Effects of temper condition on high strain-rate deformation of AA 2017 aluminum alloy in compression

    Effects of temper condition on high strain-rate deformation of AA 2017 aluminum alloy in compression A.A. Tiamiyu, A.Y. Badmos, A.G. Odeshi PII: DOI: Reference:

S0264-1275(15)30625-0 doi: 10.1016/j.matdes.2015.10.047 JMADE 786

To appear in: Received date: Revised date: Accepted date:

4 March 2015 11 October 2015 12 October 2015

Please cite this article as: A.A. Tiamiyu, A.Y. Badmos, A.G. Odeshi, Effects of temper condition on high strain-rate deformation of AA 2017 aluminum alloy in compression, (2015), doi: 10.1016/j.matdes.2015.10.047

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Effects of temper condition on high strain-rate deformation of AA 2017 aluminum alloy in compression

Department of Mechanical Engineering, University of Saskatchewan Saskatoon, SK, Canada 2College of Engineering, Department of Engineering Technology, Black Hawk College, Moline, Illinois, US

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A. A. Tiamiyu*1, A. Y. Badmos2, A. G. Odeshi1

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Abstract

Cylindrical specimens of AA 2017 alloy in T451, T651 and O temper conditions were subjected

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to quasi-static compressive loading at a strain rate of 3.2 x 10-3 s-1 and dynamic impact loading at strain rates between 3000 s-1 and 8000 s-1. The effects of strain rates and temper condition on

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damage evolution were investigated. Deformation under quasi-static loading was dominated by strain hardening while high strain rates deformation was characterized by two-peak stress-strain

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curves, indicating simultaneous occurrence of hardening and softening controlling deformation. The degree of softening after the first peak is highest in the annealed sample while the age-

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hardened specimens showed higher degree of softening beyond the second peak. Both deformed

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and transformed shear bands were observed in the impacted age-hardened specimens, whereas only deformed bands developed in the annealed specimens. The critical strains for the onset of the thermomechanical instability leading to the formation of adiabatic shear bands are about the same for both T451 and T651 specimens, but much higher for the annealed (O) specimens. The age-hardened specimens fractured in a manner consistent with ductile fracture mode that is characterized by the sequential nucleation, growth and coalescence of micro-voids inside transformed shear bands.

Keywords: AA 2017 aluminum alloy; dynamic shock loading; plastic deformation; adiabatic shear bands; fracture.

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Introduction

As a result of their high specific strength and excellent formability, aluminum alloys are choice materials for aerospace and automobile applications. Like other metallic materials, aluminum alloys can experience heterogeneous deformation leading to the development of adiabatic shear

*Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT bands (ASBs) under dynamic shock loading [1-3]. ASBs are regions of intense shear strain localization brought about by thermo-mechanical instabilities induced by localized adiabatic heating during shock loading [4-6]. They are classified into two types: deformed bands

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consisting of elongated grains and transformed bands consisting of equiaxed ultrafine grains [7]. Both types of shear bands are observed in aluminum alloys depending on the severity of loading

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and strain localization [2,8]. Several aluminum alloys that are generally known to be ductile, do

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suffer catastrophic fracture under dynamic shock loading due to adiabatic heating and strain localization [9-12]. Localization of deformation usually starts with the formation of deformed

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bands, which develop into transformed bands once a critical strain is attained [2,13,14]. Formation of transformed bands initiates dynamic failure of materials at high strain rates

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[6,15,16]. The tendency of an aluminum alloy to fragment under dynamic impact load is determined by its susceptibility to the formation of transformed bands [11,17]. The susceptibility of AA 2099 and AA 6061 aluminum alloys to the formation of transformed bands is influenced

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by the temper condition [11]. The tendency of quenched and tempered AISI 1340 steel to

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develop these bands also depends on the tempering condition [13].

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A systematic investigation of the damage evolution in aerospace aluminum alloys under dynamic shock loading is very important considering the increasing incidences of damage to aircraft fuselage due to bird strikes during take-off and landing [18]. The technological importance of scientific studies focusing on the formation of ASB in different metallic alloys is to mitigate adiabatic shear failure by suitable control of materials and process variables influencing the occurrence of ASBs in each alloy. This can be achieved by detailed studies to provide in-depth understanding of how various material variables affect damage evolution. In addition to understanding the condition promoting or mitigating the occurrence of ASBs, the kinetics and dynamics of shear band propagation and failure should be well understood. Such knowledge base is important for incorporation into new alloy design for an enhanced resistance to impact failure. Some of the factors identified to influence the occurrence of ASB in metallic alloys include strain rate sensitivity [19], microstructure [7,13], specimen geometry and dimension [19,20], local defects [19], notches [21], stacking fault energy [22], deformation temperature and grain size [23]. The presence of oxide dispersions or second phase particles has also been reported to influence the capability of tungsten heavy alloys (WHA) to resist ASB formation under dynamic shock loading [24,25]. *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT AA 2017 is an aluminum-copper-magnesium alloy with very useful applications in the aerospace industries. The dominant second phase particles, which contribute to the strengthening of Al-CuMg alloys are  (Al2Cu) and S (Al2CuMg) phases while the metastable phases / and S/S

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precipitates are responsible for high strength in the peak-aged condition [26,27]. Most of the previous works on AA 2017 aluminum alloys focused on precipitation kinetics during aging

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[28], mechanical behavior under quasi-static loading [29] or fatigue loading [30], corrosion

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behavior [31] and joining [26,32]. Ambroziak et al. [29] investigated the effects of loading direction and strain rates on yield strength and elastic modulus of the alloy under tensile loading,

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and developed constitutive models to describe the alloy’s behavior under quasi-static loading. The mechanical behavior of AA 2017 aluminum alloy at high strain-rates (>103 s-1) is not well

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investigated, although this aerospace alloy can experience high velocity impact from bird strikes when used in aircraft’s fuselage or wing. With the increasing cost of bird strikes in the aviation industry, it is important to understand the dynamic impact behavior of this alloy in order to

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effectively improve or predict their resistance to failure under this severe loading condition. This

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study aims to bridge this gap and provide information on the microstructural evolution that leads

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to failure in the alloy under dynamic impact loading. Experimental Methods

The composition of the investigated AA 2017 aluminum alloy is presented in Table 1. The alloy was received as a rolled rod in T451 temper condition. The T451 temper of the as-received alloy indicates that it was subjected to solution heat treatment process and stress relief by stretching to about 1-3% before natural aging [33]. The as-received alloy was machined into cylindrical test specimens with diameter and length of 9.5 mm and 10.5 mm, respectively. A third of the machined specimens were tested in the as-received T451 temper, a third were further artificially aged at 100C for 10 h to obtain T651 temper, while the rest were annealed to O temper. The annealed condition was achieved by solutionizing at 413 C for 3 h followed by a two-stage cooling process. The first stage involved slow cooling at 30 °C per hour to 260 °C, followed by air cooling to room temperature. Characterization of the mechanical behavior of the alloy under quasi-static compressive loading was done using Instron R5500 mechanical testing system. The specimens were subjected to *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT compressive load at a crosshead speed of 2 mm/min and a maximum load of 100 kN. The crosshead speed generated a strain rate of 3.2 x 10-3 s-1 in the specimens. Hardness measurement was taken using Mitutoyo MVK-H1 microhardness tester. High strain-rate deformation was

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carried out using split Hopkinson pressure bar (SHPB) system (Fig. 1). The specimens were sandwiched between the input and output bars while a blunt projectile fired by a light gun was

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made to strike one end of the input bar and rapidly compress the specimens. The strain gages

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mounted on the input and output bars captured the elastic wave data that were used in calculating the stress, strain and strain rate values [34]. The mechanical tests were conducted at room

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temperature. The specimens for microstructural evaluation were cold-mounted using acrylic resin, ground and polished to mirror surface finish. The polished specimens were etched using a

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solution consisting of 25 ml Methanol, 25 ml HNO3, 25 ml HCl, and 1drop of HF. Etching times for AA 2017-T451 and AA 2017-T651 specimens were about 50 seconds while AA 2017-O specimen only took 10-15 seconds to properly etch. Nikon MA100 inverted metallographic

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microscope with Pax-it! image analysis software and JEOL-JSM-6010LV scanning electron

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microscope were used for the microscopic investigations. Results and Discussion

3.1

Microstructure before Mechanical Loading

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Optical and SEM microstructures of the samples along a cross section perpendicular to the rolling direction before mechanical loading are presented in Fig. 2. The microstructures consist of dispersed second phase particles embedded in a continuous aluminum phase. The second phase particles are irregular in shape and are of varying sizes. The size range for second phase particles observed in the T451, T651 and O alloy are 8 - 25µm, 10 - 32µm and 83-140 µm respectively. This suggests that the second phase particles grew to larger sizes during the annealing treatment. In Al-Cu alloys, aluminum combines with copper to form Al2Cu precipitates. AA 2017 is an Al-Cu alloy containing a significant amount of magnesium. The presence of a substantial amount of magnesium in an Al-Cu alloy leads to formation of orthorhombic Al2CuMg particles which further increases the strength of the alloy and makes them suitable for high performance aerospace applications [35]. EDX analyses on the second phase particles in an AA 2017 alloy in a previous study showed the presence of iron rich precipitates such as Al3Fe or Al12(Fe,Mn)3Si in addition to Al2Cu () and Al2CuMg (S) *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT equilibrium phases [26]. The average grain sizes of AA 2017-T451, AA 2017-T651 and AA T2017-O alloys are about 24µm, 30µm and 58 m, respectively. Considerable grain growth thus

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occurred during annealing of the specimens. Mechanical Behavior and Microstructural Evolution during Quasi-Static Loading

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The average hardness of the investigated AA 2017 aluminum alloy was measured to be 131 HV, 114 HV, and 54 HV for the T451, T651, and O tempers respectively. These are average of ten

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hardness measurements. Artificial aging and annealing of the as-received T451 specimens resulted in about 13% and 59% decrease in hardness respectively. The typical stress-strain curves

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obtained from the quasi-static compressive loading of the specimens up to the maximum load of 100 kN are presented in Fig. 3a. Although the artificial aging of the as-received T451 specimens

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led to about 8.4% decrease in compressive elastic modulus, annealing treatment led to about 80% reduction in elastic modulus. The total engineering strains developed due to the applied

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maximum compressive load of 100 kN were determined to be 0.65, 0.64 and 0.78 respectively.

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The difference in the total strain for the T451 and T651 specimens is negligible. The higher hardness and yield strength observed in the as-received T451 specimen suggest that

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the peak strength of the alloy was already attained before the additional artificial aging. Two precipitation sequences are suggested to occur in AA 2017 alloy: SSSGP zones and SSSGPB zonesSSS) [36]. SSS is the supersaturated solid solution obtained from the solution heat treatment. The precipitation sequence and kinetics in Al-Cu-Mg alloys are discussed in detail by other researchers [37–41]. The strengthening equilibrium precipitates in Al-Cu-Mg alloys are the  (Al2Cu) and S (Al2CuMg) [42]. It is suggested that artificial aging of the as-received alloy led to the loss of coherency between the precipitates and the continuous phase. Loss of coherency will normally translate to a reduced resistance to slip since dislocations will rather loop around incoherent precipitates than to cut through it. The strengthening effect of dislocations looping around incoherent precipitates is less than that of dislocations cutting through coherent precipitates [43]. Ostwald ripening could also have occurred during the artificial aging, leading to increase in the distance between precipitates. According to Orowan mechanism, the shear stress required for dislocation to loop around incoherent precipitates is inversely proportional to the distance between them [43]. There is a critical particle size, below which the mechanism of particle shearing by dislocation operates and the critical resolved shear *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT stress (CRSS) is directly proportional to particle size. Above the critical particle size, the precipitates become incoherent and the mechanism of dislocation looping around the particles operates (Orowan mechanism). At this instance, the CRSS becomes inversely proportional to the

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precipitates’ size. Reduced coherency and/or Ostwald ripening can therefore account for the lower yield strength of the artificially aged specimens compared to the naturally aged ones. The

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that are effective barriers to the motion of dislocation (slip).

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yield strength of the annealed specimens is very low due to coarsening of the fine precipitates

It can be noted that the stress-strain curve of the T651 specimen, which is slightly below that of

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the T451 specimen at the onset of plastic deformation overlap one another when the strain exceeded 0.5. This implies that strain hardening rate is higher for the artificially aged specimen

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at higher strain resulting in the elimination of the effect of the reduction in yield strength due to artificial aging. The strain hardening curves obtained from quasi-static loading of specimens are

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presented in Fig. 3b. Strain hardening rate are very close for both the T451 and T651 specimens but consistently lower for the O-specimen up to a strain of about 0.65. At higher strains some

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spikes can be observed in the strain hardening curves for the age-hardened samples. These spikes

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correspond to the slight drop observed in the stress-strain curves of the age-hardened specimens (Fig. 3a inset). The downward spikes in the strain hardening curves may be due to recovery, whereby dislocation gains sufficient energy to break through barrier and become mobile again. It is also not unlikely that the spikes signal the onset of cracking in the age-hardened samples, more so that the spikes occur close to the end of the curves and this is not found in the annealed sample where no crack was observed (Fig. 4). If this later explanation is true, it is reasonable to suggest that the T451-specimens fractured earlier than the T651 specimens as indicated in the strain hardening curves, where the start of spikes occurred at a higher strain for the T651 sample. 3.3

Mechanical Behavior under Dynamic Mechanical Loading

The stress-strain and strain hardening curves for the specimens tested under dynamic impact loading at 5000 s-1 are presented in Fig. 5. At this strain rate, the maximum flow stresses of AA 2017-T451, AA 2017-T651 and AA 2017-O are estimated to be 532 MPa, 496 MPa and 305 MPa. Despite the T451 alloy showing a higher peak flow stress than the T651 alloy, the differences observed in their total strain are less than 3%. The annealed specimens exhibited the lowest peak flow stress and a total strain that is about 15% higher than that of AA 2017-T451 *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT specimen. Thermal softening and strain hardening compete during the plastic deformation process [44][45][46]. They both contributed to the deformation and determined the maximum peak stress. Thermal softening occurs as a result of conversion of impact energy into thermal

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energy which raises the temperature of the specimens. About 90 % of the kinetic energy of

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projectile is estimated to convert to thermal energy [34].

The stress-strain curves indicate an initial elastic deformation that is quickly overshadowed by a

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plastic deformation regime with two major peaks. Flow stress oscillation resulting in multiple peaks in stress-strain curves has been reported to be a combined effect of dislocation

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accumulation and annihilation due to dynamic recovery or dynamic recrystallization [47]. In this study, the first peak occurred not long after the yield point and the drop in stress after this peak

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could be attributed to an initial flow softening as the temperature of the specimen began to increase under the impact loading. Dislocation climb and cross slip will be enhanced, leading to

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an increased dislocation mobility. Eventually, strain hardening begins to dominate again leading to an increase in stress up to the maximum flow stress. The strain hardening after the first peak in

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the stress-strain curves can be attributed to dislocation multiplication that eventually results in

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the formation of dislocation pile-ups at barriers such as grain boundaries, Cottrell-Lomer locks and precipitates, among others [48,49]. Dislocations pile-up will usually exert back stress that opposes the motion of on-coming dislocations. Higher stress will therefore be required for further deformation. Thermal softening dominated the plastic deformation process beyond the maximum flow stress leading to a thermally induced mechanical instability and to a sharp drop in flow stress. The start of the mechanical instability corresponds to the point where excessive thermal softening occurs in narrow regions of the impacted specimen, leading to strain localization [5]. The strain hardening curves in Fig. 5b indicate strain hardening rate dropping to below zero, which indicate the dominance of thermal softening over strain hardening. The earliest drop in strain hardening rate to below zero was observed at strain values of between 0.05 and 0.15. This corresponds to the post-yield dislocation annihilation and rearrangement due to the initial increase in temperature. The critical strain for instability at higher strain is approximately the same for the AA 2017-T451 and AA 2017-T651 specimens, but much higher for the AA 2017-O specimens.

*Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT The effects of strain-rates on the dynamic impact response of the AA 2017 aluminum alloy in the different temper conditions are presented in Table 2 and Fig. 6. As the impact momentum increased, the strain rates and the total engineering strain increased for each of the investigated

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temper. Though the peak flow stress of T451 specimen increased with impact load, T651 specimen initially showed a decrease in peak flow stress, which subsequently increased as the

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impact momentum was further increased. The rather non-linear relationship between strain rates

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and the peak flow stress is attributed to the complex interaction caused by the competition between thermal softening and strain hardening during deformation [5]. The temperature rise in

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specimens under the quasi-static loading condition is not high enough to induce significant thermal softening or mechanical instabilities as observed in the specimens subjected to dynamic

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impact load. The AA 2017-O alloy specimen exhibited an enhanced plasticity and formability under both quasi-static and dynamic shock loading conditions. The relative amount of thermal softening in the annealed specimens after the first peak is greater than that for the precipitation

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hardened alloys for all the investigated strain rates.

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Occurrence of thermal softening immediately after yielding has been reported in the stress-strain

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curves of aluminum alloys during hot deformation [50,51] as observed in the dynamic impact stress-strain curves of AA 2017 alloy in the current study. As mentioned earlier, a significant increase in temperature occurs in specimens subjected to dynamic shock loading making high strain deformation somewhat similar to hot deformation. The major difference is in the fact that stress-strain curves for hot deformation decrease continuously after yielding until failure. On the other hand, the stress-strain curves obtained in this impact study show a subsequent domination of strain hardening leading to a second peak stress. Beyond the second peak stress, thermal softening dominated again leading to a continuous decrease in stress. Flow softening in a stressstrain curve has been adduced to a number of phenomena that include thermo-mechanical instability, strain induced phase transformation, localized temperature rise and texture softening [52]. In this study, relative thermal softening (Sr) beyond the first and second peak was calculated using the following equation;

 p    *100   p  

Relative Softening Sr (%)  

*Corresponding author: [email protected]

(1)

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ACCEPTED MANUSCRIPT where p is the peak stress and  is the stress (MPa) at various strain levels beyond the peak stress. Other researchers have used Eq. 1 to quantitatively estimate the amount of softening that occurs after the peak stress is attained during hot deformation [50,51,53]. Since strain hardening

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curves could not give much information on the difference in the softening of the three investigated materials after yielding, this equation was used to calculate the relative softening

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after the two peak stresses observed in the dynamic stress-strain curves.

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The plots of the relative softening against the respective strain for both peaks are presented in Fig. 7. Whereas the relative softening of the annealed specimen (O temper) occurring

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immediately after the first peak is higher than those of the precipitation hardened specimens as the strain increases beyond 0.14, reverse is the case for the softening occurring after the second

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peak. From a study on hot deformation of AA 6082 aluminum alloy, Zhang and Baker [53] observed the amount of post-yield softening to be higher in T4 temper than in annealed

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(O temper) specimens. They attributed the higher post yield softening in the AA 6082-T4 to

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higher tendency of the tempered alloy to experience dynamic recrystallization compared to the annealed samples in which the level of supersaturation is low. Unlike in hot deformation, where

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the temperature of the specimen is assumed to be relatively uniform throughout the deformation period (ignoring any temperature increase due to deformation), temperature of the impacted AA 2017 alloy specimens increased significantly during impact loading. From the experimental data, yielding was observed to occur within the first 75 microseconds after the start of deformation in the three temper conditions. The temperature and the strain reached at this stage rule out the possibility of dynamic recrystallization and particle coarsening. Dynamic recovery is therefore suggested to account for the post yield softening observed in the impact stress-strain curves of the investigated AA 2017 alloy. 3.4

Microstructural Evolution during Impact Loading

Microstructural investigations of the impacted AA 2017 aluminum alloy specimens indicated heterogeneous deformation leading to the occurrence of ASBs. Both deformed and transformed bands were observed (Fig. 8 and 9). The second phase particles were observed to align along the plastic flow direction in the deformed bands. It has been reported that transformed bands develop from deformed bands when the intensity of strain localization reaches a critical value [8,54] and this was also observed in this alloy. Table 3 provides the strains and strain rates at which *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT deformed bands were first observed in the AA 2017 alloy under different tempered conditions. Considering the required impact momentum, the critical strain and the critical strain rates for the formation of ASBs in the alloy, the annealed specimens showed the least susceptibility to the

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formation of ASBs and to fracture under impact loading. Deformed bands that are not fully developed were observed in the annealed specimens, neither was transformed band nor cracking

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observed at the maximum applied impact momentum of 43.6 kg.m/s. Transformed bands at the

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early stage of development observed in T451 and T651 specimens impacted at 41.5 kg.m/s are presented in Fig. 8. This suggests a comparable susceptibility to occurrence of transformed bands

the fracture of AA 2017-T651 specimen.

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for the two temper conditions. However, a slightly higher impact momentum was required for

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Adiabatic shear bands formed a ring on the compression plane in both T451 and T651 specimens (Fig. 9). Shear band bifurcation was observed in T651 alloy, whereas none was observed in the T451 specimens. Yang et al. [55] speculated that shear band bifurcations occurred in AA 7075

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aluminum alloy when the propagation of a “mother” shear band along its original direction is

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hindered by the presence of a barrier such as impurity or second phase particles. None of such

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barriers was observed in the region where ASBs split in the current alloy. The region ahead of the bifurcation point is relatively free of second phase particles. It is however possible that the barrier was pushed aside by the plastic flow of the soft matrix under pressure leaving a second phase particles free zone (SPFZ) at the point of bifurcation. It could also be that the SPFZ observed at the point of bifurcation was originally present before the impact. No definite explanation for the cause of the bifurcation is possible at this point and some further work will be required in order to provide the answer. However, these results show that the occurrence of bifurcation during ASB propagation in AA 2017 alloy is determined by its temper condition. The ASB on the longitudinal section of the impacted AA 2017-T451 alloy is parabolic in shape (Fig. 10a). Considering the ASB geometry on the transverse and longitudinal sections of the impacted cylindrical specimens, it can be deduced that shear band formed a thin-walled cone in the specimens. Conical shaped ASB was also reported in AISI 4340 steel subjected to dynamic impact loading [56], except that two inverted cones were observed in the steel instead of a singular cone observed in AA 2017 aluminum alloy in this study. It was reported that shear band tends to propagate along slip systems such that shear plane become parallel to the slip plane while the shear direction also becomes aligned with the slip direction [57]. It is however not *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT clear if the difference in the number of cones formed by ASB has anything to do with the difference in the slip systems of AISI 4340 steel and aluminum alloy.

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Figure 10a shows that the intensity of shearing in the parabolic ASB on the longitudinal section of an impacted specimen (41.5 kg.m/s) decreases from the side in contact with the output bar of

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the SHPB to the other side in contact with the input bar (struck by the projectile). This is schematically illustrated in Fig. 10b. Continuous solid lines A-E and A’-E’ represent alignment

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of second phase particles in the rolling direction of an un-deformed specimen before impact. The continuous dash line represents the center line of the sample. It is evident from Fig. 10a that

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elongated grains and second phase particles become discontinuous at specific regions and these are indicated by discontinuous solid lines A1-E5 and A1’-E5’ in Fig. 10b. The regions of

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discontinuity correspond to the parabolic ASBs on the longitudinal section of the impacted specimen. The shear band presented in Fig. 10a is largely a deformed band. The intensity of

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shear localization decreases towards the vertex of the cone formed by the shear band (Fig. 10). For the specimen impacted at a higher momentum of 42.6 kg.m/s, both transformed and

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deformed bands were observed along the parabolic ASB (Fig. 11) but with an increased intensity

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of transformed band. The deformed bands were observed at the vertex of the cone near the surface in contact with the input bar. The observed variation in the intensity of shearing along the shear band could be due to the variation in the intensity of adiabatic heating in the specimen. The intensity of shear strain localization in a region during impact will depend on the degree of softening there.

Results of the scanning electron microscopic investigation of the transformed bands observed in the T451 and T651 specimens are presented in Figs. 12-14. Although the grain structure of the transformed bands could not be effectively resolved with the SEM used in this study, they appear to consist of dissolved or crushed second phase particles with islands of SPFZs observed within the transformed bands. Previous investigations by other researchers showed that transformed bands consist of ultrafine equiaxed grains in aluminum alloys such as AA 6061-T6 [11], AA 2099-T8 [58], AA 7075 [8] and in other alloys such as zirconium alloy [57], titanium alloy [59], AM60B magnesium alloy [54] and AISI 1340 steel [13]. Higher resolution transmission electron microscopy (TEM) or orientation imaging techniques (OIM) using the electron back scattered

*Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT diffraction (EBSD) will be needed to adequately resolve the grain structure of the transformed bands to provide information on the size and distribution of the grains.

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The fracture of AA 2017-T451 and AA 2017-T651 was initiated by cracks which originated and propagated along the transformed bands. There was no evidence of cracking in the regions

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outside the transformed bands nor inside the deformed bands. The facture mode along transformed bands in the AA 2017-T651 specimen can be observed in Fig. 14. Similar fracture

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mode was found in the AA 2017-T451 specimen. Elliptically shaped micro-voids appear to first form inside the transformed bands in the regions where second phase particles exist. No micro-

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voids were observed to nucleate in the particle free zones of the transformed bands. This suggests that interfacial debonding between second phase particles and the continuous aluminum

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phase plays a role in nucleation of the micro-voids. An investigation by Ma et. al [60] on failure of 5A02-O aluminum alloy under quasi-static and dynamic tensile load also indicated that micro-

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voids were nucleated at second phase particles. Micro-voids can also evolve due to the existence of stress gradients in ASBs. The existence of the stress gradient tends to generate tensile stresses

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which open voids in the shear band as deformation proceeds under the impact loading [13].

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Other suggested mechanisms for micro-voids nucleation in metals include, vacancy accumulation at a high stress region, grain boundary sliding, and void nucleation at the head of dislocation pile-ups [56].

No transformed band was observed in the annealed specimens impacted at the momentum where T451 and T651 specimens fractured. The annealed specimens exhibited large plasticity and did not fracture. SEM investigation (Fig. 15) shows substantial amount of precipitations along the grain boundaries of the annealed specimens compared to the precipitation hardened specimens. The grain boundary precipitation in the annealed specimen is suggested to be aided by the prolonged exposure to high temperatures promoting diffusion during the annealing process. The quench rate, solution heat treatment temperature and aging treatment have been reported to be the important parameters influencing grain boundary precipitation in aluminum alloys [61,62]. Such grain boundary precipitations were not observed in the precipitation hardened AA 2017 specimens in this study. Steele et al. [63] however observed grain boundary segregation in AA 6111 aluminum alloy subjected to solution heat treatment followed by water quenching. Grain

*Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT boundary segregation has been shown to have effect on mechanical behavior of aluminum alloys and might contribute to the low strength of the AA 2017-O specimens in this study [64]. Conclusions

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4.

Plastic deformation and failure of AA 2017 aluminum alloy under both quasi-static and dynamic

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shock loading were investigated. Deformation was relatively homogeneous under quasi-static

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load with crack initiating at the edge of the cylindrical test specimens and propagating radially inwards. Thermal softening leading to mechanical instability causes strain localization and shear

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failure of the alloy at high strain rates. The critical strain for the instability is approximately the same for the alloy in T451 and T651 temper conditions, but much higher in the annealed

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specimens (O-Temper). Cracks were initiated inside the impacted specimens in region containing clusters of second phase particles within the transformed bands. Two peaks were observed in the dynamic stress-strain curves of the alloy in both the precipitation hardened and annealed

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conditions. The relative amount of strain softening in the annealed specimens after the first peak

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is greater than those for T451 and T651 specimens. However, the relative softening of the impacted annealed specimens beyond the second peak was observed to be the lowest of the three

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investigated temper conditions. Although the lower susceptibility of the annealed AA 2017 to ASB formation compared to that of the age-hardened ones favors their application for structures that will be exposed to dynamic impact loading, its compromised strength poses a setback. Acknowledgment

The authors wish to acknowledge the financial support of Natural Sciences and Engineering Research Council of Canada (NSERC) for the financial support of this study. References [1]

Y. Yang, H. G. Zheng, Z. J. Shi, and Q. M. Zhang, “Effect of orientation on selforganization of shear bands in 7075 aluminum alloy,” Mater. Sci. Eng. A, vol. 528, no. 6, pp. 2446–2453, Mar. 2011.

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Y. B. Xu, W. L. Zhong, Y. J. Chen, L. T. Shen, Q. Liu, Y. L. Bai, and M. A. Meyers, “Shear localization and recrystallization in dynamic deformation of 8090 Al–Li alloy,” Mater. Sci. Eng. A, vol. 299, no. 1–2, pp. 287–295, Feb. 2001.

*Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT G. M. Owolabi, A. G. Odeshi, M. N. K. Singh, and M. N. Bassim, “Dynamic shear band formation in Aluminum 6061-T6 and Aluminum 6061-T6/Al2O3 composites,” Mater. Sci. Eng. A, vol. 457, no. 1–2, pp. 114–119, May 2007.

[4]

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ACCEPTED MANUSCRIPT [41] S. C. Wang, M. J. Starink, and N. Gao, “Precipitation hardening in Al-Cu-Mg alloys revisited,” Scr. Mater., vol. 54, pp. 287–291, 2006.

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ACCEPTED MANUSCRIPT [53] B. Zhang and T. N. Baker, “Effect of the heat treatment on the hot deformation behaviour of AA6082 alloy,” J. Mater. Process. Technol., vol. 153–154, pp. 881–885, Nov. 2004.

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List of Figures

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Fig. 1. Schematic of the split Hopkinson pressure bar system used for dyna mic impact test.

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Fig. 2. Optical and scanning electron micrographs showing dispersed second phase particles in the test specimens before mechanical loading (a) AA 2017-T451, (b), AA 2017-T651 (c) and AA

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2017-O aluminum alloy.

Fig. 3. (a) Engineering stress-strain curves and (b) strain hardening curves of AA 2017-T451,

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AA 2017-T651 and AA 2017-O aluminum alloy under quasi-static compressive loading. Inset in (a) is the magnified portion of the stress-strain curve showing a slight drop in stress at higher

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strain for the age-hardened specimens.

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Fig. 4. SEM micrographs showing edge cracks on the compression plane of (a) AA 2017-T451 and (b) AA 2017-T651 specimens under quasi-static loading. No edge crack was observed in (c)

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AA 2017-O specimen.

Fig. 5. (a) Dynamic stress-strain curves and (b) strain hardening curves of AA 2017 aluminum alloy under impact load at approximately 5000 s-1. Fig. 6. Dynamic engineering stress-strain curves of (a) AA 2017-T451, (b) AA 2017-T651 and (c) AA 2017-O at increasing impact momentum. Fig. 7. Relative flow softening-strain plot beyond the (a) yield peak (first peak) and (b) second peak of stress-strain curve of AA 2017-T451, AA 2017-T651 and AA 2017-O at approximately 5000 s-1. Fig. 8. Optical micrographs showing the onset of (a) deformed band in AA 2017-T451, (b) transformed band in AA 2017-T451 and (c) transformed band in AA 2017-T651. Fig. 9. Geometry of ASB on the compression plane (transverse section) of (a) AA 2017-T451 alloy with no bifurcation and (b) AA 2017-T651 alloy with bifurcation. *Corresponding author: [email protected]

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ACCEPTED MANUSCRIPT Fig. 10. (a) Optical micrograph showing the parabolic ASB geometry on the longitudinal section of AA 2017-T451 and (b) schematic of formation of parabolic ASB geometry.

section of AA 2017-T451 specimen impacted at 42.6 kg.m/s.

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Fig. 11. (a) Optical micrograph showing transformed and deformed bands on the longitudinal

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Fig. 12. SEM micrographs of AA 2017-T451 (a) inside ASB and (b) outside ASB after been

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subjected to impact momentum of 42.6 kg.m/s.

subjected to impact momentum of 43.6kg.m/s.

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Fig. 13. SEM micrographs of AA 2017-T651 (a) inside ASB and (b) outside ASB after been

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Fig. 14. SEM micrographs showing the mode of crack nucleation and propagation along a transformed band in AA 2017-T651 aluminum alloy.

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band on the compression plane of AA 2017-O specimen.

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Figure 1

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Figure 2

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Figure 3

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Figure 4

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Figure 5

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Figure 9

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Figure 11

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Figure 12

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Figure 14

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Figure 15

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ACCEPTED MANUSCRIPT Table 1. Composition range for AA2017 aluminum alloy. Al

Cu

Mg

Mn

Si

Zn

Ti

Cr

Fe

Others

wt %

91.5 - 95.5

3.5 - 4.5

0.4 - 0.8

0.4 - 1.0

0.2 - 0.8

< 0.25

< 0.15

< 0.1

< 0.7

< 0.15

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Element

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486 446 496 494

0.28 0.35 0.41 0.48

No No Yes Yes

4870 5280 7040 7560

265 305 274 232

0.43 0.47 0.61 0.68

No No No Yes

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30.0 34.3 37.6 40.0

3290 4140 4780 5550

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AA2017-T651

30.0 34.3 37.6 40.0

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Table 2. Mechanical response of AA2017-T451, AA2017-T651 and AA2017-O under dynamic shock loading. Impact Strain Maximum Engineering Adiabatic Temper momentum rate flow stress strain after shear -1 (kg.m/s) (s ) (MPa) impact bands 30.0 3080 460 0.26 No 34.3 3960 507 0.34 No AA2017-T451 37.6 4600 532 0.40 No 40.0 5340 561 0.47 Yes

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Table 3. Summary of critical values of ASB evolution. AA2017-T451

Critical strain rate, s-1 for DB Impact Momentum, kg.m/s at the onset of TB Critical strain for TB Impact Momentum, kg.m/s at fracture

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37.6 0.41

AA2017-O 40.0 0.68

4780

7560

41.5 0.54 43.6

No TB No TB No fracture

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40.0 0.47

41.5 0.53 42.6

AA2017-T651

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Alloy Impact Momentum, kg.m/s at the onset of DB Critical strain for DB

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Graphical abstract

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ACCEPTED MANUSCRIPT Highlights 

Damage evolution in AA 2017 aluminum alloy under quasi-static and high strain-rate

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compressive loading was investigated. Deformation at low strain rates was dominated by strain hardening while high strain rates

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deformation is characterized by two-peak stress-strain curves indicating thermal 

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softening

Degree of softening after the first peak is highest in the annealed specimens while agehardened specimens show higher degree of softening beyond the second peak. Cracking occurred by sequential nucleation and coalescence of micro-voids along the

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transformed shear bands in impacted specimens.

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Microvoid nucleation along the transformed bands occurs at the second phase particles

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