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Severe plastic deformation processing of Al–Cu–Li alloy for enhancing strength while maintaining ductility
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Scripta Materialia 63 (2010) 304–307 www.elsevier.com/locate/scriptamat
Severe plastic deformation processing of Al–Cu–Li alloy for enhancing strength while maintaining ductility M.A. Mun˜oz-Morris and D.G. Morris* Department of Physical Metallurgy, CENIM-CSIC, Avenida Gregorio del Amo 8, 28040 Madrid, Spain Received 1 March 2010; revised 29 March 2010; accepted 14 April 2010 Available online 18 April 2010
An Al–Cu–Li alloy has been processed by equal channel angular pressing (ECAP) to compare the influence of strain rate and processing temperature, together with subsequent annealing treatment, on the strength and ductility of the alloy. High strength with reasonable ductility can be found under conditions where fine precipitates enhance strength but also delay strain localization and fracture. Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Equal channel angular pressing; Aluminium alloy; Mechanical properties
Aluminium–lithium base alloys are of great interest in aerospace construction because they have relatively low density and high strength, while at the same time possessing good stiffness and reasonable ductility [1,2]. There is nevertheless much interest in further improving the strength–ductility combination. Severe plastic deformation processing is a method for microstructural refinement which may lead to improved strength with improved ductility [3]. There have been relatively few studies of ECAP processing of Al–Li alloys (e.g. [4,5]), and these have not significantly enhanced such strength–ductility combinations. One study [4] carried out ECAP at high temperature and obtained some grain refinement but only low strength with good ductility, although subsequent heat treatments improved the strength somewhat. Another study [5] performed ECAP at 200–250 °C, producing very fine grains with high strength but essentially no ductility, while annealing improved ductility but led to a major loss of strength. A previous study by the present authors [6] examined hot rolling of an Al–Cu–Li alloy, with processing above 200 °C achieving some microstructural refinement as extensive precipitation of T1 phase (Al2CuLi) or d0 phase (Al3Li) occurred. The present study examines microstructural evolution and mechanical properties of the same alloy following higher levels of strain imposed by ECAP under a wide variety of processing conditions. The alloy, of composition Al–2.8 wt.% Cu–1.6 wt.% Li, with small additions of Zr, Mg, Zn, Mn and Fe, is designated alloy 2099, and is somewhat similar to AF/C-458 material * Corresponding author. Tel.: +34 91 533 8900x336; fax: +34 91 533 8016; e-mail: [email protected]
examined elsewhere [7–9]. The extensive formation of T1 phase during high-temperature exposure has been well documented [10,11], together with formation of other phases such as h0 (Al2Cu), with the d0 phase expected only at relatively low temperatures [10–12]. It is also well established [7–9] that straining prior to high-temperature exposure can significantly accelerate precipitation. The combination of high-temperature exposure with severe deformation thus offers a wide potential for microstructural modification. Samples for processing were given a solutionizing anneal of 1 h at 540 °C and then water quenched. ECAP was carried out from room temperature (RT) to 200 °C in a hydraulic machine using a circular cross-section die of diameter 20 mm with a die angle of 118°, producing a true strain of 0.7 per pass. The so-called route A was used [13], i.e. with no rotation of the sample between passes. The temperatures chosen were low enough such that either no precipitation or only precipitation to fine particle sizes [6] occurred. Significant precipitation hardening is thereby retained during processing or achievable during subsequent annealing. ECAP was carried out using 3–7 passes (strain level 2.1–4.9), restricting the total strain to moderate levels so as to achieve much of the grain refinement and hardening and ductility improvements possible by severe plastic deformation [3,5]. Samples were heated to die temperature in 5 min before pressing. At the standard pressing speed (20 mm min1) the total cycle time (preheating and pressing) was 10 min. Some tests were carried out with faster pressing (60 mm min1) and slower pressing (6 mm min1), leading to cycle times of 5–6 and 15 min, whereby most of the cycle time was spent in preheating or in pressing. The heated split
1359-6462/$ - see front matter Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2010.04.022
M. A. Mun˜oz-Morris, D. G. Morris / Scripta Materialia 63 (2010) 304–307
die could be opened hydraulically for rapid sample removal and water quenching. Tensile samples were machined for testing along the pressing direction. Two samples were tested in each case, confirming excellent data reproducibility. Microstructures were examined by scanning electron microscopy (SEM) in back-scattered electron (BSE) mode, electron back-scatter diffraction (EBSD) and transmission electron microscopy (TEM). Samples were electropolished with 20% nitric acid in methanol in the direction perpendicular to the inlet and exit pressing directions to reveal any elongation of microstructures. Several processing conditions were examined and those achieving significant improvements in stress– ductility combinations are reported here. All samples showed significant work hardening up to the ultimate tensile stress, with failure occurring either immediately thereafter or after prolonged straining under reducing nominal stress, as illustrated in Figure 1a. The strength and elongation values are shown in Table 1. Processing by ECAP for 3 passes (strain 2.1) at RT, 150–200 °C led to moderate values of yield strength with poor to good ductility (see Fig. 1b). Also shown here is the increase in strength and elongation after annealing the RT processed material. There are also important strength increases with little change in elongation when processing at 150 °C for 5 or 7 ECAP passes. Figure 1b also shows reference data for comparison. T6 refers to material solutionized and annealed for 5 h at 200 °C, leading to maximum strength at this temperature [6]. Data are also shown for yield strength and elongation of optimized commercial material following solutionizing, stretching and ageing [14]. Some of the ECAP-processed materials clearly show significant strength improvements over annealed or stretched-annealed materials, with lower but reasonable elongations. Materials always showed a deformation substructure of elongated dislocation cells, with no precipitation after RT processing and fine precipitate plates after sufficient hightemperature exposure. Various microstructures are illustrated in Figures 2 and 3. Figures 2a and 3a show typical elongated cell structures, imaged by EBSD and by TEM. EBSD confirmed that these structures were constituted by cell walls with very low misorientations, typically 70–80% less than 3°. Figure 2a illustrates the evolution of misorientation along a line drawn in the image showing that most 600
Stress (MPa)
5Ep150ºC
3EpRT
400
200
0
0
5
10
Elongation (%)
15
Yield Strength (MPa)
600
3EpRT + 200ºC
+anneal 200ºC
7Ep150ºC
550
A 5Ep150ºC
500
T6
3Ep200ºC A
450 3Ep150ºC 3EpRT
400 0
2
4
6
8
10
Elongation to Failure (%)
Figure 1. Tensile stress–elongation curves illustrating work hardening after yielding and different strain behaviour after ultimate tensile stress, depending on processing conditions. (b) Yield strength and elongation combinations after processing: ECAP at RT, 150 and 200 °C; 3, 5 and 7 passes. Material processed by 3 passes at RT was also annealed for 1/2 h at 200 °C. Data point T6 refers to material aged to peak strength without prior deformation. The two points A refer to Alcoa datasheet values [14].
305
boundaries have misorientations of only a few degrees, with others reaching up to 10°. Dislocation densities and cell dimensions were determined (Table 1) by counting intercepts on fiducial lines of images such as those shown in Figure 3a, covering several cell dimensions. Measured values are subject to significant error, and can be taken simply as indicative values. Dislocation cell width was approximately related to the inverse (or inverse square root) of dislocation density, and cell length approximately proportional to width. In view of such interrelationships, the following discussion will use dislocation density in any analysis of strengthening. High-temperature processing, during ECAP and annealing, led to only a slight recovery of the substructure, which remained as dislocation cells. High-temperature processing also led to precipitation, generally as fine plates of T1 phase, illustrated in Figure 2b (precipitates visible by atomic number contrast due to the high Cu content) and the TEM dark field (DF) images of Figure 3b–d (obtained using 00002Ti spot). The T1 precipitates are known to lie on {1 1 1}Al planes (i.e. {1 1 1}Al // {0 0 0 1}T1, with h1 1 0iAl // h1 0 1 0i) and appear as thin, circular plates [6,10,11]. The precipitate disc size can be measured from SEM images (Fig. 2b) when they are thick, or from TEM DF images (Fig. 3c and d). The disc thickness was often too small (of the order of 1–2 nm) to be determined accurately from such images, so precipitate thickness was determined instead from the extent of streaking (peak width at half intensity) of precipitate diffraction spots. The spot selected was the 0002T1 spot (found at 1/2 1/2 1/2Al; see Fig. 3c), obtained when the precipitates are oriented edgeon. Precipitate diameters and thicknesses, as well as precipitate densities, determined from TEM images are reported in Table 1 for all states where precipitation was detected, with the volume fraction of the precipitate phase determined from these two parameters also shown. The determination of particle density was especially difficult for very fine particle sizes, and these values are subject to considerable error. The precipitate thickness was apparently mostly determined by exposure time at high temperature while the precipitate disc diameter was greatly reduced by straining before or during heat treatment. Figure 3b shows very thin and large diameter precipitates in material annealed without prior strain (taken from Ref. [6]), of comparable thickness but much larger than similar annealing given after heavy straining (Fig. 3d). Particle number densities are much higher than expected in the absence of prior strain or when given small stretches before annealing [7], with particle sizes reduced accordingly and similar volume fractions of particle phase found. Figure 1b and Table 1 confirm a dramatic increase in both yield strength and elongation to failure on annealing material ECAP processed at RT. The ageing temperature (200 °C) was selected based on earlier studies [6] of age hardening of solutionized material where annealing for 2–20 h led to a flat hardness maximum at a level only slightly lower than after prolonged annealing at lower temperatures. Deformation before ageing is known to accelerate hardening [6,8], such that the annealing time for maximum hardening at 200 °C becomes conveniently short (1/2 h). Additional heat treatments (1/4 and 2 h at 200 °C) confirmed very similar strength and elongation values to those of Figure 1b and Table 1. The RT ECAP-processed material has moderate strength but fail-
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Table 1. Processing details, tensile mechanical properties and quantified microstructure. Processing history
Yield Ultimate Elongation at Elongation Particle Particle Particle number Volume Dislocation Dislocation cell density size (width of T1 strength tensile ultimate tensile to failure thickness diameter density phase (%) (m2) length) (lm) (MPa) strength strength (%) (%) (nm) (nm) (m3) 1022 (MPa)
3EpRT 3EpRT + ½ h200 °C 3Ep150 °C 5Ep150 °C 7Ep150 °C 3Ep150 °C slow 3Ep150 °C fast 3Ep 200 °C 5Ep200 °C
434 577
460 581
1.6 0.9
1.6 4.7
None 1.4
None 15
– 6.8
– 1.7
2 1016 2 1016
0.39 2.5 0.35 2
443 505 568 454 410 502 464
476 518 576 475 450 512 482
4.8 1.9 1.3 3.6 4.4 2.7 2.25
7.3 5.2 6.0 5.4 9.2 8.8 7.5
None None 1 <1 None 3.1 3.9
None None 7.5 <10 None 15 18
– – 10 1–10 – 2.5 2
– – 0.4 0.1 – 1.4 2.0
6 1015 1 1016 1.3 1016 5 1015 5 1015 1.5 1015 2 1015
0.5 2.7 0.4 2 0.35 1.5 0.5 3 0.45 2.5 0.65 4 0.5 3
The processing history indicates the number of ECAP passes at a given temperature, at the standard, slow or fast pressing speed, and any subsequent annealing.
ure occurs immediately after the ultimate tensile strength. Following annealing there is a significant increase in both yield strength (from 434 to 577 MPa) and ultimate tensile strength. The elongation to ultimate tensile strength is decreased somewhat, while the elongation to failure is increased from 1.6% to 4.7%. No change in dislocation substructure was detected after annealing (measured dislocation density identical), with presumably Cu and Li solute or precipitates slowing recovery. Strengthening increases due to fine precipitation during annealing (Fig. 3d), and these precipitates also delay strain concen-
Figure 2. Typical microstructures of transverse sections after ECAP. (a) Misorientation map obtained by EBSD: dislocation cell structure after 3 passes at RT, followed by annealing for 1/2 h at 200 °C; (b) SEM–BSE image: coarse T1 precipitates after 5 passes at 200 °C. (a) The point-to-point (in red) and point-to-origin (in blue) misorientation change along the line drawn in the image. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this paper.)
tration and the onset of failure. Such delayed failure is particularly dramatic in comparison with as-ECAPed material, which failed at the ultimate tensile strength. Figures 1b, 4a and Table 1 show that there is significant strengthening with little change of elongation to failure on increasing the ECAP strain at 150 °C. With increasing yield strength, the elongation to ultimate tensile strength decreases, but failure elongation hardly changes. No precipitation is found after 3–5 passes, but many fine precipitates are present after 7 passes (Fig. 3c), with a simultaneous slight increase in dislocation density. The strength increase is associated with the fine precipitates rather than the change in dislocation density. Even though these fine precipitates may be easily sheared, it is clear that they do not encourage strain concentration and early failure. Figure 1b and Table 1 also compare the role of processing temperature (150–200 °C) when giving 3–5 passes.
Figure 3. Microstructures obtained by TEM on transverse sections after ECAP processing. (a) Dislocation cell structure after 3 passes at RT with 1/2 h annealing at 200 °C; (b) large, thin T1 precipitates, with corresponding diffraction pattern, after annealing undeformed material for 1/2 h at 200 °C (taken from Ref. [6]); (c) fine T1 precipitates, with corresponding diffraction pattern, after 7 passes at 150 °C; (d) fine T1 precipitates after 3 passes at RT and annealing for 1/2 h at 200 °C.
M. A. Mun˜oz-Morris, D. G. Morris / Scripta Materialia 63 (2010) 304–307
Higher-temperature processing increases strength (from 443 to 502 MPa) after 3 passes but decreases it (from 505 to 464 MPa) after 5 passes, always with improved elongation to failure. Elongation at ultimate tensile strength seems generally to be inversely related to the strength level. Dislocation density is reduced somewhat after higher-temperature processing but, most importantly, precipitates appear (see Table 1). The reduced dislocation density can lower strength levels, but this is compensated by precipitation hardening when these are very fine. These precipitates appear, however, to be responsible for delaying strain localization and fracture.Figure 4b shows changes of mechanical properties for 3 ECAP passes at 150 °C at various speeds. Increasing speed reduces stress levels but increases ductility. Elongation to failure increases substantially, with elongation to ultimate tensile strength changing little (Table 1). There is, however, no clear change of dislocation substructure and only weak evidence of precipitation after processing slowly. The precipitates in such material were extremely thin discs, with diffraction spots too weak to allow reliable confirmation of structure (T1 precipitates show characteristic strong 1 0 1 0T1 diffraction, found at positions corresponding to 4/3 4/3 0Al [6,10,11], which were not detected). The precipitates here may be T1 phase or perhaps d0 phase [6,10,11]. Such incipient precipitation produces higher strength and smaller elongation to failure. The lower strength level and increased elongation when processing quickly remain difficult to understand, since no difference in dislocation structure or precipitation is found compared with the reference processing speed. Incipient precipitation possibly begins (undetected) after processing at the reference speed, which is avoided after fast processing. The corresponding lower yield strength then allows more homogeneous strain than whenever fine/incipient precipitation occurs. The present experiments have shown several ways to improve strength while retaining ductility of Al–Cu–Li. The role of dislocations and precipitates can be understood through examination of the microstructural data in Table 1: (i) after deforming to work harden, suitable annealing ensures fine and homogeneous precipitation of T1 phase, with the material strong and resistant to failure. (ii) Increasing the amount of strain at 150 °C leads to increased dislocation hardening, but it is the onset of precipitation that causes significant strengthening, while good ductility is retained. (iii) The correct ECAP temperature controls dislocation density and precipitate size. Thicker precipitates lower the strength but can further delay fracture. (iv) Controlling the speed of ECAP processing at 150 °C 9 8
Yield Strength
500
7
450
Elongation
6 5
400
ECAP 150ºC
2
3
4
5
6
Nº ECAP Passes
7
8
4
3p-ECAP 150ºC
9
Elongation
Stress (MPa)
Stress (MPa)
550
10
600 550
8
500 Ultimate Tensile Strength
7 6
450
5 Yield Strength
400 0
10
20
V
30
40
50
60
4 70
again balances greater precipitation and strengthening with lower ductility, against avoiding precipitation for better ductility but lower strength. Comparison of these strength–ductility combinations with combinations obtained by conventional processing (Fig. 1b) shows significant strength increases (from 500 to 570 MPa) but with some loss of ductility (from 9% to 6%). The elongation obtained, however, remains interesting for aerospace applications [8]. A complete analysis of the relative strengthening contributions of dislocations and cells, grain boundaries, solute and precipitates, as well as the effect on ductility, is, however, beyond the scope of the present publication. Further improvements of strength–ductility combinations are likely by optimizing annealing times and temperatures (see e.g. Ref. [9]). The present study, however, has aimed simply to demonstrate the potential interest of severe plastic deformation processing. In all cases here, the many fine precipitates cause a major strength increase above the baseline work hardening due to dislocations. Deformation to high strain levels leaves the material with an elongated dislocation cell substructure with little reduction in the grain size itself. Such deformation before or during annealing ensures much finer and more homogeneous precipitation than without deformation or after smaller strains. Homogeneously distributed fine precipitates appear to improve elongation, both to ultimate tensile strength and to failure in most cases. Clearly ECAP processing of the Al– Cu–Li alloys requires careful control of the microstructure, especially the number and size of precipitates, and it is possible that the high strength with poor ductility of Al–Li alloy reported previously [5] was induced by insufficient control of such precipitation. Financial support by the Spanish Ministry of Education, under project numbers MAT2006-01827 and MAT2009-07342, is gratefully acknowledged. The authors thank Alcoa Europe of Birmingham, UK, for material supply. [1] [2] [3] [4] [5] [6] [7] [8]
Elongation to Failure (%)
10 Ultimate Tensile Strength
Elongation to Failure (%)
600
(mm/min)
ECAP
Figure 4. Yield and ultimate tensile strength, and failure elongation, as: (a) number of ECAP passes (150 °C) increases and (b) processing speed increases (3 ECAP passes at 150 °C).
307
[9] [10] [11] [12] [13] [14]
R.J. Rioja, Mater. Sci. Eng. A257 (1998) 100. J.C. Williams, E.A. Starke, Acta Mater. 51 (2003) 5775. R.Z. Valiev, T.G. Langdon, Prog. Mater. Sci. 51 (2006) 881. M. Furukawa, P.B. Berbon, Z. Horita, M. Nemoto, N.K. Tsenev, R.Z. Valiev, T.G. Langdon, Metall. Mater. Trans. 29A (1998) 169. Z.C. Wang, P.B. Prangnell, Mater. Sci. Eng. 328A (2002) 87. M.A. Mun˜oz-Morris, I. Gutierrez-Urrutia, N. Calderon, D.G. Morris, Mater. Sci. Eng. 492A (2008) 268. B.M. Gable, A.W. Zhu, A.A. Csontros, E.A. Starke, J. Light Met. 1 (2001) 1. B.M. Gable, A.A. Csontros, E.A. Starke, J. Light Met. 2 (2002) 65. A.A. Csontros, E.A. Starke, Int. J. Plast. 21 (2005) 1097. R. Yoshimura, T.J. Konno, E. Abe, K. Hiraga, Acta Mater. 51 (2003) 2891. R. Yoshimura, T.J. Konno, E. Abe, K. Hiraga, Acta Mater. 51 (2003) 4251. B. Noble, S.E. Bray, Acta Mater. 46 (1998) 6163. J. Wang, Y. Iwahashi, Z. Horita, M. Furukawa, M. Nemoto, R.Z. Valiev, T.G. Langdon, Acta Mater. 44 (1996) 2973. Alcoa product specification: Alloy 2099-T83 and 2099T8E67, 2005.
Scripta Materialia 63 (2010) 304–307 www.elsevier.com/locate/scriptamat
Severe plastic deformation processing of Al–Cu–Li alloy for enhancing strength while maintaining ductility M.A. Mun˜oz-Morris and D.G. Morris* Department of Physical Metallurgy, CENIM-CSIC, Avenida Gregorio del Amo 8, 28040 Madrid, Spain Received 1 March 2010; revised 29 March 2010; accepted 14 April 2010 Available online 18 April 2010
An Al–Cu–Li alloy has been processed by equal channel angular pressing (ECAP) to compare the influence of strain rate and processing temperature, together with subsequent annealing treatment, on the strength and ductility of the alloy. High strength with reasonable ductility can be found under conditions where fine precipitates enhance strength but also delay strain localization and fracture. Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Equal channel angular pressing; Aluminium alloy; Mechanical properties
Aluminium–lithium base alloys are of great interest in aerospace construction because they have relatively low density and high strength, while at the same time possessing good stiffness and reasonable ductility [1,2]. There is nevertheless much interest in further improving the strength–ductility combination. Severe plastic deformation processing is a method for microstructural refinement which may lead to improved strength with improved ductility [3]. There have been relatively few studies of ECAP processing of Al–Li alloys (e.g. [4,5]), and these have not significantly enhanced such strength–ductility combinations. One study [4] carried out ECAP at high temperature and obtained some grain refinement but only low strength with good ductility, although subsequent heat treatments improved the strength somewhat. Another study [5] performed ECAP at 200–250 °C, producing very fine grains with high strength but essentially no ductility, while annealing improved ductility but led to a major loss of strength. A previous study by the present authors [6] examined hot rolling of an Al–Cu–Li alloy, with processing above 200 °C achieving some microstructural refinement as extensive precipitation of T1 phase (Al2CuLi) or d0 phase (Al3Li) occurred. The present study examines microstructural evolution and mechanical properties of the same alloy following higher levels of strain imposed by ECAP under a wide variety of processing conditions. The alloy, of composition Al–2.8 wt.% Cu–1.6 wt.% Li, with small additions of Zr, Mg, Zn, Mn and Fe, is designated alloy 2099, and is somewhat similar to AF/C-458 material * Corresponding author. Tel.: +34 91 533 8900x336; fax: +34 91 533 8016; e-mail: [email protected]
examined elsewhere [7–9]. The extensive formation of T1 phase during high-temperature exposure has been well documented [10,11], together with formation of other phases such as h0 (Al2Cu), with the d0 phase expected only at relatively low temperatures [10–12]. It is also well established [7–9] that straining prior to high-temperature exposure can significantly accelerate precipitation. The combination of high-temperature exposure with severe deformation thus offers a wide potential for microstructural modification. Samples for processing were given a solutionizing anneal of 1 h at 540 °C and then water quenched. ECAP was carried out from room temperature (RT) to 200 °C in a hydraulic machine using a circular cross-section die of diameter 20 mm with a die angle of 118°, producing a true strain of 0.7 per pass. The so-called route A was used [13], i.e. with no rotation of the sample between passes. The temperatures chosen were low enough such that either no precipitation or only precipitation to fine particle sizes [6] occurred. Significant precipitation hardening is thereby retained during processing or achievable during subsequent annealing. ECAP was carried out using 3–7 passes (strain level 2.1–4.9), restricting the total strain to moderate levels so as to achieve much of the grain refinement and hardening and ductility improvements possible by severe plastic deformation [3,5]. Samples were heated to die temperature in 5 min before pressing. At the standard pressing speed (20 mm min1) the total cycle time (preheating and pressing) was 10 min. Some tests were carried out with faster pressing (60 mm min1) and slower pressing (6 mm min1), leading to cycle times of 5–6 and 15 min, whereby most of the cycle time was spent in preheating or in pressing. The heated split
1359-6462/$ - see front matter Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2010.04.022
M. A. Mun˜oz-Morris, D. G. Morris / Scripta Materialia 63 (2010) 304–307
die could be opened hydraulically for rapid sample removal and water quenching. Tensile samples were machined for testing along the pressing direction. Two samples were tested in each case, confirming excellent data reproducibility. Microstructures were examined by scanning electron microscopy (SEM) in back-scattered electron (BSE) mode, electron back-scatter diffraction (EBSD) and transmission electron microscopy (TEM). Samples were electropolished with 20% nitric acid in methanol in the direction perpendicular to the inlet and exit pressing directions to reveal any elongation of microstructures. Several processing conditions were examined and those achieving significant improvements in stress– ductility combinations are reported here. All samples showed significant work hardening up to the ultimate tensile stress, with failure occurring either immediately thereafter or after prolonged straining under reducing nominal stress, as illustrated in Figure 1a. The strength and elongation values are shown in Table 1. Processing by ECAP for 3 passes (strain 2.1) at RT, 150–200 °C led to moderate values of yield strength with poor to good ductility (see Fig. 1b). Also shown here is the increase in strength and elongation after annealing the RT processed material. There are also important strength increases with little change in elongation when processing at 150 °C for 5 or 7 ECAP passes. Figure 1b also shows reference data for comparison. T6 refers to material solutionized and annealed for 5 h at 200 °C, leading to maximum strength at this temperature [6]. Data are also shown for yield strength and elongation of optimized commercial material following solutionizing, stretching and ageing [14]. Some of the ECAP-processed materials clearly show significant strength improvements over annealed or stretched-annealed materials, with lower but reasonable elongations. Materials always showed a deformation substructure of elongated dislocation cells, with no precipitation after RT processing and fine precipitate plates after sufficient hightemperature exposure. Various microstructures are illustrated in Figures 2 and 3. Figures 2a and 3a show typical elongated cell structures, imaged by EBSD and by TEM. EBSD confirmed that these structures were constituted by cell walls with very low misorientations, typically 70–80% less than 3°. Figure 2a illustrates the evolution of misorientation along a line drawn in the image showing that most 600
Stress (MPa)
5Ep150ºC
3EpRT
400
200
0
0
5
10
Elongation (%)
15
Yield Strength (MPa)
600
3EpRT + 200ºC
+anneal 200ºC
7Ep150ºC
550
A 5Ep150ºC
500
T6
3Ep200ºC A
450 3Ep150ºC 3EpRT
400 0
2
4
6
8
10
Elongation to Failure (%)
Figure 1. Tensile stress–elongation curves illustrating work hardening after yielding and different strain behaviour after ultimate tensile stress, depending on processing conditions. (b) Yield strength and elongation combinations after processing: ECAP at RT, 150 and 200 °C; 3, 5 and 7 passes. Material processed by 3 passes at RT was also annealed for 1/2 h at 200 °C. Data point T6 refers to material aged to peak strength without prior deformation. The two points A refer to Alcoa datasheet values [14].
305
boundaries have misorientations of only a few degrees, with others reaching up to 10°. Dislocation densities and cell dimensions were determined (Table 1) by counting intercepts on fiducial lines of images such as those shown in Figure 3a, covering several cell dimensions. Measured values are subject to significant error, and can be taken simply as indicative values. Dislocation cell width was approximately related to the inverse (or inverse square root) of dislocation density, and cell length approximately proportional to width. In view of such interrelationships, the following discussion will use dislocation density in any analysis of strengthening. High-temperature processing, during ECAP and annealing, led to only a slight recovery of the substructure, which remained as dislocation cells. High-temperature processing also led to precipitation, generally as fine plates of T1 phase, illustrated in Figure 2b (precipitates visible by atomic number contrast due to the high Cu content) and the TEM dark field (DF) images of Figure 3b–d (obtained using 00002Ti spot). The T1 precipitates are known to lie on {1 1 1}Al planes (i.e. {1 1 1}Al // {0 0 0 1}T1, with h1 1 0iAl // h1 0 1 0i) and appear as thin, circular plates [6,10,11]. The precipitate disc size can be measured from SEM images (Fig. 2b) when they are thick, or from TEM DF images (Fig. 3c and d). The disc thickness was often too small (of the order of 1–2 nm) to be determined accurately from such images, so precipitate thickness was determined instead from the extent of streaking (peak width at half intensity) of precipitate diffraction spots. The spot selected was the 0002T1 spot (found at 1/2 1/2 1/2Al; see Fig. 3c), obtained when the precipitates are oriented edgeon. Precipitate diameters and thicknesses, as well as precipitate densities, determined from TEM images are reported in Table 1 for all states where precipitation was detected, with the volume fraction of the precipitate phase determined from these two parameters also shown. The determination of particle density was especially difficult for very fine particle sizes, and these values are subject to considerable error. The precipitate thickness was apparently mostly determined by exposure time at high temperature while the precipitate disc diameter was greatly reduced by straining before or during heat treatment. Figure 3b shows very thin and large diameter precipitates in material annealed without prior strain (taken from Ref. [6]), of comparable thickness but much larger than similar annealing given after heavy straining (Fig. 3d). Particle number densities are much higher than expected in the absence of prior strain or when given small stretches before annealing [7], with particle sizes reduced accordingly and similar volume fractions of particle phase found. Figure 1b and Table 1 confirm a dramatic increase in both yield strength and elongation to failure on annealing material ECAP processed at RT. The ageing temperature (200 °C) was selected based on earlier studies [6] of age hardening of solutionized material where annealing for 2–20 h led to a flat hardness maximum at a level only slightly lower than after prolonged annealing at lower temperatures. Deformation before ageing is known to accelerate hardening [6,8], such that the annealing time for maximum hardening at 200 °C becomes conveniently short (1/2 h). Additional heat treatments (1/4 and 2 h at 200 °C) confirmed very similar strength and elongation values to those of Figure 1b and Table 1. The RT ECAP-processed material has moderate strength but fail-
306
M. A. Mun˜oz-Morris, D. G. Morris / Scripta Materialia 63 (2010) 304–307
Table 1. Processing details, tensile mechanical properties and quantified microstructure. Processing history
Yield Ultimate Elongation at Elongation Particle Particle Particle number Volume Dislocation Dislocation cell density size (width of T1 strength tensile ultimate tensile to failure thickness diameter density phase (%) (m2) length) (lm) (MPa) strength strength (%) (%) (nm) (nm) (m3) 1022 (MPa)
3EpRT 3EpRT + ½ h200 °C 3Ep150 °C 5Ep150 °C 7Ep150 °C 3Ep150 °C slow 3Ep150 °C fast 3Ep 200 °C 5Ep200 °C
434 577
460 581
1.6 0.9
1.6 4.7
None 1.4
None 15
– 6.8
– 1.7
2 1016 2 1016
0.39 2.5 0.35 2
443 505 568 454 410 502 464
476 518 576 475 450 512 482
4.8 1.9 1.3 3.6 4.4 2.7 2.25
7.3 5.2 6.0 5.4 9.2 8.8 7.5
None None 1 <1 None 3.1 3.9
None None 7.5 <10 None 15 18
– – 10 1–10 – 2.5 2
– – 0.4 0.1 – 1.4 2.0
6 1015 1 1016 1.3 1016 5 1015 5 1015 1.5 1015 2 1015
0.5 2.7 0.4 2 0.35 1.5 0.5 3 0.45 2.5 0.65 4 0.5 3
The processing history indicates the number of ECAP passes at a given temperature, at the standard, slow or fast pressing speed, and any subsequent annealing.
ure occurs immediately after the ultimate tensile strength. Following annealing there is a significant increase in both yield strength (from 434 to 577 MPa) and ultimate tensile strength. The elongation to ultimate tensile strength is decreased somewhat, while the elongation to failure is increased from 1.6% to 4.7%. No change in dislocation substructure was detected after annealing (measured dislocation density identical), with presumably Cu and Li solute or precipitates slowing recovery. Strengthening increases due to fine precipitation during annealing (Fig. 3d), and these precipitates also delay strain concen-
Figure 2. Typical microstructures of transverse sections after ECAP. (a) Misorientation map obtained by EBSD: dislocation cell structure after 3 passes at RT, followed by annealing for 1/2 h at 200 °C; (b) SEM–BSE image: coarse T1 precipitates after 5 passes at 200 °C. (a) The point-to-point (in red) and point-to-origin (in blue) misorientation change along the line drawn in the image. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this paper.)
tration and the onset of failure. Such delayed failure is particularly dramatic in comparison with as-ECAPed material, which failed at the ultimate tensile strength. Figures 1b, 4a and Table 1 show that there is significant strengthening with little change of elongation to failure on increasing the ECAP strain at 150 °C. With increasing yield strength, the elongation to ultimate tensile strength decreases, but failure elongation hardly changes. No precipitation is found after 3–5 passes, but many fine precipitates are present after 7 passes (Fig. 3c), with a simultaneous slight increase in dislocation density. The strength increase is associated with the fine precipitates rather than the change in dislocation density. Even though these fine precipitates may be easily sheared, it is clear that they do not encourage strain concentration and early failure. Figure 1b and Table 1 also compare the role of processing temperature (150–200 °C) when giving 3–5 passes.
Figure 3. Microstructures obtained by TEM on transverse sections after ECAP processing. (a) Dislocation cell structure after 3 passes at RT with 1/2 h annealing at 200 °C; (b) large, thin T1 precipitates, with corresponding diffraction pattern, after annealing undeformed material for 1/2 h at 200 °C (taken from Ref. [6]); (c) fine T1 precipitates, with corresponding diffraction pattern, after 7 passes at 150 °C; (d) fine T1 precipitates after 3 passes at RT and annealing for 1/2 h at 200 °C.
M. A. Mun˜oz-Morris, D. G. Morris / Scripta Materialia 63 (2010) 304–307
Higher-temperature processing increases strength (from 443 to 502 MPa) after 3 passes but decreases it (from 505 to 464 MPa) after 5 passes, always with improved elongation to failure. Elongation at ultimate tensile strength seems generally to be inversely related to the strength level. Dislocation density is reduced somewhat after higher-temperature processing but, most importantly, precipitates appear (see Table 1). The reduced dislocation density can lower strength levels, but this is compensated by precipitation hardening when these are very fine. These precipitates appear, however, to be responsible for delaying strain localization and fracture.Figure 4b shows changes of mechanical properties for 3 ECAP passes at 150 °C at various speeds. Increasing speed reduces stress levels but increases ductility. Elongation to failure increases substantially, with elongation to ultimate tensile strength changing little (Table 1). There is, however, no clear change of dislocation substructure and only weak evidence of precipitation after processing slowly. The precipitates in such material were extremely thin discs, with diffraction spots too weak to allow reliable confirmation of structure (T1 precipitates show characteristic strong 1 0 1 0T1 diffraction, found at positions corresponding to 4/3 4/3 0Al [6,10,11], which were not detected). The precipitates here may be T1 phase or perhaps d0 phase [6,10,11]. Such incipient precipitation produces higher strength and smaller elongation to failure. The lower strength level and increased elongation when processing quickly remain difficult to understand, since no difference in dislocation structure or precipitation is found compared with the reference processing speed. Incipient precipitation possibly begins (undetected) after processing at the reference speed, which is avoided after fast processing. The corresponding lower yield strength then allows more homogeneous strain than whenever fine/incipient precipitation occurs. The present experiments have shown several ways to improve strength while retaining ductility of Al–Cu–Li. The role of dislocations and precipitates can be understood through examination of the microstructural data in Table 1: (i) after deforming to work harden, suitable annealing ensures fine and homogeneous precipitation of T1 phase, with the material strong and resistant to failure. (ii) Increasing the amount of strain at 150 °C leads to increased dislocation hardening, but it is the onset of precipitation that causes significant strengthening, while good ductility is retained. (iii) The correct ECAP temperature controls dislocation density and precipitate size. Thicker precipitates lower the strength but can further delay fracture. (iv) Controlling the speed of ECAP processing at 150 °C 9 8
Yield Strength
500
7
450
Elongation
6 5
400
ECAP 150ºC
2
3
4
5
6
Nº ECAP Passes
7
8
4
3p-ECAP 150ºC
9
Elongation
Stress (MPa)
Stress (MPa)
550
10
600 550
8
500 Ultimate Tensile Strength
7 6
450
5 Yield Strength
400 0
10
20
V
30
40
50
60
4 70
again balances greater precipitation and strengthening with lower ductility, against avoiding precipitation for better ductility but lower strength. Comparison of these strength–ductility combinations with combinations obtained by conventional processing (Fig. 1b) shows significant strength increases (from 500 to 570 MPa) but with some loss of ductility (from 9% to 6%). The elongation obtained, however, remains interesting for aerospace applications [8]. A complete analysis of the relative strengthening contributions of dislocations and cells, grain boundaries, solute and precipitates, as well as the effect on ductility, is, however, beyond the scope of the present publication. Further improvements of strength–ductility combinations are likely by optimizing annealing times and temperatures (see e.g. Ref. [9]). The present study, however, has aimed simply to demonstrate the potential interest of severe plastic deformation processing. In all cases here, the many fine precipitates cause a major strength increase above the baseline work hardening due to dislocations. Deformation to high strain levels leaves the material with an elongated dislocation cell substructure with little reduction in the grain size itself. Such deformation before or during annealing ensures much finer and more homogeneous precipitation than without deformation or after smaller strains. Homogeneously distributed fine precipitates appear to improve elongation, both to ultimate tensile strength and to failure in most cases. Clearly ECAP processing of the Al– Cu–Li alloys requires careful control of the microstructure, especially the number and size of precipitates, and it is possible that the high strength with poor ductility of Al–Li alloy reported previously [5] was induced by insufficient control of such precipitation. Financial support by the Spanish Ministry of Education, under project numbers MAT2006-01827 and MAT2009-07342, is gratefully acknowledged. The authors thank Alcoa Europe of Birmingham, UK, for material supply. [1] [2] [3] [4] [5] [6] [7] [8]
Elongation to Failure (%)
10 Ultimate Tensile Strength
Elongation to Failure (%)
600
(mm/min)
ECAP
Figure 4. Yield and ultimate tensile strength, and failure elongation, as: (a) number of ECAP passes (150 °C) increases and (b) processing speed increases (3 ECAP passes at 150 °C).
307
[9] [10] [11] [12] [13] [14]
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