Materials and Design 39 (2012) 140–150

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Materials and Design

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High strain rate superplasticity in a nano-structured Al–Mg/SiCP compositeseverely deformed by equal channel angular extrusion

Amir Hassani ⇑, Majed ZabihiDepartment of Materials Science and Engineering, Faculty of Engineering, Semnan University, Semnan 35131-19111, Iran

a r t i c l e i n f o

Article history:Received 17 November 2011Accepted 4 February 2012Available online 28 February 2012

Keywords:A. CompositeC. ExtrusionC. Mechanical alloyingF. MicrostructureG. Scanning electron microscopy

0261-3069/$ - see front matter � 2012 Elsevier Ltd. Adoi:10.1016/j.matdes.2012.02.009

⇑ Corresponding author. Tel./fax: +98 231 3354119E-mail address: ahasani@semnan.ac.ir (A. Hassani)

a b s t r a c t

High strain rate superplastic deformation potential of an Al–4.5%Mg matrix composite reinforced with10% SiC particles of 3 lm nominal size was investigated. The material was manufactured using powdermetallurgical route and mechanical alloying which was then processed by equal channel angular extru-sion (ECAE). The composite showed a high resistance to static recrystallization. The manufacturing oper-ations atomized SiC particles to nanoscale particles and the severe plastic deformation process resulted ina dynamically recrystallized microstructure with oxide dispersoids distributed homogeneously through-out the matrix. These particles stabilized the ultra-fine grained microstructure during superplastic (SP)deformation. Testing under optimum conditions at constant strain rates led to tensile elongations>360%, but it could be further increased by control of the strain rate path. Transmission electron micro-scope (TEM) studies showed that the low angle boundary sub-grain structure obtained on heating to theSP deformation temperature developed on straining into a microstructure containing high angle bound-aries capable of sustaining grain boundary sliding.

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1. Introduction

High strain rate superplasticity (HSRS) has been characterized inultra-fine grained (�50 nm to �3 lm) aluminum alloys and com-posites developed by powder metallurgical routes [1,2] and theother processing methods such as mechanical alloying (MA) [3],friction stir processing [4], spray casting [5] and severe plastic defor-mation (SPD) [6–9]. In addition, to a large number of metallic mate-rials which exhibit superplasticity at high strain rates (P10�2 s�1), avariety of metal matrix composites can also be superplastically de-formed at relatively high strain rates [10]. High strain rate super-plasticity has attracted many metallurgists from a commercialviewpoint, because probably the main restriction in conventionalsuperplastic forming technology has always been the necessity ofapplication of very low strain rates of610�3 s�1 to give enough timefor accommodation of grain boundary sliding or grain rotationthrough diffusional and/or dislocation mechanisms during super-plastic flow. Such a restriction is obviously a central drawback formaterials like metal matrix composites, with a low capability ofbeing shaped or machined through the conventional techniques.Nieh et al. [11,12] showed that HSRS occurred at temperatures closeto, or even slightly above, the solidus temperature of the matrixalloy. These authors also revealed the important role that a liquidphase, formed at the grain boundaries and at the matrix/reinforce-

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ment interfaces, played in high strain rate superplasticity. It is wellknown that in most of particulate metal matrix composites (MMCs),dynamic recrystallization and dynamic precipitation stimulated byfine reinforcement particles (through particle stimulated nucleation(PSN) mechanism) play an important role in grain refinement[13,14] which is essential in superplastic flow to occur.

Kyung et al. obtained HSRS in a commercial 5083 Al alloy byintroducing a ultrafine grained structure of 0.3 mm through severeplastic deformation and by adding a dilute amount of scandium (Sc)as a microstructure stabilizer [6]. Al–4Mg–1Zr extruded bar wassubjected to friction stir processing (FSP) by Ma et al. [4], resultingin generation of a fine microstructure of 1.5 mm grain size. Theyinvestigated superplastic deformation behavior of FSP Al–4Mg–1Zr alloy in strain rate range of 1 � 10�3 to 1 s�1 and temperaturerange of 350–550 �C and compared with that of as-rolled one. Itwas indicated that the FSP alloy exhibited significantly enhancedsuperplasticity at a high strain rate of 1 � 10�1 s�1, and a maximumelongation of 1280% was obtained at 525 �C and 1 � 10�1 s�1.

Possibility of combination of high-strain-rate superplasticity andlow-temperature superplasticity was experimentally confirmedusing extremely fine-grained magnesium alloy by Watanabe et al.[15]. Superplastic behavior was examined using powder metallurgyprocessed Mg–Zn–Zr alloy (ZK6) through tensile testing. Thesuperplastic deformation characteristics and microstructure evolu-tion of the rolled AZ91 magnesium alloys at high strain rates wereinvestigated [16]. The dominant deformation mechanism in highstrain rate superplasticity was still grain boundary sliding (GBS).

A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150 141

The dislocation creep controlled by grain boundary diffusion wasconsidered the main accommodation mechanism observed.

The tensile superplasticity of a Ni/Si3N4(w) composite synthe-sized by electrodeposition was characterized by Chan et al. [10].A maximum elongation of 635% was obtained at a temperatureof 713 �K and a strain rate of 1 � 10�2 s�1.

Magnesium alloys processed by equal-channel angular pressingoften exhibit superplastic elongations at low strain rates. Twostrategies are available for achieving high strain rate superplastic-ity: (i) by pressing the alloys through a reduced number of passesin order to increase the thermal stability of the microstructure; and(ii) by increasing the processing temperature to permit the occur-rence of superplastic flow at higher testing temperatures [7].

Mahmudi et al. investigated the superplasticity of a fine-grainedSn–5 wt%Sb alloy, processed by equal-channel angular pressing(ECAP), by impression testing in the temperature range of 298–370 K. The deformation response of the ECAPed material with agrain size of 2.5 lm conformed to regions I, II and III, typical ofsuperplastic behavior [8]. A 5083 Al–Mg was processed by equalchannel angular extrusion (ECAE) at 150 �C according to variousroutes by Dupuy and Blandin [17]. The resulting microstructureswere characterized and tested in tensile conditions at high temper-ature. A particular attention was given to strain induced damageand to the effect of ECAE on the population of second phase parti-cles initially present in the alloy.

Low-temperature superplasticity (LTSP) and internal friction inmicrocrystalline ZK60 and AZ91 magnesium alloys processed byequal channel angular pressing (ECAP) were reported byChuvil’deev and his co-workers [18]. Low-temperature superplas-ticity was observed in ECAP-processed materials. Internal frictioncould be practically used to determine the optimum temperaturefor superplasticity. A commercial Al-7034 alloy, produced byspray-casting and with an initial grain size of �2.1 lm, was sub-jected to equal channel angular pressing (ECAP) by Cheng and hiscolleagues [5] at 473 K to a total of six passes, equivalent to an im-posed strain of �6. These authors found that in the as-pressed con-dition, there was a reasonably homogeneous microstructure ofequiaxed grains with an average grain size of �0.3 lm and MgZn2

and Al3Zr precipitates served to inhibit grain growth so that anultrafine submicrometer grain size was retained at temperaturesup to and above 600 K.

An elongation of 194% could be obtained in a mechanical alloyed(MA) Al–8 wt%Ti alloy at high temperature and high strain rate. Themajor deformation mechanism of the MA Al–Ti alloy was shown tobe grain boundary sliding at the testing condition. It was suggestedthat the presence of Al3Ti particles in MA Al–Ti alloy, could promotethe formation of high-angle grain boundaries during MA and extru-sion, which is important for easy grain boundary sliding [3].

In the present study, ultra-fine grained (UFG) microstructureobtained from straining an Al–4.5%Mg/SiCP composite manufac-tured by powder metallurgy and mechanical alloying which wasthen severely deformed by equal channel angular extrusion, sug-gested that the material might exhibit a significant superplasticbehavior if tested under optimum conditions of temperature andstrain rate. The present investigation was then carried out to deter-mine whether the material was able to exhibit superplastic behav-ior in general, or even had the ability of being superplasticallydeformed at high strain rates and, if so, to identify the optimumdeformation conditions, and the way in which the microstructureevolved during SP deformation.

2. Experimental details

The composite was manufactured by a powder route andmechanical alloying procedure through mixing an adequate

proportion of Al powder of 50 lm size with micro-scale powdersof Mg, Mn, Ti etc. in a high energy planetary ball mill. To reinforcethe produced powder alloy, 10 volume percent SiC particles of3 lm nominal size was added and blended to a homogenous mix-ture. The mixture was cold-compacted into cans, degassed andsealed in vacuum, and then subjected to hot isothermal pressing(HIP) to 100% density at temperature of 400 �C. This route pro-duced the composite with a matrix having a mean grain size of�3 lm and homogeneously dispersed SiC particles of �500 nm.The chemical analysis (wt%) of the matrix alloy was Al–4.6Mg–0.2Mn–0.2Si–0.11Cu–0.03Ti–0.02Zn–0.01Ni, which was roughlysimilar to the chemical composition of AA 5356 (4.5–5.5Mg–0.05–0.2Mn–0.06–0.2Ti–0.05–0.2Cr).

Before extruding, the annealing treatment of the composite wascarried out at a temperature of 200 �C for 1 h in a vacuum furnace.The ECAE die had a square inlet of 15 � 15 mm and, to provide aback pressure on the sample at the end of the die, a smaller outletof 14.5 � 14.5 mm. The internal angle between the two parts of thechannel was 90� with an outer arc of curvature 20�, where thesetwo parts intersected. The composite samples were cut to14.5 � 14.5 � 60 mm billets and then subjected to ECAE at roomtemperature at a rate of 5 mm s�1. During the extrusion operation,each sample was rotated by 90� in the same direction betweenconsecutive passes through the die (route BC) [19]. The detailedmanufacturing procedures of the composite and ECAE operationswill be explained in a separate article.

To investigate the HSRS deformation potential of the composite,flat tensile specimens of 10 mm gauge length, 5 mm gauge widthand a thickness of 2 mm were prepared by machining the as-extruded billets parallel to the extrusion direction, according toASTM E8M-89b [20].

Some tensile specimens were annealed at 200 �C for 1 h for thepurpose of microstructural studies. High strain rate superplasticitytesting was carried out [21,22] using a three-zone split furnace at-tached to cross-head of an Instron tensile machine which wasinterfaced with a computer program for the strain rate to be keptconstant during testing. Tensile testing to failure was carried outfor a range of constant strain rates, and the strain hardening coef-ficient, ‘‘n’’, was measured [23] as a function of strain from the truestress-true strain data obtained from these tests.

To identify the strain rate at which the material shows the opti-mum superplastic behavior, the strain rate sensitivity index ‘‘m’’was obtained from strain rate jump tests. The different experimen-tal procedures which have been suggested for determining ‘‘m’’values are based on the strain rate change techniques developedfirst by Backofen et al. [24]. In the present work uniaxial tensilespecimens were initially deformed up to an elongation of 40% ata cross-head velocity of 0.05 cm/min. This was to obtain a uniformmicrostructure before the first velocity change was imposed. Thevelocity was then rapidly reduced to a value of 0.0005 cm/minand as soon as the load was stabilized, an increased velocity wasapplied causing a jump in load to be recorded, followed by a loadsteady state. These upwards steps were continued to a cross-headvelocity of 2 cm/min.

From the load-time chart, true stresses and strain rates werecalculated and the relevant curve was logarithmically plottedassuming that the sample neither had any changes in volume dur-ing deformation nor suffered from necking, and that the structureremained stable during the test. The slope of the curve at any pointrepresented the strain rate sensitivity ‘‘m’’ for corresponding strainrate [25]:

m ¼ LogðP1=P2Þ=LogðV1=V2Þ ð1Þ

where Pi is the load at cross-head velocity Vi.To study the effect of strain rate path on superplastic ductility

and flow stress, various rapid strain rate tests to pre-determined

Fig. 1. Schematic of a spot transmission electron diffraction pattern of a [111] zoneaxis.

Fig. 2. Spot transmission electron diffraction patterns for a typical fcc structure. Thezone axis is anti-beam direction, as indicated [27].

142 A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150

strains were also carried out prior to deformation at optimumrates. To supplement the studies of microstructural evolution dur-ing tensile testing, some compression tests were carried out [26]on cylindrical specimens of 25 mm length and 12.5 mm diameterwhich were machined from the as-extruded material. These weredeformed at 280 �C, 320 �C and 380 �C to a 50% reduction in height(e = �0.69) at a strain rate of 1.5 � 10�2 s�1, and then waterquenched from the temperature prior to microstructural examina-tion. The results obtained from compression testing enabled com-paring them with those obtained from tensile testing.

Microstructural studies were carried out using optical micros-copy, scanning electron microscopy (SEM) and transmission elec-tron microscopy (TEM). SEM studies using Philips SEM 525 and505 microscopes were carried out to investigate the effect strainingon the microstructure. The latter was equipped with EDAX analysisfacilities which were used to identify the second phase particles. AJEOL-JSM 6300 scanning electron microscope was also used in back-scattered electron mode for the measurement of misorientationsbetween grain boundaries employing the EBSP technique.

To characterize the microstructural evolution resulting fromstraining, transmission electron microscopy was carried out usingHitachi H700H operating at 200 kV and Philips EM400 operatingat 120 kV both equipped with EDAX. The samples were preparedby spark-machining discs of 3 mm diameter and grinding to about200 lm thickness on a 1200 grit SiC paper. The discs were jet pol-ished in a Struers machine operating at 20 volts with a current of0.2–0.3 amps at a temperature of �30 �C using a solution of 33% ni-tric acid in methanol.

Misorientations between adjacent grains and sub-grains weremeasured in two ways. First, the SEM utilizing EBSP facilities wasused to measure misorientations between grain boundaries auto-matically. Second, the spot patterns obtained from TEM weremanipulated manually, and then calculations were performed byusing a computer program to process the measurements.

(a) Using electron back scattered pattern (EBSP)method: The samplepreparation for this purpose was conducted by electropolish-ing the specimens of 5 mm� 10 mm at�30 �C in 90% metha-nol and 10% perchloric acid after grinding and polishing. In thismethod, performed using a JEOL-JSM 6300 scanning electronmicroscope, a certain computer program used the Kikuchi dif-fraction patterns derived from each grain/sub-grain by theelectrons backscattered from the specimen. A camera installedinside the sample chamber of SEM visualized the Kikuchi pat-terns. This method is much faster and more accurate than thesecond one. Moreover, an advantage of this method is that itenables the microtextures to be studied simultaneously. How-ever, probably, due to presence of abundant sub-grains accom-panying high angle boundary grains in the material underinvestigation, it was revealed that this method had two disad-vantages; firstly, the program was unable to calculate the lowangles misorientations less than 0.5� and secondly, which wasmore important, for a heavily dislocated or an un-recrystal-lized microstructure (which was the case in the present work),revealing the exact Kikuchi patterns of the adjacent grains wasdifficult and sometimes impossible.

(b) TEM spot pattern technique: This technique, which was a longand cumbersome way, was inevitably used for those (sub)-grains that the EBSP program was unable to measure the mis-orientations between them. A variety of quantitativemisorientation measurements were manually carried outby using TEM spot patterns, utilizing a computer programwritten by the authors for this purpose to solve the matricesequations. From those equations the directions of a group ofadjacent grains, and therefore the angle between those direc-tions (i.e. the angle between the grain boundaries) could be

calculated. TEM spot patterns of a group of grains combinedby a selected grain as the reference grain and its surroundingneighbors were first photographed. In Fig. 1 which is a sche-matic of typical spot pattern photograph of a [111] zone axis,the reference direction, the two directions of [abc] and [def]and the angles between them and the reference direction, hand U, are defined. In this method [27], each time, three pairsof the main directions from either of the selected referencegrains and one of its neighbors are considered. These threepairs of directions are two reference directions of [hRkRlR]and [h0Rk0Rl0R], two anti-beam directions (the zone axes) of[hBkBlB] and [h0Bk0Bl0B], and two orthogonal directions, [hOkOlO]and [h0Ok0Ol0O]. An anti-beam direction is read from an fccpre-determined spot patterns map, which can be obtainedfrom related books (e.g. Fig. 2), by superimposing a developedfilm of a grain’s spot pattern (e.g. Fig. 1) on the map. If thesuperimposed spot pattern perfectly matches the map, thenthe indices for the zone axis i.e. [hBkBlB] of that grain will bethe same as that of the map. So, the directions [abc] and[def] are also found from the map and angles h and U aremanually measured. The reference direction (the directionperpendicular to the upper edge of the film, as shown in

A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150 143

Fig. 1) then can be calculated and eventually, the true angle ofmisorientation, c, is calculated. Detailed explanations on thistechnique can be found in Ref. [27].

To identify the grain size variation, the linear intercept method[28] was used to determine the grain size of the alloy examinedusing both optical microscope and TEM.

3. Results and discussion

3.1. Mechanical behavior

Fig. 3 shows the measured mean grain size of different ECAEedsamples in terms of number of passes they were subjected to. Asseen the smallest (sub)grain size of �250 nm was brought aboutafter 8th pass of ECAE, while the as-annealed sample had the larg-est grain size of �7 lm. With increasing the number of ECAEpasses up to 8, grain size of the composite decreases but further in-crease in passes results in increasing the grain size. This is in con-trast with the research reports published by many authors [29–36].Such continuous grain refinement is normally explicated based onthe argument that a decrease in the grain diameter (D) is related tothe strain (e) and initial grain size (D0), and for a plain strain geom-etry [34], e = ln(D0/D). For dynamic recrystallization to take place,the grain diameter (D) has to decrease to a critical value which isof the order of sub-grain size (d) [35,36]. The critical strain (ec) atwhich the high angle grain boundaries (HAGBs) spacing (D) be-comes equal to the sub-grain size (d) is approximately given byabove equation when D � d, i.e. ec = ln(D0/d). This relationship rep-resents an upper limit of strain because the grain diameter may befurther reduced by grain subdivision during deformation [37–39].

Instead, plastic deformation induced grain growth has beenwidely reported in various severe plastic deformation processes.Shear band formation during ECEA may lead to grain coarsening.Special nature of nanostructured materials which is somehow dif-ferent from normal micron-scale materials [40] and strain inducedgrain growth mechanism which is predominant when high amountof accumulated strains (i.e. higher cycles of SPD) are applied,should be taken into account. For ultra-fine grained materials pro-duced by more appropriate ECAP processes like route C and Bc,macroscopic shear banding can be more easily attributed to astrain path change or to grain coarsening and grain rotation. Thereis also the possibility that the shear deformation during ECAP leadsto non-homogeneous band-like shaped grains. Hence, it can beconcluded that firstly, the shear band starts at a locally coarsenedgrains or at a patch which is much more easily deformed. Triggeredby the localized plastic deformation, the coarsened patch spreads

Fig. 3. Measured mean (sub)grain size of different ECAEed samples against numberof ECAE passes.

out and the macroscopic shear band forms. Secondly, the extendedshear band is caused by a strain path change and catastrophicshear instability will be the consequence [40]. The main coarseningmechanism seems to be that some grains grow at the expense ofothers, leading to a coarser grain structure with fewer, largergrains. Occasionally, low angle grain boundaries between grainswith similar orientation simply vanish as the grains rotate intoidentical orientation. During the deformation, a large amount ofmechanical work is done on the system. Most of this energy isdeposited in the grain boundaries, and one could speculate thatthe observed coarsening is simply due to heating of the grainboundaries. However, the thermal conductivity is relatively large,and the heat is quickly spread to the entire grains and then re-moved by the Langevin dynamics [41]. Simulation of the procedureindicates that during the deformation, the temperature of the grainboundaries does not rise significantly over the average tempera-ture of the sample, so it does not seem likely that the coarseningis due to local heating. To check this assumption, a simulationwas performed without any deformation. Instead, the sameamount of energy was deposited directly into the grain boundariesas thermal energy. This was done by increasing the fluctuatingforce in the Langevin dynamics, but only for the atoms in the grainboundaries. No grain coarsening was seen in this case, providingfurther evidence that the coarsening is caused by the applied strain[42].

Zhang and his co-workers [43] who examined grain growth innanocrystalline Cu under the microhardness indenter, found thatgrain coarsening was even faster at cryogenic temperatures thanat room temperature. The results of the work carried out by Liaoet al. [44] on deformation induced grain growth in electrodepositednano-crystalline Ni during high pressure torsion (HPT) indicatedthat high stress and severe strain were required for inducing graingrowth, and the upper limit of grain size was determined bydeformation mode and parameters. Strain induced grain coarseningwas observed during uniaxial compression of an as-deposited bulknanocrystalline Ni–Fe at ambient temperature [45]. SPD-inducedgrain growth has also been reported during uniaxial tension [46].Wang et al. [47] who reported grain coarsening during high pressuretorsion (HPT) suggested that SPD-induced grain growth wasachieved primarily via grain rotation. The grain growth changesthe structure of nanocrystalline materials and therefore will affecttheir mechanical properties. Molecular dynamics simulations andtheoretical analysis have been carried out to understand the mech-anisms of deformation induced grain growth and contradictory re-sults have been reported. Sansoz and Dupont [48] showed thatstress driven nanocrystalline grain growth resulted mainly fromrotation of nano-grains and propagation of shear bands, while Far-kas et al. [49] suggested that, for grain having the same size as thosepresented in Ref. [48] grain boundary migration is responsible forthe nano-crystalline grain growth. Experimental results obtainedvia in situ TEM investigations suggested that deformation inducedgrain growth can be realized via grain rotation and grain boundarysliding [50,51]. Nanograined materials have some unusual proper-ties. Li [52] found that high purity and non-equilibrium structureare necessary conditions for mechanical grain growth. He suggestedthat the material must be pure enough so that free dislocations areavailable to move out of the boundary. But the boundary should notbe in the lowest-energy state so that extra dislocations are availableto be emitted by stress. Based on these conditions, he devised somemethods to avoid low temperature grain growth. Legros et al. [53]strained free-standing nanocrystalline Al thin films in situ in TEMat room temperature and observed extensive grain-boundarymigration accompanied the in situ loading. This grain growth pre-ceded dislocation activity, and measured boundary velocities weregreater than could be explained by diffusive processes. The unam-biguous observations of stress-assisted grain growth were compat-

Fig. 4. (a) Flow stress and (b) m values versus strain rate.

Fig. 5. Elongation to failure as a function of temperature and strain rate.

Fig. 6. True stress vs true strain at various temperature at e9 = 1.5 � 10�2 s�1.

144 A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150

ible with recently proposed models for stress-coupled grain-bound-ary migration.

Probably, grain growth mechanisms during ECAE processes, inparticular, have not been studied yet. Therefore, the mechanismsinvolved in grain coarsening at higher plastic strains of Al–Mg/SiCP

composite will be focused by the present authors in the next paper.All tensile and compression specimens were prepared from thefine-grained 8-pass extruded samples.

The relationship between flow stress and strain rate was exam-ined using cross-head velocity cycling tests [24] for temperaturesof 280 �C, 320 �C, 380 �C and 450 �C. It can be seen in Fig. 4a thatthere is a general decrease in flow stress level with increasing tem-perature, and that at each temperature the logarithmic plot showsthe sigmoidal form usually associated with SP materials. The slopesof the curves which are equal to the strain rate sensitivities, or mvalue, are shown in Fig. 4b. Maximum m-values were obtainedfor strain rate range 10�2 s�1 – 10�3 s�1 for the four temperatures.However, ‘‘m’’ values of <0.3 at 280 �C were marginal for SP behav-ior, but ‘‘m’’ increased to >0.55 at 380 �C. As seen, with increasingthe temperature ‘‘m’’ decreases again and at a high temperatureof 450 �C it is lower than 0.3. Therefore, the optimum temperaturefor superplastic forming was determined to be 380 �C.

Constant strain rate to failure tests were carried out at temper-atures ranging from 220 �C to 460 �C, and the results are summa-rized in Fig. 5. It can be seen for each strain rate that the tensileelongation to failure passes through a maximum at �380 �C, withthe highest elongation of �360% being observed for a high strainrate examined (1.5 � 10�2 s�1). True stress-true strain relation-ships are shown in Fig. 6 for the samples subjected to various ECAEpasses at the optimum strain rate of 1.5 � 10�2 s�1 and tempera-ture of 380 �C. The tests results are in agreement with those indi-cated in Fig. 3. The 12-pass ECAEed sample which initially had alarger grain size and as a result, a higher strength, showed a lower

elongation coupled with a higher flow stress when subjected totensile deformation. But the highest elongation with lowest flowstress was obtained from the ultra-fine grained specimen ECAEedfor 8 passes. Indeed, this specimen which consists of an initial ul-tra-fine grain size stabilized at superplastic temperatures with col-laboration of nano-scale ceramic particles and oxide dispersoids,was capable of sustaining grain boundary sliding. This reduces flowstress and increases elongation to failure, as seen in Fig. 6. For thestates examined, peak stress values were recorded for true strainsof 0.2–0.4. This is probably due to the onset of necking, but despitethis the specimens, except for the as-annealed sample, were able tosustain appreciable strains to failure. Strain rate jump tests werecarried out at 320 �C and 380 �C to examine the variation of m asa function of strain. It can be seen in Fig. 7 that at both tempera-tures, the value of m is too low in the early stages of deformationto maintain substantial tensile stability.

It was first shown by Ash and Hamilton [54] that strain harden-ing could make a significant contribution to tensile stability for al-loys in which the superplastic microstructure evolved by dynamicor strain-enhanced recrystallization. For the present work, mea-sured values of the strain hardening index, n, are superimposedon those of m in Fig. 7. It can be seen that at the start of deforma-tion, although the m values are low, the values of n are relativelyhigh, and strain hardening will make a contribution to tensile flowstability. Nevertheless, at neither temperature is this sufficient toprevent the onset of necking (Fig. 6). However, with increasingstrain rate at 380 �C the values of m is increasing and at the stagewhen n ? 0, it is sufficiently high (>0.3) to inhibit rapid neck prop-agation so leading to a significant tensile strains to failure. Fig. 7shows that the value of m continues to increase with increasing

Fig. 8. Variation of Instability Parameter, ‘‘I’’, with strain at e9 = 1.5 � 10�2 s�1 at thetemperature of 320 �C and 380 �C.

Fig. 7. Variation of m and n with strain.

A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150 145

strain almost to the point of failure. Similar behavior is apparent at320 �C, but at the stage when n ? 0, the value of m is still <0.3. Sur-prisingly, this is sufficiently high to inhibit catastrophic failure andm continues to rise slowly throughout deformation to a value >0.3.

The ability to avoid necking depends on the increase in the flowstress (dr) caused by strain hardening at any cross-sectional posi-tion on the tensile specimen, where the deformation tends to behigher than in the rest of it. If dr is high enough to stop deforma-tion in that cross-section and thus to transfer strain to some otherregion, deformation will be stable and necking will be prevented.According to Hart’s law [55] the deformation is stable if:

dÅ=dA � 0 ð2Þ

where A is the instantaneous cross-section area, dÅ is the growthrate of an instability, dA is the difference in cross-sectional area be-tween the neck region and the normal section.

Caceres and Wilkinson [56] have summarized the combined ef-fects of strain and strain rate hardening, based on Hart’s criterion,by use of an Instability parameter, I:

I ¼ ð1� c�mÞ=m ð3Þ

where c is the hardening parameter defined by Hart [55]:

c ¼ 1=r� ðdr=deÞ ð4Þ

then the rate of development of a neck, dÅ, can be related to m andc. Instability parameter, I, which considers both c and m, can beused to determine the relative stability of the deformation fordifferent conditions. When the value of I is greater than 0, deforma-tion is unstable with rapid neck growth. Conversely, the greatest

stability is indicated by the smallest value of I. By plotting I as afunction of strain for a strain rate in the three superplastic regionsthe stability of a deforming specimen can be predicted. In both Re-gions (I) and (II) low values of I represent delay of instability up tolarge strains. In Region (III) the low m leads to unstable flow and theneck grows rapidly. For the values of I � 0 development of the uni-form strains, eu, is predicted when:

eu ¼ n=ð1�mÞ ð5Þ

The instability parameter, I, is shown in Fig. 8 as a function ofstrain for two temperatures of 320 �C and 380 �C at strain rate of1.5 � 10�2 s�1. It can be seen that for deformation at the lowertemperature, I is maintained less than zero to a strain of �0.6(�85%) after which the instability parameter is kept about zeroto a strain of �1.00 (�170%), indicating development of the uni-form straining (I � 0), and then a rapid increase in the instabilityparameter occurs leading to unstable flow and hence, growingthe neck rapidly.

For the optimum temperature of 380 �C, the tensile stability ismaintained with the values of instability parameter, I, less than zeroto a strain of �0.75 (�110%) after which a larger uniform strainingdevelops (I � 0) to a strain of approximately 1.45 (�320%) and thenthe tensile stability decreases where I increases rapidly. These indi-cate that at an optimum strain rate, the stability of the deformationincreases as the temperature is raised and the larger uniform strainsare expected to develop as the temperature is increased to theoptimum temperature.

Previous studies have shown for alloys in which SP microstruc-ture evolves by strain-enhanced recrystallization that the applica-tion of a rapid pre-strain rate followed by deformation at a lower,near optimum strain rate can lead to a significant enhancement ofsuperplastic ductility [57]. In the present study, the effect of vari-ous rapid strain rates, applied for pre-strains of 0.3 (35% elonga-tion) and 0.4 (50% elongation) was examined and the results areshown in Fig. 9. It can be seen for the range of pre-strain ratesexamined that there are two important effects: the tensile strainto failure increases with increasing pre-strain rate, and flow stresslevels progressively decrease after pre-straining. Highest elonga-tion to failure for the specimen subjected to the highest pre-strainrate can be explained as to the initial ultra fine microstructure, sta-bilized by nano-scale particles, which has been provided by ECAEprocessing, remains almost unchanged throughout superplasticflow. This is due to the rapid deformation that does not give time

Fig. 11. Profiles of tensile specimens (a) undeformed, (b) strained to failure ate9 = 1.5 � 10�2 s�1 and 380 �C, (c) pre-strained to 0.4 at e9 = 8.3 � 10�1 s�1 andstrained to failure at e9 = 1.5 � 10�2 s�1 and 380 �C.

Fig. 9. Effect of pre-strain rate on flow stress and strain to failure at e9 = 1.5 � 10�2

s�1.

Fig. 10. TEM micrograph of an almost recrystallized sub-structure after thespecimen was pulled to 50% at T = 380 �C and 8.3 � 10�1 s�1.

146 A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150

for diffusive mechanisms to occur. Thus, grain growth and the sub-sequent premature tensile failure are prevented. The microstruc-tural studies revealed that a specimen pulled to a strain of 0.4 ata fast strain rate of 8.3 � 10�1 s�1 had maintained its recrystallizedmicrostructure (Fig. 10), hence capable of sustaining grain bound-ary sliding. These microstructural studies support the view that theenhancement of superplastic ductility in this composite material,brought about by rapid pre-strain rates, is related to the micro-structural evolution rather than mechanical behavior such as thestability of deformation based on Hart’s criterion [55]. The effectof pre-strain rate on flow stress and elongation to failure is tabu-

Table 1Effect of pre- _e on flow stress and elongation to failure at optimum strain rate.

Pre- _e (s�1) Pre-e r At pre-eapplied(MPa)

r At e = 0.4 at_e ¼ 1:5� 10�2

(MPa)

Elongationto failure(%)

– – – 24 3603.3 � 10�1 s�1 0.3 26.5 14 3855 � 10–1 s�1 0.3 28.5 16 4058.3 � 10�1 s�1 0.4 39.5 18 510

lated in Table 1. Fig. 11 shows the profiles of tensile specimenpulled to failure, with and without the application of a rapid pre-strain. Hence the benefits of controlling the strain rate path areclearly demonstrated for this material.

3.2. Microstructural examinations

Fig. 12a shows SEM micrograph of the as-extruded compositeindicating a homogeneous distribution of both larger and smallernano-scale particles. Fig. 12b and c display EDS patterns for theseparticles revealing that they are SiC and Al2O3, respectively. SiCparticles were atomized to nano-scale particles by PM and MAprocessing and ECAE operations, and smaller particles i.e. Al2O3

dispersoids were produced mainly during the powder processingand mechanical alloying.

To evaluate the effect of strain rate path at high temperatures onparticle coarsening, a series of SEM examinations were carried out(Fig. 13) on the above mentioned specimens which were subjectedto pre-strain rate tensile testing. Fig. 13a shows both types of parti-cles in the as-extruded material before tensile testing. SEM analysisrevealed that the sample pulled to 50% at highest pre-strain rate of8.3 � 10�1 s�1 show lowest particle coarsening (Fig. 13b) due to lackof enough time for grain boundary migration/rotation leading toparticle coarsening [35], while the specimens deformed at loweststrain rate of 1.5 � 10�2 s�1 experiences higher particle coarseningduring SP deformation (Fig. 13e) because of appropriate time forthe particles to grow at SP temperature. It can be inferred fromFig. 13b–e that the higher strain rate the lower particle growth isachieved.

TEM examination revealed a heavily dislocated sub-structurewhich showed a high resistance to static recrystallization (Fig. 14aand b). In general, small particles, normally less than 1 lm in diam-eter retard recrystallization by pinning sub-grain boundaries andrestraining sub-grain coalescence [58]. The pinning effect of the par-ticles on sub-grain boundaries is clearly seen in Fig. 14b. It has beenproposed by Humphreys [59] that recrystallization is retarded if

Fv=d � 0:1 ð6Þ

where Fv is the particle volume fraction and d the particle size. Forthe SiC particles in the present material Fv � 0.1 and d � 0.08 lm,thus fulfilling the conditions predicted for the inhibition of recrys-tallization. In addition to these particles, presence of fine oxide dis-persoids would also enhance the pinning of sub-grain boundaries. Adislocated sub-structure would also tend to be retained if the driv-ing force for recrystallization was reduced by hot working [60]which is consistent with the method of production of the compositebeing investigated.

Fig. 13. SEM; coarsening of SiC and Al2O3 particles during superplastic deformation; (a) as-extruded sample and (b) pulled sample to 50% at pre-strain rate of 8.3 � 10�1 s�1,(c) specimen pulled to 30% at pre-strain rate of 5 � 10�1 s�1, (d) specimen elongated to 30% at pre-strain of 3.3 � 10�1 s�1 and (e) specimen pulled to 30% at strain rate of1.5 � 10�2 s�1.

Fig. 12. (a) SEM micrograph of as-extruded composite, (b and c) EDS patterns for SiC and Al2O3, respectively.

A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150 147

Fig. 15. Effect of tensile deformation on boundary misorientation in sheet material.

Fig. 16. Mean values of boundary misorientation for specimens deformed intension and compression.

Fig. 14. TEM; (a) as-extruded material showing dislocated structure, (b) annealed at 380 �C for 60 min showing dislocations and sub-boundaries pinned by dispersoids.

148 A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150

On straining at the optimum superplastic temperature of380 �C, it was noted that the strain rate sensitivity, m, was initiallylow, and this would be expected if relatively few large angleboundaries, capable of sustaining grain boundary sliding, werepresent. If the material is to exhibit superplastic behavior it isnecessary for the low angle boundaries to be converted to higherangle boundaries. The observation that m increases with increasingstrain indicates that it is occurring, particularly at 380 �C (Fig. 7).However, in the early stages, deformation must occur by intragran-ular slip involving the generation and movement of dislocations,and this would lead to high initial rate of strain hardening (n)which is observed (Fig. 7). If the dislocations are absorbed bysub-grain boundaries this would result in increasing misorienta-tion across the boundaries, and eventually to a change indeformation mechanism from slip to grain boundary sliding. Thischange is consistent with increasing in m and decrease in n ob-served (Fig. 7).

Measurements of grain boundary misorientation have beenmade for the as-extruded material and for specimens deformedin both tensile and compression and the results are shown in Figs.15 and 16. Fig. 15 shows the distributions of misorientation in theas-extruded material and in the tensile specimens deformed at380 �C to elongations of 50% and 100% (e = 0.4 and 0.69, respec-tively). The as-extruded composite shows mainly low angle bound-aries with relatively few high angle boundaries. On deformation at380 �C there is a progressive increase in the number of high angleboundaries. The corresponding TEM microstructures are seen inFigs. 17a and 15b. At a strain of 0.4 there is still evidence of dislo-cations and sub-grains in the microstructure (Fig. 17a), but after astrain of 0.69 the material appears to be substantially recrystal-lized (Fig. 17b).

The mean misorientations of boundaries are shown in Fig. 16,both for the data in Fig. 15 and also for the specimens deformedin compression. It can be seen that for the compressive deforma-tion at 320 �C there is only a relatively small change in misorienta-tion. At 380 �C the changes are more marked, although still less

Fig. 17. TEM of specimens deformed in tension showing (a) evidence of sub-grains and dislocations (e = 0.4), (b) an almost fully recrystallized microstructure (e = 0.69).

A. Hassani, M. Zabihi / Materials and Design 39 (2012) 140–150 149

than those for tensile deformation at the optimum deformationtemperature. A work by Salishchev et al. [61] showed a similarmode of boundary misorientation evolution for titanium duringdeformation.

4. Conclusions

An ultra-fine Al–4.5%Mg/SiCP composite was produced by pow-der metallurgy route and mechanical alloying procedure and thenwas severely deformed by equal channel angular extrusion (ECAE)for 8 passes. With increasing the number of ECAE passes up to 8,grain size of the composite decreased but further increase in passesresulted in increasing the grain size which was proven to be basedon plastic deformation induced grain growth mechanism. Thismaterial showed a significant potential for superplastic deforma-tion. Testing at constant strain rate under optimum conditions forSP tensile flow led to elongations >510%. The as extruded materialhad a microstructure consisting of heavily dislocated sub-grainswhich was reinforced with a relatively large volume fraction ofnanoscale SiC particles. The composite had a high resistance to sta-tic recrystallization. On deformation at the optimum temperature,strain hardening made a significant contribution to the initial ten-sile stability while the strain rate hardening parameter, m, wasincreasing from its initially low value. The application of a relativelyrapid pre-strain rate led to an enhancement of tensile ductility, andto a fall in flow stress. Measurement of boundary misorientationsfor specimens deformed in tension or compression showed thatthe initial low angle boundary sub-structure developed on straininginto a fine microstructure with a high proportion of large angleboundaries capable of sustaining SP flow.

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