Materials Science & Engineering A 640 (2015) 42–50

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Microstructure and mechanical properties of an Al–Mg–Si tube pro-cessed by severe plastic deformation and subsequent annealing

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

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

a r t i c l e i n f o

Article history:Received 17 March 2015Received in revised form26 May 2015Accepted 28 May 2015Available online 28 May 2015

Keywords:Severe plastic deformationRecrystallizationAnnealingTubeAl–Mg–Si

x.doi.org/10.1016/j.msea.2015.05.09993/& 2015 Elsevier B.V. All rights reserved.

esponding author. Fax: þ98 5138806057.ail address: farshidi@um.ac.ir (M.H. Farshidi).

a b s t r a c t

This study is aimed to realize evolution of microstructure and mechanical properties of aluminum 6061alloy tube subjected to Severe Plastic Deformation (SPD) and subsequent annealing. For this purpose, thetube is initially processed by different passes of an SPD process called Tube Channel Pressing (TCP) andthen subjected to a subsequent annealing at 473 °K for 2 h. Afterwards, tension test is used for theevaluation of mechanical properties while Electron Back-Scattered Diffraction (EBSD) equipped ScanningElectron Microscopy (SEM) and Transmission Electron Microscopy (TEM) are utilized for the micro-structural characterizations. Results show that the Continuous Static Recrystallization (CSRX) is the mainrestoration phenomenon during annealing of aluminum 6061 alloy, even after imposing a moderateplastic strain. For instance, CSRX has been observed during annealing treatment after imposing anequivalent plastic strain as low as 1. However, the used annealing treatment causes different micro-structural variations in specimens depending on the pass number of TCP. As an illustration, while theaverage grain size impressively decreases due to annealing of 1 pass TCPed specimen, it moderatelyincreases after annealing of 5 passes TCPed specimen. This is due to development of a bimodal micro-structure after 5 pass of TCP which leads to a different evolution of microstructure during successiveannealing. It is also notable that TCPed and annealed specimens show higher strength and ductilitycompared with as TCPed specimens which is attributed to the occurrence of precipitation hardeningbesides restoration phenomenon during the annealing treatment.

& 2015 Elsevier B.V. All rights reserved.

1. Introduction

One of the most common methods for grain refinement ofmaterials is application of Severe Plastic Deformation (SPD) pro-cesses in which reduction of the grain size takes place by imposingintense plastic strains (e.g. equivalent strain of 1-100) without asignificant change of outer dimensions. For instance, Very-FineGrained (VFG) and Ultra-Fine Grained (UFG) materials can beproduced using SPD method which their grain sizes are between1 and 10 mm and less than 1 mm, respectively. Usually, a decreaseof deformation temperature causes an increase of the strength-ening rate and a decrease of the grain size reduction saturationlimit. Therefore, SPD processes are often applied in cold and warmregimes [1–3].

As shown before, the grain refinement of metals subjected to acold/warm SPD process (Tdeformationo0.5Tm) occurs due to

Continuous Dynamic Recrystallization (CDRX) [4]. In CDRX, thegrain refinement takes place by propagation of Low AngleBoundaries (LABs) and following evolution of LABs to High AngleBoundaries (HABs) inside “parent grains” due to imposing intenseplastic strain. Different mechanisms have been presented for thisphenomenon as thoroughly discussed before. As an illustration, itis believed that CDRX in materials bearing high Stacking FaultEnergy (SFE), such as pure and lowly alloyed aluminum, is mainlytaken place by two mechanisms [3–5]:

(1)

Multiplication and Migration of Dislocations (MMD): homo-genous multiplication of dislocation as a result of deformation,migration and rearrangement of dislocations to form homo-shaped sub-grains (or LABs), rotation of these sub-grains toform main grains characterized by HABs.

(2)

Intersection of Micro-Shear Bands (IMSB): appearance of LABsdue to propagation of Micro-Shear Bands (MSBs) as a result ofdeformation, evolution of LABs to HABs due to intersection ofMSBs by further deformation.

Fig. 1. Schematic illustrations of different CDRX mechanisms: (a) MMD mechanism and (b) IMSB mechanism.

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It is noteworthy that the IMSB usually results in a bimodalnecklace shaped microstructure while the MMD results in a rela-tively homo-shaped microstructure as illustrated in Fig. 1.

Although a cold/warm SPD treatment can result in a reductionof the grain size, the kinetic of this phenomenon is relatively slow.For example, high strains in range of 8–12 are needed to developan UFGed material by a cold/warm SPD treatment. In addition, thestrain hardened materials obtained after a cold/warm SPD treat-ment reveal relatively low ductility [1,4]. Therefore, the cold/warmSPD of materials are often combined with subsequent annealingtreatments which usually lead to occurrence of Static Re-crystallization (SRX) [4,6]. SRX can occur either as DiscontinuousSRX (DSRX) or Continuous SRX (CSRX) which are schematicallyillustrated in Fig. 2. As can be seen, DSRX occurs by nucleation andgrowth of new grains. In this phenomenon, a nucleus grows up bylong range migration of its grain boundaries and forms a new grainbearing little dislocation density. Comparatively, CSRX takes placeby initiation of LABs and their evolution to HABs to form newgrains instead of old strain hardened grains. This occurs by mi-gration and rearrangement of dislocations to form LABs and fol-lowing evolution of LABs to HABs by further accumulation ofdislocations in LABs [6]. Compared with those developed by DSRX,the dislocation density in grains developed by CSRX is relativelyhigher. This leads to higher inside misorientation of grains de-veloped by CSRX. In addition, while DSRX results in equilibriumand sharp grain boundaries, CSRX results in ill-defined and wavygrain boundaries [4]. It is believed that CSRX takes place by an-nealing treatment after imposing intense strains while DSRX oc-curs by annealing after imposing relatively lower strains [4,6].However, the transient strain in which CSRX prevails to DSRX, isless studied. Similarly, little studies have considered influence ofgrain refinement mechanisms (i.e. MMD and IMSB) – activatedduring a cold/warm SPD – on the following occurrence of SRX

phenomenon during post-annealing. For example, it is not clearthat whether new grains developed during SPD can perform asnucleus for occurrence of DSRX during annealing or not. In addi-tion, the evolution of bimodal microstructures developed by SPDprocesses during subsequent annealing treatments has not studiedyet.

It is estimated that about 90% of aluminum extruded productsare made from Al–Mg–Si alloys since these alloys have attractivecharacteristics such as the precipitation hardening capability andthe excellent workability. Therefore, multiple studies have focusedon SPD of these alloys [7–19]. However, the majority of thesestudies have considered Al–Mg–Si alloys in simple geometries (i.e.Billets and Sheets) and SPD of extruded Al–Mg–Si products, suchas tubes, has been less studied. Recently, a new SPD process calledTube Channel Pressing (TCP) has been developed to process tub-ular geometries [9,10]. Since precipitation hardening can be acti-vated during the annealing treatment of Al–Mg–Si alloys [7,12–15], the combination of TCP and post annealing of these alloys canbe a good solution for development of tubes bearing an excellentstrength and a good ductility. On the other hand, a few studiesconcentrated on the evolution of mechanical properties duringannealing/aging treatments after SPD processes suggest that theeffects of these treatments are strongly dependent on the imposedstrain by SPD [15,20]. Therefore, this work aimed to study evolu-tions of microstructure and mechanical properties of an Al–Mg–Sitube processed by TCP and subsequent annealing.

2. Process, materials and methods

The schema of TCP is shown in Fig. 3. As can be seen, the tubepasses in a bottleneck channel which results in a consecutive de-crease and increase of its diameter. Additionally, the direction ofcurvature of tube wall is changed consequently during TCP.

Fig. 2. Schematic illustrations of static recrystallization phenomena: (a) Discontinuous Static Recrystallization (DSRX) and (b) Continuous Static Recrystallization (CSRX).

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Therefore, consecutive shear strains also impose during TCP. Theaverage equivalent plastic strain imposed by each pass of TCPusing the die characterized by Rdie of 7.5 mm and θdie of 26° hasbeen calculated about 1 by an FEM analysis [10].

An Al–Mg–Si alloy tube was received in extruded form. Theinner and outer diameters of tube were 19 mm and 26 mm, re-spectively. The chemical composition of tube was Al–1.01 Mg–0.49Si–0.31Cu–0.24Fe–0.06Cr wt% which is identical to the alu-minum 6061 alloy. 75 mm long specimens were cut from as re-ceived tube and subjected to solution treatment at 803 K for 1 h.This solution treatment causes dissolution of almost all pre-cipitations since after this treatment, no evidence of any pre-cipitation was detected by X-Ray diffraction [7,21,22]. After solu-tion treatment, different specimens were subjected to 1, 3 and

Fig. 3. Schematic illust

5 passes of TCP at the room temperature. These specimens arecalled 1TCP, 3TCP and 5TCP, respectively. After processing by TCP, agroup of specimens were subjected to an annealing treatment at473 K (about 0.5Tm) for 2 h and then immediately quenched inroom temperature water. These specimens are called 1TCPA,3TCPA and 5TCPA, respectively. Tension samples were machinedparallel to longitudinal direction of specimen as illustrated pre-viously [21,23]. The strength of specimens was obtained by ten-sion tests at the room temperature using the strain rate of5�10�4 s�1.

Transmission Electron Microscopy (TEM) samples were pre-pared by mechanical polishing to 100 mm thickness and sub-sequent jet-polishing in a 75% CH3OH–25% HNO3 solution at�30 °C for 5–15 min. The TEM studies were accomplished by at

ration of TCP [10].

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least three different samples from each studied specimen usingJEM-2100F machine adapted in the acceleration voltage of 200 kV.Scanning Electron Microscopy (SEM) samples were prepared byion polishing using the IB-9010 cross section polisher to obtainElectron Back Scattering Diffraction (EBSD) maps. These mapswere achieved by JSM-7001F machine adapted at the accelerationvoltage of 15 kV and the INCA 4.09 software. The mapping wasrepeated at least three times for each specimen to obtain relativelyaccurate results. The step sizes of EBSD measurement were se-lected between 0.5 and 0.1 μm. The cut off angle for definition ofHABs in EBSD maps was considered as 5° since it has been shown

Fig. 4. EBSD maps of different specimens: (a) As received specimen, (b) 1TCP, (c) 3TCP, (dgrains.

that the appearance of grain boundaries with angle higher than 5°in aluminum results in Hall–Pitch strengthening [24,25].

3. Results

Fig. 4(a)–(d) compares EBSD maps of the specimens subjectedto different passes of TCP. As can be seen in Fig. 4(a), the as re-ceived tube has equiaxed Coarse Grains (CG) with an average grainsize of 35 mm. As shown in Fig. 4(b), few UFGs have appeared afterthe first pass of TCP which implies a limited alteration of the

) 5TCP, (e) 1TCPA, (f) 3TCPA, (g) 5TCPA and (h) the color code for the orientation of

Fig. 5. Comparison of: (a) average grain size and (b) HABs fraction after TCP and subsequent annealing; Comparison of distributions of grain size in: (c) 1TCP vs. 1TCPA,(d) 3TCP vs. 3TCPA and (c) 5TCP vs. 5TCPA.

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microstructure. Imposing of 3 passes of TCP results in a mixedmicrostructure which consists of UFGs, VFGs and CGs as illustratedin Fig. 4(c). After 5 passes of TCP, the microstructure mainly con-sists of UFGs and VFGs as shown in Fig. 4(d). Comparing Fig. 4(a)–(d), it is obvious that the processing by TCP causes a developmentof a bimodal necklace shaped microstructure as discussed later.

Annealing of 1TCP specimen results in arise of finer grains asshown in Fig. 4(e). In 1TCPA specimen, two types of grains can betraced: UFGs appeared during TCP and VFGs developed duringsubsequent annealing treatment. These VFGs substitute initialcoarse strain hardened grains and are characterized by ill-definedwavy boundaries as can be seen in Fig. 4(e). As shown here, thereis a relatively high misorientation inside these grains which in-dicates their high dislocation densities. Similar microstructurescan be seen for 3TCPA and 5TCPA specimens as illustrated in Fig. 4(f) and (g). Nonetheless, comparing Fig. 4(d) and (g), it is obvious

that the annealing treatment has impressively lowered fraction offiner grains (e.g. UFGs) in 5TCPA specimen. In addition, VFGs ap-peared in 5TCPA specimen are relatively coarser than VFGs of 5TCPspecimen. This indicates the occurrence of grain growth in thisspecimen as discussed later.

Fig. 5(a) compares the average grain size of specimens beforeand after annealing treatment. As illustrated here, the annealingtreatment after 1 pass of TCP causes a remarkable decrease of thegrain size. Despite so, the grain size is moderately increased afterannealing of 5TCP specimen. Fig. 5(b) shows the HABs fraction forspecimens before and after annealing treatment. As can be seen,the HABs fraction is impressively increased after annealing of 1TCPspecimen. Despite so, a notable decrease of HABs fraction can beobserved after annealing of 5TCP specimen. Fig. 5(c)–(e) comparesthe distribution of grain size in xTCP and xTCPA specimens. As canbe seen in Fig. 5(c), the annealing treatment after 1 pass of TCP

Fig. 6. Microstructures of different specimens: (a) 1TCP, (b) 5TCP, (c) 1TCPA and (d) 5TCPA; Appearance of precipitations in: (e) 1TCPA specimen and (f) 5TCPA specimen.

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causes a relative decrease in fraction area of coarser grains (e.g.grains greater than 30 mm) in benefit of increase of fraction area offiner grains (i.e. grains in range of 0–3 mm). This is due to occur-rence of recrystallization phenomenon. Comparatively, the an-nealing treatment after 5 pass of TCP causes a decrease of fractionarea of finer grains and an increase of fraction area of coarsergrains which indicates occurrence of the grain growth duringannealing.

Fig. 6(a)–(f) compares the microstructures of specimens pro-cessed by TCP and subsequent annealing. Collating Fig. 6(a) and(c), it is obvious that the annealing treatment after 1 pass of TCP

causes development of finer grains which are characterized by ill-defined wavy boundaries. In comparison, the evolution of micro-structure during annealing of 5TCP specimen is relatively differentas shown in Fig. 6(b) and (d). At first, the grain boundaries aresharper in 5TCPA specimen in comparison of 1TCPA specimen. Thisis due to imposing of a larger strain which promotes re-crystallization phenomenon. Secondly, while UFGs developed byTCP can be easily found in 5TCP specimen, they can hardly betraced in 5TCPA specimen. This illustrates growth of grains duringannealing of 5TCP specimen which is mentioned before. Besidesevolution of grains, the annealing treatment results in

Fig. 7. Comparison of: (a) strength and (b) elongation after TCP and subsequent annealing treatment; and (c) comparison of the engineering stress–strain curves for TCPAspecimens, as solution treated specimen and solution treated-annealed/aged specimen.

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development of precipitations as shown in Fig. 6(e) and (f). Asthoroughly discussed before [7,22,26], it is believed that the con-sequence of precipitation in an Al–Mg–Si alloy is:

GP I Zone GP II Zone− → − (β′′) → β′ → βIt is well-known that the morphology of GP-I zone is spherical

while the morphology of β′′, β′ and β are needle-like (elliptical),rod-like and platelet-like, respectively. As can be seen, precipita-tions appeared in 1TCPA and 5TCPA specimens have an ellipticalmorphology which has been related to β′′ precipitations [7,22,26].In addition, the size of precipitation appeared in 5TCPA specimenis coarser than that appeared in 1TCPA specimen. This is attributedto the remarkable increase of diffusion rate by increase of imposedplastic strain in SPDed materials as discussed before [27–30]. Theemergence of precipitations during annealing treatment results inan improvement of strength as discussed later.

Fig. 7 compares evolutions of the strength and ductility ofspecimens after processing by the TCP and subsequent annealing.As can be seen, yield strength of 310–380 MPa and fracture strainof 17–25% is obtained by the used treatment depending on the TCPpass number. Comparatively, the strength and the ductility of si-milar Al–Mg–Si alloys treated by other SPD and heat treatmentprocedures are reported in ranges of 240–420 MPa and 5–20%,respectively [7–19]. This denotes that the treatment used in thiswork has shown a significant capability to improve the strengthand the ductility of tube.

4. Discussion

From what told above, it can be inferred that TCP results indevelopment of a bimodal necklace-shaped microstructure similarto what illustrated in Fig. 1. As discussed before, this

microstructure arises due to occurrence of IMSB promoted CDRX[3,4,9]. In general, the IMSB mechanism leads to development oftwo groups of grains: finer grains (i.e. UFGs) and coarser grains (i.e.VFGs and CGs) similar to what shown in Fig. 4(a)–(d). In com-parison, the microstructures of TCPed and annealed specimensconsist of finer grains developed during TCP and coarser grainsdeveloped during TCP and/or annealing as can be seen in Fig. 4(e)–(g). As shown here, coarser grains developed during annealing arecharacterized by ill-defined wavy boundaries and bear a highdislocation density. These trends are similar to what shown inFig. 2(b) for the evolution of microstructure by CSRX which in-dicates occurrence of this phenomenon during applied annealingtreatment for used alloy. It is notable that UFGs developed duringan SPD process bear low inside dislocation density [4]. Regardingthis, one may propose that these UFGs may act as nucleus for in-itiation of DSRX during subsequent annealing of Al–Mg–Si alloyssubjected to SPD. Despite this, results of this work decline this ideasince the microstructure of TCPA specimens represent occurrenceof CSRX rather than DSRX as explained in discussions for Fig. 4(e)–(g). Therefore, it can be demonstrated that the CSRX is the pre-dominant restoration phenomenon of an Al–Mg–Si alloy subjectedto a moderate plastic strain – as low as 1 – during subsequentannealing. Comparatively, occurrence of CSRX is reported duringannealing of pure copper and stainless steels after imposingstrains greater than 3 [4]. This difference can be explained by highSFE of aluminum compared with those materials which accel-erates its dislocation mobility [6,13,14].

As mentioned in explanations of Fig. 5, while the decrease ofaverage grain size and increase of HABs fraction after annealing of1TCP specimen occurs due to CSRX, the increase of average grainsize and decrease of HABs fraction after annealing of 5TCP speci-men occurs due to grain growth phenomenon. This is related to

M.H. Farshidi et al. / Materials Science & Engineering A 640 (2015) 42–50 49

higher progression of CDRX in 5TCP specimen during TCP whichmakes it susceptible for subsequent grain growth during anneal-ing. Compared with 1TCP specimen, 5TCP specimen has been af-fected by more severe CDRX which causes much progression ofrecrystallization and higher fraction of UFGs. Besides this, 5TCPspecimen bears greater stored energy and higher rate of diffusiondue to more intense plastic deformation [27–30]. Therefore, theaffinity for grain growth during subsequent annealing of 5TCPspecimen is higher than that of 1TCP specimen which causes rapidelimination of finer grains (i.e. UFGs) in benefit of coarser grains. Itis also notable that the upper bound of grain size in 5TCPA spe-cimen is 30 mm as shown in Fig. 5(d) which is much greater than9 mm related to 5TCP specimen. This represents an extensivelyaccelerated and abnormal grain growth in a group of grains whichis not expectable at the used annealing temperature. In fact, lim-ited growth of microstructure has been reported during annealingof the used alloy in this temperature (473 K) in previous studies[13,31,32]. In comparison, the abnormal grain growth during an-nealing after 5 passes of TCP is related to development of a bi-modal microstructure during TCP which accelerates the growth ofgrains due to a great driving force for elimination of finer grains(i.e. UFGs) in the microstructure. In addition, since the fragmen-tation of initial microstructure through TCP has been occurred byIMSB mechanism, the elimination of finer grains developed byIMSB results in vicinity of coarser grains to each other (see Fig. 1).These coarser grains may have a similar crystallographic orienta-tion since they may be derived from one parent grain. Therefore,the elimination of finer grains causes a reduction of HABs andabnormal increase of grain size due to a probable coalescence ofcoarser grains as illustrated in Figs. 4 and 5.

As illustrated in Fig. 7(a), 2 h of annealing of TCPed specimensat 473 K results in a moderate increase of their yield strength. Thiscan be explained by activation of age hardening during annealingas shown before [7,12–15]. However, little changes have beenpreviously obtained for strength of TCPed specimens after 20 min.of annealing in similar temperature [21]. This has been explainedby decrease of dislocation density besides arise of pre-β′′ pre-cipitations which results in a balance between strengthening ef-fect of precipitations and softening effect of restoration [21]. Col-lating with those results, the moderate strength increase of TCPedspecimens after 2 h of annealing is related to incidence of semi-coherent β′′ precipitations which causes more strengthening effectin comparison of coherent pre-β′′ precipitations [19,26]. In addi-tion, the annealing strengthening effect is decreased by increase ofTCP pass numbers as shown in Fig. 7(a). This is related to incidenceof more coarse precipitations during annealing of a specimensubjected to higher TCP pass numbers as previously illustrated inFig. 6(e) and (f). Note that coarsening of precipitations reducestheir strengthening effect due to increase of their free distance[14]. Beside moderate strength increase, the elongation is im-pressively increased by annealing treatment after TCP as illu-strated in Fig. 7(b) which can be explained by restoration effect ofannealing. It is also noteworthy that the fracture strains of TCPAspecimens are impressively higher than that of the solution trea-ted-annealed/aged specimen. This is mainly due to increase of thestrain accommodated after necking point in TCPA specimens asillustrated in Fig. 7(c). As can be seen, the necking strains of TCPAspecimens are almost equal to necking strain of solution treated-annealed/aged specimen. However, TCPA specimens bear moreextensive strain after necking point which results in their higherductility. This can be explained by the lower grain size of TCPAspecimens which increases their strain rate sensitivity and de-celerates the occurrence of fracture after necking point [33–35].

5. Conclusions

Considering the presented results, it can be concluded that:

1.

The development of a bimodal microstructure during TCP re-sults in occurrence of different trends during subsequent an-nealing of TCPed specimens.

2.

Continuous static recrystallization is the main restorationphenomenon during annealing of the aluminum 6061 alloy at473 K after imposing a moderate plastic strain as low as 1.

3.

The average grain size is impressively decreased by annealingafter 1 pass of TCP due to occurrence of continuous static re-crystallization. Despite so, annealing after 5 passes of TCPcauses a moderate increase of the average grain size which isrelated to elimination of finer grains by occurrence of graingrowth during annealing.

4.

Since the restoration and the precipitation hardening occurtogether, both of the strength and ductility of the aluminum6061 alloy improve by 2 h of annealing at 473 K after TCP.

Acknowledgments

The authors wish to thank the research boards of FerdowsiUniversity of Mashhad, Sharif University of Technology andDoshisha University for the financial support and the provision ofresearch facilities used in this work.

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