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INFLUENCE OF GRAIN MORPHOLOGY AND TEXTURE ON THE DEFORMATION AND FRACTURE OF EXTRUDED CP276 Al-Li ALLOY
F. Broussaud, C. Diot
To cite this version:
F. Broussaud, C. Diot. INFLUENCE OF GRAIN MORPHOLOGY AND TEXTURE ON THE DEFORMATION AND FRACTURE OF EXTRUDED CP276 Al-Li ALLOY. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-597-C3-603. �10.1051/jphyscol:1987369�. �jpa-00226600�
JOURNAL DE PHYSIQUE
CoLloque C3, supplement au n09, Tome 48, septembre 1987
INFLUENCE OF GRAIN MORPHOLOGY AND TEXTURE ON THE DEFORMATION AND FRACTURE OF EXTRUDED CP276 A1-Li ALLOY
F. BROUSSAUD and C. DIOT
ONERA, 29, Avenue de la Division Leclerc, B.P. 7 2 , F-92322 ChStillon Cedex, France
Abstract
The mechanical properties of a C P 276 extruded f l a t bar have been evaluated in t h e T351 and T851 conditions. The tensile properties in t h e extrusion direction (L) exhibit a pronounced profile along t h e long-traverse direction (T). Changes in slip line pattern and in f r a c t u r e mode have been studied a s a function of t h e position of L specimens along t h e T direction. Rupture modes depend on grain morphology. Deformation characteristics appear t o be related t o variations of preferred local crystallographic orientation. Schmid and orientation factors have been estimated for t h e edge and t h e c e n t r e of t h e f l a t bar.
INTRODUCTION
Conventional processing of Al-Li-Cu-Mg-Zr alloy gives rise t o a n unrecrystaIIised grain structure with a pronounced texture, which results in a strong anisotropy of mechanical properties [1,21. P e t e r s and a1 [3! also demonstrated t h a t mechanical properties varied through t h e thickness of a plate material due t o variation of texture. A similar evolution is observed along t h e long-transverse direction of C P 276 extruded f l a t bar. This paper details t h e evaluation of different metallurgical parameters in order t o explain t h e profile of tensile properties.
EXPERIMENTAL
The C P 276 alloy (2.25 wt% Li-2.85-wt% Cu-0.55 wt% Mg-0.10 wt% Zr) was supplied by CEGEDUR PECHINEY (CRV) in t h e form of a n extruded f l a t bar (100 x 13 mm2) in t h e T351 temper (2% stretch). The alloy was studied for three conditions : T351, T8x51 (underaged) and T85 1 (peak-aged).
Round specimens for tensile testing were c u t off parallel t o t h e extrusion direction (L) in t h e mid-thickness plane a t different locations (relative t o t h e extrusion axis) along t h e long- traverse direction (T). Each specimen location will be indicated by i t s distance (mm) from the bar edge.
Characterisation of t h e microstructure was carried o u t using both optical and scanning electron microscopy (SEM). Samples were prepared by standard techniques. SEM f r a c t u r e examination was carried o u t on t h e fractured tensile samples. Furthermore, optical microscopy examinations were performed on polished and etched cross-sections.
The d d o r m a t i o n behaviour was examined on f l a t test-pieces polished on L x S and L x T planes . Slip t r a c e s were observed by optical microscopy using a Nomarski interference contrast amplifier.
Thin foils f o r transmission electron microscopy (TEM) investigations were double-jet electrothinned in 30% nitric acid-70% methanol electrolyte a t -30°C (voltage :
Texture approach was performed on a n X-ray 0120 goniometer (ONERA). The l 2
vO1tsk
I220 pole fig- ures were obtained by using t h e Schultz reflexion technique (CRV). Intensities have been corrected for defocalization. The X-ray source was a Cu X-ray tube operating at 30 kV and 30 mA. Pole figures w e r e determined a t 5mm and 45 mrn from t h e edge of t h e bar. Samples were examined on t h e L x S plane, and not L x T, in order t o avoid t h e e f f e c t of t h e texture evolution along T-direction.RESULTS
Preliminary Vickers hardness measurements performed all over t h e section of t h e f l a t bar ( ~ 3 5 1 condition) revealed a strong heterogeneity, as illustrated by t h e isohardness map in fig. 1.
These results a r e confirmed by tensile t e s t s in L-direction : variation of tensile characteristics
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through the T-direction is shown in figure 2. Tensile properties a r e symmetrical relative t o the extrusion axis and their evolution is continuous from the edge t o the center of the bar. Yield stress and rupture strength a r e minimum a t the centre of the bar whereas elongation is maximum. The difference of elongation values between the centre and the edge of the bar is maximum in t h e T351 condition and decreases after prolonged aging treatments. The stress- strain curves exhibit two salient features :
- Serrated yielding which may indicate a Portevin-Le Chatelier e f f e c t is observed in the T351 condition, but not a f t e r artificial aging.
- The shape of stress-strain curves is dependent on the location of the specimen within the bar relative to t h e extrusion axis (fig. 3). Work hardening is higher for specimens originating from the centre of the bar.
Fig. I : Isohardness map. Plane T.S. -
.. -
Fig. 2 : Proof and tensile strength profiles : A T351 , m T8x51, T851
MPa
4 Fig. 3 : Stress-strain curves (T8x51) a, 5 mm b, 17.5 mm c, 45 mm
300
In order t o explain the evolution of tensile properties along the T-axis, composition, microstructure, slip morphology, fracture modes and texture were successively investigated.
Composition and microstructure
No significant gradient of addition elements (Li, Cu, ~ g ) has been detected along the T direction. Li content was determined by emission spectrometry, Cu and Mg by chemical analysis.
The grain size distribution observed in the T x S plane (Table I) appears t o be inhomogeneous and t o vary from a fine grain structure near t h e edge of t h e bar t o a coarse grain structure a t the centre (fig. 4). At the periphery of t h e bar, a recrystallised layer 100 um thick can be seen.
At t h e centre of t h e bar, a typically uniecrystallised pancake grain structure is observed exhibiting elongated grains parallel t o the extrusion direction with large zones constituted of well-developed sub-grains. The ratio between the size of the grains along the T and S directions increases from 1 a t a distance of t h e bar edge of 5 mm t o 7 a t t h e centre of the bar.
Fig. 4 : Macrographs of the grain shape (L.T. plane) a, 5 mm b, 45 mm Table I : Evolution of grain size
Deformation
Slip line patterns, observed on f l a t polished specimens in the T351 condition after straining up t o 1% or 4%, depend on the location of the specimens in t h e bar, relative t o the extrusion axis. For centre specimens, slip lines a r e straight and well outlined, especially within large elongated grains. A dominant slip line system with angles between t h e slip lines and the stress axis of 24' in the L x S plane and 55" in the L x T plane was evidenced. For specimens
Grain sizes p
Etched samples revealed the existence of two distinct insoluble particles. The larger ones with an average size of 10 Um contain an appreciable amount of Cu. Their distribution is hazardous and recrystallisation areas a r e located around these particles. The second type of particles, 1Pm in dimension, a r e aligned parallel t o the extrusion axis ; their composition is close t o A17Cu~Fe.
TEM samples of aged alloy indicated the presence of homogeneously distributed 6 ' precipitates. The TI and S9 phases were also observed for the two artificial aging treatments.
In short, concerning the size and distribution of insoluble particles, dispersoids (A132r) and precipitates, no significant evolution of the microstructure has been detected along T-direction.
Locations (mm)
T direction S direction
1 30 45
4 110 170
5 190 180
735 250
10 280 110
45 700
80
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Fig. 5 : Optical micrographs of t h e shear bands observed on t h e LS plane in longitudinal t e s t e d specimens a, a t 2,5 mm b,at 45 mm
Some TEM examinations of strained samples (T8x51 and T851) were carried o u t : long slip bands (parallel t o (1 11) planes) extend across several grains.
Rupture behavior
In t h e T351 condition, t h e rupture surface is identical for specimens originating from t h e edge or t h e c e n t r e of t h e bar. The f r a c t u r e surface is dominated by large shear facets.
Macroscopically, t h e s u r f a c e is inclined 45" t o t h e stress axis while t h e elongated nature of t h e grains produced a stepped appearance, especially f o r specimens c u t off a t t h e centre. Local intergranular rupture may be observed. Examination of longitudinal sections of f r a c t u r e d specimens clearly evidenced t h a t this intergranular rupture coincides with recrystallised areas.
In underaged and peak-aged conditions, t h e rupture mode appears different between t h e edge and c e n t r e specimens. This difference i s particularly pronounced in t h e T851 t e m p e r (fig.
6). For specimens originating from t h e e d g e of t h e bar, t h e f r a c t u r e is predominantly intergranular (fig. 7). For c e n t r e specimens, t h e rupture mode is more complex, a s f r a c t u r e surface showed both intergranular failure (including delamination of longitudinal grains boundaries) and shear failure.
Fig. 6 : Scanning electron micrographs showing t h e f r a c t u r e surface a , at 5mm b, a t 45mm
4 Fig. 7 : Intergranular f r a c t u r e (5mm T851)
Texture
Examinations with @/@ goniometer showed t h a t t h e intensities of preferred crystallographic orientation varied along T-direction. The planes T x S, L x S and L x T were examined :
- T x S plane : t h e (220), (331) and (420) planes a r e absent on t h e diagram. A strong (111) component is observed near t h e edge of t h e bar. The (422) intensities increases around t h e c e n t r e of t h e bar.
- L x S plane : all t h e planes a r e present. An evolution is noted from edge t o c e n t r e along t h e T direction. A (220) component is observed near t h e edge of t h e bar. A (11 1) component is present around t h e c e n t r e of t h e bar.
- L x T plane : all t h e planes a r e observed. The ( I 11) and (331) intensities a r e weak.
Diffraction d a t a s for th pl nes L x T and T x S a r e given for comparison on figure 8.
The figure 9 shows t h e f 22 C f pole figures. At 5 mm away from t h e edge of the bar, there is a very strong t e x t u r e : t h e plane L x S i s of (110) type, t h e longitudin_al orientation is < I l l > . This orientation is accompanied by weaker components : { 110) <001>, { 11?) < 11 1,. A t 45 mm, several t e x t u r e components a r e present :{ I10)< 001>,1 1 0 0 ) ~ 001>,{ 11 1)'<112>.
T.S p l a n e L . T plane
Fig. 8 : Diffraction d a t a s for t h e L.T and T.S planes
JOURNAL DE PHYSIQUE
DISCUSSION Fracture
SEM examinations showed that the grain shape has a large influence on the fracture modes.
For specimens originating from t h e centre of t h e bar, the grain aspect ratio results in shear rupture and delamination. When the aging treatment reaches the peak-aged temper, intergranular rupture becomes easier : thus, a competition establishes between shear and intergranular rupture. For edge specimens, due t o the small grain size, intergranular rupture is favoured. Taking in account the difference in rupture stress level between edge and centre, no fracture mode seems t o be more catastrophic than the other one.
strengthening
1. Grain size contribution
The most commonly used relationship t o rely the proof stress and the grain size is that of Hall-Petch : 0 = 0 + kd-112 where d is the grain size, k a constant and (J the flow stress of the grain interior. In t h e A1 alloys, the grain size hardening is known t o be low. K is in the range 0.06-0.25 M P ~ J ~ t41. This factor is 0.1 for AI-3Li in t h e as quenched condition [51, 0.23 for peak- aged A1-2Li-2Mg alloy 161. The experimental difference between the proof stresses of samples c u t off a t 7.5 mm away from the bar edge is 80 Mpa for the T851 temper. The maximum difference between these two specimens by applying t h e Hall-Petch relationship with parameter established by Dinsdale t61 and by taking the grain size in the T-direction which overestimates their contribution to proof stress is 8MPa. So the contribution of the grain size hardening is too small t o explain the difference between the proof stresses of the edge and centre specimens.
2. Texture contribution
0 is related t o texture by t h e relationship 0 = Mar where T is the CRSS on < I 1 I> planes and M is an orientation factor averaging the distribution of grain orientations. When this distribution is isotropic, M ' 3 (Taylor factor).
The study of variations of texture along the T axis allows M t o be estimated a t the edge and a t the centre of the bar. At the bar edge, the texture is very strong, with < I l l > along the extrusion direction. M can be approximated by t h e reciprocal of maximum Schmid factor (Table 2). With this orientation, six slip systems experience the same stress with a Schmid factor value of 0.272 (giving therefore M = 3.67 a t the bar edge). This result is consistent with the observation of multiple slip line systems. At the bar centre, the evaluation of M appears t o be more complex:
indeed, it remains some incertainties concerning the dominant texture component. Examination of the (220) pole figure does not allow t o find a major component. Assuming that this resu_lt is correct, M can be considered a s close t o 3 (Taylor factor). But, 8/28 diagramms give <112> a s preferred longitudinal orientation. This result is confirmed by analysis of the dominant slip line system observed on t h e f l a t test-bars. The theoretical traces on the L x S and L x T planes a r e inclined respectively 30" and 54" t o the L-direction, in agreement with experimental observations. In this case, M is estimated by reciprocal of corresponding maximum Schrnid factor (Table 21, i.e M = 2.45.
Table 2 : Schmid factors of slip system (1 11 ) 10 >
Slip system lo I TI (111) [loll
[1103 [Ol l l (1y1) flOl1 [1101
to1 11 ( l l i ) [roil [I1 01
Longitudinal direction L a l o n g < l l l >
o
0,272 0,272 0,272 0 0,272 0,272 0,272 0
L along < 112 >
0,408 0,136 0,272 0,136 0,408 0,272 0,272 0,272 0
In conclusion, a t t h e edge of t h e bar M = 3.67, a t t h e c e n t r e of t h e bar M is assumed t o b e in t h e range 2.45-3. Thus t h e ratio M edge/M c e n t r e is in t h e range 1.2-1.5. The experimental ratio R0.2 edge/R0.2 c e n t r e i s in t h e range 1.13-1.29". So t h e r e i s a reasonable agreement between these t w o ratios and a s a consequence, i t c a n be concluded t h a t t e x t u r e hardening is mainly responsible for t h e strength evolution along t h e T-direction.
CONCLUSION
Grain size and grain shape have a direct influence on t h e f r a c t u r e mode.
Combined e f f e c t s of grain parameters could explain t h e profile of t h e tensile properties along t h e T direction, but actually t h e strength evolution observed in t h e f l a t bar is mainly due t o t e x t u r e contribution.
Acknowledgements
The authors acknowledge STPA for financial support. They wish t o thank M. Doudeau, P.
Sainfort, F. L e Gouic from PechnineyICRV for t h e pole figures obtention, and also G. Lapasset f r o m ONERA for helpful discussions.
References
[I] FOX S., McDarmaid D.S., Flower H.M. (1986) in Aluminium technology186, Institute of Metal, 327.
[2] FOX S., Flower H.M., McDarmaid D.S. (1986) in Aluminium alloys : their physical and mech- anical properties, t h e Institute of Metal, 939.
[3] P e t e r s M., Eschweiler J., Welprnann K. (1986) Scripta Met. 20, 259.
[k] Evensen S.D., Ryum N., Embury J.D. (1975) Mat. Sc. Eng. 18, 221.
[5] Jensrud 0. (1985) in ALuminium Lithium Alloys 111 edited by Baker, Gregson, Harris and Peel, Institute of Metal (Oxford), 411.
[6] Dinsdale D., Harris S.J., Noble B. (1981) in Aluminium Lithium Alloys I, 101.
