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On the evolution of heterogeneous microstructure and microtexture in impacted aluminum–lithium alloy
Journal of Alloys and Compounds 578 (2013) 183–187
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
On the evolution of heterogeneous microstructure and microtexture in impacted aluminum–lithium alloy N.P. Gurao ⇑, A.O. Adesola, A.G. Odeshi, J.A. Szpunar Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, Canada S7N 5A9
a r t i c l e
i n f o
Article history: Received 24 February 2013 Received in revised form 24 April 2013 Accepted 25 April 2013 Available online 6 May 2013 Keywords: Metals and alloys Anisotropy X-ray diffraction
a b s t r a c t The microstructure and texture evolution in an aluminum–lithium alloy showing adiabatic shear band after dynamic deformation has been carried out using electron backscatter diffraction and bulk texture measurement using X-ray diffraction. The uniquely heterogeneous microstructure is characterized by fine sub-micron grains of the matrix phase and small spherical precipitates while the region away from the shear band shows deformed elongated grains and plate shaped precipitates. The interface between the shear band and rest of the sample shows unique deformation features indicating the existence of both continuous as well as geometric dynamic recrystallization in the impacted sample. It is inferred that the coarse grain size accompanied with heterogeneous deformation conditions during dynamic loading leads to the coexistence of competing recrystallization mechanisms in aluminum–lithium alloy. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Light weight aluminum–lithium (Al–Li) alloys offer an exciting opportunity for reducing the fuel consumption in the aerospace and automobile industry [1]. New generation Al–Li 2099 alloy developed by Alcoa Inc. in 1997 overcomes the problems pertaining to planar anisotropy, lower transverse ductility and lower toughness associated with conventional Al–Li alloys [1–3]. This alloy has high specific modulus and excellent fatigue and cryogenic toughness compared to other aerospace aluminum alloys. The properties of the alloy can be modulated significantly by altering the heat treatment prior to application. The most common heat treatment is the T8 temper, wherein, the alloy is solution treated, subjected to 6% stretching and then aged. The dislocations introduced during pre-age stretching act as nucleation sites for the T1 (Al2CuLi) precipitates that constitute the major strengthening phase along with the coherent d0 (Al3Li) phase [4–5]. There have been few investigations on these alloys to study the deformation behavior in different loading conditions and strain rates [6]. However, the impact behavior of this alloy in various aging conditions and the subsequent evolution of microstructure and microtexture have not been explored. Of particular importance is the need to understand the mechanism of formation of adiabatic shear bands in this material so as to avoid failure during impact. It is therefore imperative to study the micro-mechanisms operative during the dynamic deformation of AA 2099 aluminum alloy. The present ⇑ Corresponding author. Present address: Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kalyanpur, Kanpur 208 016, India. Tel.: +91 512 259 6688; fax: +91 512 259 7505. E-mail address: [email protected] (N.P. Gurao). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.04.176
investigation aims at deciphering the operative micro-mechanisms that contribute to heterogeneous deformation during impact in the AA 2099-T8 alloy using Electron Back Scatter Diffraction and bulk X-ray texture measurement. 2. Experimental AA 2099 aluminum alloy with chemical composition of 2.69 Cu, 1.8 Li, 0.6 Zn, 0.3 Mg, 0.3 Mn and 0.08 Zr in weight percent was provided by Alcoa Inc. in the form of hot rolled plate. Cylindrical specimens with 10.5 mm height and 9.5 mm diameter were machined in the rolling direction and subjected to dynamic impact loading using an instrumented direct impact Hopkinson bar. The cylindrical specimens were impacted with a blunt projectile fired by a light gun at impact momentum ranging between 28 and 39 kg m/s. On impact, elastic waves were produced, which propagated through it onto the output bar. The elastic waves, captured by strain gage attached to the bar, were used to generate the dynamic stress–strain curves for the impacted specimens. A detailed list of samples and the strain rate as well as strain achieved during dynamic loading are presented in Table 1. The deformed samples were ground, polished and prepared metallographically to one micron diamond polishing. Bulk texture measurements were carried out on a Bruker D8 Advance diffractometer with Cu Ka radiation and a 2D Hi-star detector. Incomplete (1 1 1), (2 0 0), (2 2 0) and (1 1 3) pole figures were measured. Resmat software [7] was used to calculate the complete Orientation Distribution Function and determine the complete pole figures as well as volume fraction of important fibers. The sample impacted at highest momentum was then subjected to colloidal silica polishing using Vibromet for 12 h. Orientation Imaging Microscopy was carried out on a field emission gun scanning electron microscope Hitachi SU6600 fitted with Oxford Instruments electron backscatter diffraction. Data analysis was carried out using hkl Channel 5 data acquisition and analysis software.
3. Results and discussion The initial material in the T8 aged condition was characterized by presence of elongated grains and second phase particles which
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Table 1 Experimental conditions and volume fraction of important texture fibers for the impacted AA 2099-T8 aluminum alloy. Strain rate (s1)
Impact momentum (kg m/s)
0 28 31 33 39
Maximum flow stress (MPa)
As received material 4878 7494 8706 8903
Engineering strain
390 370 350 360
Fiber component (%)
0.4257 0.6540 0.7598 0.7770
(a)
h1 0 1i
h1 1 1i
8.4 10.9 7.2 8.2 6.6
27.1 20.9 25.2 29.1 39.9
5.0 19.8 15.5 10.2 9.2
(b)
400
Stress (MPa)
h1 0 0i
300
200
4878 s-1 7494 s-1 8706 s-1 8903 s-1
100
500 µm 0 0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
Strain
I
II
III
(c)
500 µm
(d)
25 µm
Fig. 1. (a) Optical micrograph of the initial AA 2099-T8 alloy (b) stress–strain response of the alloy impacted at different impact momentum and evolution of shear band in sample impacted at 8903 s1 at (c) low magnification and (d) high magnification.
are distributed uniformly throughout the matrix (Fig. 1a). In the peak-aged condition the alloy will also contain T1 (Al2CuLi) and d0 (Al3Li) precipitates [4,5]. In addition, other phases like T2 (Al6CuLi3), S0 (Al2CuMg), b0 (Al3Zr) and d (AlLi) are also present in AA 2099-T8 alloy. The stress–strain curves of the samples impacted at different momentum are shown in Fig. 1b. The stress–strain curves are characterized by the presence of multiple peaks for all the samples. The sample deformed at the highest strain rate shows three distinct peaks compared to two peaks for the remaining samples. The bulk texture measurement shows an increase in h1 0 1i component that is a characteristic FCC compression texture component [8] with increase in impact pressure (Table 1). The microstructure of the deformed samples is characterized by the presence of multiple shear bands devoid of large second phase particles (Fig. 1c). However, SEM image at higher magnification (Fig. 1d) indicated the presence of spherical second phase particles of varying sizes within the shear bands. The second phase particles inside the shear bands are much smaller than those observed outside the shear band region. However, no information was obtained about the microstructure of the matrix phase from the SEM studies
and hence Electron Back Scatter Diffraction was used to characterize the microstructure. In order to decipher the heterogeneity in deformation of impacted 2099 alloy, we monitor the evolution of microstructure and microtexture away from the shear band (region I), near the shear band (region II) and within the shear band (region III) for the sample impacted at the highest pressure (Fig. 1c). Fig. 2 shows the pattern quality, inverse pole figure (IPF) map and inverse pole figure for the sample impacted at highest momentum in the aforementioned three regions. The microstructure of region I show elongated grains with significant orientation gradients within individual grain. The black region in the IPF map represents the location of the second phase particles that came out during colloidal silica polishing. The IPF map of a region III (Fig. 2) indicates presence of elongated grains as well as equiaxed grains and the grains are an order of magnitude smaller than the region far away from the shear band. The equiaxed grains show lower intragranular misorientation and form in clusters. The IPF map of region II shows a combination of very large elongated grains and bands of fine grains. Another important observation is the wavy nature of boundaries in this region. In addition, a
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185
Fig. 2. Pattern quality (top) and inverse pole figure (IPF) map (bottom) in the three different region of the sample deformed at the highest impact momentum. The inverse pole figure in compression direction (going into the plane of the paper) is also shown.
heterogeneous evolution of microtexture is observed in the sample. Region I shows a characteristic h1 0 1i fiber component and region III shows a h1 0 0i component while clustering of orientation between h1 1 1i and h1 0 1i line is observed for region II. The misorientation distribution obtained from the three different regions is shown in Fig. 3a. The region far away from shear band is characterized by high fraction of low angle grain boundaries while the region within shear band shows high fraction of high angle boundaries. The EBSD data was used to describe substructure evolution in terms of Local Average Misorientation (LAM) parameter in the three regions (Fig. 3b). Region I shows very high fraction of low misorientation and a higher value of LAM value. Region II shows lower fraction of low angle misorientation as well as lower average LAM value than Region I. Region III shows very low fraction of low angle LAM and the lowest average LAM. The partitioning of LAM into low and high misorientation in Region I is attributed to evolution of heterogeneous dislocation structure comprising of low and high dislocation density regions in the microstructure [9]. This heterogeneity reduces gradually from regions I to II and the region within the shear band shows minimum LAM and hence minimum dislocation density. Thus, the microstructure as well as microtexture evolution indicates the presence of recrystallization within the shear band. In order to understand the heterogeneity of deformation process at high strain rate, the origin of adiabatic shear bands has to be understood. During dynamic deformation heat gets localized
in particular region of the sample leading to formation of adiabatic shear bands [10–12]. The absence of large precipitates in the shear bands clearly indicates that the temperature reached the level above their dissolution temperature [13]. However, the presence of smaller spherical particles in this region indicates that the dissolution of the second phase particles was not complete during deformation and subsequent cooling. The dissolution of the second phase particles occurs via diffusion of copper atoms. The diffusion rate is higher in regions with sharp curvature and hence higher chemical potential according to Gibbs Thompson equation [14]. Hence there is change in shape of the second phase particles from plate like to spherical shape inside the shear band. The increase in Cu content of the matrix phase from 2.38 wt% in region I to 2.76 wt% in region III as determined from Energy Dispersive Spectroscopy corroborates the enrichment of the matrix phase by precipitate dissolution. Higher temperature that is expected within the shear band, also aids in recrystallization of the microstructure. However, new grains formed within the shear band are not expected to form by classical discontinuous dynamic recrystallization (dDRX) as the time available for recrystallization is inadequate for nucleation and growth of the recrystallized grains. Instead the material is expected to undergo continuous or geometric/rotational dynamic recrystallization [15]. The structure of various coherent precipitates and nanoscale dispersoids within the shear band requires extensive transmission electron microscopy investigation and is beyond the scope of this study. Coherent precipitates
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(b) 0.12 RI RII RIII
0.8
RI RII RIII
0.10 0.08
0.6
Fraction
Cumulative fraction
(a) 1.0
0.4
0.06 0.04
0.2
0.02 0.00
0.0 0
10
20
30
40
50
0
60
Misorientation°
1
2
3
4
5
Local Average Misorientation°
Fig. 3. (a) Cumulative distribution of misorientation and (b) local average misorientation distribution in the three distinct region of the AA 2099-T8 aluminum specimen impacted at the highest momentum.
Region I
Region II
100µ µm
100µm
1.5
3.7
40
60
35
50
Misorientation°
Misorientation°
30 25 20 15
40 30 20
10 10
5 0
0 0
10
20
30
40
50
60
70
Distance (µm)
0
5
10
15
20
25
30
35
40
Distance (µm)
Fig. 4. Taylor factor map for the impacted AA 2099 sample in different regions and misorientation distribution within large grains near the interface.
like T1 and T2 with dimensions of the order of few hundred of nanometer in one dimension are neither expected to dissolve nor play any role in particle stimulated nucleation during high strain rate deformation. The presence of large second phase particles in the artificially aged alloy can be expected to act as initiators for dynamic recrystallization (particle stimulated nucleation PSN DRX). However, the h1 0 0i texture observed within the shear band negates the occurrence of the same as it leads to randomization of texture. The presence of multiple peaks in the stress–strain curve (Fig. 1b) can be attributed to the occurrence of multiple dynamic
recrystallization processes in the high strain high strain rate process [16]. High strain rate deformation is characterized by grain fragmentation and higher number of geometrically necessary dislocations to maintain strain compatibility. During the course of deformation, these dislocations arrange themselves to form elongated sub grain boundaries. The misorientation of the sub-grain boundaries increases with the amount of deformation and new geometrically necessary boundaries are generated within these elongated sub-grains due to different regions of the same grain rotating in different direction to maintain strain compatibility. Gradually these sub-grain boundaries can pinch off to form newly
N.P. Gurao et al. / Journal of Alloys and Compounds 578 (2013) 183–187
recrystallized grains with lower dislocation density. The elimination of sub-grain boundaries and limited possibility of formation of high misorientation mobile grain boundary leads to smaller grain size at location of the shear band. The aforementioned mechanism is termed as geometric dynamic recrystallization (gDRX) [17]. The multiple serrations in the stress–strain curve as well as the IPF map of region I with some flattened grains (blue) and IPF map of region III (Fig. 2) with cluster of grains having slight misorientation substantiate this argument. Thus the multiple peaks in the stress–strain curve indicate towards multiple gDRX events. However, a closer look at the IPF map of the interface region II indicates very high misorientation forming in some grains. These regions are formed by grain boundaries swiping sub-grain boundaries during deformation and indicate the presence of continuous dynamic recrystallization (cDRX) [18]. Fig. 4a shows the Taylor factor map of regions I and II. Grains with high Taylor factor (hard orientations) show wavy boundaries in both the region and are expected to act as precursor for gDRX. However, misorientation profile drawn in large grains near the interface in region II clearly show low point to point misorientation accompanied with grain subdivision by geometrically necessary boundaries. Such grain subdivision is due to local lattice rotation along initial grain boundaries aided by formation of geometrically necessary boundaries to accommodate heterogeneous deformation by cDRX [21]. We therefore, believe that the initial reduction of grain size is by continuous dynamic recrystallization that is followed by geometric dynamic recrystallization in the newly formed grains. Another possibility is that cDRX operates in coarse grains with low aspect ratio or in geometrically softer grains while gDRX occurs in smaller grains with high aspect ratio or higher Taylor factor grains. However, the former mechanism wherein, gDRX is preceded by cDRX seems to be dominant as conventional recrystallization texture component h1 0 0i is dominant in the shear band area. The texture is not expected to change qualitatively during cDRX [19–20], however, a characteristically different texture from the deformation texture indicates the presence of gDRX or dDRX. Since the time scale associated with high strain rate deformation is very small, the latter is not expected to occur in the present case. Etter et al. [22] have shown the presence of cDRX and gDRX in 5251 aluminum alloy in the cold rolled and annealed condition respectively. Therefore, it is proposed that due to the difference in deformation state in the impacted sample in terms of difference in Taylor factor and grain size/aspect ratio, different mechanisms can be operative in different regions of the same sample.
187
4. Conclusions High strain rate deformation of AA 2099 aluminum alloy in T8 temper condition leads to profuse shear banding that contributes to heterogeneous microstructure and microtexture evolution. The presence of micron sized grains with h1 0 0i texture within the shear band indicate the presence of geometric dynamic recrystallization while the interface between the shear band and sample shows wavy boundaries indicating continuous dynamic recrystallization. Thus two competing dynamic recrystallization mechanisms namely the cDRX and gDRX occur concomitantly during dynamic deformation of aluminum–lithium alloy. References [1] R.J. Rioja, J. Liu, Metall. Mater. Trans. A 43 (2012) 3325–3337. [2] V.K. Jain, K.V. Jata, R.J. Rioja, J.T. Morgan, A.K. Hopkins, J. Mater. Proc. Technol. 73 (1998) 108–118. [3] A.A. Csontos, E.A. Starke Jr., Metall. Mater. Trans. A 31 (2000) 1965. [4] S.C. Huang, M.J. Starink, Int. Mater. Rev. 50 (2005) 193. [5] R. Yoshimura, T.J. Konno, E. Abe, K. Hiraga, Transmission electron microscopy study of the evolution of precipitates in aged Al–Li–Cu alloys: the h0 and T1, Acta Mater. 51 (2003) 4251–4266. [6] R. J. McDonald, Characterization of delamination in 2099-T861 Aluminum– Lithium, PhD thesis, University of Illinois at Urbana-Champaign (2009). [7] TexTools Software, Resmat Corporation, Montreal, QC, Canada. [8] U.F. Kocks, C.N. Tome, H.-R. Wenk, Texture and Anisotropy: Preferred Orientation in Polycrystals and their Effect on Materials Properties, Cambridge University Press, 1998. [9] Y. Zhong, F. Yin, T. Sakaguchi, K. Nagai, K. Yang, Acta Mater. 55 (2007) 2747– 2756. [10] S.E. Schoenfeld, T.W. Wright, Int. J. Solids Struct. 40 (2003) 3021–3037. [11] C. Zener, J.H. Hollomon, J. Appl. Phys. 15 (1944) 22–32. [12] H.C. Rogers, Annu. Rev. Mater. Sci. 9 (1979) 283–311. [13] R.W. Chen, K.S. Vecchio, J. De Phys. IV 4 (1994) 459–464. [14] D.A. Porter, K.E. Esterling, Phase transformations in metals and alloys, London, Chapman and Hall, p. 514. [15] F.J. Humphreys, M. Hutchinson, Recrystallization and Related Annealing Phenomena, Elsevier Ltd., p. 451. [16] U. Andrade, M.A. Meyers, K.S. Vecchio, A.H. Chokshi, Acta Metall. Mater. 42 (1994) 3183–3195. [17] M.A. Meyers, V.F. Nesterenko, J.C. LaSalvia, Q. Xue, Mater. Sci. Eng. A 317 (2001) 204–225. [18] T.G. Nieh, L.M. Hsiung, J. Wadsworth, R. Kaibyshev, Acta Mater. 46 (1998) 2789–2800. [19] W. Skrotzki, N. Scheerbaum, C. Oertel, R.A. Massion, S. Suwas, L.S. Toth, Acta Mater. 55 (2007) 2013–2024. [20] N.P. Gurao, R. Kapoor, S. Suwas, Metall. Mater. Trans. A 41 (2010) 2794– 2804. [21] R. Kaibyshev, O. Stdikov, A. Goloborodko, T. Sakai, Mater. Sci. Eng. A 344 (2003) 348–356. [22] A.L. Etter, T. Baudin, N. Fredj, R. Penelle, Mater. Sci. Eng. A 445–446 (2007) 94– 99.
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
On the evolution of heterogeneous microstructure and microtexture in impacted aluminum–lithium alloy N.P. Gurao ⇑, A.O. Adesola, A.G. Odeshi, J.A. Szpunar Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, Canada S7N 5A9
a r t i c l e
i n f o
Article history: Received 24 February 2013 Received in revised form 24 April 2013 Accepted 25 April 2013 Available online 6 May 2013 Keywords: Metals and alloys Anisotropy X-ray diffraction
a b s t r a c t The microstructure and texture evolution in an aluminum–lithium alloy showing adiabatic shear band after dynamic deformation has been carried out using electron backscatter diffraction and bulk texture measurement using X-ray diffraction. The uniquely heterogeneous microstructure is characterized by fine sub-micron grains of the matrix phase and small spherical precipitates while the region away from the shear band shows deformed elongated grains and plate shaped precipitates. The interface between the shear band and rest of the sample shows unique deformation features indicating the existence of both continuous as well as geometric dynamic recrystallization in the impacted sample. It is inferred that the coarse grain size accompanied with heterogeneous deformation conditions during dynamic loading leads to the coexistence of competing recrystallization mechanisms in aluminum–lithium alloy. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Light weight aluminum–lithium (Al–Li) alloys offer an exciting opportunity for reducing the fuel consumption in the aerospace and automobile industry [1]. New generation Al–Li 2099 alloy developed by Alcoa Inc. in 1997 overcomes the problems pertaining to planar anisotropy, lower transverse ductility and lower toughness associated with conventional Al–Li alloys [1–3]. This alloy has high specific modulus and excellent fatigue and cryogenic toughness compared to other aerospace aluminum alloys. The properties of the alloy can be modulated significantly by altering the heat treatment prior to application. The most common heat treatment is the T8 temper, wherein, the alloy is solution treated, subjected to 6% stretching and then aged. The dislocations introduced during pre-age stretching act as nucleation sites for the T1 (Al2CuLi) precipitates that constitute the major strengthening phase along with the coherent d0 (Al3Li) phase [4–5]. There have been few investigations on these alloys to study the deformation behavior in different loading conditions and strain rates [6]. However, the impact behavior of this alloy in various aging conditions and the subsequent evolution of microstructure and microtexture have not been explored. Of particular importance is the need to understand the mechanism of formation of adiabatic shear bands in this material so as to avoid failure during impact. It is therefore imperative to study the micro-mechanisms operative during the dynamic deformation of AA 2099 aluminum alloy. The present ⇑ Corresponding author. Present address: Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kalyanpur, Kanpur 208 016, India. Tel.: +91 512 259 6688; fax: +91 512 259 7505. E-mail address: [email protected] (N.P. Gurao). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.04.176
investigation aims at deciphering the operative micro-mechanisms that contribute to heterogeneous deformation during impact in the AA 2099-T8 alloy using Electron Back Scatter Diffraction and bulk X-ray texture measurement. 2. Experimental AA 2099 aluminum alloy with chemical composition of 2.69 Cu, 1.8 Li, 0.6 Zn, 0.3 Mg, 0.3 Mn and 0.08 Zr in weight percent was provided by Alcoa Inc. in the form of hot rolled plate. Cylindrical specimens with 10.5 mm height and 9.5 mm diameter were machined in the rolling direction and subjected to dynamic impact loading using an instrumented direct impact Hopkinson bar. The cylindrical specimens were impacted with a blunt projectile fired by a light gun at impact momentum ranging between 28 and 39 kg m/s. On impact, elastic waves were produced, which propagated through it onto the output bar. The elastic waves, captured by strain gage attached to the bar, were used to generate the dynamic stress–strain curves for the impacted specimens. A detailed list of samples and the strain rate as well as strain achieved during dynamic loading are presented in Table 1. The deformed samples were ground, polished and prepared metallographically to one micron diamond polishing. Bulk texture measurements were carried out on a Bruker D8 Advance diffractometer with Cu Ka radiation and a 2D Hi-star detector. Incomplete (1 1 1), (2 0 0), (2 2 0) and (1 1 3) pole figures were measured. Resmat software [7] was used to calculate the complete Orientation Distribution Function and determine the complete pole figures as well as volume fraction of important fibers. The sample impacted at highest momentum was then subjected to colloidal silica polishing using Vibromet for 12 h. Orientation Imaging Microscopy was carried out on a field emission gun scanning electron microscope Hitachi SU6600 fitted with Oxford Instruments electron backscatter diffraction. Data analysis was carried out using hkl Channel 5 data acquisition and analysis software.
3. Results and discussion The initial material in the T8 aged condition was characterized by presence of elongated grains and second phase particles which
184
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Table 1 Experimental conditions and volume fraction of important texture fibers for the impacted AA 2099-T8 aluminum alloy. Strain rate (s1)
Impact momentum (kg m/s)
0 28 31 33 39
Maximum flow stress (MPa)
As received material 4878 7494 8706 8903
Engineering strain
390 370 350 360
Fiber component (%)
0.4257 0.6540 0.7598 0.7770
(a)
h1 0 1i
h1 1 1i
8.4 10.9 7.2 8.2 6.6
27.1 20.9 25.2 29.1 39.9
5.0 19.8 15.5 10.2 9.2
(b)
400
Stress (MPa)
h1 0 0i
300
200
4878 s-1 7494 s-1 8706 s-1 8903 s-1
100
500 µm 0 0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
Strain
I
II
III
(c)
500 µm
(d)
25 µm
Fig. 1. (a) Optical micrograph of the initial AA 2099-T8 alloy (b) stress–strain response of the alloy impacted at different impact momentum and evolution of shear band in sample impacted at 8903 s1 at (c) low magnification and (d) high magnification.
are distributed uniformly throughout the matrix (Fig. 1a). In the peak-aged condition the alloy will also contain T1 (Al2CuLi) and d0 (Al3Li) precipitates [4,5]. In addition, other phases like T2 (Al6CuLi3), S0 (Al2CuMg), b0 (Al3Zr) and d (AlLi) are also present in AA 2099-T8 alloy. The stress–strain curves of the samples impacted at different momentum are shown in Fig. 1b. The stress–strain curves are characterized by the presence of multiple peaks for all the samples. The sample deformed at the highest strain rate shows three distinct peaks compared to two peaks for the remaining samples. The bulk texture measurement shows an increase in h1 0 1i component that is a characteristic FCC compression texture component [8] with increase in impact pressure (Table 1). The microstructure of the deformed samples is characterized by the presence of multiple shear bands devoid of large second phase particles (Fig. 1c). However, SEM image at higher magnification (Fig. 1d) indicated the presence of spherical second phase particles of varying sizes within the shear bands. The second phase particles inside the shear bands are much smaller than those observed outside the shear band region. However, no information was obtained about the microstructure of the matrix phase from the SEM studies
and hence Electron Back Scatter Diffraction was used to characterize the microstructure. In order to decipher the heterogeneity in deformation of impacted 2099 alloy, we monitor the evolution of microstructure and microtexture away from the shear band (region I), near the shear band (region II) and within the shear band (region III) for the sample impacted at the highest pressure (Fig. 1c). Fig. 2 shows the pattern quality, inverse pole figure (IPF) map and inverse pole figure for the sample impacted at highest momentum in the aforementioned three regions. The microstructure of region I show elongated grains with significant orientation gradients within individual grain. The black region in the IPF map represents the location of the second phase particles that came out during colloidal silica polishing. The IPF map of a region III (Fig. 2) indicates presence of elongated grains as well as equiaxed grains and the grains are an order of magnitude smaller than the region far away from the shear band. The equiaxed grains show lower intragranular misorientation and form in clusters. The IPF map of region II shows a combination of very large elongated grains and bands of fine grains. Another important observation is the wavy nature of boundaries in this region. In addition, a
N.P. Gurao et al. / Journal of Alloys and Compounds 578 (2013) 183–187
185
Fig. 2. Pattern quality (top) and inverse pole figure (IPF) map (bottom) in the three different region of the sample deformed at the highest impact momentum. The inverse pole figure in compression direction (going into the plane of the paper) is also shown.
heterogeneous evolution of microtexture is observed in the sample. Region I shows a characteristic h1 0 1i fiber component and region III shows a h1 0 0i component while clustering of orientation between h1 1 1i and h1 0 1i line is observed for region II. The misorientation distribution obtained from the three different regions is shown in Fig. 3a. The region far away from shear band is characterized by high fraction of low angle grain boundaries while the region within shear band shows high fraction of high angle boundaries. The EBSD data was used to describe substructure evolution in terms of Local Average Misorientation (LAM) parameter in the three regions (Fig. 3b). Region I shows very high fraction of low misorientation and a higher value of LAM value. Region II shows lower fraction of low angle misorientation as well as lower average LAM value than Region I. Region III shows very low fraction of low angle LAM and the lowest average LAM. The partitioning of LAM into low and high misorientation in Region I is attributed to evolution of heterogeneous dislocation structure comprising of low and high dislocation density regions in the microstructure [9]. This heterogeneity reduces gradually from regions I to II and the region within the shear band shows minimum LAM and hence minimum dislocation density. Thus, the microstructure as well as microtexture evolution indicates the presence of recrystallization within the shear band. In order to understand the heterogeneity of deformation process at high strain rate, the origin of adiabatic shear bands has to be understood. During dynamic deformation heat gets localized
in particular region of the sample leading to formation of adiabatic shear bands [10–12]. The absence of large precipitates in the shear bands clearly indicates that the temperature reached the level above their dissolution temperature [13]. However, the presence of smaller spherical particles in this region indicates that the dissolution of the second phase particles was not complete during deformation and subsequent cooling. The dissolution of the second phase particles occurs via diffusion of copper atoms. The diffusion rate is higher in regions with sharp curvature and hence higher chemical potential according to Gibbs Thompson equation [14]. Hence there is change in shape of the second phase particles from plate like to spherical shape inside the shear band. The increase in Cu content of the matrix phase from 2.38 wt% in region I to 2.76 wt% in region III as determined from Energy Dispersive Spectroscopy corroborates the enrichment of the matrix phase by precipitate dissolution. Higher temperature that is expected within the shear band, also aids in recrystallization of the microstructure. However, new grains formed within the shear band are not expected to form by classical discontinuous dynamic recrystallization (dDRX) as the time available for recrystallization is inadequate for nucleation and growth of the recrystallized grains. Instead the material is expected to undergo continuous or geometric/rotational dynamic recrystallization [15]. The structure of various coherent precipitates and nanoscale dispersoids within the shear band requires extensive transmission electron microscopy investigation and is beyond the scope of this study. Coherent precipitates
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(b) 0.12 RI RII RIII
0.8
RI RII RIII
0.10 0.08
0.6
Fraction
Cumulative fraction
(a) 1.0
0.4
0.06 0.04
0.2
0.02 0.00
0.0 0
10
20
30
40
50
0
60
Misorientation°
1
2
3
4
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Fig. 3. (a) Cumulative distribution of misorientation and (b) local average misorientation distribution in the three distinct region of the AA 2099-T8 aluminum specimen impacted at the highest momentum.
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Fig. 4. Taylor factor map for the impacted AA 2099 sample in different regions and misorientation distribution within large grains near the interface.
like T1 and T2 with dimensions of the order of few hundred of nanometer in one dimension are neither expected to dissolve nor play any role in particle stimulated nucleation during high strain rate deformation. The presence of large second phase particles in the artificially aged alloy can be expected to act as initiators for dynamic recrystallization (particle stimulated nucleation PSN DRX). However, the h1 0 0i texture observed within the shear band negates the occurrence of the same as it leads to randomization of texture. The presence of multiple peaks in the stress–strain curve (Fig. 1b) can be attributed to the occurrence of multiple dynamic
recrystallization processes in the high strain high strain rate process [16]. High strain rate deformation is characterized by grain fragmentation and higher number of geometrically necessary dislocations to maintain strain compatibility. During the course of deformation, these dislocations arrange themselves to form elongated sub grain boundaries. The misorientation of the sub-grain boundaries increases with the amount of deformation and new geometrically necessary boundaries are generated within these elongated sub-grains due to different regions of the same grain rotating in different direction to maintain strain compatibility. Gradually these sub-grain boundaries can pinch off to form newly
N.P. Gurao et al. / Journal of Alloys and Compounds 578 (2013) 183–187
recrystallized grains with lower dislocation density. The elimination of sub-grain boundaries and limited possibility of formation of high misorientation mobile grain boundary leads to smaller grain size at location of the shear band. The aforementioned mechanism is termed as geometric dynamic recrystallization (gDRX) [17]. The multiple serrations in the stress–strain curve as well as the IPF map of region I with some flattened grains (blue) and IPF map of region III (Fig. 2) with cluster of grains having slight misorientation substantiate this argument. Thus the multiple peaks in the stress–strain curve indicate towards multiple gDRX events. However, a closer look at the IPF map of the interface region II indicates very high misorientation forming in some grains. These regions are formed by grain boundaries swiping sub-grain boundaries during deformation and indicate the presence of continuous dynamic recrystallization (cDRX) [18]. Fig. 4a shows the Taylor factor map of regions I and II. Grains with high Taylor factor (hard orientations) show wavy boundaries in both the region and are expected to act as precursor for gDRX. However, misorientation profile drawn in large grains near the interface in region II clearly show low point to point misorientation accompanied with grain subdivision by geometrically necessary boundaries. Such grain subdivision is due to local lattice rotation along initial grain boundaries aided by formation of geometrically necessary boundaries to accommodate heterogeneous deformation by cDRX [21]. We therefore, believe that the initial reduction of grain size is by continuous dynamic recrystallization that is followed by geometric dynamic recrystallization in the newly formed grains. Another possibility is that cDRX operates in coarse grains with low aspect ratio or in geometrically softer grains while gDRX occurs in smaller grains with high aspect ratio or higher Taylor factor grains. However, the former mechanism wherein, gDRX is preceded by cDRX seems to be dominant as conventional recrystallization texture component h1 0 0i is dominant in the shear band area. The texture is not expected to change qualitatively during cDRX [19–20], however, a characteristically different texture from the deformation texture indicates the presence of gDRX or dDRX. Since the time scale associated with high strain rate deformation is very small, the latter is not expected to occur in the present case. Etter et al. [22] have shown the presence of cDRX and gDRX in 5251 aluminum alloy in the cold rolled and annealed condition respectively. Therefore, it is proposed that due to the difference in deformation state in the impacted sample in terms of difference in Taylor factor and grain size/aspect ratio, different mechanisms can be operative in different regions of the same sample.
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4. Conclusions High strain rate deformation of AA 2099 aluminum alloy in T8 temper condition leads to profuse shear banding that contributes to heterogeneous microstructure and microtexture evolution. The presence of micron sized grains with h1 0 0i texture within the shear band indicate the presence of geometric dynamic recrystallization while the interface between the shear band and sample shows wavy boundaries indicating continuous dynamic recrystallization. Thus two competing dynamic recrystallization mechanisms namely the cDRX and gDRX occur concomitantly during dynamic deformation of aluminum–lithium alloy. References [1] R.J. Rioja, J. Liu, Metall. Mater. Trans. A 43 (2012) 3325–3337. [2] V.K. Jain, K.V. Jata, R.J. Rioja, J.T. Morgan, A.K. Hopkins, J. Mater. Proc. Technol. 73 (1998) 108–118. [3] A.A. Csontos, E.A. Starke Jr., Metall. Mater. Trans. A 31 (2000) 1965. [4] S.C. Huang, M.J. Starink, Int. Mater. Rev. 50 (2005) 193. [5] R. Yoshimura, T.J. Konno, E. Abe, K. Hiraga, Transmission electron microscopy study of the evolution of precipitates in aged Al–Li–Cu alloys: the h0 and T1, Acta Mater. 51 (2003) 4251–4266. [6] R. J. McDonald, Characterization of delamination in 2099-T861 Aluminum– Lithium, PhD thesis, University of Illinois at Urbana-Champaign (2009). [7] TexTools Software, Resmat Corporation, Montreal, QC, Canada. [8] U.F. Kocks, C.N. Tome, H.-R. Wenk, Texture and Anisotropy: Preferred Orientation in Polycrystals and their Effect on Materials Properties, Cambridge University Press, 1998. [9] Y. Zhong, F. Yin, T. Sakaguchi, K. Nagai, K. Yang, Acta Mater. 55 (2007) 2747– 2756. [10] S.E. Schoenfeld, T.W. Wright, Int. J. Solids Struct. 40 (2003) 3021–3037. [11] C. Zener, J.H. Hollomon, J. Appl. Phys. 15 (1944) 22–32. [12] H.C. Rogers, Annu. Rev. Mater. Sci. 9 (1979) 283–311. [13] R.W. Chen, K.S. Vecchio, J. De Phys. IV 4 (1994) 459–464. [14] D.A. Porter, K.E. Esterling, Phase transformations in metals and alloys, London, Chapman and Hall, p. 514. [15] F.J. Humphreys, M. Hutchinson, Recrystallization and Related Annealing Phenomena, Elsevier Ltd., p. 451. [16] U. Andrade, M.A. Meyers, K.S. Vecchio, A.H. Chokshi, Acta Metall. Mater. 42 (1994) 3183–3195. [17] M.A. Meyers, V.F. Nesterenko, J.C. LaSalvia, Q. Xue, Mater. Sci. Eng. A 317 (2001) 204–225. [18] T.G. Nieh, L.M. Hsiung, J. Wadsworth, R. Kaibyshev, Acta Mater. 46 (1998) 2789–2800. [19] W. Skrotzki, N. Scheerbaum, C. Oertel, R.A. Massion, S. Suwas, L.S. Toth, Acta Mater. 55 (2007) 2013–2024. [20] N.P. Gurao, R. Kapoor, S. Suwas, Metall. Mater. Trans. A 41 (2010) 2794– 2804. [21] R. Kaibyshev, O. Stdikov, A. Goloborodko, T. Sakai, Mater. Sci. Eng. A 344 (2003) 348–356. [22] A.L. Etter, T. Baudin, N. Fredj, R. Penelle, Mater. Sci. Eng. A 445–446 (2007) 94– 99.