enhance this transformation during the initial stage of loading. Therefore the subgrain structure before loading is critical. As the test t e m p e r a t u r e was raised to 450 oC, the larger driving force at this higher test t e m p e r a t u r e m a d e the grain e v o l u t i o n effect d u r i n g reheat time between rolling passes less pronounced, leading to a more similar behavior at 450 °C. The superplasticity behavior of the route E material at 350 oC was studied in more detail, and the resulting elongations are listed in Table 3. The d e p e n d e n c e of uniaxial superplastic elongation as a function of strain rate is shown in Fig. 5. The o p t i m u m initial strain rate was 8x10 -4 s -1 and the highest elongation achieved was 645% at this temperature, well above the set goal. Although the d e t e r m i n e d m-values s u g g e s t e d that a lower rate m i g h t correspond to a higher elon~;ation, the result over the strain rates 2x10 -4 to 2x10 -3 s -1 did not reveal the this trend. This is postulated to be partly d u e to the small variation of m-value over this strain rate region (-0.35 in Fig. 1). In a d d i t i o n , the p r e v e n t i o n from necking can also be i m p a r t e d by the contribution from w o r k hardening (n-value). The lower initial n-value for tests at 2x10 -4 s -1, compared with that at 8x10 -4 s1, might result in a lower elongation. Minor effects from grain growth should also be accounted for, thus making the reasoning rather complicated. TABLE 3 The Superplastic Behavior of the Route E Mateiial Tested at 350 °C Initial ¢ (s-~)
2x10 -4
Second-step e (s-~)
. . 300+25%
Elongation
8x10 -4 .
. . . 600-L-_50%
2x10 -3
250+25%
2x10 -3 (for 2ram) 2x10 -4 360%*
5x10 -3 (for 2mm) 8x10 -4 460%*
* Single Test Since two-step straining has been widely recognized to be beneficial to superplastic a l u m i n u m alloys m a d e via d y n a m i c recrystallization, the specimens were also tested using an initial higher rate (2 or 5x10 -~ s -1) to 2 m m of tensile displacement followed by a loading rate at 10 -4 ranges. The result did not suggest any improvement. Tfiis might be because the flow stress at 350 o c a n ~ 10 -3 s -1 has been considerably high (-40 MPa); thereby micro-cavitation tended to develop d u r i n g the early stage, resulting in p r e m a t u r e fracture. Due to the fact the stored energy before loading was higher than normal, the samples did not s h o w s u p e r p l a s t i c i t y at conventional s u p e r p l a s t i c t e m p e r a t u r e , i.e, 500 oC or above. At such temperatures the T2 phase (A16Li3Cu) will be completely dissolved and cannot act in the role as A13Zr to suppress grain growth. It follows that the grains in such a high strained specimen will g r o w substantially and thus the superplasticity will not be retained. It was seen that if the specimen was first loaded at 350-400 oC for a short period of time followed by continuous loading at'525 oC, the superplastic property was recovered. The sub~;rain size before superplastic loading seen by TEM was of the order of 0.5 ~tm. After s u p e r p l a s t i c l o a d i n g at 8x10 -4 s -1 and 350 o c to 500% elongation, the grain Rrew to -1 ~tm (350 oC) to 5 ~tm ~450 oC, in Fig. 6), as compared with -18 ~tm of the purchased 8090 SP~ sheet tested at 8x10 -4 s -1 and 525 oC. The cavitation can also be seen in Fig. 6 to be much better i m p r o v e d when tested at low temperatures. Meanwhile, the width of Li- a n d / o r Mg-depletion layer was reduced from -200 lain at 525 °C to 10 ~tm at 350 °C or 45 ~tm at 450 °C. In view of the i m p r o v e d properties of grain size, cavitation and solute depletion, the post-form mechanical p r o p e r h e s are expected to be i m p r o v e d as well.
1128
LOW T SUPERPLASTICITY
IN AI-Li
Vol.
28, No. 9
..It has long. been recognized .that the . dislocation . . substructure . . formed during superplastic. straining at high temperatures is dlfhcu|t to retain for TEM examination. The dislocation morphology at l~igh temperatures will be relaxed or even annealed during slow furnace cooling to room temperature. The current loading temperature is lowered down to 350 °C and the specimen can be . rapidly . cooled . to. below 200 OCby simply openin the furnace. .and blowng cold air to the specimen gauge section. It is considered that more o~ the d~slocatlon arrangement could be maintained. This might be one of the other benefits from studying low temperature superplasticity. Figure 7 shows the examples of the TEM micrographs showing various dislocation morphologies near grain boundaries. The long and well-arranged dislocation arrays seem to be associatedwith the accommodation mechanism for grain boundary sliding and migration. In most cases, only one set of dislocations were seen, suggesting that these dislocations were mostly originated from one single process of accommodation. Meanwhile, the large T2 particles were not observed to be a major obstacle for the superplastic flow at 350-450 °C. Systematic microstructural characterization on grain boundary sliding anddislocation accommodation will be continuously examined in the future. Summary The normal forming temperature for aircraft-used superplastic aluminum lithium base alloys usually needs to be withl~n the range of 510-530 oC. Rapid recrystailization, grain growth, and soluteloss from surface layers cannot be prevented from at this high temperature. Thus, the development for low temperature superplastic A1-Li base alloys becomes meaningful. Through a series of thermomechanical treatments, the 8090 AI-Li base alloy becames superplastic at temperatures from 350 to 450 oC. The TMT-processed sheets exhibit uniaxlal tensile elongations over 600% while tested at 350 °C and 8x10 -4 s -1, without adding back pressure. The m-value within the test conditions usually ranged from 0.3 to 0.4, and the flow stress was much higher than normal cases. The grain growth, cavitation and solute depletion effects were seen to be improved greatly. Finally, the microstructure of the superplastically deformed specimens was interesting and is worth detailed characterization in order to explain the superplastic mechanism. Acknowledgements The partial sponsorship by the China Steels Corp and Aero-Industry Development Center in Taiwan, R.O.C. is gratefully acknowledged. References 1. 2.
J. Wadsworth, A. R. Pelton and R. E. Lewis, Metall. Trans., 16A, 2319 (1985). P.J. Winkler, T. Heinrich, R. Keyte, G. J. Mahon and R. A. Ricks, Aluminum-Lithium VI, DGM, Germany, 1069 (1992). 3. R. Grimes, W. S. Miller and R, G. Bulter, J. de Physeque, C3, 239 (1987). 4. N. Ridley, D. W. Livesey and J. Pilling, J. de Physeque, C3, 251 (1987). 5. J.M. Papazian, G. G. Bott and P. Shaw, Mater. Sci. Eng, 94, 219 (1987). 6. T.R. McNellev, E. -W. Lee and M. E. Mills, Metall. Trans., 17A, 1035 (1986). 7. E.-W. Lee, T. R. McNelley and A. F. Stengel, Metall. Trans., 17A, 1043 (1986). 8. E.-W. Lee and T. R. McNelley, Mater. Sci. Eng., 93, 45 (1987). 9. S.J. Hales and T. R. McNelley, Acta Metall., 36, 1229 (1988). 10. S.J. Hales, T. R. McNelley and H. J. McQueen, Metall. Trans., 22A, 1037 (1991).
Vol.
28, No.
9
LOW T SUPERPLASTICITY
IN AI-Li
1129
0.8
100
[] • o • •
0.7 0.6 0.5
Route A at 450 Route C at 450 Route C at 400 Route C at 350 RouteEat 350
[]
!
0.4
1o
•
0.3 I ~ .
8 ........
1
Route Route Route Route ! 10 -3
........
i
10 -4
10 -5
A at 450 C at 450 C at 400 C at 350 .......
"C "C "C "C
b
0.0 10 -5
10 -2
10-4
10 -3
10 -2
Strain rate (s-l)
Illustration of the variation of (a) flow stress and (b) strain rate sensitivity (m) as a function of strain rate when tested at 350-450 °C.
100
90
8
'
400 "C 350 *C 300 "C
7O
b
2,1o
s1
8x 10--4 s-I 2x10-3 s-1
80 70
~" 60
60
~_. 50
50
~4o
40
o~ 30
,a
20
30 20
10
1
O! 0.0
FIG. 2
I
i
Slxain rate (s-l)
FIG. 1
•
|
0.2 0.1
'C "C *C *C °C
i
i
0.5
1.0 Strain
•
i
1.5
•
2.0
0.0
0.5
1.0 Strain
1.5
2.0
True stress a n d strain curves for sheets processed by route E tested at (a) 8x10 -4 s -1 and various temperatures, and (b) 350 °C and various initial strain rates.
800 • [ •
R ~ m _ .
600
Route A Route C
_ outeE
°= 400 t~ 200
0 250
FIG. 3
Tensile specimens of route A and E after loading. 1,2 and 3 represent 350, 400 and 450 oC respectively.
•
FIG. 4
i
300
•
i
350
•
i
.
i
•
i
400 450 500 Temperature (°C)
•
i
550
.
600
The v a r i a t i o n of s u p e r p l a s t i c elongation as a f u n c t i o n of test temperature at an initial strain rate of 8x10 -4 s -1.
i130
LOW T SUPERPLASTICITY
800
IN AI-Li
Vol.
28, No.
] [] Route E at 350 "C
/L t
7110 600
3110 200 1110 ~ 10-4
FIG. 5
FIG. 6
10 -3 Strain rate (s-l)
10 -2
The superplastic elongation variation of sheets processed by route E as a function of test strain rate at 350 °C.
Optical micrographs of the sheets after loading to -500% elongation: (a) route E material tested at 350 °C, (b) route A material tested at 450 °C, and (c) purchased SPF sheet tested at 525 °C.
FIG. 7 TEM micrographs showing the dislocation activity near grain boundaries in sheets processed by route E tested at 350 OCand 2x10-4 s -1 to 300% elongation.
9
Vol. 28, pp. 1125-1130, 1993 Printed in the U.S.A.
Pergamon Press Ltd. All rights reserved
LOW TEMPERATURE SUPERPLASTICITY IN 8090 A1-Li ALLOYS H-P Pu and J. C. H u a n g Institute of Materials Science and Engineering National Sun Yat-Sen University Kaohsiung, Taiwan, R.O.C. (Received January ii, 1993) (Revised February 23, 1993)
Introduction There have been a number of aluminum alloys which , after being subjected to appropriate thermomechanical treatments (TMTs), exhibit high strain-rate sensitive superplasticity while deformed at a temperature near 0.9T m [1]. Examples include the AI-Cu series (e.g. Supral 100 and 150), the AI-Mg series (e.g. 5083 A1), the A1-Zn-Mg series (e.g. 7075 and 7475 A1), and the A1-Li series (e.g. 2090 and 8090 A1). A m o n g these alloys, superplastic aluminum-lithium alloys are still in the early developing stage. Aluminum-lithium alloys have been recognized as attractive candidates for aircraft structure materials. The outstanding superplasticity behavior makes them even more competitive. Current research [2] has been directed toward exploring the superplastic forming (SPF) and diffusion bonding (DB) capability, so as to develop a complete SPF/DB process for engineering applications. The current optimum superplastic temperature for A1-Li alloys usually needs to be within the range of 510-530 oC [3,4]. Since the grain size in A1-Li base alloys is mainly stabilized by the small A13Zr particles, which usually do not act as highly effective grain growth inhibitors, at least not as effective as is the case in the the two-phase (x/[~ Ti base alloys, the grain growth problem cannot be prevented at such high temperatures. Meanwhile, a Li-depletion zone (and also Mg-depletion in the A1-Li-Cu-Mg base alloys)has been observed [5] near the surface of the SPF thin sheet and was considered to be detrimental to post-form properties. Thus, the development of lower-temperature superplastic A1-Li base alloys becomes interesting and meaningful. Research along this line has been reported in a number of papers by McNelley's group [6-10] for A1-Mg and A1-Mg-Li alloys. Although the result m a y not be of practical impact due to the much higher stress involved, it might be of scientific interest in understanding the superplasticity behavior. The current study was intended to develop a series of TMT processes which will allow the 8090 A1-Li base alloy to have superplastic capability (a) at temperatures below 450 oC so as to suppress to a certain degree of the extensive grain growth and Li-loss; (b) at temperatures above 300 °C so that the flow stress will not be too high to induce intensive cavitation; and(c) in excess of 400% elongation, thus reaching the level for certain commercial applications, if possible. Experimental The Alcan 8090 aluminum-lithium alloys in the form of thick plates (50 mm) were selected for TMT processes in this study. The as-received plates were in the T73 condition and did not possess any superplasticity. The composition was A1-2.4wt%Li-1.15%Cu-0.67%Mg-0.11%Zr-0.05%Fe. The trial TMT processes for each route varied in a fewparameters, such as the cold work amount before aging, overaging temperature and time, rolling reduction for each pass, reheat time between rolling passes, and the apphcation of cross-rolling or not. Rolling was performed mainly in the direction parallel to the rolling direction of the as-received plate. The as-received 8090-T73 thick plates were tirst hot rolled at 520°C to -15 m m thickness, followed by solution treatment and 0, 5 or 10% cold rolling. The overaging treatment was done at 350-400 °C for 4-20 h. Warm rolling was performed at 260-300 °C to 85-90% reduction. The controlling factor appeared to be the stored energy before tensile loading. No static recrystallization heat treatment was carried out before loading. Table -
- -
1125 0956-716X/93 $6.00 + .00 Copyright (c) 1993 Pergamon Press Ltd.
"
1
1126
LOW T SUPERPLASTICITY
IN AI-Li
Vol.
28, No.
summarizes the TMT routes listed in decreasing sequence of the stored energy. For comparison, a commercial 8090 SPF sheet (1.6 mm), purchased from Superform company and possessing normal high temperature superplasticity, was also evaluated. TABLE 1 Summary of the TMT Routes Tried Route Reduction per pass Reheat time Cross rollin[~
A high short no
B high medium no
C low medium no
D high medium yes
E high long no
F high long yes
Tensile samples were m a c h i n e d from the as-rolled sheets with tensile axes parallel to the rolling direction; gauge length was set to be 10 m m throughout the study. Tensile testing was conducted mainly at constant cross-head speed (except for m-value measurement). Initial strain rates in the range 10-4 to 10-3 s "1 were used. Test temperatures ranged from 300 to 525 °C. To d e t e r m i n e the strain rate sensitivity, m, sti'ain rate jump tests were c o n d u c t e d . Optical metallography and transmission electron microscopy (TEM) were used to back up the results. Thin foil examinatxon was performed on a JEOL 200CX STEM. Results and Discussion The strain rate sensitivity, m-value, was first evaluated over the strain rates from 8x10 -5 to 8x10 -j s -~ and temperatures from 350 to 450 oC. The results are plotted in Fig. 1. The m-value was determined to be within 0.3-0.4 over the strain rates of 2x10 -4 to 2x10 -3 s -1, and it reached 0.5 or 0.2 at lower or hiRher rates. It is evident that the m-value tends to increase with decreasin~ strain rate, suggesting the o p t i m u m rate to be 8x10-5 s-1 or less. This result is not surprising since tile lower test temperature w i H b e compensated b~ allower strain rate (thus correspondfng to a lower flow stress). ~mce strain rates lower than 2x10- s- appear to be impractical, the majority of current tests were thereby conducted at 8x10 -4 s -1. The uniaxial tensile results at 350, 400 and 450 °C and an initial strain rate of 8x10 -4 s -1 are resented in Table 2. No hydrostatic back pressure was imposed during loading. The flow stress ig. 2) was seen to be significantly higher than that of conventional superplastic a l u m i n u m alloys, which will be unfavorable for SPF practices. In comparison, the commercial 8090 SPF sheet (1.6 mm) purchased from Superform c o m p a n y was also tested u n d e r the same condition. This purchased material, which exhibited superplastic elongation over 500% at temperatures a r o u n d 525 °C, did not show similar superplastic behavior at 350 oC. TABLE 2 The Uniaxial Tensile Test Results Performed at an Initial Strain Rate of 8x 104 s1. The Test Results Typically Showed a Scatter Band of +50% Route
A
300 oC 350 oc 400 oc 450 oc 510 - 525 oc
. . 270% 350% 710% 50%
B .
. . 280% 410% 365% 120%
C . . . 270% 360% 505% 100%
D
E
230% 290% 320% 120%
130% 645% 552% 550% 90%
F
Purchased SPF 8090 AI . . . . 170% 120% 190% -200% -110% 550%
It can be seen from Table 2 that the elongation data at 350 °C are in the neighborhood of 250650%. Cross rolling d i d not seem to improve the superplastic properties, b u t actually degraded slightly the elongation while using the same total rolling reductxon ratio. As the test temperature
9
Vol.
28, No.
9
LOW T S U P E R P L A S T I C I T Y
IN AI-Li
1127
increased to 400 oC, flow stresses became lower and the elongation was within the range of 350550%. At 450 oC, the stress level d r o p p e d d o w n to 15 MPa o r l e s s and the superplastic elongation changed from 310% of path A to 710% of Path E. Figure 3 shows the tested specimens processed via path A a n d E. It can be seen that the deformation at 350-450 oC was sufficiently uniform. The variation of elongation as a function of test temperature can be seen in Fig. 4. Based on the data in Table 2, it seems a longer reheat time will render the rolled structure to better-defined subgrains, which m a y be transferred into highly-misoriented grains u p o n loading throu~gh recovery and partial recrystallization. This effect was seen to be especially a p p a r e n t for samples tested at 350 oC, where the test temperature itself was not high enough to enhance this transformation during the initial stage of loading. Therefore the subgrain structure before loading is critical. As the test t e m p e r a t u r e was raised to 450 oC, the larger driving force at this higher test t e m p e r a t u r e m a d e the grain e v o l u t i o n effect d u r i n g reheat time between rolling passes less pronounced, leading to a more similar behavior at 450 °C. The superplasticity behavior of the route E material at 350 oC was studied in more detail, and the resulting elongations are listed in Table 3. The d e p e n d e n c e of uniaxial superplastic elongation as a function of strain rate is shown in Fig. 5. The o p t i m u m initial strain rate was 8x10 -4 s -1 and the highest elongation achieved was 645% at this temperature, well above the set goal. Although the d e t e r m i n e d m-values s u g g e s t e d that a lower rate m i g h t correspond to a higher elon~;ation, the result over the strain rates 2x10 -4 to 2x10 -3 s -1 did not reveal the this trend. This is postulated to be partly d u e to the small variation of m-value over this strain rate region (-0.35 in Fig. 1). In a d d i t i o n , the p r e v e n t i o n from necking can also be i m p a r t e d by the contribution from w o r k hardening (n-value). The lower initial n-value for tests at 2x10 -4 s -1, compared with that at 8x10 -4 s1, might result in a lower elongation. Minor effects from grain growth should also be accounted for, thus making the reasoning rather complicated. TABLE 3 The Superplastic Behavior of the Route E Mateiial Tested at 350 °C Initial ¢ (s-~)
2x10 -4
Second-step e (s-~)
. . 300+25%
Elongation
8x10 -4 .
. . . 600-L-_50%
2x10 -3
250+25%
2x10 -3 (for 2ram) 2x10 -4 360%*
5x10 -3 (for 2mm) 8x10 -4 460%*
* Single Test Since two-step straining has been widely recognized to be beneficial to superplastic a l u m i n u m alloys m a d e via d y n a m i c recrystallization, the specimens were also tested using an initial higher rate (2 or 5x10 -~ s -1) to 2 m m of tensile displacement followed by a loading rate at 10 -4 ranges. The result did not suggest any improvement. Tfiis might be because the flow stress at 350 o c a n ~ 10 -3 s -1 has been considerably high (-40 MPa); thereby micro-cavitation tended to develop d u r i n g the early stage, resulting in p r e m a t u r e fracture. Due to the fact the stored energy before loading was higher than normal, the samples did not s h o w s u p e r p l a s t i c i t y at conventional s u p e r p l a s t i c t e m p e r a t u r e , i.e, 500 oC or above. At such temperatures the T2 phase (A16Li3Cu) will be completely dissolved and cannot act in the role as A13Zr to suppress grain growth. It follows that the grains in such a high strained specimen will g r o w substantially and thus the superplasticity will not be retained. It was seen that if the specimen was first loaded at 350-400 oC for a short period of time followed by continuous loading at'525 oC, the superplastic property was recovered. The sub~;rain size before superplastic loading seen by TEM was of the order of 0.5 ~tm. After s u p e r p l a s t i c l o a d i n g at 8x10 -4 s -1 and 350 o c to 500% elongation, the grain Rrew to -1 ~tm (350 oC) to 5 ~tm ~450 oC, in Fig. 6), as compared with -18 ~tm of the purchased 8090 SP~ sheet tested at 8x10 -4 s -1 and 525 oC. The cavitation can also be seen in Fig. 6 to be much better i m p r o v e d when tested at low temperatures. Meanwhile, the width of Li- a n d / o r Mg-depletion layer was reduced from -200 lain at 525 °C to 10 ~tm at 350 °C or 45 ~tm at 450 °C. In view of the i m p r o v e d properties of grain size, cavitation and solute depletion, the post-form mechanical p r o p e r h e s are expected to be i m p r o v e d as well.
1128
LOW T SUPERPLASTICITY
IN AI-Li
Vol.
28, No. 9
..It has long. been recognized .that the . dislocation . . substructure . . formed during superplastic. straining at high temperatures is dlfhcu|t to retain for TEM examination. The dislocation morphology at l~igh temperatures will be relaxed or even annealed during slow furnace cooling to room temperature. The current loading temperature is lowered down to 350 °C and the specimen can be . rapidly . cooled . to. below 200 OCby simply openin the furnace. .and blowng cold air to the specimen gauge section. It is considered that more o~ the d~slocatlon arrangement could be maintained. This might be one of the other benefits from studying low temperature superplasticity. Figure 7 shows the examples of the TEM micrographs showing various dislocation morphologies near grain boundaries. The long and well-arranged dislocation arrays seem to be associatedwith the accommodation mechanism for grain boundary sliding and migration. In most cases, only one set of dislocations were seen, suggesting that these dislocations were mostly originated from one single process of accommodation. Meanwhile, the large T2 particles were not observed to be a major obstacle for the superplastic flow at 350-450 °C. Systematic microstructural characterization on grain boundary sliding anddislocation accommodation will be continuously examined in the future. Summary The normal forming temperature for aircraft-used superplastic aluminum lithium base alloys usually needs to be withl~n the range of 510-530 oC. Rapid recrystailization, grain growth, and soluteloss from surface layers cannot be prevented from at this high temperature. Thus, the development for low temperature superplastic A1-Li base alloys becomes meaningful. Through a series of thermomechanical treatments, the 8090 AI-Li base alloy becames superplastic at temperatures from 350 to 450 oC. The TMT-processed sheets exhibit uniaxlal tensile elongations over 600% while tested at 350 °C and 8x10 -4 s -1, without adding back pressure. The m-value within the test conditions usually ranged from 0.3 to 0.4, and the flow stress was much higher than normal cases. The grain growth, cavitation and solute depletion effects were seen to be improved greatly. Finally, the microstructure of the superplastically deformed specimens was interesting and is worth detailed characterization in order to explain the superplastic mechanism. Acknowledgements The partial sponsorship by the China Steels Corp and Aero-Industry Development Center in Taiwan, R.O.C. is gratefully acknowledged. References 1. 2.
J. Wadsworth, A. R. Pelton and R. E. Lewis, Metall. Trans., 16A, 2319 (1985). P.J. Winkler, T. Heinrich, R. Keyte, G. J. Mahon and R. A. Ricks, Aluminum-Lithium VI, DGM, Germany, 1069 (1992). 3. R. Grimes, W. S. Miller and R, G. Bulter, J. de Physeque, C3, 239 (1987). 4. N. Ridley, D. W. Livesey and J. Pilling, J. de Physeque, C3, 251 (1987). 5. J.M. Papazian, G. G. Bott and P. Shaw, Mater. Sci. Eng, 94, 219 (1987). 6. T.R. McNellev, E. -W. Lee and M. E. Mills, Metall. Trans., 17A, 1035 (1986). 7. E.-W. Lee, T. R. McNelley and A. F. Stengel, Metall. Trans., 17A, 1043 (1986). 8. E.-W. Lee and T. R. McNelley, Mater. Sci. Eng., 93, 45 (1987). 9. S.J. Hales and T. R. McNelley, Acta Metall., 36, 1229 (1988). 10. S.J. Hales, T. R. McNelley and H. J. McQueen, Metall. Trans., 22A, 1037 (1991).
Vol.
28, No.
9
LOW T SUPERPLASTICITY
IN AI-Li
1129
0.8
100
[] • o • •
0.7 0.6 0.5
Route A at 450 Route C at 450 Route C at 400 Route C at 350 RouteEat 350
[]
!
0.4
1o
•
0.3 I ~ .
8 ........
1
Route Route Route Route ! 10 -3
........
i
10 -4
10 -5
A at 450 C at 450 C at 400 C at 350 .......
"C "C "C "C
b
0.0 10 -5
10 -2
10-4
10 -3
10 -2
Strain rate (s-l)
Illustration of the variation of (a) flow stress and (b) strain rate sensitivity (m) as a function of strain rate when tested at 350-450 °C.
100
90
8
'
400 "C 350 *C 300 "C
7O
b
2,1o
s1
8x 10--4 s-I 2x10-3 s-1
80 70
~" 60
60
~_. 50
50
~4o
40
o~ 30
,a
20
30 20
10
1
O! 0.0
FIG. 2
I
i
Slxain rate (s-l)
FIG. 1
•
|
0.2 0.1
'C "C *C *C °C
i
i
0.5
1.0 Strain
•
i
1.5
•
2.0
0.0
0.5
1.0 Strain
1.5
2.0
True stress a n d strain curves for sheets processed by route E tested at (a) 8x10 -4 s -1 and various temperatures, and (b) 350 °C and various initial strain rates.
800 • [ •
R ~ m _ .
600
Route A Route C
_ outeE
°= 400 t~ 200
0 250
FIG. 3
Tensile specimens of route A and E after loading. 1,2 and 3 represent 350, 400 and 450 oC respectively.
•
FIG. 4
i
300
•
i
350
•
i
.
i
•
i
400 450 500 Temperature (°C)
•
i
550
.
600
The v a r i a t i o n of s u p e r p l a s t i c elongation as a f u n c t i o n of test temperature at an initial strain rate of 8x10 -4 s -1.
i130
LOW T SUPERPLASTICITY
800
IN AI-Li
Vol.
28, No.
] [] Route E at 350 "C
/L t
7110 600
3110 200 1110 ~ 10-4
FIG. 5
FIG. 6
10 -3 Strain rate (s-l)
10 -2
The superplastic elongation variation of sheets processed by route E as a function of test strain rate at 350 °C.
Optical micrographs of the sheets after loading to -500% elongation: (a) route E material tested at 350 °C, (b) route A material tested at 450 °C, and (c) purchased SPF sheet tested at 525 °C.
FIG. 7 TEM micrographs showing the dislocation activity near grain boundaries in sheets processed by route E tested at 350 OCand 2x10-4 s -1 to 300% elongation.
9