Materials and Design 50 (2013) 620–634

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Materials and Design

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The optimum combination of tool rotation rate and traveling speedfor obtaining the preferable corrosion behavior and mechanicalproperties of friction stir welded AA5052 aluminum alloy

0261-3069/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.matdes.2013.03.027

⇑ Corresponding author. Tel.: +98 919 7642980; fax: +98 21 66165717.E-mail addresses: Mohibagheri@yahoo.com (M. Bagheri Hariri), S.gholami85@g-

mail.com (S. Gholami Shiri), Yaghoubinezhad@mehr.sharif.ir (Y. Yaghoubinezhad),M.rahvard@yahoo.com (M. Mohammadi Rahvard).

Mohiedin Bagheri Hariri a, Sajad Gholami Shiri b, Yadollah Yaghoubinezhad a,⇑,Masoud Mohammadi Rahvard c

a Materials Science and Engineering Department, Sharif University of Technology, Azadi Ave., P.O. Box 11365-9466, Tehran, Iranb Materials Science and Engineering Department, Shiraz University, Zand Ave., Shiraz, Iranc Materials Science and Engineering Department, Iran University of Science and Engineering, Narmak Ave., Tehran, Iran

a r t i c l e i n f o

Article history:Received 28 November 2012Accepted 9 March 2013Available online 21 March 2013

Keywords:Aluminum alloysFriction stir weldingCorrosionMicrostructureMechanical

a b s t r a c t

This study attempts to find an optimum combination of the welding tool rotation rate (x) and travelingspeed (t), concerning the corrosion and mechanical properties of Friction Stir Welded (FSWed) AA5052Aluminum alloy. The effect of the tool speeds on the FSWed AA5052 are investigated via potentiodynamicpolarization, open circuit potential (OCP) monitoring, test of the susceptibility to intergranular corrosion,weight loss, tension and micro-hardness tests. Optical microscope and Scanning Electron Microscopy(SEM) were employed for studying the morphology and analyzing the probable intergranular attacks.It was found that by increasing t up to 200 mm/min at x = 400 rpm, the microstructural evolution isin a way that the finer grain structure intensifies the anodic reactivity of the Stir Zone (SZ). At faster t(about 250 mm/min), further grain reinforcement resulted in a predominant effect of passive film forma-tion and thereby an unexpected high corrosion resistant SZ with a proper mechanical characteristics wasattained.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Due to the fact that Al alloys have specific properties such aslight weight, corrosion resistance, specific strength, low price,excellent formability and super plasticity, the inclinations to usethese alloys are developed [1–3]. For Al alloys, bonding and weld-ing technologies should also be established apart from alloy design,microstructure control, plastic forming, casting, surface treatment,etc. [4]. The welding of Al and its alloys has always represented agreat challenge for designers and technologists [5]. The conven-tional joining processes for this alloy include gas metal arc welding(GMAW), gas tungsten arc welding (GTAW), plasma arc welding(PAW) and electron beam welding (EBW) [6,7]. But the strengthof joints through the conventional fusion welding can only bereached 50–70% of the parent metal [8]. Besides, it has beenclaimed that recrystallization, grain growth, crack and residualstress adjacent to the welded regions cause the weakening of theweld; consequently, the welded domains of the most joints are cor-rosion susceptible [9]. Because of the segregation, porosities and

the loss of alloying elements during a conventional fusion weldingprocess, Al alloys are grouped as ‘difficult to weld’ metals. FSWprocess since its inception in the year 1991 by The Welding Insti-tute (TWI) in the UK, suggests an alternative through solid-statebinding that remove the problems emanating from nugget solidifi-cation [10,11]. FSW as a solid-state welding process can com-pletely omit the melting defects e.g. liquation cracking, porositiesand distortions in the weld. This technique offers excellentmechanical characteristics up to the point that the weld joints pro-duced by FSW have been reported to have the tensile properties of80–100% of the parent metal. In fact, FSW combines frictional heat-ing and stirring motion to soften and mix the interface betweentwo metal sheets to produce fully consolidated welds [12–13]. Be-cause of the various grain microstructure and dynamically grainsrecrystallization phenomena occurring during the heat treatmentand FSW process, 5xxx series of stir welded Al alloy have a differ-ent corrosion susceptibility in sub-regions of the jointed zone.AA5052 shows a fairly higher corrosion resistance against localizedor general attack, compared to Al–Si (4xxx) and Al–Si–Mg (6xxx)grades which are more sensitive to pitting corrosion [14]. Lately,Galvao and coworkers have examined three tool geometries, (i.e.flat, conical and scrolled) for FSW of Cu plates. In the same weldingconditions, the scrolled shoulder contributes to the better mechan-ical properties and the finer grain microstructure [15]. Depending

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 621

on various welding parameters, weld zones can include severalsub-domains with different microstructure, containing variousgrain sizes (G.S.) structures and residual stress/strains. So variousstructural zones, i.e., SZ, Heat Affected Zone (HAZ), Thermo-Mechanical Affecter Zone (TMAZ) and Base Metal (BM) displayindividual sensitivity to corrosion [16]. Previous studies haveexhibited more susceptibility of the SZ to local corrosion in com-parison to the BM of AA2024, AA5454 and AA7075 [17]. Cu deple-tion along grain boundaries for AA7075 is claimed to beresponsible for higher susceptibility of the SZ to intergranular at-tack. Additionally, the SZ in AA1100, AA2024, AA6061, 7010-T7651 and Al–Cu–Mg–Ag alloys also show vulnerability to local-ized corrosion that has attributed to the wide precipitate-free re-gions and coarse precipitates in hardenable Al alloys [18]. ForFSWed AA2219, Srinivasan et al. correlated the improved corrosionresistant in the SZ to the dissolution/coarsening of the strengthen-ing precipitates in the zone and subsequently the reduction in thegalvanic driving forces [9]. Dissolution effects of the second phasesof CuCl2 in AA2219 during the multi/single pass friction stir pro-cessing and fusion based laser melting have been studied else-where. These methods develop the corrosion resistant of thealloy due to reducing the number of sites available for galvaniccoupling [19,20]. Jariyaboon et al. have comprehensively discussedhow corrosion characteristics of FSWed AA2024-T351 can be de-pended on microstructure during the FSW process [21]. Shenet al. investigated comparatively the microstructure and electro-chemical corrosion behavior of dissimilar FSWed 5058/6082. Atthe given welding conditions, the corrosion resistant of the SZwas higher than those of two parent alloys [22]. Many studies haveattempted to find out the effects of welding parameters on themechanical characteristics of FSWed alloys [23–25]. In fact, theinfluence of FSW parameters mainly x and t on the microstructureand mechanical properties of AA5052 have been studied in the lit-eratures without taking into account the electrochemical behaviorof the joint [1,10]. To the author’s best knowledge, it is the first re-port attempting to find an optimum welding condition for FSWedAA5052 at a fixed tool axis of 2.5� and plunge force of 800 kg thatconsiders both corrosion and mechanical properties.

There is a limited understanding of the relationship betweenmicrostructure and corrosion characteristics of FSWed Al alloys[21]. The aim of this work is to survey the effects of x and t onthe corrosion behavior of FSWed AA5052. Eventually, an optimumcondition, concerning the corrosion vulnerability and mechanicalcharacteristics of the welded regions will be introduced.

2. Experimental details

2.1. Material

The base metal was a rolled sheet of AA5052 with a thickness of4 mm and the chemical composition, which is given in Table 1.

2.2. Welding procedure

The welding plate samples with dimensions of 4 � 70 �150 mm, were longitudinally butt-welded parallel to the roll direc-tion using an automatic vertical milling machine. A backing plateof carbon steel was placed under the two-butted base metal tobe welded. The H13 hardened steel was utilized as a welding tool

Table 1Chemical composition (in wt.%) of AA5052 alloy.

Alloy Si Fe Cu Mn Mg Cr Al

AA5052 0.115 0.372 0.017 0.055 2.363 0.191 Balance

and a typical shoulder with 20 mm diameter and a tapper cylindri-cal pin were used that had a base and head diameters of 5 and4 mm, respectively. The tool axis was tilted 2.5� from the verticalaxis to facilitate the weld consolidation at the plunging force of800 kg.

2.3. Tensile strength and microhardness evaluations

The transverse tensile specimen was prepared according to theASTM: E8 standard. To obtain a reproducible data, tensile strengthevaluation was generally accomplished by three samples and Vick-ers microhardness was also implemented under the loading about50 g for 10 s.

2.4. Microstructural observations

Optical microscopy (Olympus model PME3) and SEM were uti-lized to analyze the morphology and other probable localized at-tacks in various regions of FSWed samples. The transverse cross-section of joint was carefully prepared applying standard metallo-graphic techniques consisting of grinding with emery 1200 and3000 abrasive paper, followed by polishing and etching in a suit-able etchant solution composed of hydrofluoric (40 vol.%), hydro-chloric (37 vol.%), nitric (65 vol.%) acids and potassium chromatein distilled water till the various regions of the welded sampleswere roughly observable during optical microscopy observations.

2.5. Evaluation of the corrosion behavior of FSWed AA5052

2.5.1. Conventional cyclic polarization test together with OCPmonitoring

According to the ASTM: G69, the potentiodynamic polarizationmeasurements were carried out in a typical three-electrode cellcontaining a corrosive electrolyte (60 g/L NaCl) and a stainless steelplate with dimensions of 0.3 � 2 � 2 cm was operated as an auxil-iary electrode. The potentiodynamic experiments were carried outfrom �1 V to 1 V with respect to the OCP at a scan rate of 5 mV/s.The OCP was monitored for 3000 s in 3.5 wt.% NaCl solution. Beforethe experiment, an applied cathodic potential of �0.1 V vs. SCE for45 s provided a reproducible homogeneous surface conditions. Thepotentiodynamic experiments and OCP monitoring were per-formed by a Potentiostat/Galvanostat Autolab model PGSTAT302Nequipped with Nova 1.5 software. A Saturated Calomel Electrode(SCE) was utilized as a reference electrode to which all potentialsare referred. For minimizing the IR drop effect, the capillary bridgetip of SCE was very close to the surface of working electrode. Theexperiments were repeated for several times to ensure the repro-ducibility of the results.

2.5.2. Susceptibility to intergranular corrosion attackIn order to determine the susceptibility to the intergranular cor-

rosion attack, the immersion test in a concentrated 70 wt.% HNO3

solution at 30 �C for 24 h was carried out by a similar methodaccording to references [26,27]. The prepared specimens wereweighed to an accuracy of about 0.0001 g. After weight loss testin mentioned electrolyte, the samples were washed in distilledwater and acetone, dried at 65 �C, and weighed again.

2.5.3. Static immersion testIn accordance with the ASTM: G31, immersion test of prepared

samples with dimensions of about 0.35 � 0.5 � 1 cm were per-formed in 3.5 wt.% NaCl solution for about 17 days. Three out ofeach welded specimen were immersed in the mentioned solution.Before immersion, all samples were mechanically polished withemery 1200 and 3000 abrasive paper and ultrasonically cleanedfor 3 min in distilled water and acetone, and finally dried in warm

622 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

air. All weight measurements were carried out with a METTLERTOLEDO model XS204, scale with an accuracy of 0.0001 g. All elec-trolytes were prepared from analytical grade using double distilledwater.

3. Results and discussions

3.1. Mechanical properties of FSWed AA5052

3.1.1. Macro and micro-metallographyIn Fig. 1a, b and c, the macrograph cross-sections of FSWed

plates at x = 800, 1600 and 2500 rpm were represented, respec-tively (t = 50 mm/min).

At the tool speeds faster than 400 rpm the SZ, HAZ, TMAZ andBM demonstrate the presence of probable defects like voids andcracks. It was clearly found that an increase in x caused to a moreheterogeneity in the weld zones and subsequently the hazardousdefects such as tunneling and porosities were remarkably ap-peared (as demonstrated in Fig. 1c).

Fig. 2 schematically illustrates the FSW process in a three-dimensional cartesian/cylindrical space. The symbols rp, rs, hp, a,h and r denote the tool pin radius, shoulder radius, tool pin height,tool shoulder cone angle, the angle between radial vector and thewelding direction and the position along tool radius, respectively.

In FSW, due to the effects of the shear failure of the weldingmaterials, the increase of the friction stress at the tool–metal inter-face is restricted. So, it is deduced that the classical Coulomb fric-

Fig. 1. Macrographs cross-sections of the FSWed AA5052 obtained at the toolspeeds of (a) 400, (b) 1600, and (c) 2500 rpm; t = 50 mm/min.

Fig. 2. Schematic representation of the FSW

tion law cannot be employed in high levels of x where thedynamic effects of the welding tool is increasing. However, themodified Coulomb model includes these effects and thereby isapplicable at faster x as well. The friction heat generation, Qfric,can be evaluated by [28]:

Qfric ¼ gs c�

ð1Þ

where g, s and c�

are the heat efficiency, friction stress and slippingrate, respectively. c

�is defined as xr(1 � d) in which d is introduced

as the extent of slip [29]. Moreover, the tangential rate of the shoul-der with respect to the workpiece is given by xr – t sinh [30]. So,Eq. (1) can be rewritten as follows:

Qfric ¼ gsðxr � t sin hÞð1� dÞ ð2Þ

The term t sinh would be neglected at much larger values of xr.In a pure sliding situation where d comes close to 0, the frictionstress is given by lP in which l and P are the friction coefficientand the contact pressure, respectively. In contrast, when d is nearto 1, the material sticks to the welding tool and the shear failurelimits the friction stress, and consequently the heat is mostly gen-erated due to the shear deformation. For these sticking conditions,the contact shear stress becomes equal to yield shear stress(rs=

ffiffiffi3p

). rs is the yield stress and is a function of temperature[28,29]. Considering Eq. (2) and by combining the analyticalexpressions for sliding and sticking conditions, the overall frictionheat is generated from the tool shoulder and the pin is also esti-mated as follows [29]:

Q ¼ 23px

drsffiffiffi3p þ ð1� dÞlP� �

� ðr3s � r3

pÞð1þ tan aÞ þ r3p þ 3r2

php

h ið3Þ

Owing to the tool movement during FSW, the friction heat flowis evaluated by Qfric/t (i.e., by (xr – t sinh/t) regarding Eq. (2));hence, it can be concluded that the friction heat flow increasesby increasing x or decreasing t [31]. In order to exhibit the effectof t on the weld microstructure, the speeds of 50 and 350 mm/minwere selected to show the influence in Fig. 3.

Obviously, some voids, cracks and unbounded local defects areobserved in the SZ attained at higher traveling speed than350 mm/min (Fig. 3a and c). The formation of defects in the SZ de-

system and the welding parameters.

Fig. 3. (a) Macrograph, (c) micrograph cross-sections of the FSWed AA5052 at t = 350 mm/min, (b) macrograph, and (d) micrograph of the FSWed AA5052 at t = 50 mm/min;x = 400 rpm.

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 623

pends on both x and t which affect the material flow and heat gen-eration during FSW process. As a matter of fact, at faster t, the toolrotation speed should be increased simultaneously to prevent anyprobable welding defect such as voids and cracks [32]. RegardingEqs. (2) and (3), due to a more intense Qfric at extra fast x (or slowt), pasty or even melt material appears in the SZ; and melting defectssuch as shrinkage, voids, tunnel, cracks and residual stress were re-sulted. On the other hand, by decreasing Qfric, the flow of materialwill reduce; consequently, the metal cannot completely fill thenotch produced by pin. Therefore, defects such as tunnel, un-bounded zone and void were appeared in the weld zones. So, a prop-er combination of x and t can lead to a more perfect weld duringFSW process [33]. Under different values ofx and t, the temperatureand material flow pattern are studied by many researchers. Both theincrease of x and decrease of t result in the increment of the stirringeffects which enhance the welding quality [34]. For several Al alloys,it was speculated that the maximum welding temperature in FSWprocess can be described as [16]:

TTm¼ k

x2

t� 104

� �a

ð4Þ

in which a is reported to be in a range from 0.04 to 0.06, k is be-tween 0.65 and 0.75 and Tm is the melting point (in �C). So, it is ex-pected that the maximum temperature in FSW is increased withincreasing x and decreasing t (increasing the (x2/t) ratio). Further-more, considering Eq. (2), it can be seen that Qfric and the frictionheat flow can be decreased with the increase of t. For these reasonsand based on the principles of grain growth, the finer grain struc-ture is observed at faster t (see Fig. 3). The average grain size evo-lution due to the continuous dynamic recrystallization phenomena(CDRX) during FSW of Al alloys is numerically modeled by Fratiniand Buffa. Although, an intense plastic deformation in the SZ resultsin a fine equiaxed grain microstructure, in TMAZ the inadequatedeformation is occurred without any recrystallization phenomenon[34]. Therefore, the average grain size in the SZ is smaller than thatin TMAZ (as it can be seen in Fig. 3d). The macrograph and micro-graph images of the perfect joint which was obtained at(x = 400 rpm, t = 50 mm/min) are shown in Fig. 3b and d. The Ze-ner–Hollomon parameter expresses the temperature compensatedeffective strain rate which is given by the following equation [35]:

Z ¼ e�

expðQ=RTÞ ð5Þ

where Q, T and R, are the activation energy, the absolute tempera-ture and the gas constant, respectively. e

�is the effective strain rate

and according to the Von Mises criteria it can be calculated byP 23 ðe�ijÞ

2h i1=2

. In order to obtain a more precise relationship for e�

in Eq. (5), the effective strain rate is expressed as a function of toolspeeds (x,t) in a Cartesian coordinate system as [30,36]:

e�¼ 2

34p2r2

px2

3600� F1 þ

2prpxt60

� F2 þ t2 @d@x

� �2

þ 12

@d@y

� �2

þ 12

@d@z

� �2" #( )" #1=2

ð6Þ

where F1 and F2 are the functions of x, y, d, @d@x,

@d@y and @d

@z. It was as-sumed that the material movement in z-direction was insignificantduring FSW of AA5052. According to the works of McQeen et al., thefinal recrystallized grain size in Al alloys can be estimated by Eq. (7)[37]:

G:S: ¼ ð�0:60þ 0:08 log ZÞ�1 ð7Þ

With decreasing t, the effective strain rate is decreased (withregard to Eq. (6)), and subsequently the Zener–Hollomon parame-ter is reduced which leads to the grain coarsening in accordancewith Eq. (7). Moreover, considering Eqs. (2) and (3), the grain sizeenlargement takes place at slower t owing to the higher frictionheat flow and the occurrence of the grain growth phenomenon(compare Fig. 3c and d). Some additional static grain growth canalso take place in SZ during the cooling conditions and after a se-vere plastic deformation [16]. Fig. 4 relates to the SEM observationof SZ at the welding condition of (x = 400 rpm, t = 250 mm/min).

The Clemex Vision PE as a powerful image processing softwarewas employed for determining the grain size. The microstructureof the SZ at different x is shown in Fig. 5.

By comparing the grain size in Fig. 5a (G.S. = 2.1 lm) and Fig. 4(G.S. = 0.96 lm), it will be deduced again that an increase of t re-sults in grain reinforcement. At the fast tool speeds of about350 mm/min it is suggested that x should be increased simulta-neously to avoid the weld defects [32]; however, we ignoredmicrostructural investigations at faster t than 300 and 315 mm/min; since our joints under such a condition were completelyinappropriate.

Apparently in Fig. 5, by increasing x the coarse grains will begenerated as a result of the grain growth phenomenon which hap-pens at higher Qfric (considering Eqs. (2) and (3)). Additionally,accompanying with x increasing (or x2/t ratio increment), themaximum temperature is increased in the weld region which im-proves the grain growth occurrence. On the other hand, by consid-ering Eqs. (5)–(7), the degree of deformation (e and/or e

�) increases

by increasing x (or x2/t) and thereby the generation of a fine grain

Fig. 4. SEM micrograph of the SZ for FSWed AA5052 at (x = 400 rpm, t = 250 mm/min).

Fig. 5. Optical microstructure of the SZ in different tool speeds of

45

47

49

51

53

55

57

59

61

-10 -5 0 5 10

Mic

roha

rdne

ss (

HV

)

Distance from joint line (mm)

(a)

Base

(b)SZ SZ

Base Base

100 µm

Fig. 6. (a) Profile of microhardness and (b) optical micrograph of the transve

624 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

structure is expected based on the models ascribing the CDRX phe-nomenon [34]. However, the final recrystallized grain size duringFSW of AA5052 depends on the factor that predominantly deter-mines the process. During FSW, it is claimed that the friction heatand peak temperature are the dominant factors controlling therecrystallized grain size, and generally the grain size increases withincreasing x or the ratio of x2/t [16,32]. Actually, the finest grainsize was observed at the tool speeds of (x = 400 rpm, t = 250 mm/min). It must be noted that, the size of the grain proportionally isconcluded from the microstructural observation.

3.1.2. Micro-hardnessFig. 6a illustrates the micro-hardness profile of the weld ac-

quired at (x = 400 rpm, t = 100 mm/min). It can be seen that thehardness around the center of keyhole increases gradually fromthe BM through the TMAZ and abruptly increases in the SZ.

As shown in Fig. 6b, the minimum grain size value by a nominaldimension of few micrometers is observed in the SZ due to an in-tense degree of deformation in this region. Although the TMAZundergoes the plastic deformation, the recrystallization cannot oc-cur in this zone because of an insufficient deformation strains [34].Therefore, in accordance with the Hall–Petch relationship, the fin-est grain in SZ associated with a higher Vickers hardness value. Al

(a) 400, (b) 800, (c) 1600, and (d) 2500 rpm; t = 50 mm/min.

TMAZBase

Joint line

SZ SZ

TMAZ TMAZBase

rse cross-sections for FSWed AA5052 at (x = 400 rpm, t = 100 mm/min).

45

50

55

60

65

70

0 50 100 150 200 250 300

Ave

rage

Mic

roha

rdne

ss (

HV

)

Traverse Speed (mm/min)

Fig. 7. (a) Hardness of the FSWed AA5052 at different tool traverse speeds;x = 400 rpm.

Fig. 9. Tensile specimen after fracture; the FSW accomplished at (x = 400 rpm,t = 50 mm/min).

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 625

alloys are classified into heat-treatable (precipitation-hardenable)and non-heat-treatable (solid-solution-hardened) alloys. Someinvestigations revealed that the hardness profiles for these twocategories of Al alloys are not analogous during FSW. In a numberof precipitation-hardened Al alloys, FSW creates a soft regionaround the weld center. It was suggested that such a softening iscaused by coarsening and dissolution of strengthening precipitatesduring the thermal cycle of FSW [38–40]. For pure Al without anysecond-phase particles, the hardness in the nugget zone wasslightly higher than that in the base material, and the maximumhardness was located in the TMAZ. In general, for the solid-solu-tion-hardened Al alloys, FSW does not result in weld softening.Microstructural factors which influence the hardness of FSWedspecimens for solid-solution-hardened Al alloys were suggestedby various investigators [41,42]. As denoted in Fig. 7 there is anincreasing trend of hardness with increasing t (at a constantx = 400 rpm).

As it was mentioned in Section 3.1.1, the grain size normally de-creases at faster t as a result of the synergistic effects of the frictionheat flow diminishing (Eq. (2)) and increasing of the degree ofdeformation (Eqs. (5)–(7)). Hence, the faster t contributes to a finergrain structure (see Fig. 3) and thereby the increment of micro-hardness [43].

100

120

140

160

180

200

220

240

260

0 500 1000 1500 2000 2500 3000

Ten

sile

Str

engt

h (M

Pa)

Rotation Rate (rpm)

(a)

Fig. 8. (a) Tensile strength of FSWed AA5052 specimens at different tool rotation ratedifferent traverse speeds; x = 400 rpm.

3.1.3. Tensile strengthFig. 8a and b show the ultimate tensile strength variations of

the welded samples as a function of x and t, respectively.As it can be seen in Fig. 8a, the speeds of 400 and 800 rpm (at

t = 50 mm/min) lead to a good junction where the tensile strengthsof the SZ are close to that for the BM. From Fig. 8b, it can be con-cluded that the tensile strength of the samples which welded att of 25–250 mm/min are close to the strength of the BM(249 MPa). Two factors are important to get a relatively higher ten-sile strength. The presence of a fine microstructure in the SZ and aperfect SZ without any defects are essential to enhance thestrength of the weld zone. By increasing t up to 300 mm/min,the strength of the SZ decreases, because of the presence of theweld deficiencies.

3.1.4. Analysis of fracture surfaceIn order to understand the nature of fracture for the SZ speci-

mens, a few samples were examined. Fig. 9 shows an example ofthe tensile specimen which was welded at (x = 400 rpm,t = 50 mm/min).

The relevant SEM micrographs of the fracture surface can beseen in Fig. 10.

It is evident that the welding parameters substantially affect thenature of the fracture and the mode of failure is clearly classified asductile fracture. More ductile fractures with a less void coalescenceare evident in Fig. 10b which is a result of the finer grain structureachieved under condition of (x = 400 rpm, t = 250 mm/min) [10].Compared to Fig. 10b, the dimples in the Fig. 10a are deeper andtearing edges are thicker containing many micro-pores. The voidsize in the SZ at (x = 400 rpm, t = 250 mm/min) is smaller thanthat in the same rotation rate and at t = 50 mm/min; which impliesthat a smaller stretched zone was present at the crack tip in the

0

50

100

150

200

250

300

0 50 100 150 200 250 300 350

Ten

sile

Str

engt

h (M

Pa)

Traverse Speed (mm/min)

(b)

s; t = 50 mm/min and (b) ultimate tensile strength of FSWed AA5052 samples at

Fig. 10. SEM fractographs of the SZ samples after FSW of AA5052 at tool speeds of (a) 50 and (b) 250 mm/min; x = 400 rpm.

626 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

condition of (x = 400 rpm, t = 50 mm/min) that leads to a smallplastic zone and rather low ductility.

3.2. Corrosion behavior of FSWed AA5052 specimens

The electrochemical corrosion investigations were separatedinto two different sections. In the first section the potentiodynamicpolarization studies together with OCP monitoring and the test ofintergranular corrosion susceptibility were performed in order tofind the valuable data about the electrochemical behavior of theFSWed AA5052. Afterward, the static immersion test was accom-plished so as to support and confirm the polarization results.

3.2.1. Polarization study and the susceptibility to localized corrosionFirst of all, we attempted to find out the effect of x variations as

the first important factor on the microstructure and corrosionbehavior of the weld zones. We will specify the best x duringFSW of AA5052 at a fixed t of 50 mm/min, tool axis of 2.5� andplunge force of 800 kg. At the second stage, the influence of t vari-ations will be investigated and finally an optimum combination ofx and t will be introduced.

3.2.1.1. The effect of rotation speed variations on corrosion character-istics of FSWed AA5052. Fig. 11a and b illustrate the reproduciblecathodic and anodic cyclic polarization curves of SZ and HAZ ac-quired in various x and at a constant t = 50 mm/min, in 60 g/LNaCl, while Table 2 summarizes the relevant data of Ecorr. and icorr.

concluded using the Tafel extrapolation method.

-2.5

-2

-1.5

-1

-0.5

0

0.5

-3 -2 -1 0 1 2 3 4 5 6

E v

s. S

CE

(V

)

Log (i) (µA/cm2)

SZ: 400 rpm

SZ: 800 rpm

SZ: 1600 rpm

SZ: 2500 rpm

BM

-1.5

-1.3

-1.1

-0.5

-0.7

-0.9

0 1 2 3 4 5

Eprot.

EPit. Hysteresis loop

(c)

(a)

Fig. 11. Potentiodynamic polarization curves of FSWed AA5052 for (a) SZ, (b) HAZ inx = 800 rpm; Scan rate 5 mV/s, t = 50 mm/min.

All cases have a similar cathodic polarization behavior, whereasin anodic polarization branches immediately after OCP, there is astrong rising in dissolution current density with no significant pas-sivation area, which corresponds to the pitting breakdown poten-tial of duplex protective oxide film on AA5052 [44]. Such abehavior was also reported elsewhere for scandium AA5052 [44],Al–4Zn–2Mg alloy [45] and AA2024-T3 in 3.5 wt.% NaCl [46]. Somephenomena decelerating the passivity and encouraging the local-ized corrosion. (e.g., intermetallic phases such as Al3Mg2 have lessnoble dissolution potential relative to Al matrix and act as anodicspots [47,48]). For a better analogy, the outcome values of Ecorr.

and icorr. have been illustrated in Fig. 12a and b.It is evident that Ecorr. in the SZ and HAZ is approximately 0.1–

0.625 V and 0.01–0.14 V more negative than the BM, respectively.Moreover, icorr. of BM is nearly 4.5–65 and 2.5–45 order of magni-tude greater than those in the SZ and HAZ, respectively. Fig. 13 rep-resents the optical microscopy images of the BM, HAZ and SZ forFSWed AA5052 at x = 400, 800, 1600 and 2500 rpm, after 24 himmersion in 70 wt.% HNO3 solution. The mass losses per unit areaof the specimens are listed in Table 3.

The highest degree of weight loss for the SZ obtained at x = 800and 2500 rpm might be due to the networks dissolution of Al–Mgintermetallic compound around the grain boundaries [26,27].Regarding Table 3, the base alloy in all x experiences a mass lossabout 7.6 to 13.3 mg/cm2 that confirms the AA5052 is an inher-ently intergranular-resistant alloy [48]. Choi and coworkers haverecently investigated the intergranular corrosion properties ofFSWed AA5083 by the nitric acid weight loss test [26]. The mass

Log (i) (µA/cm2)

-2

-1.8

-1.6

-1.4

-1.2

-1

-0.8

-0.6

-0.4

-0.2

-1 0 1 2 3 4 5 6

E v

s. S

CE

(V

)

HAZ: 400 rpm

HAZ: 800 rpm

HAZ: 1600 rpm

HAZ: 2500 rpm

BM

Active corrosion ratesLow to passive rates

(b)

60 g/L NaCl, in various x and (c) cyclic polarization curve of the SZ obtained at

Table 2Polarization diagram parameters in 60 g/L NaCl for weld zones of FSWed AA5052 atdifferent x; t = 50 mm/min.

Rotation speed (rpm) Ecorr. (V vs. SCE) icorr. (lA/cm2)

SZ HAZ BM SZ HAZ BM

400 �0.844 �0.858 �0.751 0.96 1.41 63.45800 �1.374 �0.884 �0.749 70.8 6.31 64.56

1600 �0.888 �0.842 �0.755 9.54 26.3 70.562500 �1.26 �0.842 �0.754 14.13 25.11 63.87

0

10

20

30

40

50

60

70

80

0 500 1000 1500 2000 2500 3000

i corr

.(µA

/cm

2 )

Rotation Speed (rpm)

SZ

HAZ

Traverse speed: 50 mm/min

-1.5 -1.4 -1.3 -1.2 -1.1

-1 -0.9 -0.8 -0.7 -0.6 -0.5

0 500 1000 1500 2000 2500 3000

Eco

rr. v

s. S

CE

(V

)

Rotation Speed (rpm)

SZ

HAZ

(a)

(b)

Fig. 12. Effect of tool rotation rate variations on (a) icorr. and (b) Ecorr. of the weldzones; t = 50 mm/min.

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 627

loss for the HAZ at all x is a result of the preferential corrosion ofrandomly distributed Al–Mg intermetallic compound along grainboundaries [26,27]. Obviously, the intergranular resistant (e.g.,for the BM), randomly (e.g., for the HAZ) or network attack (e.g.,for the SZ at x = 800 and 2500 rpm) along grain boundaries canbe seen in the metallographic examination in Fig. 13. The highestanodic activity in the SZ occurred at x = 800 and 2500 rpm withthe most negative dissolution potential of �1.37 and �1.26 V,respectively. It can be referred to more sensitivity of the weld tointergranular corrosion in these speeds (as shown in Fig. 13c andg); however, a lower intergranular attack in the SZ can obviouslybe seen at x = 400 and 1600 rpm (Fig. 13a and e). Almost in allx, a more negative Ecorr. in the SZ with respect to that in theHAZ, demonstrates a preferential attack in the nugget region. Themicrograph images confirm the former claim. As shown inFig. 13, the main corrosion attack is appeared as localized corrosionin the SZ as a result of the dominant dissolution of Al–Mg interme-tallic networks around grains in nitric acid [26,48]. In all speeds ex-cept 800 rpm, the higher corrosion current in the HAZ than that inthe SZ does not oppose more positive Ecorr. in the HAZ; while inmany previous literature corrosion resistance improvement issimultaneously concluded both from Ecorr. decrement and icorr.

increment [2,19]. But we also take account of the different notionsof corrosion potential and corrosion current (as a thermodynamicand kinetic concepts, respectively). Afterwards, we will discuss thisclaim in more details. As seen in Fig. 13 however, x does notchange appreciably the intergranular corrosion attack in a regularmanner, but it directly influences the grain structure and passivityof the weld zones [49–51]. As illustrated in micrograph images inFig. 13b, d, f and h, the intergranular attack in the HAZ does notchange conspicuously by increasing x; however, the width ofintergranular attack in the SZ slightly increases at faster x (com-paring Fig. 13a, e and g). Regarding Table 2, and according to Ral-ston criteria (icorr. of tens to hundreds of lA/cm2 for corrosivemetal–solution combination [50]), in all x the corrosion current

of the HAZ is almost classified as low to passive corrosion rates.All micrograph observations are in a good consistency with theoutcome values of Ecorr. presented in Table 2. The intergranular cor-rosion is a outcome of the higher intermetallic density of Al–Mgphases aligned in grain boundaries (comparing Fig. 13c, d and i),i.e. the intermetallic phase of Al–Mg with less nobility has moretendencies to segregate in the grain boundaries which decreasethe dissolution potential [21,48]. As it can be seen from metallo-graphic images in Fig. 13, the parent metal generally experiencesa lower intergranular corrosion than both the SZ and the HAZ. Con-trary to Ecorr., corrosion current density of the SZ and HAZ havemore significant variations at different x. Additionally, in all x,the corrosion current of the SZ and HAZ increase by increasingx. The corrosion potential is a thermodynamic parameter that justcorrelates to the familiar chemical potential of the alloy substanceand expresses nothing about corrosion kinetic (icorr.) as a moreprominent concept in deterioration behavior of materials, particu-larly for inherently passive metals such as AA5052 [52,53]. For thepassive metal-solution combinations, the kinetic aspects of thecorrosion behavior predominate the thermodynamic aspects ofcorrosion, i.e. the influence of passivation as a result of the grainreinforcement dominates the effect of the intermetallic phases dis-solution that segregated in grain boundaries (e.g., at x = 400 or1600 rpm) [48]. Clearly, the worst weld obtained at(x = 800 rpm, t = 50 mm/min) where the SZ experiences the mostnegative anodic dissolution potential of �1.38 V vs. SCE and thehighest deterioration current density around 71 lA/cm2. The sig-nificant decline in the corrosion resistance at x = 800 rpm (as illus-trated in Fig. 13c and d) can be attributed to the cumulative effectof grain growth, low strain rate, high temperature deformation andweld micro-defects such as cracks, voids, etc. [47,51]. However, abroad grain size distribution from BM toward HAZ and SZ, or con-versely from SZ to BM, might encourage the galvanic effects be-tween various weld regions [49]. During the more complexmultipass processes compared to the FSW (e.g., equal channelangular pressing, etc.), the effect of texture variations (in additionto the grain size) on corrosion behavior should be taken into ac-count [54]. But during FSW, the main phenomenon is severe shearplastic deformation, leading to a dynamic recrystallized fine-grained microstructure in the SZ [16,21]. As shown in microscopicobservations (comparing Fig. 13a, c, e and g), for all x, the grainsize in SZ is smaller than that in the BM (it was disclosed earlierin Figs. 3 and 6b). Hitherto, there has not developed a definitivecomprehending of the influence of grain size on the corrosion rateof metallic alloys. Ralston et al. reviewed the dependency of icorr. tograin size with a diverse range of metals and suggest the followingrelationship [50]:

icorr: ¼ aþ bðG:S:Þ�12 ð8Þ

where the constant a has been claimed to be a function of environ-ment and b is a material dependant constant. Ralston et al. foundthat b assumes a negative value in low corrosive or passivatingenvironments (icorr. 6 10 lA/cm2). This critical value has been di-vided the metal–electrolyte combinations into two different classes,namely low corrosive to passive electrolytes (icorr. 6 10 lA/cm2 andb < 0) and highly active metal-environment interactions which re-veal icorr. of tens to hundreds of lA/cm2 (b > 0) [50]. However, dur-ing FSW the grain microstructure development has been claimed tobe a complex phenomenon [55]; but it was stated that a finer grainsize in SZ is usually observed by decreasing x. It was demonstratedthat during FSW of alloyed Al plates, the peak temperature in theweld region is a strong function of the pseudo-heat index (x2/t).In addition, there is a decline in the mentioned peak with increasingthe distance from the SZ centerline [16]. Regarding the values oficorr. in Table 2, almost for all x, the FSW process changes the grain

network attack pit

pit

pit

pit

(a) (b)

(c) (d)

(f)(e)

100 µm

100 µm

100 µm

100 µm

50 µm

50 µm

50 µm

50 µm100 µm

(g) (h)

(i)

Fig. 13. Optical micrographs showing corrosion attack after immersion of FSWed AA5052 in 70 wt.% HNO3 for 24 h; the values of x-zones are (a) 400 rpm-SZ, (b) 400 rpm-HAZ, (c) 800 rpm-SZ, (d) 800 rpm-HAZ, (e) 1600 rpm-SZ, (f) 1600 rpm-HAZ, (g) 2500 rpm-SZ, (h) 2500 rpm-HAZ, and (i) BM; t = 50 mm/min. Arrows indicate intergranularattack, pitting or network attacks along grain boundaries.

628 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

microstructure of the BM in such a way that the SZ–electrolyteinteraction classifies as passive interactions. Considering Eq. (8),the reduction in grain size would encourage the higher corrosion

resistance and decrease the icorr.. As a matter of fact, during FSWthe grains in SZ experience intense plastic deformation that leadsto a fine recrystallized grain structure with a higher density of

Table 3Summary of the mass loss per unit area for the SZ specimens at different x afterimmersion in 70 wt.% HNO3 for 24 h; t = 50 mm/min.

Rotation speed (rpm) Mass loss per unit area (mgr/cm2)

SZ HAZ BM

400 11.36 16.94 11.15800 31.23 17.63 7.58

1600 16.91 18.35 7.922500 27.82 13.34 13.26

-0.95

-0.93

-0.91

-0.89

-0.87

-0.85

-0.83

-0.81

0 500 1000 1500 2000 2500 3000

EO

CPvs

. SC

E (

V)

t (sec)

1. 400 rpm

2. 800

3. 1600

4. 2500

-0.86

-0.85

-0.84

-0.83

-0.82

-0.81

500 1000 1500 2000 2500

OC

Pav

e. v

s. S

CE

(V

)

t (sec)

1

2

3

4

3

4

2

1

Fig. 14. Evaluation of OCP vs. time for the SZ in 3.5 wt.% NaCl after FSW of AA5052at different x; t = 50 mm/min. Inset shows the smoothed plots of OCP vs. time.

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 629

nucleation sites for passive film formation [56]. Consequently, thefine equiaxed grains in SZ will probably accelerate the passive filmcreation and, therefore, diminish the icorr. (as shown in Fig. 11a)[57]. As noted in Section 3.1.1, by increasing the ratio of (x/t) themaximum Qfric in the SZ increases (Eqs. (2) and (3)), and accordingto the principles for recrystallization, the coarse recrystallizedgrains will be generated in the SZ [16], i.e., the corrosion currentdensity of SZ in FSWed AA5052 increases at faster x (as shown inFig. 11a). As illustrated in Fig. 12, in all x the corrosion potentialsof HAZ have nearly the same value and increasing x leads to a slightincrement of icorr. During the FSW, the grains in SZ experience an in-tense plastic deformation and thermal exposure leading to a fine-grain microstructure, while the HAZ only undergoes friction heatingand subsequently grain coarsening [55]. As it can be seen in Table 2,the icorr. of the HAZ in all x except 800 rpm are greater than thosefor the SZ. At faster x the corrosion rate of the weld zones irregu-larly increase due to the grain coarsening. Based upon the Ecorr. fac-tor, the main corrosion attack is appeared in the SZ in the form ofintergranular corrosion; while with regard to the variations of icorr.

the higher corrosion current is a result of the grain coarsening in theHAZ. Considering the corrosion properties of the weld, both param-eters have the commensurate values at the speeds of 400 and1600 rpm.

The cyclic polarization studies were conducted to investigatethe sensitivity to the pitting corrosion for FSWed AA5052. As illus-trated in Fig. 11c, the potential scanning was reversed after reach-ing the diffusion-limiting anodic current and finally intersects theforward anodic branch at the protection potential below which pitsdo not initiate (Eprot). There is a higher current density upon revers-ing the potential sweeping toward cathodic direction just for theSZ acquired at x = 800 rpm, 2500 rpm and for BM which is corre-sponding to the continued pits growing even after scan reversaland susceptibility to pitting attacks that preferentially initiates atthe surface defects, e.g. void, cracks, etc. [51,58]. No hysteresisloops were seen in the SZ and HAZ at other tool speeds which re-veal that the SZ obtained at x = 800 and 2500 rpm (see Fig. 13cand g) are merely disposed to pitting attack. Evidently, at thesespeeds a higher density of localized attacks can be seen in the SZcompared to the tool speed of 400 rpm (see Fig. 13). The most po-sitive pitting potential for the SZ and HAZ of FSWed AA5052 atx = 400 rpm exhibits a better pitting corrosion resistance of theweld. So, generally the joint demonstrated an improved corrosionproperties compared to the BM except at x = 800 and 2500 rpm;while in many of prior literature, the SZ has similar or poorer cor-rosion resistance than BM.

Fig. 14 shows the OCP variation as a function of time in 3.5 wt.%NaCl solution for the SZ which was achieved at different x.

The potential rapidly moves in the positive direction and onlyafter about 300 s reaches to a relatively stable value in view ofthe spontaneous generation of pre-layer of Al(OH)3 [46]. As itwas expected, the stable OCP as an inherently substance-correlatedparameter is not influenced during the FSW process and the varia-tions of x does not change the stable OCP value appreciably. Thereis a meaningful increase in the OCP fluctuations in the SZ atx = 800 and 2500 rpm that indicates more susceptibility to pitting

corrosion. This increment in the noise level is attributed to local-ized breakdown due to the chloride ion attack and repassivationof the protective oxide film [45]. The SZ at x = 800 rpm exhibitsthe most negative OCP that indicates its tendency to localized cor-rosions. These observations are consistent with the polarization re-sults where the SZ experiences a nobler Ecorr. at x = 400 and1600 rpm compared to those at the rates of 800 and 2500 rpm. Evi-dently, for the SZ at the x = 400 rpm, the kinetic of the passivationphenomenon is so rapid that the OCP reaches to its steady-state va-lue at the first moments of immersion in the corrosive electrolyte.This significant development of the passivation behavior is mainlydue to the finer grains and consequently the higher density ofnucleation sites for passive film formation, as was previously men-tioned [57]. The observations are all in a good accordance withthose gotten by our previous polarization investigations whilethe corrosion properties of the weld regions decline towards thetool speeds of 400, 1600, 2500 and 800 rpm, respectively. At thetool speeds extremely slower than 400 rpm, an unsuitable jointwill be attained due to an insufficient Qfric while x faster than2500 rpm cause a defective weld with inappropriate corrosionproperties in both of the HAZ and SZ [16,21]. Therefore, att = 50 mm/min, the rate of 400 rpm is obviously a preferable xto obtain a superior corrosion resistant FSWed AA5052.

3.2.1.2. The effect of the tool traveling speed alteration on thecorrosion properties of FSWed AA5052. Fig. 15a and b shows thepolarization curves of the SZ and HAZ for FSWed AA5052 at differ-ent t and at a constant x of 400 rpm in 60 g/L NaCl. It is worthmentioning that the potentiodynamic curves are the average ofseveral repeatable measurements. Table 4 summarizes the out-come data of icorr. and Ecorr.. The higher cathodic current densityfor the SZ at t = 250 mm/min indicates the acceleration of oxygenand hydrogen reduction reactions. All the anodic branches show asimilar response of the weld region to the potential scanning in60 g/L NaCl solution.

It has been elsewhere demonstrated that at a constant x, thereis no clarified alteration in the grain size by changing t, but it caninfluence the anodically-active area and decreases the reactive re-gion in FSWed zones [21]. Moreover, t and temperature of the tool-piece affect the pattern of metal flow and the grain structure. Maet al. reported that a higher ratio of (t/x), diminishes the grain sizein the nugget zone of friction stir processed AA7575 [59]. As it wasmentioned in Section 3.1.1 by comparing the macrographs inFig. 3c and d, the finer grain structure is attained at faster t as a re-

-2 -1 0 1 2 3 4 5 6

E v

s. S

CE

(V

)

Log (i) (µA/cm2) Log (i) (µA/cm2)

SZ: 50 mm/min

SZ: 80 mm/min

SZ: 100 mm/min

SZ: 200 mm/min

SZ: 250 mm/min

BM

0

-1 0 1 2 3 4 5 6

E v

s. S

CE

(V

)

HAZ: 50 mm/min

HAZ: 80 mm/min

HAZ: 100 mm/min

HAZ: 200 mm/min

HAZ: 250 mm/min

BM

(a)

Active corrosion ratesLow to passive rates

(b)

-0.2

-0.2

-0.4

-0.6

-0.8

-1

-1.2

-1.4

-1.6

-1.8

-0.4

-2

-0.6

-0.8

-1

-1.4

-1.6

-1.2

Fig. 15. Polarization curves of (a) SZ, (b) HAZ in 60 g/L NaCl, for FSWed AA5052 in various t; scan rate 5 mV/s, x = 400 rpm.

Table 4Polarization diagram parameters in 60 g/L NaCl for weld zones of FSWed AA5052 in different t; x = 400 rpm.

Traveling speed (mm/min) Ecorr. (V vs. SCE) icorr. (lA/cm2)

SZ HAZ BM SZ HAZ BM

50 �0.844 �0.858 �0.751 2.51 2.69 62.9780 �0.837 �0.825 �0.756 16.59 14.12 63.71

100 �0.84 �0.812 �0.749 67.61 40.73 66.58200 �0.831 �0.813 �0.75 63.09 64.56 70.92250 �0.74 �0.813 �0.73 0.63 77.43 65.18

630 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

sult of the synergistic effects of more intense strain rate and smal-ler friction heat flow generation. The parameters of x and t havedifferent effects on the pattern of material flow and thereby onthe microstructure of the weld zones [49,59]. In the previous sec-tion, it was seen that by increasing x the effect of the Qfric and graingrowth phenomenon in nugget region, dominates the influence ofsuperplastic deformation and recrystallization phenomena. Mean-while, the grain boundary slipping leads to the migration of bound-aries and consequently the strain-induced grain growth will occur.In addition, it is notable that the grain growth is remarkably morerapid than the formal static grain growth during superplasticdeformation [59]. The increase of t contributes mainly to the grainreinforcement as a result of Qfric diminution and strain rate incre-ment up to the point that an extremely fine grain structure inthe SZ achieved at t = 250 mm/min causes a sudden decrementin the corrosion current (compare G.S. reported in Fig. 4 with thatin Fig. 5a) [45,59]. We schematically illustrate in Fig. 16 that how

(a) (b) (c)

Fig. 16. Schematic of the competitive effect of passivation and intergranular or granular(c) 200 mm/min and (d) 250 mm/min; x = 400 rpm.

the increment of t and subsequently the grain reinforcement inSZ leads to a dominant effect of intergranular corrosion due tothe grain reinforcement (up to 200 mm/min) or predominant effectof passivation (at 250 mm/min).

Almost for all t other than 50 mm/min, the corrosion currenthigher than 10 lA/cm2 for the weld zones, reveals an active corro-sion rates where the grain reinforcement leads to the corrosionaccelerating [49,50]. Accordingly, the increment of t up to200 mm/min, causes a finer grain structure in the SZ and the higherdensity of grain boundaries improves the overall surface reactivityand opposite to the effect of x, the grain reinforcement results inicorr. increment [50,60]. But the corrosion rate retrogression of thenugget zone at t = 250 mm/min chiefly attributes to a soundmicrograins, absence of the rough anodically active intermetalliccompounds along grain boundary and finally, facileness of passivefilm formation [45,57]. The noblest Ecorr. in SZ at t = 250 mm/min isin a good agreement with the passivation behavior of the weld.

(d)

Passivated areas Unpassivated areas Intergranular or granular corrosion

corrosion for the SZ after FSW at the tool speeds of (a) 80 mm/min, (b) 100 mm/min,

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 631

Hence, the discussion on the optimal mix of x and t for attainingthe weld zone with improved corrosion properties is a significantsubject. At higher t both the effects of the faster superplastic defor-mation and lesser welding peak temperature led to an extra finegrain structure in the SZ. While at these speeds the Qfric diminish-ing is the only reason for the finer grain morphology in the HAZ.The SZ at t = 250 mm/min, experiences an extra fine grain struc-ture (as shown in photomicrograph examination in Fig. 4 and sche-

network attack

pit

pit

pit

50 µm

50 µm

50 µm

50 µm

(a)

(c)

(e)

(g)

Fig. 17. Micrographs showing corrosion attack after immersion of FSWed AA5052 in 70 w(c) 100 mm/min-SZ, (d) 100 mm/min-HAZ, (e) 200 mm/min-SZ, (f) 200 mm/min-Hintergranular attack, pitting or network corrosions along grain boundaries.

matically in Fig. 16d) which accelerate the passivation kineticswhile a smaller grain size in HAZ just activates the corrosion ratewithout any acceleration of the passive film nucleation. The SZ att = 250 mm/min has a corrosion potential of �0.74 V vs. SCE (no-bler than BM) and a corrosion rate about 0.63 lA/cm2 and the val-ues of icorr. for the HAZ and BM are approximately equal. Thespeeds higher than 250 mm/min under our experimental condi-tions are not recommended, since it can overshadow the joint from

pit

pit

100 µm

100 µm

100 µm

100 µm

(b)

(d)

(f)

(h)

t.% HNO3 for 24 h; the values of t-zones are (a) 80 mm/min-SZ, (b) 80 mm/min-HAZ,AZ,(g) 250 mm/min-SZ, and (h) 250 mm/min-HAZ; x = 400 rpm. Arrows show

632 M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634

a perspective of the insufficient Qfric; moreover, it cannot signifi-cantly change the HAZ microstustructre [16]. The width of the re-gions susceptible to the intergranular corrosion depends on thewelding parameters especially on x and t [21]. We investigatethe susceptibility to the intergranular attack in the SZ and HAZ ofAA5052 specimens which were welded at various t by exposingthem to 70 wt.% HNO3 solution at 30 �C for 24 h. Fig. 17 showsthe optical micrographs of the weld zones obtained at different tafter immersion.

As shown in Fig. 17, in the speeds of 80, 100 and 200 mm/min,the main attack occurs in the SZ and further extends to the HAZ; Inaddition, the intergranular attack is widening at faster t. The inter-granular corrosion intensity in the SZ at lower t (about 50 to80 mm/min) is slower than that of medium t (between 100 to200 mm/min). The intergranular attack in the HAZ at faster t (be-tween 100 to 250 mm/min) is more accelerated than those inslower t (about 50 to 80 mm/min). In the SZ at t = 200 mm/min(Fig. 17e), a relatively continuous destroyed networks along grainboundaries suggests the preferential attack around grain due tothe more active boundaries [10,16]. As shown in Fig. 17e and f,there are an intense intergranular or transgranular attacks at bothSZ and HAZ at t = 200 mm/min; furthermore, as it can be seen, ahigher density of preferential local attacks around grain bound-aries indicates highest susceptibility to the intergranular attack.So such a defective weld may be attributed to the collective effectsof higher density of active boundaries in the SZ and HAZ, intergran-ular attack and localized anodic dissolution of intermetallic phases[21,27,48]. As illustrated in Fig. 17g and h, at t = 250 mm/min themain attack obviously is seen in the HAZ (nearly uniform corro-sion) and the SZ undergoes a very slight and nearly insignificantintergranular attack. As mentioned earlier, the reason for inter-granular corrosion resistance of the SZ at t = 250 mm/min is re-lated to the further grain reinforcement and the absence ofcoarse grain that retards the passivation kinetics [47,50]. In otherword, at t = 250 mm/min, the grain size becomes extremely fineso that the passivation prevails the overall surface dissolution inthis microstructure with large number of grain boundaries[47,49,57]. The observations are consistent with the potentiody-namic polarization results. So considering all the above featuresduring FSW of AA5052 x = 400 rpm, the best corrosion resistant

0 500 1000 1500 2000 2500 3000

Wei

ght

loss

(m

gr)

Rotation Speed (rpm)

SZ

HAZ

0

10

20

30

40

50

60

70

80

0 1000 2000 3000

i cor

r. ( µ

A/c

m2 )

Rotation Speed (rpm)

SZ

HAZ

Base Metal

3.2

2.7

2.2

1.7

1.2

0.7

0.2

(a)

Fig. 18. Comparison between the variations of weight loss and corrosion current densi(immersion was carried out for about 410 h in 3.5 wt.% NaCl solution at 30 �C).

weld is acquired at t = 250 mm/min. It implies that during theoptimization efforts in the previous section, the fixed t of50 mm/min was not the best speed for FSW of AA5052.

3.2.2. Static immersion test in 3.5 wt.% NaCl solutionFinally, we compare the results of potentiodynamic polarization

test with weight loss experiment just to determine the validity ofthe icorr. variation trend for FSWed samples. The weight loss ofspecimens after immersion in 3.5 wt.% NaCl solution for about410 h are presented in Fig. 18a and b, respectively as a functionof x and t. For a better comparison in each case, the results of icorr.

for the SZ and HAZ versus x and t are presented. The points of thecurves refer to the average weight loss and the bars indicate thestandard deviation of the three parallel immersed samples. All datafor the weight loss are listed in Tables 5 and 6. After 410 h ofimmersion, the samples were covered by a white adhering driedAl(OH)3 gel [44]. We made use of a distilled chromic acid and stiffplastic brush to remove all adhering layers.

After immersion, the highest weight loss for the SZ atx = 800 rpm is corresponds to our previous polarization results.As mentioned earlier a smaller mass loss for the HAZ at slower xcould be related to the nucleation and growth of the passive film.The variation of x does not remarkably change the mass loss ofthe specimens that was cut off from the HAZ region. Moreover,by increasing t at a constant x of 400 rpm, up to 200 mm/minthe weight loss increases and then suddenly decreases att = 250 mm/min. This observation is consistent with our discussionabout the improved corrosion resistance due to an extra fine grainstructure obtained at the speeds of (x = 400 rpm, t = 250 mm/min). The highest weight losses at t = 100 and 200 rpm confirmthe worst corrosion properties of the SZ at these rates as shownpreviously. The trend of the weight loss variation for HAZ speci-mens as a function of t differs slightly from that for corrosion cur-rent density. But generally, they approximately follow the sametendencies. Apart from the rates of 50 and 250 mm/min, for all tthe SZ specimens indicate higher weight loss than the HAZ speci-mens. During FSW process of AA5052, the highest weight losseswere obtained under the following condition: (x = 800 rpm,t = 50 mm/min) and (x = 400 rpm, t = 100 or 200 mm/min).Undoubtedly, the results of the weight loss test together with

0 100 200 300

Traveling Speed (mm/min)

SZ

HAZ

i cor

r. ( µ

A/c

m2 )

0

20

40

60

80

40 90 140 190 240Traveling Speed (mm/min)

SZ

HAZ

(b)

Wei

ght

loss

(m

gr)

3.2

2.7

2.2

1.7

1.2

0.7

0.2

ties for the SZ and HAZ of FSWed AA5052 (a) at different x and (b) at different t.

Table 5The mass of the SZ and HAZ specimens acquired at different x, before and after immersion in 3.5 wt.% NaCl solution for 410 h and the relevant results of weight losses;t = 50 mm/min.

Rotation speed (rpm) Wo (gr) W1 (gr) DW (mgr) DWave. (mgr)

SZ HAZ SZ HAZ t (hr.) SZ HAZ SZ HAZ

400 0.4155 0.307 0.4153 0.3062 410 0.2 0.8 0.621 0.9340.3605 0.3502 0.3595 0.3494 410 1 0.80.3843 0.4056 0.3836 0.4044 410 0.7 1.2

800 0.3496 0.3026 0.3471 0.3011 410 2.5 1.5 2.1 1.2340.4048 0.3338 0.4035 0.3327 410 1.3 1.10.4471 0.3651 0.4446 0.364 410 2.5 1.1

1600 0.3614 0.3468 0.3603 0.3452 410 1.1 1.6 0.967 1.440.3347 0.3293 0.3342 0.3277 410 0.5 1.60.3245 0.2631 0.3232 0.262 410 1.3 1.1

2500 0.378 0.3615 0.3776 0.3601 410 0.4 1.4 0.4 1.40.2995 0.3757 0.299 0.3744 410 0.5 1.30.3409 0.3869 0.3406 0.3854 410 0.3 1.5

Table 6The mass of the SZ and HAZ specimens obtained at different t, before and after immersion in 3.5 wt.% NaCl solution for 410 h and the relevant results of weight losses;x = 400 rpm.

Traveling speed (mm/min) Wo (gr) W1 (gr) DW (mrg) DWave. (mgr)

SZ HAZ SZ HAZ t (h) SZ HAZ SZ HAZ

50 0.4155 0.307 0.4153 0.3062 410 0.2 0.8 0.621 0.9340.3605 0.3502 0.3595 0.3494 410 1 0.80.3843 0.4056 0.3836 0.4044 410 0.7 1.2

80 0.4301 0.3718 0.4293 0.3708 410 0.8 1 1.3334 1.03340.4115 0.3448 0.4101 0.3437 410 1.4 1.10.3642 0.3862 0.3624 0.3852 410 1.8 1

100 0.3171 0.3503 0.315 0.3488 410 2.1 1.5 1.9363 1.46670.2832 0.376 0.2813 0.3746 410 1.9 1.40.396 0.3416 0.3941 0.3401 410 1.9 1.5

200 0.4426 0.4352 0.4409 0.4336 410 1.7 1.6 2.054 1.36670.4361 0.29 0.4342 0.2888 410 1.9 1.20.4631 0.553 0.4611 0.5516 410 2 1.4

250 0.4418 0.3207 0.441 0.3199 410 0.8 0.8 0.6 1.0010.3735 0.2573 0.373 0.2566 410 0.5 0.70.3815 0.3145 0.381 0.313 410 0.5 1.5

M. Bagheri Hariri et al. / Materials and Design 50 (2013) 620–634 633

our prior observations for AA5052 show an improvement of corro-sion properties of the BM after FSW under conditions of(x = 400 rpm, t = 250 mm/min), (400 rpm, 50 mm/min) and(1600 rpm, 50 mm/min).

4. Conclusions

The effects of the consequential parameters of FSW (x and t) onthe corrosion and mechanical properties of AA5052 have beencarefully studied. At different conditions, the final grains dimen-sions are mainly influenced by the combination effects of the fric-tion heat flow and the degree of deformation. The finalrecrystallized grain size depends on the factor that dominantlydetermines the FSW process. The most intense intergranular attackin the form of network attacks around grain boundaries was ob-served for the SZ which obtained at x = 800 and 2500 rpm. Thewidth of intergranular attack in the SZ slightly increases at fasterx. The variations of x between 400 to 2500 rpm at a fixedt = 50 mm/min, transform the microstructure in a way that theSZ–electrolyte interaction was classified as passive interactionswhere the finer grain size leads to a lower corrosion activity. Inspite of x, the alteration of t between 80 to 200 mm/min at a con-stant t = 400 rpm, change the structure so that the SZ–electrolytecombination can be classified as active corrosion rates interactionswhere the grain reinforcement leads to the corrosion development.

The remarkable decline of the SZ corrosion rate was seen at(x = 400 rpm, t = 250 mm/min) and is mainly attributed to the ex-tra fine grain structure that accelerates the passive film nucleation.At t = 80, 100 and 200 mm/min, the major intergranular attack wasseen in the SZ, while the main corrosion was obviously occurred inthe HAZ at (x = 400 rpm and t = 250 mm/min). It was found thatthe optimum corrosion and mechanical properties of the weldzones were obtained at (x = 400 rpm, t = 250 mm/min). At thenext level, the combinations of (400,50) and then (1600,50) wereintroduced as the preferential welding conditions, respectively.

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