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Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading
Engineering Failure Analysis xxx (2013) xxx–xxx
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Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading A.G. Odeshi a,⇑, A.O. Adesola a, A.Y. Badmos b a b
University of Saskatchewan, Saskatoon, Canada Black Hawk College, Moline, USA
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Aluminum alloys Precipitation hardening Dynamic shock loading Adiabatic shear bands Fracture
a b s t r a c t Microstructural aspects of the deformation and failure of AA 6061 and AA 2099 aluminum alloys under dynamic impact loading are investigated and compared with their responses to quasi-static mechanical loading in compression. Cylindrical specimens of the alloys, heat-treated to T4, T6 and T8 tempers, were subjected to dynamic compressive loading at strain rates of between 2800 and 9200 s1 and quasi-static compressive loading at a strain rate of 0.0032 s1. Plastic deformation under the dynamic impact loading is dominated by thermal softening leading to formation of adiabatic shear bands. Both deformed and transformed shear bands were observed in the two alloys. The shear bands offer preferential crack initiation site and crack propagation path in the alloys during impact loading leading to ductile shear fracture. While cracks propagate along the central region of transformed bands in AA 6061 alloy, the AA 2099 alloy failed by cracks that propagate preferentially along the boundary region between the transformed shear bands and the bulk material. Whereas the AA 2099 alloy shows the greatest propensity for adiabatic shear banding and failure in the T8 temper condition, AA 6061 alloy is most susceptible to formation of adiabatic shear bands and failure in the T4 temper. Deformation under quasi-static loading is dominated by strain hardening in the two alloys. Rate of strain hardening is higher for naturally aged AA 6061 than the artificially aged alloy, while the strain hardening rate for the AA 2099 alloy is independent of the temper condition. The AA 2099 alloy shows a superior mechanical behaviour under quasi-static compressive loading whereas the AA 6061 shows a higher resistance to impact damage. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Aluminum alloys are very popular for automotive and aerospace applications as a result of their high specific strength, good fracture toughness, excellent corrosion resistance and good formability. These alloys can encounter dynamic shock loading as in the case of automobile crash or bird strike on aircraft’s fuselage or wings. Heat generation and retention along narrow paths during such severe and rapid loading will lead to local thermal softening and mechanical instability that result in intense strain localization along narrow bands called adiabatic shear bands [1–3]. Deformed shear bands (DSBs) consisting of elongated grains and transformed shear bands (TSBs) consisting of ultrafine equi-axed grains are the two types of adiabatic shear bands commonly observed in metals subjected to dynamic shock loading. In most cases, strain localization begins with the formation of deformed bands which change into transformed bands as strain intensity increases [4].
⇑ Corresponding author. Tel.: +1 306 966 5118; fax: +1 306 966 5427. E-mail address: [email protected] (A.G. Odeshi). 1350-6307/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Although shear bands are typically reported to occur in materials subjected to dynamic shock loadings, however, region of local weakening leading to deformation localization and formation of shear bands was reported by Xu et al. [5] in AA 8090 aluminum–lithium alloy during low cycle fatigue loading. The tendency for deformation localization along shear bands was reported to depend on the degree of softening occurring as a result of shearing of d0 precipitates in the alloy by moving dislocations. Fatigue cracks eventually propagate more rapidly along the shear bands leading to a shortened fatigue life. Shear bands have also been reported to occur during fatigue damage of Zr52.5Cu17.9Ni14.6Al10.0Ti5.0 bulk metallic glass [6]. The difference in the strength of the very soft shear bands and its rigid surrounding will generate tensile stresses inside the shear bands, which can open up microvoids that coalesce and initiate cracking along the shear bands [7,8]. FCC metals such as aluminum with high stacking fault energies are reported to show lower susceptibility to adiabatic shear banding during cold rolling than those with low stacking fault energy such as copper [9–12]. A copper–silicon alloy that contained 5.7 atom% silicon showed lower tendency to formation of shear band than the alloy containing 8.8 atom% Si with lower stacking fault energy [10]. Kamijo and Fujiwara [13] observed that shear band did not occur in cold-rolled high purity aluminum until the percentage reduction was above 90%. The occurrence of the shear bands in the high purity aluminum was reported to be strongly dependent on rolling parameters. When annealed, recrystallization was observed to occur preferentially along the shear bands, underscoring the high level of strain intensity inside the shear bands. Textural and geometrical softening was suggested by Asaro and Needleman [14] as the cause of adiabatic shear banding in FCC metals during cold rolling. Despite the high stacking fault energy of aluminum, its alloys have shown considerable susceptibility to strain localization and formation of adiabatic shear bands when subjected to dynamic shock loading [15–19]. The propensity of aluminum alloys to form adiabatic shear bands and experience adiabatic shear failure at high strain rates is determined by the alloy composition, temper condition, and crystal orientation [15–19]. Jia et al. [15] reported that the formation of shear bands in FCC single crystals depends on the initial orientation of the crystals. Morii et al. [16] observed their occurrence in single crystals of Al–Mg alloys to be influenced by the magnesium content and crystal orientation. Shear 1 1 and not in ð0 1 1Þ½1 0 0 Al–Mg crystals. Pérez-Bergquist et al. [17] compared strain localbands were observed in ð2 1 1Þ½1 ization in AA 5059, 5083 and 7039 alloys, which are used in light armour vehicles, and observed susceptibility to shear strain localization and formation of shear bands to be highest in the 7039 alloy. The 7039 alloy has significantly higher zinc and lower magnesium content compared to the other two alloys. TEM investigation of the cores and the outer regions of the shear bands in the three alloys showed that they consist of fine equi-axed and highly elongated dislocation cells, respectively. In another study by Leech [18], it was observed that the properties of the shear band in AA 7039 depend on the original microstructure of the alloy before dynamic shock loading. A comparative study of adiabatic shear banding in AA 2090 and weldalite 049™ Al–Li alloys with that in Al–Mg–Zn (AA 7039) by Lee et al. [19] showed the 7039 to be more susceptible to adiabatic shear bands. Another comparative study on strain localization in an Al–Mg–Si–Cu (AA 6061-T6) alloy and an Al– Mg–Mn (5083-H131) alloy by Odeshi et al. [20] showed the AA 5083 to have a lower tendency for strain localization and adiabatic shear failure under the same impact conditions. In the current study, a comparative evaluation of the effects of the original microstructure and strain rates on adiabatic shear failure of AA 6061 and 2099 aluminum alloys is done. The AA 6061 is an Al–Mg–Si–Cu alloy which finds wide-ranging structural applications owing to its high specific strength, good fracture toughness and excellent corrosion resistance. The AA 2099 is a commercial Al–Li alloy containing significant amounts of Cu and Zn. This alloy was developed for high performance aerospace application and is characterised by high specific modulus, excellent fatigue and cryogenic properties. The effects of temper condition on the microstructural evolution associated with the occurrence of adiabatic shear bands and failure in the alloys are discussed. The mechanical responses of the two alloys under quasi-static and dynamic shock loading are compared.
2. Materials and methods Both AA 6061 and AA 2099 aluminum alloys investigated in this study contain Mg and Cu alloy additions. AA 6061 contains silicon and no lithium while AA 2099 contains Li and no silicon. The composition ranges for the two alloys are presented in Table 1. The AA 6061 alloy was received in T6 temper condition while AA 2099 was supplied in T8 temper. The
Table 1 Chemical composition range for AA 6061 and AA 2099 aluminum alloys. Element
AA 6061 (wt.%)
AA 2099 (wt.%)
Al Li Cu Mg Si Zn Mn Zr Cr
>95 – 0.15–0.40 0.8–1.2 0.4–0.8 <0.25 <0.15 – 0.04–0.35
>95 1.6–2.0 2.4–3.0 0.1–0.5 – 0.4–1.0 0.1–0.5 0.05–0.12 –
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as-received wrought alloys were machined into cylindrical test specimens, 9.5 mm in diameter and 10.5 mm long. A third of the AA 6061 specimens were left in the as-received T6 temper while a third each were heat treated to T4 and T8 temper conditions. The T4 temper specimens were solution heat-treated at 540 °C for 2 h, followed by water quenching and natural aging at room temperature. The T8 temper was obtained by solution treatment at 540 °C for 2 h and water-quenching, followed by room temperature aging for 48 h, 17% cold work, and finally, aging at 160 °C for 18 h. Similarly, a third of AA 2099 specimens were tested in the as-received T8 temper while a third each were heat treated to have T4 and T6 temper. Optimum mechanical properties are achievable in AA 2099 alloys when subjected to two-step aging [21]. Therefore, the precipitation heat treatment for the AA 2099 to obtain the T6 temper was carried out in two steps as in the as-received T8 sample. The treatments involved the solution treatment at 540 °C for 2 h, followed by quenching, natural aging for 48 h and two steps artificial aging. The first step of the artificial aging treatment involved soaking at 120 °C for 12 h followed immediately by a second-step aging at 160 °C for 24 h. Some of the test specimens were subjected to quasi-static compressive loading at a strain rate of 0.0032 s1 using Instron R5500 series load frame up to a maximum load of 100 kN. The dynamic mechanical test was conducted using direct impact Hopkinson bar, in which the test specimens were impacted by a blunt steel projectile at momentum ranging between 28 and 44 kg m/s. This test system is described elsewhere [22]. Detailed microstructural studies of the specimen before and after the various mechanical tests were carried out using optical and scanning electron microscopy. Fig. 1 shows typical optical and scanning electron micrographs of the alloys in the as-received conditions before impact. Second phase particles in the AA 6061 alloy are shown to be randomly distributed while those in the AA 2099 alloys are aligned along the rolling direction of the plate. 3. Results and discussion Fig. 2 represents typical stress strain curves obtained for the alloys under impact and quasi-static compressive loading. Under the quasi-static loading, plastic deformation is controlled by strain hardening, and stress increased continually with strain as the applied load increased. Under the impact loading, the plastic deformation is controlled by simultaneous actions of thermal softening and strain hardening. Thermal softening is a consequence of conversion of about 90% of the kinetic energy of the projectile into heat energy while the remaining 10% is expended on deforming the material [23]. Strain hardening which initially dominates the plastic deformation is quickly overshadowed by thermal softening, resulting in stress decrease with further increase in strain. Eventually adiabatic heating occurs, leading to intense localized thermal softening and mechanical instabilities that triggers the occurrence of adiabatic shear bands. Figs. 3 and 4 show the effects of temper
Fig. 1. Optical and SEM micrographs of the as-received (a) AA 2099-T8 and (b) 6061-T6 aluminum alloys showing the morphology and distribution of the second phase particles within the continuous aluminum-rich phase.
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Fig. 2. Stress–strain curves for 2099-T4 alloy under quasi-static and dynamic impact loading.
Fig. 3. Stress–strain curves (a and b) and strain hardening rate vs. strain curves (c and d) for AA 6061 and AA 2099 aluminum alloys under quasi-static loading at a strain rate of 3.2 103/s.
condition on the stress–strain responses of the alloys under quasi-static loading. For both alloys, the T8 specimens have the highest yield strength while the T4 alloys have the lowest yield strength. The AA 2099 specimens showed higher yield strength and lower strain compared to the AA 6061 specimens under quasi-static compressive loading. The higher yield strength of the artificially aged specimens with or without cold work (T6 and T8) compared to the naturally aged alloys (T4) is due to the formation of fine precipitates during artificial aging, which hinder dislocation motion and result in higher resistance to yielding. The precipitates contributing majorly to strengthening in AA 6061 alloy in peak-aged
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Fig. 4. Effect of temper condition on yield strength and the engineering strain (corresponding to a compressive load of 100 kN) of AA 6061 and 2099 aluminum alloys.
Fig. 5. Typical dynamic impact stress–strain curves for AA 6061 and 2099 aluminum alloys.
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condition is b00 (Mg5Si6), which is needle-like and coherent with the continuous phase along the needle axis [24,25]. In a slightly over-aged condition, semi-coherent rod-like b0 (Mg2Si) precipitates form and contribute to higher strength of the artificially aged 6061 alloys [25,26]. Other precipitates that can contribute to higher strength of AA 6061 aluminum alloy in T6 and T8 temper have been identified to include B0 (Mg9Al3Si7), U1 (MgAl2Si2) and U2 (MgAlSi) and h0 (CuAl2) [24–26]. Aluminum 2099 belongs to the Al–Li–Cu–Mg system, in which d0 (Al3Li) particles are precipitated as the strengthening phase. This phase can remain coherent even after extensive aging [27,28]. Coarsening of the d0 phase with time at the artificial aging temperature leads to the formation of d (AlLi) equilibrium phase [27]. The presence of magnesium, copper and zirconium results in precipitation of other phases in addition to the d0 phase. These phases include T1 (Al2CuLi), T2 (Al6CuLi3), h0 (Al2Cu), b0 (Al3Zr) and S (Al2CuMg) [29]. The formation of these coherent or semi-coherent precipitates during artificial aging will impose high resistance to dislocation motion, resulting in the observed higher yield strength of the artificially aged AA 2099 alloy. The dislocations generated during strain hardening before precipitation hardening of AA 2099-T8 alloy act as nucleation sites for T1 and S precipitates [27,30]. The higher strength of 2099-T8, compared to AA 2099-T6 under quasi-static loading, can be attributed to nucleation of more precipitates in the T8 alloy resulting from high dislocation density at the onset of artificial aging in addition to the strain hardening effect of the intermediate cold-work. Although the yield stress of the AA 6061-T6 alloy is greater than that of the AA 6061-T4 alloy, the strain corresponding to the applied compressive load of 100 kN is unexpectedly higher for the AA 6061-T6 alloy. This is attributable to a higher strain hardening rate in the T4 alloy under the quasi-static loading in compression as shown in Fig. 3c. The same observation was reported in a similar Al–Mg–Si alloy (AA 6022) in which strain hardening rate is higher in the naturally aged (T4 temper) than in the artificially aged conditions [31]. On the contrary, there is no difference in the strain hardening rate for the AA 2099 alloy in the three temper conditions (Fig. 3d). Consequently, the total strain of the AA 2099-T6 alloy with higher yield strength is lower than that of the weaker AA 2099-T4 alloy, as expected. The cause of the higher strain hardening rate in the AA 6061-T4 alloy is not very clear, but may not be unconnected with the high concentration of as-quenched vacancies in Al– Mg–Si alloys, which become mobile during plastic deformation [32,33]. These vacancies can collapse forming loops which generate new dislocations during plastic deformation [33,34] of the as-quenched AA 6061 alloy (T4) and enhance its strain hardening capability as observed in this study. Fig. 5 shows the effects of aging condition on the dynamic stress–strain curves of the AA 6061 and AA 2099 alloys under impact loading. Although the AA 2099 has higher strength under quasi-static load, it experiences more thermal softening and consequently lower deformation resistance under impact loading. Comparing the maximum flow stress attained in the dynamic stress–strain curves and the critical strains for stress collapse, the AA 6061 shows a lower susceptibility to strain localization and adiabatic shear banding than the AA 2099 alloy. Whereas the dynamic stress–strain curves suggest that the AA 6061 alloy will have higher susceptibility to thermal softening and adiabatic shear banding in the T4 condition than in the T8 temper, the reverse appears to be the case for the AA 2099 alloy in which T4 alloy showed the least susceptibility. The results of the dynamic impact response of the alloys, as influenced by their temper conditions, are summarised in Table 2. Deformed shear bands (DSBs) and/or transformed shear bands (TSBs) developed in the alloys under impact load depending on the temper condition and impact momentum. DSBs develop at low impact momentums, and consist of second phase particles that
Table 2 Dynamic impact response of the investigated AA 2099 and AA 6061 aluminum alloys. Alloy
Impact momentum (kg m/s)
Total strain
Peak flow stress (MPa)
Strain rate (s1)
ASB
2099-T4
28 31 33 39
0.32 0.50 0.54 0.69
315 350 372 380
3630 5723 6238 7945
None None Deformed Transformed
2099-T6
28 31 33 39
0.33 0.44 0.68 0.74
378 380 330 410
3812 4984 7773 8497
Deformed Deformed Transformed Transformed
2099-T8
28 31 33 39
0.25 0.50 0.68 0.73
420 445 448 442
2889 5723 7820 8390
Transformed Transformed Transformed Transformed
6061-T4
33 39 44
0.61 0.69 0.79
390 560 336
6973 7905 9096
Transformed Transformed Transformed
6061-T6
33 39 44
0.59 0.70 0.80
430 571 336
6754 7971 9062
Deformed Transformed Transformed
6061-T8
33 39 44
0.54 0.65 0.75
401 482 502
6164 7426 8540
Deformed Transformed Transformed
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align in shear flow directions (Fig. 6). They were also observed ahead of TSB propagating paths in specimens impacted at high momentums. This suggests that the TSBs develop from the initially formed DSBs. Fig. 7 shows the impact momentum at which TSB was first observed in the impacted specimens. The minimum impact momentum for the formation of transformed band is lower for AA 2099 than for AA 6061. This confirms a higher propensity for formation of TSB in the AA 2099 alloy. Optical macrographs showing the impacted surface and side-view of some impacted specimens (Fig. 8) also confirms higher propensity of AA 2099 alloy to form transformed shear bands. The development of TSBs in the alloys during impact caused a change in the geometry of the transverse section of the impacted specimens from circular to elliptical as the impact momentum was increased. The elliptical and V-shape trajectory of the TSBs on the respective transverse and longitudinal sections of the impacted specimens are observable by visual inspection (Fig. 8). Visual inspections of the impacted specimens also indicate that the temper conditions that are most susceptible to adiabatic shear failure in AA 6061 and AA 2099 alloys are T4 and T8, respectively. In order to show the differences in the degree of adiabatic heating inside the shear bands that formed in the two investigated alloys, the temperature increase inside the shear bands was estimated using the following equation [1]:
DT ¼
b qC v
Z
rde
ð1Þ
where DT is temperature rise inside the transformed shear bands; b is fraction of the impact energy that is converted to thermal energy (taken as 0.9) [23]; q and Cv are densities and specific heat capacities of the alloys, respectively. Fig. 9 shows the
Fig. 6. Optical micrographs showing (a) deformed shear bands (DSBs) adjacent to transformed shear bands (TSBs) in an impacted AA 6061 alloy and (b) deformed shear bands in an impacted AA 2099 alloy.
Fig. 7. The impact momentum at which transformed shear bands were first observed in the alloys as the impact momentum of the projectile was increased from 28 kg m/s.
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Fig. 8. Optical macrographs showing changes in the geometry of the impacted specimens with increase in impact momentum.
a
250
b 250
AA 2099
Temperature (°C)
Temperature (°C)
150
T8 T6 T4
100
50
0
AA 6061
200
200
150
T4 T6 T8
100
50
0
26
28
30
32
34
36
Impact momentum (kg m/s)
38
40
0
30 0
32
34
36
38
40
42
44
46
Impact Momentum (kg m/s)
Fig. 9. Estimated temperature inside the transformed shear bands as a function of the impact momentum and temper condition for (a) AA 2099 and (b) AA 6061 aluminum alloys.
estimated relative temperature rise inside the transformed shear bands for various impacted specimens. For the AA 6061 alloy, the temperature rise inside adiabatic shear bands is highest for the alloy in T4 condition and lowest in the T8 condition. In AA 2099, the temperature rise is highest in the T8 condition and least in the T4 condition. This shows that the susceptibility of aluminum alloys has a direct correlation with the temperature rise leading to localized thermal softening. This temperature rise is in turn dependent on the applied momentum, alloy composition and the temper condition of the aluminum alloys. The temperature rise was observed to increase with the applied momentum for all the temper conditions. Typical transformed shear bands (TSBs) observed in the optical micrographs of the specimens subjected to the maximum applied impact momentum are presented in Fig. 10. The TSBs are characteristically featureless under an optical microscope. They appear white or dark depending on the intensity of etching during preparation. Heavily etched TSBs appear darker. The results of the scanning electron microscopic investigations of the impacted specimens reveal more detailed information on the difference in the morphology of the microstructures in the bulk materials and the TSBs. The results of the SEM study show that the second phase particles in the AA 6061 and AA 2099 alloys outside the TSBs have rectangular and elliptical shapes respectively (Figs. 11 and 12). As in optical microscopy, the microstructures of the TSBs in the AA 2099 alloy could not be resolved under the SEM except that the SEM images indicate that dissolution of the second phase particles occurred inside the TSBs. Fine spherical particles with average sizes of less than 2 lm were observed at the central regions of the TSBs in the AA 2099-T8 alloy (Fig. 11). The TSBs are devoid of the much bigger elliptical second phase particles that are observed in the bulk materials. They apparently dissolve as a result of the high intensity of adiabatic heating and severe strains within the TSBs during impact. Carbide dissolution was similarly reported to occur in TSBs during shear strain localization in AISI 4340 steel [7]. Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Fig. 10. Optical micrographs showing the transformed shear bands observed in the AA 6061 and AA 2099 impacted at the maximum applied momentum of 44 and 39 kg m/s respectively.
Fig. 11. SEM micrographs showing transformed shear bands (TSBs) in impacted AA 2099-T8 aluminum alloy.
The fine spherical particles at the centre of the TSBs in AA 2099-T8 could be the remaining portion of incompletely dissolved second phase particles or newly recrystallized grains. Because these grains are observed at the centre of the transformed bands where the peak temperature is expected to be highest, we believe these grains to be new dynamically recrystallized grains. Moreover, these grains do not appear to be as brittle as the second phase particles observed in the bulk materials. The recrystallization occurring simultaneously with deformation at elevated temperatures is termed dynamic recrystallization. Beside the few new equi-axed grains observed at the centre of the transformed bands in the AA 2099T8 alloy, the microstructure of the continuous phase inside the transformed bands could not be resolved under the scanning electron microscope used in this study. Transformed shear bands which appear featureless under an optical microscope and low resolution SEM have been reported by several authors [1–3] to consist of nano-sized grains when observed under a Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Fig. 12. SEM micrographs showing transformed shear bands (TSBs) in impacted AA 6061-T8 aluminum alloy.
transmission electron microscope. We believe that the size of such recrystallized grains in the predominantly featureless areas of the transformed bands in the AA 2099 alloy is beyond the resolution limit of the SEM. Unlike in the case of the AA 2099 alloy, the microstructure of the TSBs in the AA 6061 alloy was resolvable using the scanning electron microscope, and they consists of fine equi-axed grains with an average size of 0.6 lm (Fig. 12). For the AA 6061-T8 alloy, the equi-axed particles are located at the centre of the transformed bands while the peripheral regions consist of a continuous phase with minimal or no second phase particles or recrystallized grains. Such second phase particle-free regions adjacent to TSBs are not noticeable in the T4 and T6 alloys. There are two possible mechanisms for the formation of the equi-axed grains inside the TSBs. One possibility is the fragmentation of the second phase particles occurring inside the shear bands as a result of these particles crushing against one another due to hydrostatic pressure from the rigid walls of the shear bands during impact. Grain elongation followed by fragmentation of pearlite lamellae followed by spheroidization has equally been suggested for the formation of fine equi-axed grains in pearlitic steels [35,36]. Fragmentation of the pearlite lamellae was reported to occur as a result of partitioning into transverse cell by dislocation movement under high hydrostatic pressure and shear stress. The plastic flow of the softened surrounding matrix phase could enhance the crushing of the second phase particles against one another while the soft continuous phase flows away from the centre of the transformed bands. However, in an earlier study on adiabatic shear banding in alumina particles reinforced AA 6061-T6 aluminum matrix composite, these hydrostatic pressure only caused a denser packing of the reinforcing alumina particles along the shear bands and no crushing of the alumina particles was reported inside the shear bands [37]. Another possible mechanism for the formation of the new fine equi-axed grain in the impacted AA 6061 alloy, which is more probable, is the dynamic recrystallization as earlier suggested for the fine spherical grains observed in the central region of the TSBs in the AA 2099 alloy. There are two types of recrystallization; static and dynamic recrystallization. While static recrystallization occurs subsequent to plastic deformation, dynamic recrystallization occurs simultaneously with deformation as a material is subjected to mechanical load at elevated temperatures. Derby [38] suggested that dynamic recrystallization can occur either by grain boundary rotational mechanism or grain boundaries migration through the deforming microstructure. The grain size obtained during dynamic recrystallization by migration of high angle grain boundaries during hot working of metals is reported to have stronger dependence on the deformation stress than temperature [38,39]. Dynamic recrystallization by grain boundary rotation is reported to account for formation of ultrafine equi-axed grains inside transformed bands in several metals [1,39,40]. The deformation time during dynamic shock loading is too short to allow grain boundary migration. It is generally believed that the severe strain and adiabatic heating trigger the dynamic recrystallization and the formation of shear bands in metals. A contrary opinion has recently been given by Rittel et al. [41,42], who suggested that strain localization is initiated by local softening caused by these dynamically recrystallized grains. They showed evidence of dynamically recrystallized grains in samples in which TSBs have not been fully developed. Figs. 13 and 14 show cracks propagation along TSBs in impacted specimens of the AA 2099-T8 and 6061-T4, respectively. These are the temper conditions which show the highest propensities for TSB’s formation and cracking for the respective alloys. Cracks initiated at and propagated along the interface between the TSBs and the bulk material in the AA 2099 alloy
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Fig. 13. Crack propagation at the interface between transformed shear band (TSB) and bulk material in AA 2099-T8 alloy.
Fig. 14. Crack propagation inside a transformed shear band (TSB) in AA 6061-T4 alloy.
Fig. 15. SEM micrographs showing the fracture surface of AA 2099-T8 aluminum alloy that fractured under impact loading. Fracture (a) inside the bulk material and (b) along the interface between TSB and the bulk material.
(Fig. 13). In AA 6061 alloy cracks are initiated inside the TSBs and they propagated along the central region of the shear bands (Fig. 14). The TSBs formed in the AA 2099 alloy are more susceptible to cracking than those in AA 6061 alloy. The difference in the visco-plasticity of the soft shear bands and that of its rigid surrounding material during impact loading generated the tensile stresses that opened up voids that coalesced to initiate crack inside the TSBs of the AA 6061 alloy [7,8]. The interfacial cracks between the TSBs and the bulk materials in AA 2099 may be due to a weak interface, which shears easily as the soft material in the TSB undergoes plastic flow under impact loading. The results of the SEM examination of the fracture surface of a fragmented AA 2099 alloy are shown in Fig. 15. The specimen shows a very smooth round fracture surface. These confirm shearing of the specimens around the transformed shear bands. Cleavage fracture of the second phase particles and ridges showing ductile fracture of the continuous aluminum-rich phase are observed on the fracture surface. As indicated in the figure, two regions can be identified on the fractured surface. At the centre of the fracture surface, irregular topography
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showing ridges or striations in the continuous phase can be observed adjacent to the TSBs (Fig. 15b). A less ductile fracture is observed in the bulk material farther away from the TSBs (Fig. 15a). None of the AA 6061 specimens fractured under the investigated impact condition despite the higher impact momentum, to which they were subjected. 4. Conclusions The effects of strain rate and temper condition on strain localization and adiabatic shear failure in AA 6061 and AA 2099 aluminum alloys are investigated and compared. The following conclusions can be drawn from the results obtained: 1. AA 2099 has higher deformation resistance under quasi-static loading whereas AA 6061 alloy is more resistant to deformation under dynamic impact loading. 2. Failure of the alloys under impact loading is preceded by strain localization along transformed shear bands which initially occurred as deformed bands. 3. AA 6061 is most susceptible to adiabatic shear failure in the T4 temper, whereas AA 2099 is most susceptible in the T8 temper condition. 4. Transformed bands in AA 6061 alloy consist of fine recrystallized equi-axed grains with an average size of 0.6 lm while the microstructure of the transformed bands in AA 2099 is not resolvable under the same scanning electron microscope. 5. Cracks initiate and propagate along the interface between the adiabatic shear bands and bulk material in the AA 2099 but along the central region of the shear bands in the AA 6061 alloy. 6. Fragmentation of specimen occurs in AA 2099 under the impact loading while none of the AA 6061 specimens fractured, even at higher impact momentum, suggesting higher cracking susceptibility of AA 2099 under impact loading.
Acknowledgments The authors wish to acknowledge the support of Alcoa incorporation in providing the investigated AA 2099 alloy and Natural Sciences and Engineering Research Council of Canada for their financial support of this study. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
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Contents lists available at SciVerse ScienceDirect
Engineering Failure Analysis journal homepage: www.elsevier.com/locate/engfailanal
Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading A.G. Odeshi a,⇑, A.O. Adesola a, A.Y. Badmos b a b
University of Saskatchewan, Saskatoon, Canada Black Hawk College, Moline, USA
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Aluminum alloys Precipitation hardening Dynamic shock loading Adiabatic shear bands Fracture
a b s t r a c t Microstructural aspects of the deformation and failure of AA 6061 and AA 2099 aluminum alloys under dynamic impact loading are investigated and compared with their responses to quasi-static mechanical loading in compression. Cylindrical specimens of the alloys, heat-treated to T4, T6 and T8 tempers, were subjected to dynamic compressive loading at strain rates of between 2800 and 9200 s1 and quasi-static compressive loading at a strain rate of 0.0032 s1. Plastic deformation under the dynamic impact loading is dominated by thermal softening leading to formation of adiabatic shear bands. Both deformed and transformed shear bands were observed in the two alloys. The shear bands offer preferential crack initiation site and crack propagation path in the alloys during impact loading leading to ductile shear fracture. While cracks propagate along the central region of transformed bands in AA 6061 alloy, the AA 2099 alloy failed by cracks that propagate preferentially along the boundary region between the transformed shear bands and the bulk material. Whereas the AA 2099 alloy shows the greatest propensity for adiabatic shear banding and failure in the T8 temper condition, AA 6061 alloy is most susceptible to formation of adiabatic shear bands and failure in the T4 temper. Deformation under quasi-static loading is dominated by strain hardening in the two alloys. Rate of strain hardening is higher for naturally aged AA 6061 than the artificially aged alloy, while the strain hardening rate for the AA 2099 alloy is independent of the temper condition. The AA 2099 alloy shows a superior mechanical behaviour under quasi-static compressive loading whereas the AA 6061 shows a higher resistance to impact damage. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Aluminum alloys are very popular for automotive and aerospace applications as a result of their high specific strength, good fracture toughness, excellent corrosion resistance and good formability. These alloys can encounter dynamic shock loading as in the case of automobile crash or bird strike on aircraft’s fuselage or wings. Heat generation and retention along narrow paths during such severe and rapid loading will lead to local thermal softening and mechanical instability that result in intense strain localization along narrow bands called adiabatic shear bands [1–3]. Deformed shear bands (DSBs) consisting of elongated grains and transformed shear bands (TSBs) consisting of ultrafine equi-axed grains are the two types of adiabatic shear bands commonly observed in metals subjected to dynamic shock loading. In most cases, strain localization begins with the formation of deformed bands which change into transformed bands as strain intensity increases [4].
⇑ Corresponding author. Tel.: +1 306 966 5118; fax: +1 306 966 5427. E-mail address: [email protected] (A.G. Odeshi). 1350-6307/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Although shear bands are typically reported to occur in materials subjected to dynamic shock loadings, however, region of local weakening leading to deformation localization and formation of shear bands was reported by Xu et al. [5] in AA 8090 aluminum–lithium alloy during low cycle fatigue loading. The tendency for deformation localization along shear bands was reported to depend on the degree of softening occurring as a result of shearing of d0 precipitates in the alloy by moving dislocations. Fatigue cracks eventually propagate more rapidly along the shear bands leading to a shortened fatigue life. Shear bands have also been reported to occur during fatigue damage of Zr52.5Cu17.9Ni14.6Al10.0Ti5.0 bulk metallic glass [6]. The difference in the strength of the very soft shear bands and its rigid surrounding will generate tensile stresses inside the shear bands, which can open up microvoids that coalesce and initiate cracking along the shear bands [7,8]. FCC metals such as aluminum with high stacking fault energies are reported to show lower susceptibility to adiabatic shear banding during cold rolling than those with low stacking fault energy such as copper [9–12]. A copper–silicon alloy that contained 5.7 atom% silicon showed lower tendency to formation of shear band than the alloy containing 8.8 atom% Si with lower stacking fault energy [10]. Kamijo and Fujiwara [13] observed that shear band did not occur in cold-rolled high purity aluminum until the percentage reduction was above 90%. The occurrence of the shear bands in the high purity aluminum was reported to be strongly dependent on rolling parameters. When annealed, recrystallization was observed to occur preferentially along the shear bands, underscoring the high level of strain intensity inside the shear bands. Textural and geometrical softening was suggested by Asaro and Needleman [14] as the cause of adiabatic shear banding in FCC metals during cold rolling. Despite the high stacking fault energy of aluminum, its alloys have shown considerable susceptibility to strain localization and formation of adiabatic shear bands when subjected to dynamic shock loading [15–19]. The propensity of aluminum alloys to form adiabatic shear bands and experience adiabatic shear failure at high strain rates is determined by the alloy composition, temper condition, and crystal orientation [15–19]. Jia et al. [15] reported that the formation of shear bands in FCC single crystals depends on the initial orientation of the crystals. Morii et al. [16] observed their occurrence in single crystals of Al–Mg alloys to be influenced by the magnesium content and crystal orientation. Shear 1 1 and not in ð0 1 1Þ½1 0 0 Al–Mg crystals. Pérez-Bergquist et al. [17] compared strain localbands were observed in ð2 1 1Þ½1 ization in AA 5059, 5083 and 7039 alloys, which are used in light armour vehicles, and observed susceptibility to shear strain localization and formation of shear bands to be highest in the 7039 alloy. The 7039 alloy has significantly higher zinc and lower magnesium content compared to the other two alloys. TEM investigation of the cores and the outer regions of the shear bands in the three alloys showed that they consist of fine equi-axed and highly elongated dislocation cells, respectively. In another study by Leech [18], it was observed that the properties of the shear band in AA 7039 depend on the original microstructure of the alloy before dynamic shock loading. A comparative study of adiabatic shear banding in AA 2090 and weldalite 049™ Al–Li alloys with that in Al–Mg–Zn (AA 7039) by Lee et al. [19] showed the 7039 to be more susceptible to adiabatic shear bands. Another comparative study on strain localization in an Al–Mg–Si–Cu (AA 6061-T6) alloy and an Al– Mg–Mn (5083-H131) alloy by Odeshi et al. [20] showed the AA 5083 to have a lower tendency for strain localization and adiabatic shear failure under the same impact conditions. In the current study, a comparative evaluation of the effects of the original microstructure and strain rates on adiabatic shear failure of AA 6061 and 2099 aluminum alloys is done. The AA 6061 is an Al–Mg–Si–Cu alloy which finds wide-ranging structural applications owing to its high specific strength, good fracture toughness and excellent corrosion resistance. The AA 2099 is a commercial Al–Li alloy containing significant amounts of Cu and Zn. This alloy was developed for high performance aerospace application and is characterised by high specific modulus, excellent fatigue and cryogenic properties. The effects of temper condition on the microstructural evolution associated with the occurrence of adiabatic shear bands and failure in the alloys are discussed. The mechanical responses of the two alloys under quasi-static and dynamic shock loading are compared.
2. Materials and methods Both AA 6061 and AA 2099 aluminum alloys investigated in this study contain Mg and Cu alloy additions. AA 6061 contains silicon and no lithium while AA 2099 contains Li and no silicon. The composition ranges for the two alloys are presented in Table 1. The AA 6061 alloy was received in T6 temper condition while AA 2099 was supplied in T8 temper. The
Table 1 Chemical composition range for AA 6061 and AA 2099 aluminum alloys. Element
AA 6061 (wt.%)
AA 2099 (wt.%)
Al Li Cu Mg Si Zn Mn Zr Cr
>95 – 0.15–0.40 0.8–1.2 0.4–0.8 <0.25 <0.15 – 0.04–0.35
>95 1.6–2.0 2.4–3.0 0.1–0.5 – 0.4–1.0 0.1–0.5 0.05–0.12 –
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as-received wrought alloys were machined into cylindrical test specimens, 9.5 mm in diameter and 10.5 mm long. A third of the AA 6061 specimens were left in the as-received T6 temper while a third each were heat treated to T4 and T8 temper conditions. The T4 temper specimens were solution heat-treated at 540 °C for 2 h, followed by water quenching and natural aging at room temperature. The T8 temper was obtained by solution treatment at 540 °C for 2 h and water-quenching, followed by room temperature aging for 48 h, 17% cold work, and finally, aging at 160 °C for 18 h. Similarly, a third of AA 2099 specimens were tested in the as-received T8 temper while a third each were heat treated to have T4 and T6 temper. Optimum mechanical properties are achievable in AA 2099 alloys when subjected to two-step aging [21]. Therefore, the precipitation heat treatment for the AA 2099 to obtain the T6 temper was carried out in two steps as in the as-received T8 sample. The treatments involved the solution treatment at 540 °C for 2 h, followed by quenching, natural aging for 48 h and two steps artificial aging. The first step of the artificial aging treatment involved soaking at 120 °C for 12 h followed immediately by a second-step aging at 160 °C for 24 h. Some of the test specimens were subjected to quasi-static compressive loading at a strain rate of 0.0032 s1 using Instron R5500 series load frame up to a maximum load of 100 kN. The dynamic mechanical test was conducted using direct impact Hopkinson bar, in which the test specimens were impacted by a blunt steel projectile at momentum ranging between 28 and 44 kg m/s. This test system is described elsewhere [22]. Detailed microstructural studies of the specimen before and after the various mechanical tests were carried out using optical and scanning electron microscopy. Fig. 1 shows typical optical and scanning electron micrographs of the alloys in the as-received conditions before impact. Second phase particles in the AA 6061 alloy are shown to be randomly distributed while those in the AA 2099 alloys are aligned along the rolling direction of the plate. 3. Results and discussion Fig. 2 represents typical stress strain curves obtained for the alloys under impact and quasi-static compressive loading. Under the quasi-static loading, plastic deformation is controlled by strain hardening, and stress increased continually with strain as the applied load increased. Under the impact loading, the plastic deformation is controlled by simultaneous actions of thermal softening and strain hardening. Thermal softening is a consequence of conversion of about 90% of the kinetic energy of the projectile into heat energy while the remaining 10% is expended on deforming the material [23]. Strain hardening which initially dominates the plastic deformation is quickly overshadowed by thermal softening, resulting in stress decrease with further increase in strain. Eventually adiabatic heating occurs, leading to intense localized thermal softening and mechanical instabilities that triggers the occurrence of adiabatic shear bands. Figs. 3 and 4 show the effects of temper
Fig. 1. Optical and SEM micrographs of the as-received (a) AA 2099-T8 and (b) 6061-T6 aluminum alloys showing the morphology and distribution of the second phase particles within the continuous aluminum-rich phase.
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Fig. 2. Stress–strain curves for 2099-T4 alloy under quasi-static and dynamic impact loading.
Fig. 3. Stress–strain curves (a and b) and strain hardening rate vs. strain curves (c and d) for AA 6061 and AA 2099 aluminum alloys under quasi-static loading at a strain rate of 3.2 103/s.
condition on the stress–strain responses of the alloys under quasi-static loading. For both alloys, the T8 specimens have the highest yield strength while the T4 alloys have the lowest yield strength. The AA 2099 specimens showed higher yield strength and lower strain compared to the AA 6061 specimens under quasi-static compressive loading. The higher yield strength of the artificially aged specimens with or without cold work (T6 and T8) compared to the naturally aged alloys (T4) is due to the formation of fine precipitates during artificial aging, which hinder dislocation motion and result in higher resistance to yielding. The precipitates contributing majorly to strengthening in AA 6061 alloy in peak-aged
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Fig. 4. Effect of temper condition on yield strength and the engineering strain (corresponding to a compressive load of 100 kN) of AA 6061 and 2099 aluminum alloys.
Fig. 5. Typical dynamic impact stress–strain curves for AA 6061 and 2099 aluminum alloys.
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condition is b00 (Mg5Si6), which is needle-like and coherent with the continuous phase along the needle axis [24,25]. In a slightly over-aged condition, semi-coherent rod-like b0 (Mg2Si) precipitates form and contribute to higher strength of the artificially aged 6061 alloys [25,26]. Other precipitates that can contribute to higher strength of AA 6061 aluminum alloy in T6 and T8 temper have been identified to include B0 (Mg9Al3Si7), U1 (MgAl2Si2) and U2 (MgAlSi) and h0 (CuAl2) [24–26]. Aluminum 2099 belongs to the Al–Li–Cu–Mg system, in which d0 (Al3Li) particles are precipitated as the strengthening phase. This phase can remain coherent even after extensive aging [27,28]. Coarsening of the d0 phase with time at the artificial aging temperature leads to the formation of d (AlLi) equilibrium phase [27]. The presence of magnesium, copper and zirconium results in precipitation of other phases in addition to the d0 phase. These phases include T1 (Al2CuLi), T2 (Al6CuLi3), h0 (Al2Cu), b0 (Al3Zr) and S (Al2CuMg) [29]. The formation of these coherent or semi-coherent precipitates during artificial aging will impose high resistance to dislocation motion, resulting in the observed higher yield strength of the artificially aged AA 2099 alloy. The dislocations generated during strain hardening before precipitation hardening of AA 2099-T8 alloy act as nucleation sites for T1 and S precipitates [27,30]. The higher strength of 2099-T8, compared to AA 2099-T6 under quasi-static loading, can be attributed to nucleation of more precipitates in the T8 alloy resulting from high dislocation density at the onset of artificial aging in addition to the strain hardening effect of the intermediate cold-work. Although the yield stress of the AA 6061-T6 alloy is greater than that of the AA 6061-T4 alloy, the strain corresponding to the applied compressive load of 100 kN is unexpectedly higher for the AA 6061-T6 alloy. This is attributable to a higher strain hardening rate in the T4 alloy under the quasi-static loading in compression as shown in Fig. 3c. The same observation was reported in a similar Al–Mg–Si alloy (AA 6022) in which strain hardening rate is higher in the naturally aged (T4 temper) than in the artificially aged conditions [31]. On the contrary, there is no difference in the strain hardening rate for the AA 2099 alloy in the three temper conditions (Fig. 3d). Consequently, the total strain of the AA 2099-T6 alloy with higher yield strength is lower than that of the weaker AA 2099-T4 alloy, as expected. The cause of the higher strain hardening rate in the AA 6061-T4 alloy is not very clear, but may not be unconnected with the high concentration of as-quenched vacancies in Al– Mg–Si alloys, which become mobile during plastic deformation [32,33]. These vacancies can collapse forming loops which generate new dislocations during plastic deformation [33,34] of the as-quenched AA 6061 alloy (T4) and enhance its strain hardening capability as observed in this study. Fig. 5 shows the effects of aging condition on the dynamic stress–strain curves of the AA 6061 and AA 2099 alloys under impact loading. Although the AA 2099 has higher strength under quasi-static load, it experiences more thermal softening and consequently lower deformation resistance under impact loading. Comparing the maximum flow stress attained in the dynamic stress–strain curves and the critical strains for stress collapse, the AA 6061 shows a lower susceptibility to strain localization and adiabatic shear banding than the AA 2099 alloy. Whereas the dynamic stress–strain curves suggest that the AA 6061 alloy will have higher susceptibility to thermal softening and adiabatic shear banding in the T4 condition than in the T8 temper, the reverse appears to be the case for the AA 2099 alloy in which T4 alloy showed the least susceptibility. The results of the dynamic impact response of the alloys, as influenced by their temper conditions, are summarised in Table 2. Deformed shear bands (DSBs) and/or transformed shear bands (TSBs) developed in the alloys under impact load depending on the temper condition and impact momentum. DSBs develop at low impact momentums, and consist of second phase particles that
Table 2 Dynamic impact response of the investigated AA 2099 and AA 6061 aluminum alloys. Alloy
Impact momentum (kg m/s)
Total strain
Peak flow stress (MPa)
Strain rate (s1)
ASB
2099-T4
28 31 33 39
0.32 0.50 0.54 0.69
315 350 372 380
3630 5723 6238 7945
None None Deformed Transformed
2099-T6
28 31 33 39
0.33 0.44 0.68 0.74
378 380 330 410
3812 4984 7773 8497
Deformed Deformed Transformed Transformed
2099-T8
28 31 33 39
0.25 0.50 0.68 0.73
420 445 448 442
2889 5723 7820 8390
Transformed Transformed Transformed Transformed
6061-T4
33 39 44
0.61 0.69 0.79
390 560 336
6973 7905 9096
Transformed Transformed Transformed
6061-T6
33 39 44
0.59 0.70 0.80
430 571 336
6754 7971 9062
Deformed Transformed Transformed
6061-T8
33 39 44
0.54 0.65 0.75
401 482 502
6164 7426 8540
Deformed Transformed Transformed
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align in shear flow directions (Fig. 6). They were also observed ahead of TSB propagating paths in specimens impacted at high momentums. This suggests that the TSBs develop from the initially formed DSBs. Fig. 7 shows the impact momentum at which TSB was first observed in the impacted specimens. The minimum impact momentum for the formation of transformed band is lower for AA 2099 than for AA 6061. This confirms a higher propensity for formation of TSB in the AA 2099 alloy. Optical macrographs showing the impacted surface and side-view of some impacted specimens (Fig. 8) also confirms higher propensity of AA 2099 alloy to form transformed shear bands. The development of TSBs in the alloys during impact caused a change in the geometry of the transverse section of the impacted specimens from circular to elliptical as the impact momentum was increased. The elliptical and V-shape trajectory of the TSBs on the respective transverse and longitudinal sections of the impacted specimens are observable by visual inspection (Fig. 8). Visual inspections of the impacted specimens also indicate that the temper conditions that are most susceptible to adiabatic shear failure in AA 6061 and AA 2099 alloys are T4 and T8, respectively. In order to show the differences in the degree of adiabatic heating inside the shear bands that formed in the two investigated alloys, the temperature increase inside the shear bands was estimated using the following equation [1]:
DT ¼
b qC v
Z
rde
ð1Þ
where DT is temperature rise inside the transformed shear bands; b is fraction of the impact energy that is converted to thermal energy (taken as 0.9) [23]; q and Cv are densities and specific heat capacities of the alloys, respectively. Fig. 9 shows the
Fig. 6. Optical micrographs showing (a) deformed shear bands (DSBs) adjacent to transformed shear bands (TSBs) in an impacted AA 6061 alloy and (b) deformed shear bands in an impacted AA 2099 alloy.
Fig. 7. The impact momentum at which transformed shear bands were first observed in the alloys as the impact momentum of the projectile was increased from 28 kg m/s.
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Fig. 8. Optical macrographs showing changes in the geometry of the impacted specimens with increase in impact momentum.
a
250
b 250
AA 2099
Temperature (°C)
Temperature (°C)
150
T8 T6 T4
100
50
0
AA 6061
200
200
150
T4 T6 T8
100
50
0
26
28
30
32
34
36
Impact momentum (kg m/s)
38
40
0
30 0
32
34
36
38
40
42
44
46
Impact Momentum (kg m/s)
Fig. 9. Estimated temperature inside the transformed shear bands as a function of the impact momentum and temper condition for (a) AA 2099 and (b) AA 6061 aluminum alloys.
estimated relative temperature rise inside the transformed shear bands for various impacted specimens. For the AA 6061 alloy, the temperature rise inside adiabatic shear bands is highest for the alloy in T4 condition and lowest in the T8 condition. In AA 2099, the temperature rise is highest in the T8 condition and least in the T4 condition. This shows that the susceptibility of aluminum alloys has a direct correlation with the temperature rise leading to localized thermal softening. This temperature rise is in turn dependent on the applied momentum, alloy composition and the temper condition of the aluminum alloys. The temperature rise was observed to increase with the applied momentum for all the temper conditions. Typical transformed shear bands (TSBs) observed in the optical micrographs of the specimens subjected to the maximum applied impact momentum are presented in Fig. 10. The TSBs are characteristically featureless under an optical microscope. They appear white or dark depending on the intensity of etching during preparation. Heavily etched TSBs appear darker. The results of the scanning electron microscopic investigations of the impacted specimens reveal more detailed information on the difference in the morphology of the microstructures in the bulk materials and the TSBs. The results of the SEM study show that the second phase particles in the AA 6061 and AA 2099 alloys outside the TSBs have rectangular and elliptical shapes respectively (Figs. 11 and 12). As in optical microscopy, the microstructures of the TSBs in the AA 2099 alloy could not be resolved under the SEM except that the SEM images indicate that dissolution of the second phase particles occurred inside the TSBs. Fine spherical particles with average sizes of less than 2 lm were observed at the central regions of the TSBs in the AA 2099-T8 alloy (Fig. 11). The TSBs are devoid of the much bigger elliptical second phase particles that are observed in the bulk materials. They apparently dissolve as a result of the high intensity of adiabatic heating and severe strains within the TSBs during impact. Carbide dissolution was similarly reported to occur in TSBs during shear strain localization in AISI 4340 steel [7]. Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Fig. 10. Optical micrographs showing the transformed shear bands observed in the AA 6061 and AA 2099 impacted at the maximum applied momentum of 44 and 39 kg m/s respectively.
Fig. 11. SEM micrographs showing transformed shear bands (TSBs) in impacted AA 2099-T8 aluminum alloy.
The fine spherical particles at the centre of the TSBs in AA 2099-T8 could be the remaining portion of incompletely dissolved second phase particles or newly recrystallized grains. Because these grains are observed at the centre of the transformed bands where the peak temperature is expected to be highest, we believe these grains to be new dynamically recrystallized grains. Moreover, these grains do not appear to be as brittle as the second phase particles observed in the bulk materials. The recrystallization occurring simultaneously with deformation at elevated temperatures is termed dynamic recrystallization. Beside the few new equi-axed grains observed at the centre of the transformed bands in the AA 2099T8 alloy, the microstructure of the continuous phase inside the transformed bands could not be resolved under the scanning electron microscope used in this study. Transformed shear bands which appear featureless under an optical microscope and low resolution SEM have been reported by several authors [1–3] to consist of nano-sized grains when observed under a Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015
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Fig. 12. SEM micrographs showing transformed shear bands (TSBs) in impacted AA 6061-T8 aluminum alloy.
transmission electron microscope. We believe that the size of such recrystallized grains in the predominantly featureless areas of the transformed bands in the AA 2099 alloy is beyond the resolution limit of the SEM. Unlike in the case of the AA 2099 alloy, the microstructure of the TSBs in the AA 6061 alloy was resolvable using the scanning electron microscope, and they consists of fine equi-axed grains with an average size of 0.6 lm (Fig. 12). For the AA 6061-T8 alloy, the equi-axed particles are located at the centre of the transformed bands while the peripheral regions consist of a continuous phase with minimal or no second phase particles or recrystallized grains. Such second phase particle-free regions adjacent to TSBs are not noticeable in the T4 and T6 alloys. There are two possible mechanisms for the formation of the equi-axed grains inside the TSBs. One possibility is the fragmentation of the second phase particles occurring inside the shear bands as a result of these particles crushing against one another due to hydrostatic pressure from the rigid walls of the shear bands during impact. Grain elongation followed by fragmentation of pearlite lamellae followed by spheroidization has equally been suggested for the formation of fine equi-axed grains in pearlitic steels [35,36]. Fragmentation of the pearlite lamellae was reported to occur as a result of partitioning into transverse cell by dislocation movement under high hydrostatic pressure and shear stress. The plastic flow of the softened surrounding matrix phase could enhance the crushing of the second phase particles against one another while the soft continuous phase flows away from the centre of the transformed bands. However, in an earlier study on adiabatic shear banding in alumina particles reinforced AA 6061-T6 aluminum matrix composite, these hydrostatic pressure only caused a denser packing of the reinforcing alumina particles along the shear bands and no crushing of the alumina particles was reported inside the shear bands [37]. Another possible mechanism for the formation of the new fine equi-axed grain in the impacted AA 6061 alloy, which is more probable, is the dynamic recrystallization as earlier suggested for the fine spherical grains observed in the central region of the TSBs in the AA 2099 alloy. There are two types of recrystallization; static and dynamic recrystallization. While static recrystallization occurs subsequent to plastic deformation, dynamic recrystallization occurs simultaneously with deformation as a material is subjected to mechanical load at elevated temperatures. Derby [38] suggested that dynamic recrystallization can occur either by grain boundary rotational mechanism or grain boundaries migration through the deforming microstructure. The grain size obtained during dynamic recrystallization by migration of high angle grain boundaries during hot working of metals is reported to have stronger dependence on the deformation stress than temperature [38,39]. Dynamic recrystallization by grain boundary rotation is reported to account for formation of ultrafine equi-axed grains inside transformed bands in several metals [1,39,40]. The deformation time during dynamic shock loading is too short to allow grain boundary migration. It is generally believed that the severe strain and adiabatic heating trigger the dynamic recrystallization and the formation of shear bands in metals. A contrary opinion has recently been given by Rittel et al. [41,42], who suggested that strain localization is initiated by local softening caused by these dynamically recrystallized grains. They showed evidence of dynamically recrystallized grains in samples in which TSBs have not been fully developed. Figs. 13 and 14 show cracks propagation along TSBs in impacted specimens of the AA 2099-T8 and 6061-T4, respectively. These are the temper conditions which show the highest propensities for TSB’s formation and cracking for the respective alloys. Cracks initiated at and propagated along the interface between the TSBs and the bulk material in the AA 2099 alloy
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Fig. 13. Crack propagation at the interface between transformed shear band (TSB) and bulk material in AA 2099-T8 alloy.
Fig. 14. Crack propagation inside a transformed shear band (TSB) in AA 6061-T4 alloy.
Fig. 15. SEM micrographs showing the fracture surface of AA 2099-T8 aluminum alloy that fractured under impact loading. Fracture (a) inside the bulk material and (b) along the interface between TSB and the bulk material.
(Fig. 13). In AA 6061 alloy cracks are initiated inside the TSBs and they propagated along the central region of the shear bands (Fig. 14). The TSBs formed in the AA 2099 alloy are more susceptible to cracking than those in AA 6061 alloy. The difference in the visco-plasticity of the soft shear bands and that of its rigid surrounding material during impact loading generated the tensile stresses that opened up voids that coalesced to initiate crack inside the TSBs of the AA 6061 alloy [7,8]. The interfacial cracks between the TSBs and the bulk materials in AA 2099 may be due to a weak interface, which shears easily as the soft material in the TSB undergoes plastic flow under impact loading. The results of the SEM examination of the fracture surface of a fragmented AA 2099 alloy are shown in Fig. 15. The specimen shows a very smooth round fracture surface. These confirm shearing of the specimens around the transformed shear bands. Cleavage fracture of the second phase particles and ridges showing ductile fracture of the continuous aluminum-rich phase are observed on the fracture surface. As indicated in the figure, two regions can be identified on the fractured surface. At the centre of the fracture surface, irregular topography
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showing ridges or striations in the continuous phase can be observed adjacent to the TSBs (Fig. 15b). A less ductile fracture is observed in the bulk material farther away from the TSBs (Fig. 15a). None of the AA 6061 specimens fractured under the investigated impact condition despite the higher impact momentum, to which they were subjected. 4. Conclusions The effects of strain rate and temper condition on strain localization and adiabatic shear failure in AA 6061 and AA 2099 aluminum alloys are investigated and compared. The following conclusions can be drawn from the results obtained: 1. AA 2099 has higher deformation resistance under quasi-static loading whereas AA 6061 alloy is more resistant to deformation under dynamic impact loading. 2. Failure of the alloys under impact loading is preceded by strain localization along transformed shear bands which initially occurred as deformed bands. 3. AA 6061 is most susceptible to adiabatic shear failure in the T4 temper, whereas AA 2099 is most susceptible in the T8 temper condition. 4. Transformed bands in AA 6061 alloy consist of fine recrystallized equi-axed grains with an average size of 0.6 lm while the microstructure of the transformed bands in AA 2099 is not resolvable under the same scanning electron microscope. 5. Cracks initiate and propagate along the interface between the adiabatic shear bands and bulk material in the AA 2099 but along the central region of the shear bands in the AA 6061 alloy. 6. Fragmentation of specimen occurs in AA 2099 under the impact loading while none of the AA 6061 specimens fractured, even at higher impact momentum, suggesting higher cracking susceptibility of AA 2099 under impact loading.
Acknowledgments The authors wish to acknowledge the support of Alcoa incorporation in providing the investigated AA 2099 alloy and Natural Sciences and Engineering Research Council of Canada for their financial support of this study. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
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Please cite this article in press as: Odeshi AG et al. Failure of AA 6061 and 2099 aluminum alloys under dynamic shock loading. Eng Fail Anal (2013), http://dx.doi.org/10.1016/j.engfailanal.2013.02.015