113Use of TEM Kikuchi bands for microstructure and thermal stability study of ECA pressed AA1200...

Corresponding author: M. Cabibbo, e-mail: m.cabibbo@univpm.it

Rev.Adv.Mater.Sci. 25 (2010) 113-121

© 2010 Advanced Study Center Co. Ltd.

USE OF TEM KIKUCHI BANDS FOR MICROSTRUCTUREAND THERMAL STABILITY STUDY OF ECA PRESSED

AA1200 VIA ROUTES A, C, Bc

M. Cabibbo

Department of Mechanical Engineering, Marche Polytechnic University, Via Brecce Bianche, 60131-Ancona, Italy

Received: September 17, 2009

Abstract. The evolution of the microstructure during ECAP depends upon the shearingdeformation paths introduced into the material. Some authors, using selected-area electrondiffraction, reported an effectiveness of cell (low-angle boundaries) evolution toward high-angleboundaries on the order Bc > C > A. This study focused on the cell and grain boundary evolutionof an AA1200 during ECAP, up to a true strain of =8.64, in routes A, C, and Bc. TEM Kikuchi bandswere used for the quantitative evaluation of boundary misorientation. A different hierarchy of cellevolution to grains was found, and this was as C > Bc > A; while in terms of reducing grain size,the route efficiency remained as Bc C > A. Thermal stability of the severely deformed alloy wasstudied by annealing the alloy at 0.5, 0.6, 0.7 TM for 2 hrs. It appeared that, even if route Bc is theone involving a faster microstructure grain refining, route C is likely to be the most stable uponreheating.

1. INTRODUCTION

Methods for obtaining ultra-fine-grained materialsare currently attracting a great deal of attention.This is primarily due to the superior mechanical andphysical properties the ultra-fine-grained materialsexhibit such: increased strength, hardness, wearresistance [1,2]. Severe plastic deformation (SPD)techniques are able to produce bulk materials withfine grained structure [3]. Among the SPD tech-niques, equal channel angular pressing (ECAP)under particular circumstances is able to give grainsizes in the submicrometre range of ~400-800 nm[3-6].

In ECAP, the microstructure evolution with accu-mulative strain can be easily driven and optimizedby modifying the specific shear deformation path. Infact, the microstructure evolution during ECAP isdirectly linked to a complex evolution of dislocationnetworks. That is, dislocation recombination andannihilation phenomena drive the microstructuremodification to a quick and effective grain refinement.The way the billet is repetitively introduced into theL-shaped channels of the die determines the spe-cific shear deformation path. Thus, the billet can beintroduced into the die with no rotation betweenpasses (unidirectiona] straining path – route A . A]-ternatively, billet can be introduced by rotating it of

114 M. Cabibbo

180 deg between passes (route C), or with a rota-tion of +90 deg in the same direction (route Bc) [8-12]. Other routes are known in literature, but for asake of simplicity here are mentioned only the onesused in this study.

To date, a number of researchers have studieddifferent aspects of the specific cell and grain evolu-tion under different strain paths (routes) with differ-ent techniques [8-12]. For instance, in a study ofpure aluminium, deformed at room temperature bya 90° equal-channels die, to a strain of 4, Iwahashiet al. [6,11] reported that the effectiveness of cellsto evolve most rapidly into an array of high angleboundaries (HABs) is on the order Bc > C > A. Theyalso reported route Bc to be the most effective routefor producing equiaxed fine grain structure with HABs.In route A, a large portion of high angle boundaries(HABs) is generated, route C generates cycles oflike deformation at each 2N passes, cycle of defor-mation extended to 4N passes in route Bc. Uponthese considerations, the grain refining is expectedto be less effective in routes Bc and C, compared toroute A. Route Bc appears to be preferable to routeC as it does not induce any deformation in the planenormal to the pressing direction [8]. Route Bc iseffective in refining the microstructure, because foursets of shear bands are introduced into the materialat each 4N-cycle. In route A, the shear bands,formed in the second pass, intersect the billet axisof symmetry at an angle of 45 deg, while the shearbands that formed in the first pass changed orienta-tion to an angle of about 20 deg to the billet axis.The same procedure repeats at the following passesand continuously transforms the previous shearbands into an elongated grained structure along theflow direction. Apart from the shear deformation pathsto which the material is subjected, Zhu and Li [8,12]showed that the die L-shaped corner plays a keyrole in determining which route is likely to more ef-fectively generate HAB into the material microstruc-ture. Thus they reported route Bc as the most ef-fective deformation route for grain refinement in thecase of a 90 deg L-channel die corner die, whileroute A as most effective when using a 120 deg die.

Whatever the route used, grains subdivide bythe effect of newly introduced HABs on finer andfiner length-scales, until ultimately a limit is reachedwhere the HAB spacing converges with the cell size[13]. As a consequence, the boundary evolution pro-ceeds from cells that have only cells as nearestneighbours towards cells having both grains and cellsas nearest neighbours [14]. Average spacing be-tween dislocation boundaries (cells) is a common

microstructure parameter in deformed materials dueto its inverse relationship with flow stress, as amplydocumented in the literature [15]. Studies ofmisorientation angles and of the early cell (oftencalled subgrain) formation process are less investi-gated (e.g. [1,16-18]). Among these, very few pa-pers reported TEM systematic studies on LAB andHAB misorientation evolution with strain for differentstrain paths [7,16,18-20]. These studies mainly fo-cused on the correlation between deformation pathand the resulting microstructure. The present workreports a TEM characterization of the evolution ofLAB and HAB misorientation with accumulativestrain. Kikuchi bands were used to quantitativelyestimate the boundary misorientation, and then tomeasure cell and grain size.

An important feature concerning the ultra-fine-grained structure of metals, which has always totake into consideration, is the thermal stability ofthe severely deformed microstructure. That is, for amechanical point of view, refined microstructuresmust retain their sub-micrometric structure uponheating at least up to 0.5 TM (where TM is the alloymelting temperature). Several studies have reportedthermal stability and superplasticity behaviour ofECA pressed dispersoid-containing aluminium-mag-nesium alloys, showing a rather good thermal sta-bility due to the presence of the Sc-dispersoids [21-23]. With this respect, Langdon et al., in severalstudies [23-25], reported that a refined microstruc-ture has a mechanical meaningful application onlyif it is retained (i.e. no recrystallization occurs) uponheating. It is straightforward that there will be a tem-perature threshold-like limit. This temperature willdepend on the specific alloy (in the present case,aluminium alloy), the level of severe plastic defor-mation, and the thermo-mechanical processes towhich the material is subjected.

In this study, the microstructure of the alloy sub-jected to ECAP, via different strain paths (routes A,C and Bc), up to a strain of e=8.64, was quantita-tively characterized using transmission electronmicroscopy. Thermal stability studies were carriedout for three different temperatures (0.5 TM, 0.6 TM,0.7 TM) in all the three routes: A, C, and Bc.

2. EXPERIMENTAL DETAILS ANDMETHOD

For this study a commercially pure aluminium(AA1200: 0.7Si, 0.3Fe, 0.1Zn, 0.03Cu, 0.02Mn,wt.%) was used. The material was machined to rod-shaped bars 10 mm in diameter and 100 mm inlength, and then homogenised (540 °C / 4 h). The

115Use of TEM Kikuchi bands for microstructure and thermal stability study of ECA pressed AA1200...

microstructure of the bars was homogeneous withaverage grain size of 330 m. Bars were subjectedto severe plastic deformation by ECAP at room tem-perature with pressing forces ranging 20 to 70 kNand speed of 4 mm min-1. ECAP was carried outusing an open die whose two parts were bolted togive a single internal L-channel of cylindrical shape.The two cylindrical channels intersect at an angle

= 90 deg and a curvature = 20 deg with equalcross section of 10 10 mm2. This die configurationinduces to the specimen a true strain of = 1.08 ateach pass [6]. The alloy was subjected to threedifferent routes: A, uniaxial straining, C, reversestraining of 180 deg at each pass, and Bc, +90 degrotation at each consecutive pass.

Microstructure inspections of the ECAP speci-mens were carried out using a transmission elec-tron microscope. TEM specimens were ground to athickness of about 170 m and then surfaces werechemical etched with a solution of perchloric acid(HClO4) in methanol at room temperature for fewminutes. Samples were sectioned along the Z-plane[9]. Thin discs were extracted from the TEM speci-mens, along the central zone of the Z-plane, andprepared to the last stage using electro-polishing at-40 °C, 18 V with a solution of 30 pct nitric acid(HNO3) and 75 pct methanol (CH3OH). The Z-planewas keep as inspection plane in all the three routesused in this study. Thus, in route A and C section-ing was constantly performed along the longer axisof symmetry of the billet out of a plane containingthe pressing and transverse directions. In route Bc,sectioning of the samples was performed followingthe +90 deg billet rotation at each ECAP pass. Thus,considering the literature definition, inspections werecarried out in a sequence: +Z, –Y, –Z, +Y.

TEM discs were examined in a PHILIPS CM-200TM equipped with a double-tilt specimen holder,operated at 200 kV.

The angle of boundary misorientation was mea-sured either form Kikuchi patterns [18,19,26] or fromMoiré fringes [27]. More than 250 boundaries weremeasured per strain level for each of the three routes.The orientation of the different cells were measuredby taking a cell as reference and calculating therelative misorientation of the different surroundingcrystals. The misorientation matrix is thus as RCR =RPR RCP, where RCP is the matrix of the anglesbetween the crystal and the electron beam, and RPR

the matrix of the angles between the reference crys-tal and the beam direction (which is not necessarilyat a low-indexed zone axis). Moiré fringes were evalu-ated, instead of Kikuchi analysis, for those bound-

aries showing them, and typically in cell boundarymisorientation < 2°. The angular measurementwas carried out by counting the number of Moirédark-bright pairs, D, directly on stage. The corre-sponding misorientation across the boundary is thendetermined as = 2-0.5 a/D, where a is the crystallattice spacing [18,27].

Average cell, , and grain, d, spacing betweenboundaries was determined using test lines drawnon TEM micrographs at equal distance parallel andperpendicular to the direction of cell elongation. Dueto the microstructure inhomogeneity induced by theshear deformation throughout the material, the re-gions from which the quantitative misorientationevaluation was performed were selected to be rep-resentative of the microstructure features. Thus, foreach experimental condition, at least three sets ofmeasurements were performed from within the shearbanding zones and form within zones outside theseregions.

Polarized optical microscopy (POM) was car-ried out on central sections along the Z-plane. Sur-faces were ground and etched using the same meth-ods followed for the TEM specimens. Mean grainsize was measured using parallel test lines drawnon the POM micrographs.

3. RESULTS AND DISCUSSION

Fig. 1 shows the cell structures after four and heightECAP passes. Grain subdivision starts first at thelargest length scale (i.e. with primary deformationbands), and then, moves down to the next availablelevel (grain walls and eventually cell walls) with ac-cumulative strain. A banded structure of elongatedcells / grains develops on route A (Figs. 1a and 1d).This is due to the monotonic increase of the sheardeformation induced in this route. After the first pass,the shear bands form an angle of ~45 deg with re-spect to the pressing direction. As the strain accu-mulates, the deviation of shear band inclination tothe pressing direction progressively rises. By con-trast, cells / grains morphology is essentially re-stored at every 2 and 4 passes, in route C and Bc,respectively. As a consequence, cell and grain struc-tures look much more equiaxed with accumulativestrain in these two deformation routes (Figs. 1b,1c, 1e, 1f). In route Bc, the homogeneity of the cellsis evident at both strain levels ( = 4.32 and =8.64).

Fig. 2 gives the distribution of misorientation, ,in all the three routes. After the first pass, 75 pct ofthe boundaries have low-angle character ( <15 deg).Distribution in route A (Fig. 2a) appears to evolve

116 M. Cabibbo

Fig. 1. TEM images of cell structure after 4 and 8 passes for routes A (a), (d), route C (b), (e), and route Bc(c), (f). The pressing direction (PD) is also reported.

Fig. 2. Cumulative frequency distribution ofmisorientation in the three routes: A (a), C (b),and Bc (c). Every plot is labelled with the corre-sponding number of ECAP passes. The Mackenziedistribution is reported as dotted line plot.

from low-angle boundaries to high-angle boundariesquite rapidly. After a strain of = 3.24 (3rd pass), astep at ~ 2 deg starts to develop and consolidate inthe distribution. A second step at misorientationrange of between 15 and 18 deg starts to developat a strain = 4.32 (4th pass), while a step at = 45

deg only appears at = 8.48 (8th pass). The cumu-lative misorientation slowly approach to theMackenzie distribution but remains quite distant fromit even at the maximum strain. On the other hand,route C distribution (Fig. 2b) approaches to theMackenzie distribution much quickly, and it starts

117Use of TEM Kikuchi bands for microstructure and thermal stability study of ECA pressed AA1200...

showing random distribution for misorientationangles f > 40 deg, at the 8th pass. Thus, the genera-tion of high-angle boundaries is, to some extent,slower. Distribution of the different strain levels ap-pear more separate and shows step-likemisorientation increment from a strain level of E =3.24 (3rd pass). Misorientation distribution of routeBc (Fig. 2c) is not dissimilar to the one of route C.Yet, some relevant differences can be highlighted.At the even-numbered passes 2n the distributionappears quite similar to that of the preceding odd-numbered passes 2n -1. Moreover, the rate of evo-lution in odd-numbered passes is higher than in theeven-numbered passes, where the distribution re-mains practically constant or, in some cases, itslightly increases. At the maximum strain of =8.64, the fraction of the low-angle boundaries, F(< 15 deg), is still 30 pct and, as a matter of fact, itis still higher than the expected random orientation(Mackenzie). Steps in the curves indicate the pres-ence and evolution of a pronounced texture in theobserved regions [28]. Yet, the limited area of in-vestigations makes not possible to use the existingsteps in the misorientation distributions for a reli-able interpretation of the texture evolution with ac-cumulative strain.

Fig. 3 has been derived from Fig. 2 and reportsthe fraction of boundaries grouped in 15 deg rangesof misorientation: < 15 deg - Low Angle Bound-aries (LABs), 15 < 30 deg - Low High AngleBoundaries (LHABs), 30 < 45 deg - MediumHigh Angle Boundaries (MHABs), 45 deg - VeryHigh Angle Boundaries (VHABs). A rather mono-tonic decrease of LABs is common to all the threeroutes. In particular, the shrinking of low-angle frac-tion with strain in route Bc is characterized by aclear step-like behaviour. At every 2n passes, thefraction is essentially equal to the previous 2n -1pass. In route A, a relevant increase of the LHAB,MHAB, and VHAB is only observed at a strain of = 5.40 (after 5 passes). In route C, in spite of apeak at 40 pct of the LHAB at = 3.24, LHAB andMHAB do not show a relevant increment with strain,the VHAB population continuously increases withstrain from 6 pct, at = 1.08 (1 pass) to 42 pct at = 8.64 (8 passes). In route Bc, the population re-duction of the LAB contributes to a slow, but con-tinuous, unit increment of all the HAB population(LHAB, MHAB, and VHAB) with similar rate. Theunit increment of all the three HAB ranges is rathermodest accounting in some cases to a two-fold riseafter 8 passes compared to the first pass fraction.This reflects the still relatively high fraction of theLABs at the maximum strain of = 8.64, in all the

three routes. Nevertheless, the fraction of the LHABis constantly the highest among the group of theHABs with the only exception of route C, startingfrom the 5th pass. That is, LABs evolve into HABsrather gradually and the depletion in the LAB distri-bution with strain essentially contribute to the popu-lation of the LHABs. These progressively evolve intoMHAB and eventually to VHAB. Route Bc shows a

Fig. 3. Frequency of misorientation f in different 15deg interval ranges. Open-circle datapoint: LowAngle Boundaries (LABs, < 15 deg); solid-circledatapoints: Low High Angle Boundaries (LHABs, 15deg < 30 deg); solid-triangle datapoints: Me-dium High Angle Boundaries (MHABs, 30 deg <45 deg); solid-square datapoints: Very High AngleBoundaries (VHABs, 45 deg). Route A (a), C(b), and Bc (c).

118 M. Cabibbo

linear and almost monotonic growth of the faction ofall the three HAB ranges with strain. It is likely thatthe route Bc, among the three routes here consid-ered, drives the generation of the HAB with strain inthe most progressive and effective manner, from theLAB to LHAB, MHAB and then VHAB.

Shear plane changes at each ECAP pass, in-duced in route Bc, are likely to favour shear bandintersections, which develop a near spatially uni-form network in the material, and consequently leadsto a higher efficiency in HAB generation. Furukawaet al. [29], in an Al-Mg alloy, suggested that a wideangular range of slip traces might favour the forma-tion of fine equiaxed grains. It may account for thepresent observations, that among the three routesstudied, the grains developed in route Bc have themost equiaxed shape and the smallest size.

The mean fraction of HABs, fHAB, as a function ofECAP strain is shown in Fig. 4. At the maximumstrain of = 8.64, fHAB reaches 66, 81 and 71 pct, inroute A, C and Bc, respectively. The unit incrementof fHAB shows some differences among the threeroutes. Route A is the only one clearly showing a S-shape curve, which is likely to mean that a satura-tion limit of HAB is reached at = 8.64. The plot ofroute C constantly stays above the other two curvesmeaning a faster generation of HAB since the earli-est strain levels. The plot of route Bc shows an al-ternate distribution of datapoints across the inter-cept trend curve, which reflects the different dislo-cation accumulation and annihilation phenomenaoccurring at each 2n pass respect each previous2n -1 pass. Among the three different routes, at theintermediate strain level, = 4.32, the material pro-cessed via route A and route Bc contained the low-

Fig. 4. Mean fraction of HAB, HAB, as a function of ECAP strain for routes A, C, and Bc.

est density of high angle boundaries with a high-angle grain boundary fraction of only 33 pct, whileroute C reached an almost double fraction, 63 pct.At the maximum strain, = 8.64, the HAB fractionin route A is of 66 pct, in route C is higher (81 pct),and intermediate in route Bc (71 pct). Thus, thisfigure invariably shows that the efficiency of gener-ating HAB is as: C > Bc > A.

In other studies [11] the efficiency of generationof HAB was reported to be as Bc > C > A, whichwould mean a much larger HAB fraction, fHAB, inroute Bc respect to the present findings. This dis-crepancy is believed to be attributed to a quite dif-ferent evaluation method for the HAB fraction evalu-ation out of their boundary misorientation measure-ments. In the present study, a systematic TEM studywas carried out using a quantitative method throughthe Kikuchi bands misorientation measurement.This method is far more reliable than other semi-quantitative, or even qualitative, methods. Otherquantitative techniques have also been developedand now commonly used for these purposes. Amongthese, the FEGSEM Electron-Back-Scattered Dif-fraction (EBSD) method is quite attractive and time-saving. Nevertheless, results obtained by usingEBSD suffer from an overestimation of HABs as themisorientation reso]ution ]ies within 2 – 3 deg and,more important, the boundaries showing Moiréfringes can not be detected. Moiré fringes that weremuch more abundant in route Bc compared to routeC microstructure. This feature can explain the dis-crepancy of the present findings with data publishedby Huang, Cao, and others using FEGSEM EBSD[7,19,30-32].

119Use of TEM Kikuchi bands for microstructure and thermal stability study of ECA pressed AA1200...

In order to test the thermal stability, the severelydeformed alloy was annealed at 0.5,0.6, and 0.7 TM

/ 2 h (where TM = 933K is the alloy melting tempera-ture). Figs. 5a and 5b show POM of the post-ECAPannealed alloy, after the 4th and 8th pass, respec-tively. The 4 passes-ECAP and annealed microstruc-tures show a grain coarsening with annealing tem-peratures. In particular, route A experienced the mostdramatic grain coarsening with a microstructure al-most no longer showing traces of the fine-grainstructure obtained during ECAP. The annealed mi-

Fig. 5. Representative POM images of the three routes after the post-ECAP annealing. (a): microstructureafter 4 passes and annealing at: 0.5 TM (330 °C), 0.6 TM (390 °C), and 0.7 TM (430 °C). (b): microstructureafter 8 passes and annealing at same temperatures.

crostructures of the 8 ECAP passes condition at0.5 TM show large portion of ultra-fine grained struc-ture, especially for route A and Bc. As the anneal-ing temperature rise, grain coarsen to a great ex-tent in route A, while route Bc still shows a reason-able fine grain structure, with limited coarsening.The case of route C is somewhat different. The =8.64 + 0.5 TM annealed conditions showed a grainstructure coarser in the average, but much morehomogeneous, respect route A and Bc. The grainstructure remains essentially stable with the anneal-

120 M. Cabibbo

ing temperature. That is, the POM showed grainscoarsened upon annealing at 0.5 TM, but the grainsdid not further coarsen significantly at 0.6 and 0.7TM, either in the 4 and 8 passes conditions.

Fig. 6 shows the evolution of the grain spacing,d, either for the ECAP and post-ECAP annealedconditions. Mean grain size of the ECAP threeroutes monotonically reduces from 2.7 mm, afterthe first pass ( = 1.08), down to ~700 nm in routeA, and ~400 for routes C and Bc after 8 passes ( =8.64). Upon reheating to temperatures ranging 0.5-to-0.7 TM, grain size practically raised by an orderof magnitude, irrespective of the route used. Even ifin any case the mean grain size coarsened to thelevel of the undeformed condition (d = 330 m), themicrostructure after annealing can not longer beconsidered as ultrafine or sub-micrometric. A staticrecrystallization occurred during annealing. The re-crystallization rate is directly driven by the amountof severe plastic deformation the material underwentduring ECAP and by the following annealing tem-perature. The ultra-fine grained structure instabilityupon annealing is certainly due to the absence ofany kind of dispersoids or fine precipitates in thealloy. The only significant results concerns the routeC where microstructure revealed a grain size homo-geneity over the three annealing temperatures andthe two strain levels ( = 4.32 and = 8.64).

Fig. 6. Grain size, d, after = 4.32 (4 passes) and = 8.64 (8 passes), and after post-ECAP annealing to0.5, 0.6, and 0.7 TM. mean grain size of the undeformed homogenised alloy condition is also reported.

4. CONCLUDING REMARKS

This study focused on the role of three different ECAPstrain paths, route A, C, and Bc, in the microstruc-ture evolution prior and after annealing at 0.5, 0.6,0.7 TM (TM: alloy melting temperature). An AA1200and a = 90 deg ECAP die were used. Boundarymisorientation has been successfully characterizedusing Kikuchi bands and Moiré fringes on TEM.These techniques, although time-consuming, are farmore reliable than other automated techniques, suchas the FEGSEM OIM which also use Kikuchi bandsfor the boundary misorientation measurements. Thelower boundary misorientation accuracy limit wasof = 0.1 deg. The role of the boundary showingMoiré fringes was also clarified and it is believed tobe the key aspect for the different route efficiencyhierarchy found in the present work compared tosome other literature findings.

It was found thus that:(i) in terms of generation of HABs, the route effi-

ciency is C > Bc > A;(ii) in terms of reducing grain size, the route effi-

ciency is Bc > C > A;Results of the material subjected to annealing

showed that:(iii) static recrystallization occurred in all the three

routes. In particular, route C showed rather goodgrain size uniformity over the annealing tempera-tures and the strain levels.

121Use of TEM Kikuchi bands for microstructure and thermal stability study of ECA pressed AA1200...

ACKNOWLEDGEMENTS

The author acknowledges Mr. D. Ciccarelli for thePOM specimen preparation.

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