Microstructrual evolution of aluminum 6061 alloy throughtube channel pressing

M.H. Farshidi a,b, M. Kazeminezhad a,n, H. Miyamoto b

a Department of Materials Science and Engineering, Sharif University of Technology, Azadi Avenue, Tehran, Iranb Department of Mechanical Engineering, Doshisha University, Kyotanabe City, Kyoto, Japan

a r t i c l e i n f o

Article history:Received 1 June 2014Received in revised form17 July 2014Accepted 18 July 2014Available online 29 July 2014

Keywords:Severe plastic deformationTube channel pressingGrain refinement6061 Aluminum alloyMicroshear-band

a b s t r a c t

The evolution of microstructure and the mechanical properties of an aluminum 6061 alloy tubesubjected to a novel severe plastic deformation process called tube channel pressing were studied.Results show that the tube channel pressing causes an impressive grain refinement and strengthening ofthe material. As an illustration, mostly ultrafine grained tube ðD¼ 760 nmÞ bearing the yield strength of352 MPa can be obtained after 5 passes of tube channel pressing. Additionally, the tube channel pressingcauses the heterogeneous grain refinement resulting in a bimodal distribution of size of grains after 7thpass. Therefore, it can be inferred that the main microstructure refining mechanism caused by tubechannel pressing process is intersection of microshear-bands.

& 2014 Elsevier B.V. All rights reserved.

1. Introduction

Ultrafine Grained (UFG) and Nano Grained (NG) materials areattractive due to their unique characteristics such as the excellentstrength and the special superplastic capabilities. Regularly, anUFGed material is defined as a material with average grain size ofless than 1 mmwhile a NGed material is defined as a material withaverage grain size of less than 100 nm. Among multiple methodsdeveloped for production of UFGed/NGed materials, severe plasticdeformation (SPD) is more considered due to its attractive speci-fications such as lower defects introduced by processing and acapability of production of bulk UFGed/NGed materials. Therefore,many processes have been developed to impose SPD on differentgeometries of materials. For example, Equal Channel AngularPressing (ECAP) for rods, High Pressure Torsion (HPT) for disksand Accumulative Roll Bonding (ARB) for sheets have beenintroduced by Segal et al. [1], Valiev et al. [2] and Tsuji et al. [3],respectively. Nonetheless, few studies were concentrated on SPDof tubes. Naghasekhar et al. [4] and Toth et al. [5] have focused ondevelopment of tubular ECAP and High Pressure Tube Twisting(HPTT), respectively. Nevertheless, these processes can be onlyused for special tubular geometries. Faraji et al. [6] developedTubular Channel Angular Pressing (TCAP) which can be used forSPD of different tubular geometries. Despite so, TCAP imposesrelatively high strains in each pass which increases the risk of

fracture during processing. In addition, the variation consequenceof tube diameter in TCAP may result in wall thinning of tube asshown later [7]. In response to these limitations, a new SPDprocess for tubes has been developed called “Tube ChannelPressing” (TCP) [8]. As shown before, TCP imposes moderate strainin each pass and does not cause tube wall thinning or otherunfavorable dimensional change if a suitable die design is selected[8,9]. Therefore, TCP can be an excellent solution for developmentof UFGed/NGed tubes. On the other hand, aluminum 6xxx alloysare widely used in different industries due to their excellentworkability, corrosion resistance and precipitation hardening cap-ability. It is estimated that 90% of aluminum extruded products arefabricated from these alloys [10]. As an illustration, aluminum6061 alloy is widely used to fabricate seamless tubes for construc-tion, marine and automotive industries. Considering these facts,one may guess that this alloy can be a strong candidate fordevelopment of high strength tubes through imposing of an SPDprocess. Therefore, it is proposed to study the evolutions ofmechanical properties and microstructure of aluminum 6061 alloyduring processing by TCP.

SPD processes are often used in cold and warm regimes wherethe deformation temperature is less than half of the melting point.It is believed that the main grain refining phenomenon duringcold/warm SPD of a ductile metal is continuous dynamic recrys-tallization. As shown by Sakai et al. [11] and Mishra et al. [12],continuous dynamic recrystallization can occur by differentmechanisms such as those explained below:

(a) Multiplication and Migration of Dislocations (MMD): homo-genous multiplication of dislocations in initial grains due to imposing

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Materials Science & Engineering A

http://dx.doi.org/10.1016/j.msea.2014.07.0610921-5093/& 2014 Elsevier B.V. All rights reserved.

n Corresponding author: Tel.: þ98 21 66165227; fax: þ98 21 66005717.E-mail address: mkazemi@sharif.edu (M. Kazeminezhad).

Materials Science & Engineering A 615 (2014) 139–147

of plastic strain, migration and rearrangement of dislocations to formrelatively homo-shaped subgrains with dense dislocation walls,rotation of subgrains and formation of new grains bearing sharpboundaries.

(b) Intersection of MicroShear Bands (IMSB): appearance ofLow Angle Boundaries (LABs) in initial grains due to incidence ofmicroshear-bands (MSB) by plastic deformation, evolution of theLABs to High Angle Boundaries (HABs) due to intersection of MSBswhich results in formation of new grains.

Fig. 1 shows a schematic illustration of these mechanisms. Ascan be seen, the IMSB mechanism results in a bimodal micro-structure in which more coarse grains are surrounded by finergrains. Comparatively, the MMD mechanism results in a relativelyhomogenous microstructure.

Since it needs high mobility of dislocations, the MMD mechan-ism is usually more active in materials bearing high stacking faultenergy such as aluminum alloys. Despite this, the IMSB mechan-ism can occur in different metals and alloys as explained before[11]. In addition, other parameters can affect the activity of thesemechanisms. For example, the lower deformation temperatureand the higher concentrations of alloying elements can result insuperiority of the IMSB mechanism and deactivation of the MMDmechanism as illustrated by Stidikov et al. [13]. The deformationgeometry may also affect the activity of these mechanisms. As anexplanation, the activity of IMSB mechanism is mainly reportedafter imposing of Multi-Axial Forging (MAF) which causes inci-dence of an intersecting shear plane couple [11,13]. Consideringthis, one may guess that the incidence of multi-directional shearplanes and concentration of strains in these planes appear in MAFcan accelerate IMSB mechanism. Nonetheless, no study has con-sidered the effect of deformation geometry on the activity of thismechanism.

Since imposing of SPD causes refining of grains, it is expectedthat the strength of an SPD processed material increases respect-ing the Hall–Petch relation:

σ ¼ σ0þkD�0:5 ð1Þ

Here, σ0 is the frictional strength of the alloy, k is the strengtheningcoefficient and D is the average grain size. Although this relationprovides a relatively accurate estimation of the strength, Hansen[14] and Kamikawa et al. [15] have noticed that the k value isheavily dependent on the imposed plastic strain. For example,while the k value of Al–Mg–Si alloys in annealed conditionhas been reported as 0.086 MPa m0.5, it is evaluated as 0.18–0.33 MPa m0.5 after imposing of different SPD treatments[16–18]. This difference is mainly due to dislocation strengtheningafter imposing of plastic strains. For more accurate evaluation ofthe strength after SPD, the contributions of grain boundarystrengthening and dislocation strengthening can be independentlyconsidered as shown by Kamikawa et al. [15]:

σ ¼ σ0þσDisþσGB ð2ÞHere, σDis and σGB respectively are the contributions of thedislocation strengthening and grain boundary strengtheningobtained as below:

σDis ¼ αMGbffiffiffi

ρp ð3Þ

σGB ¼ k1D�0:5 ð4Þ

where α is a constant, M is the Taylor factor, G is the shearmodulus, b is the Burgers vector, ρ is the dislocation density and k1is the Hall–Petch coefficient in annealed condition.

2. Principles of TCP

The schema of TCP process is illustrated in Fig. 2. As shown inFig. 2(a), the tube is passing through a channel which has a neckzone in the middle. After finishing each pass of TCP, the die,containing mandrel and specimen, is rotated by 1801 and ramstarts the process from the other side of tube. When neededpasses are imposed, the die will be opened and specimen contain-ing mandrel will be expelled from the die. Then, the specimen isfixed and the mandrel is pulled out by pressing.

Fig. 1. Schematic illustration of (a) MMD mechanism and (b) IMSB mechanism.

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As shown in Fig. 2(b), the neck zone can be divided to threesub-zones. As explained before, each of these sub-zones results ina shear strain [8,9]. The consequence of these shear strains issimilar to route C in the conventional ECAP process. Note that thesecond shear stage is twice of the first and third stages. As a result,the direction of any element in tube wall does not change aftereach pass of TCP. In addition, the consequent decrease andincrease of tube diameter during TCP result in accumulation ofhoop strains. Therefore, the TCP imposed strain is a mixture ofshear strains and hoop strains.

The geometries of die and mandrel control deformation beha-vior of tube during processing by TCP. As shown in Fig. 2(c), theprofile of die in neck zone consists of three tangent circles. Themandrel profile has similar geometry. The radius of tangent circlesðRdieÞ and the height of die convex ðΔγdieÞ are main geometricalparameters of a TCP die. The mandrel geometry must be setregarding to the die geometry as shown before [9]. Considering

previous experiments, in this work, Rdie and ðΔγdieÞ were consid-ered 7.5 and 1.5 mm while the tube inner and outer diameterswere 19 mm and 26 mm, respectively. As a result, an averageequivalent strain of about 1 can be obtained in each pass of TCP.

3. Materials and methods

6061 Aluminum alloy tube was received in wrought form andcut to 75 mm long specimens. The chemical composition of theused tube was Al–1.01Mg–0.49Si–0.31Cu–0.24Fe–0.06Cr wt%. Spe-cimens were solid solution treated at 530 1C for 1 h and quenchedimmediately before TCP. Then, they were subjected to 1, 3, 5 and7 passes of TCP at room temperature using MoS2 aerosol as alubricant. Tensile test samples were machined in the longitudinaldirection of TCP processed tubes. The tension tests were achievedat room temperature using strain rate of 5�10�4.

Fig. 2. Principles of TCP.

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Scanning Electron Microscopy (SEM) samples were cut in bothR–θ and R–Z planes and prepared by ion polishing using the IB-9010 cross section polisher. Electron Back Scattering Diffraction(EBSD) mappings were carried out using JSM-7001F machineusing the acceleration voltage of 15 kV. The INCA 4.09 softwarewas used to obtain and analyze EBSD maps. The EBSD mappingsfor each specimen were repeated by three different step sizes of0.5, 0.1 and 0.025 μm. For each step size, EBSD maps were achievedfor at least two different places of specimens to eliminate the effect oflocal microstructural differences. The EBSD quality maps were usedto characterize MSBs whereas darker regions represent the MSBs.Additionally, the EBSDmaps can be used to obtain dislocation densityof specimens as shown by Kapoor et al. [19]:

ρ¼ φloc:�φerr:

bΔxð5Þ

where φloc: is the average point to point misorientation inside majorgrains, φerr: is the error of misorientation measurement and Δx is themapping step. Theoretically, the dislocation density can be measuredin grains as fine as 50 nm using finest step size (0.025 μm). The cut offangle for measurement of dislocation density inside grains wereconsidered as 51.

Comparing the grain size and dislocation density obtained byEBSD, the strengths of tubes were calculated using Eqs. (2)–(4). Forthis purpose, M, α, G, b and k1 were respectively taken in accountas 3.06, 0.25, 26 GPa, 0.286 nm and 0.086 MPa m0.5. The frictionalstrength ðσ0Þ was obtained as 90 MPa by fitting the strength oftube in annealed condition.

Transmission Electron Microscopy (TEM) samples were cut inR–θ plane and mechanically polished to 100 mm thickness. Then,the TEM samples were jet polished in a solution of 75% CH3OH and25% HNO3 at �30 1C for 5–15 min. The JEM-2100F machine withacceleration voltage of 200 kV was used for TEM studies.

4. Results

Fig. 3(a)–(e) compares the EBSD maps of specimens in R–θplane. Equiaxed coarse grains can be observed before TCP asshown in Fig. 3(a). After 1 pass, although few finer grains can beobserved, no overall grain refinement can be traced as can be seenin Fig. 3(b). A significant grain refinement can be observed after3 passes of TCP where a few Microns Sized Grains (MSGs) areformed as illustrated in Fig. 3(c). In addition, the grain sizedistribution is relatively heterogeneous after 3rd pass and coarsegrains have remained in some areas. UFGs accompanying withMSGs can be observed after 5 passes of TCP as presented in Fig. 3(d) implying the obvious grain size heterogeneity of this specimen.In addition, more coarse grains are usually surrounded by finergrains. After 7th pass, the grain size range is between a fewmicrometers to about 100 nm as can be seen in Fig. 3(e). Inaddition, more coarse grains are surrounded by finer grains.Fig. 3(f)–(h) shows the microstructure of tube after imposing 0,3 and 7 passes of TCP in R–Z plane. Considering Fig. 3(a)–(h), it isclear that the trend of grain refinement is similar in both R–θ andR–Z planes.

Fig. 4(a)–(c) shows quality maps of 0, 1 and 3 passes TCPprocessed specimens in R–θ plane. As shown here, MSBs arewidely propagated and intersected after imposing 1 and 3 passesof TCP. Therefore, it is clear that the MSBs affected area spreadswith increasing TCP pass numbers.

Fig. 5(a) compares the variation of average grain size indifferent planes during processing by TCP. As can be seen, thegrain refining in R–Z plane is more rapid than that in R–θ plane.Additionally, the decrease of average grain size has a saturationlimit which is estimated about 0.5 mm. A comparable grain

refinement behavior has been reported for similar alloys subjectedto other SPD processes [20,21]. Fig. 5(b) compares the fraction ofHABs in TCP processed specimen at both R–θ and R–Z planes. Ascan be seen, the fraction of HABs is about 50% before TCP. Thisnumber is decreased after 1 pass due to increase of the LABs as aresult of deformation. By imposing of further TCP passes, thefraction of HABs is increased and approaches to about 80% after7 passes. Similarly, the fraction of HABs has been evaluated as 79%after imposing a plastic strain of 8 using ECAP [20].

Table 1 compares the dislocation density of specimens afterimposing of 0–7 passes of TCP. As can be seen, the dislocationdensity increases by imposing of 1–3 passes of TCP. Despite this,the dislocation density moderately decreases after 3rd pass. Theincrease of dislocation density to a saturation limit by imposing asevere deformation and its moderate decrease by further deforma-tion is a well-known phenomenon reported frequently. Forinstance, Gubicza et al. [22] and Rezaei et al. [23] have reportedthis phenomenon for similar aluminum alloys subjected to otherSPD processes.

Fig. 6(a)–(e) shows the distribution of normalized size of grainsin R–θ plane. As shown in Fig. 6(a), grains are ranged between0 and 2 times of the average grain size ðDÞ before TCP. Despite this,the grain size distribution tends to a bimodal distribution byimposing TCP passes as illustrated in Fig. 6(b)–(e). For example,after 7 passes of TCP, about 35 and 15% of the area fraction areoccupied by grains sized in ranges of 0–1 and 8–9 times of theaverage grain size, respectively. Similar trend can be seen in R–Zplane as shown in Fig. 6(f)–(h). This represents the development ofa bimodal microstructure in aluminum 6061 alloy subjected to TCPprocess as discussed later.

Fig. 7 compares the microstructures of specimens after impos-ing of 3–7 passes of TCP. As illustrated in Fig. 7(a), MSGs appearafter 3 passes of TCP. Selected Area Electron Diffraction (SAED)pattern also shows the existence of few grains. UFGs can be tracedbeside MSGs in the microstructure of 5 passes TCPed specimen ascan be seen in Fig. 7(b). As shown in Fig. 7(c) and (d), NGsaccompany with UFGs and MSGs in microstructure after 7 passeswhich represents the propagation of the bimodal microstructurementioned before. In addition, UFGs and NGs usually surroundmore coarse grains. On the other hand, bright rings can beobserved in the SAED patterns which represent the existence ofnumerous grains.

Table 2 compares the variation of yield strength and elongationafter imposing of 0–7 passes of TCP. As illustrated here, the yieldstrength is increased and the elongation is decreased with increas-ing pass number of TCP from 0 to 5. Despite so, an impressiveincrease of elongation and decrease of strength can be observedafter 7th pass. Comparatively, Kim et al. [24] have reported rapidincrease of strength and fall of elongation in similar alloy pro-cessed by imposing plastic strains of 0–4 using ECAP. Nonetheless,they have evaluated little variations of strength and elongationafter imposing plastic strains of 4–8.

5. Discussion

As a result of microstructural studies, it shall be demonstratedthat the processing by TCP causes the grain refinement ofaluminum 6061 alloy. Additionally, the grain refining is morerapid in R–Z plane compared with R–θ plane. Similarly, a morerapid grain refining has been reported in transversal plane ofspecimens processed by ECAP through route C [25]. This is relatedto higher shear strain in transversal plane of specimens due to theconsequence of route C. When route C is used, the shear directionis always aligned in the transversal plane of specimen which

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Fig. 3. EBSD maps in R–θ plane after (a) 0 pass, (b) 1 pass, (c) 3 passes, (d) 5 passes and (e) 7 passes of TCP; EBSD maps in R–Z plane after (f) 0 pass, (g) 3 passes, and (h) 7passes of TCP.

M.H. Farshidi et al. / Materials Science & Engineering A 615 (2014) 139–147 143

results higher accumulation of strains in this plane compared withcross sectional plane.

As shown in Fig. 4, MSBs are widely propagated duringprocessing by TCP. In addition, the microstructure of tube afterprocessing by TCP is extensively heterogeneous as illustrated inFigs. 3, 6 and 7. In this microstructure, more coarse grains aresurrounded by finer grains which illustrates that the main activegrain refinement mechanism during processing by TCP is IMSB

mechanism. However, the MMD mechanism is also active duringprocessing by TCP and causes increase of the dislocation density asillustrated in Table 1.

Since the IMSB mechanism is based on the “intersection ofmicro shear bands”, multi-directional MSBs are needed for activa-tion of this mechanism. It is noteworthy that the deformationoccurs in one shear plane during processing by conventional ECAProute C as shown by Zhu and Lowe [26]. Therefore, shear bands

Fig. 4. The EBSD's quality maps after (a) 0 pass, (b) 1 pass and (c) 3 passes of TCP.

Fig. 5. Variation of (a) the average grain size and (b) the fraction of HABs during processing by TCP.

Table 1The calculated dislocation density of aluminum 6061 alloy processed by different passes of TCP.

Pass number 0 1 3 5 7

Dislocation density (1/m2) 1012 (4.271.2)�1014 (4.571.4)�1014 (3.371.0)�1014 (2.170.7)�1014

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Fig. 6. Normalized distributions of size of grains in R–θ plane after (a) 0 pass, (b) 1 pass, (c) 3 passes, (d) 5 passes and (e) 7 passes of TCP; normalized distributions of size ofgrains in R–Z plane after (f) 0 pass, (g) 3 passes and (h) 7 passes of TCP.

M.H. Farshidi et al. / Materials Science & Engineering A 615 (2014) 139–147 145

intersection does not occur in this process and shear bandspartially revert in next pass as illustrated by Miyamoto et al.[27]. Regarding so, one may question the role of IMSB in grainrefinement caused by TCP. Nevertheless, it must be consideredthat although the shear consequence of TCP is similar to that ofECAP route C, there are big differences between these processes. Atfirst, the channel is curvy in TCP in opposite of the conventionalECAP. Therefore, shears take place not only in the direction ofchannel's inner angle bisector, but also in the tangent direction ofchannel as shown by Zhilyaev et al. [28]. This results in appearanceof three shear planes couples as illustrated in Fig. 8. Secondly, theTCP has one more strain term compared with the ECAP route Cwhich is the hoop strain. This strain term can also result inappearance of a couple of shear planes which can intersect to

Fig. 7. TEM microstructure of processed tubes after (a) 3 passes, (b) 5 passes and (c, d) 7 passes of TCP.

Table 2The yield strength and elongation of aluminum 6061 alloy processed by differentpasses of TCP.

Pass Number 0 1 3 5 7

Yield strength (MPa) 110710 265711 325710 352712 282715Elongation (%) 3073 970.8 570.7 470.6 1372

Fig. 8. Schematic illustration of shear planes during TCP. Arrows indicate the shearplane and rectangles indicate the shape of an element in each step.

M.H. Farshidi et al. / Materials Science & Engineering A 615 (2014) 139–147146

previously mentioned shear planes. This means that the proces-sing of tube by TCP results in appearance of four active shearsystems in each pass compared with one active shear system ofthe conventional ECAP route C. This is also higher than two activeshear systems of the MAF process which shows high tendency toIMSB mechanism. Therefore, the higher number of shear planescan explain the high activity of IMSB mechanism during proces-sing by TCP.

Calculated and experimentally obtained values of the strengthafter different passes of TCP are compared in Fig. 9. As can be seen,Eqs. (2)–(4) can nearly predict the trend of strength variation ofprocessed tubes. However, the predicted strengths of TCP pro-cessed tubes are usually lower than the values obtained byexperiments. This is related to segregation of alloying elements(Mg, Si and Cu) on dislocations which enhances the strength ofTCP processed specimens as reported before [18]. In addition, afterprocessing by TCP, room temperature aging is accelerated due toincrease of diffusion rate. This can also strengthen the material asdiscussed before [29]. Despite other specimens, the experimen-tally obtained strength of 7 passes TCP processed specimen islower than the calculated amount. In addition, the fall of strengthand the increase of elongation after 7th pass are in contradictionto the little variation of strength and elongation during imposingof strains between 4 to 8 reported by Kim et al. [24]. Consideringthe incidence of a bimodal NG–MSG microstructure after 7th pass,one can relate the fall of strength and the increase of elongation tothis microstructure. In a bimodal NG–MSG microstructure, thegrain boundary sliding can be activated in NGs while MSGsaccommodate cracks initiated due to grain boundary sliding bydislocation slip predominated deformation. This results in adecrease of strength and increase of elongation as discussed before[30,31].

6. Conclusions

The results of this work can be presented as follow:

1- TCP is an effective process for strengthening and grain refine-ment of tubes. Using 5 passes of TCP, an UFGed aluminum 6061alloy tube can be obtained bearing the yield strength of352 MPa.

2- Imposing of TCP process causes heterogeneity in distribution ofsize of grains. As an illustration, a bimodal microstructure

arises after imposing of 7 passes of TCP in which grains aredistributed between 100 nm to a few micrometers.

3- Intersection of microshear bands is the main mechanism ofgrain refinement during processing by TCP. The superiority ofthis mechanism is attributed to activation of four shear systemsby processing through TCP.

4- An impressive strength decrease and elongation increase canbe observed in 7 passes TCP processed specimen. This isattributed to development of a bimodal NG–MSG microstruc-ture caused by TCP.

Acknowledgments

The authors wish to thank the research board of SharifUniversity of Technology and Doshisha University for the provisionof research facilities used in this work.

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Fig. 9. Comparison of the experimentally obtained and the calculated yieldstrength after different TCP passes.

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