Friction Stir-Welded Dissimilar Aluminum

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Friction Stir-Welded Dissimilar AluminumAlloys: Microstructure, Mechanical Properties,

and Physical StateM. Ghosh, Md.M. Husain, K. Kumar, and S.V. Kailas

(Submitted March 22, 2013; in revised form July 11, 2013; published online August 8, 2013)

A356 and 6061 aluminum alloys were joined by friction stir welding at constant tool rotational rate withdifferent tool-traversing speeds. Thermomechanical data of welding showed that increment in tool speedreduced the pseudo heat index and temperature at weld nugget (WN). On the other hand, volume ofmaterial within extrusion zone, strain rate, and Zenner Hollomon parameter were reduced with decrease intool speed. Optical microstructure of WN exhibited nearly uniform dispersion of Si-rich particles, fine grainsize of 6061 Al alloy, and disappearance of second phase within 6061 Al alloy. With enhancement in weldingspeed, matrix grain size became finer, yet size of Si-rich particles did not reduce incessantly. Size of Si-richparticles was governed by interaction time between tool and substrate. Mechanical property of WN wasevaluated. It has been found that the maximum joint efficiency of 116% with respect to that of 6061 alloywas obtained at an intermediate tool-traversing speed, where matrix grain size was significantly fine and

those of Si-rich particles were substantially small.

Keywords aluminum alloys, friction stir welding, light microscopy,mechanical characterization

1. Introduction

Nowadays, friction stir welding (FSW) of versatile materials

finds widespread application in automobile and aerospaceindustries for fabricating primary and secondary components.The technique was developed and patented in the UK in 1990

by The Welding Institute for welding of plates in solid state(Ref 1). The procedure is a complex solid-state thermome-chanical process, in which a rotating tool with a shoulder and

pin moves through rigidly clamped plates, placed in butting/lap/fillet joint configuration over metallic support (Ref 2).Shoulder maintains intimate contact with top surface ofworkpiece. Heat is generated by friction at shoulder and pinsurfaces. Material gets softened under severe plastic deforma-tion, and flow occurs along welding direction with thetranslation of tool. Material is thus transported from the front

end of tool to trailing edge, where it is forged to form a joint.

Along the welding line, the side, where the direction of toolrotation is the same as that of traversing of tool, is called

advancing side, with other side being termed as retreating side.Ultimate microstructure considering grain size, second-phase

fraction, dissolution and reappearance of precipitate in differentzones like heat-affected zone (HAZ), thermomechanicallyaffected zone (TMAZ), and weld nugget (WN) along with

joint efficiency depend on total heat input, cooling rate, plasticstrain, material flow, and state of stress. The plastic strain andstrain rate are very high and substantially greater thanconventional metal-working processes like extrusion, rolling,forging, etc. (Ref3). Welding variables, for example, rotationalspeed, traversing speed, tool tilt angle, plunging depth,plunging speed, motor torque, normal load, and tool design,

have predominant effect on welding defects, residual stress, andjoint quality (Ref 1). Three types of material flow have beenidentified during FSW. First is the rotation of plasticized

material around tool, which is governed by the revolution oftool generating friction between tool and workpiece. Second isthe downward movement of material by pin nearby itself andsubsequent upward motion of an equivalent amount of materialaway from pin. Third one is the relative motion of material

between tool and workpiece (Ref2).Till date, the process has made reasonable breakthrough in

respect of copper alloys, magnesium alloys, titanium alloys,steels, nickel alloys, molybdenum alloys, Al-alloy matrixcomposites, and thermoplastics (Ref 1, 3). However, major

thrust has been given for FSW of Al and its derivatives toproduce lap/butt joints consisting of similar materials. Suchtype of studies highlighted microstructural-mechanical propertyco-relation, mechanism of weld-zone formation, heat input and

temperature rises, materials flow prediction using markertechnique, fatigue response, fracture toughness, corrosionresistance, and residual stress determination for transition joints(Ref2,4-9).

It has been established that variation in weldability of Alalloys occurs owing to the presence of various alloying

elements and heat-treatment condition. In general, precipitationbehavior in 2xxx and 7xxx series alloys have been matched

M. Ghosh and MD.M. Husain, Materials Science & TechnologyDivision, CSIRNational Metallurgical Laboratory, Jamshedpur831007, India; K. Kumar, Department of Materials Science &Engineering, University of Northern Texas, Denton, TX 76203-5017;and S.V. Kailas, Department of Mechanical Engineering, IndianInstitute of Science, Bengaluru 560012, India. Contact e-mails:ghosh_mnk@yahoo.com and mainakg@nmlindia.org.

JMEPEG (2013) 22:38903901 ASM InternationalDOI: 10.1007/s11665-013-0663-3 1059-9495/$19.00

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with precipitation hardenable 6xxx series alloys. Coarsening ofprecipitates and solutionization of needle precipitates wereobserved in HAZ of welded alloy 6063 alloy. Characterizationat WN of friction stir-welded 6061 Al alloy indicated the

presence of second phase, but their identity was not established(Ref 10). In another attempt, same AA6061 Al alloy plates,both in O and T6 temper conditions, were joined by FSW usingfour different sets of weld parameters (Ref10). Microstructuraland mechanical characterization of the joints were made bydetailed optical microscopy investigations, extensive hardness

measurements and tensile tests. The effect of temper conditionon joint performance was explored in addition to the effect ofweld parameters. For alloy AA 6082, it is found, that somecoarsening of precipitates occurred along with partial dissolu-

tion of them at high FSW temperature (Ref2).Microstructural investigation in nonheat-treatable friction

stir-welded Al alloys in cold worked,tempered condition hasdescribed the loss of hardness across the weld. FSW of Al-5086plates in H32 condition at tool rotational speed of1600 rpmand tool-traversing speeds of 175, 200, and 225 mm/minexhibited failure through stir zone during tensile test (Ref11).Microhardness measurements and bend test results were inaccordance with tensile properties. This was attributed to the

loss of cold-work hardening within stirred zone due to heat

generation during welding. No cracking phenomenon wasobserved at the time of bend testing of welds, although a littleporosity was present. FSW of 5082 Al alloy was also carriedout, which produced uniform hardness across the weld (Ref2).

For heat-treatable Al alloys, FSW promoted the formation ofrelatively softer region at WN with respect to other parts oftransition joint due to dissolution of second phases; however,strength could be regained with subsequent growth of strength-ening precipitates during thermal cycle (Ref 12, 13). For

nonheat-treatable Al alloys, FSW produced hardened region atWN because of the development of fine grain structure andhomogeneous distribution of nondissolved second phase. HAZ

became weak mainly because of decrease in dislocation density

and increase in grain size (Ref14,15). Considering the diverseresponses as mentioned above, a few attempts have been madeto study weldability of two different Al alloys. One of theexamples is friction stir spot welding (FSSW) between AA2024-T3 and 5754-H22 alloy sheets at tool rotation speed of

1500 rpm (Ref 16). Tool plunge depths were of 2.45,2.55, and2.65 mm from specimen surface with dwell timesof 2, 5, and 10 s. The maximum strength was achieved for thejoint produced by placing AA 5754-H22 sheet on the top with

the tool plunge depth of2.65 mm and dwell time of 10 s withthe characteristic of pull-out nugget fracture. The minimumstrength was obtained for the joint produced by placing AA2024-T3 sheet on the top with tool plunge depth of2.45 mmand dwell time of 2 s with typical cross-nugget failure. It has

been inferred that an increase in tool penetration depth up to acertain limit ensured the increment in joint strength.

Challenge arises when welding consists of dissimilar alloysbecause of their different physicochemical properties, precip-itation behavior, trend in defects accumulation, and phase

transformation characteristic during welding. FSW of 5083alloy to 6061 alloy under various tool rotation and traversingspeeds is one of the examples (Ref 17). For this couple,microhardness near interface showed heterogeneous distribu-

tion, and the bond strength was 63% with respect to that of6061 alloy with drop in elongation. In another endeavor,Kumbhar and Bhanumurthy (Ref 18) have pointed out the

absence of rigorous mixing during FSW of 5052-6061 alloy.They reported the highest bond strength of 71% of that of6061 alloy with elongation of3%. Effect of FSW parameterson microstructure was evaluated for transition joint consisting

of 2024-T3 and 7075-T6 Al alloys (Ref 19). For this couplestirring zone (SZ), microstructure was heterogeneous becauseof shorter welding time and decorated with onion-ring-likepattern illustrating differences in grain size and composition.Microhardness at WN was close to the hardness of 2024 Al,

and weaker region appeared at the periphery of shoulder. Most

of the weld joints failed from HAZ of 2024 Al, and in fewcases, from SZ; however, overall strength and elongation of theassemblies were lower than those of base alloy. Commerciallypure aluminum (CP Al) was also friction stir welded at different

welding speeds with 7039 aluminum alloy (Ref20). For them,onion-ring pattern was found at WN. Microhardness showed anupward trend starting from CP Al side with a little zigzag near

SZ. FSW and FSW followed by heat treatment indicated UTSvalues to the tune of 61 and 68% of that of CP Al,respectively, for the same couple. Fracture occurred either

through CP Al or CP Al-HAZ interface. Palanivel and Mathews(Ref 21) have attempted to join 6351-T6 to AA5083-H111aluminum using different tool geometry at constant rpm but

various tool-traversing speeds. The highest bond strength was

275 MPa when welding was done with square cross-sectionpin at 63 mm/min welding speed. FSW of 6061 Al to 2024 Alreported formation of three regions within nugget; a dispersionzone, stirring-induced plastically deformed zone containing

alternate lamellar structure, and equiaxed zone (Ref22). Mileset al. (Ref23) have made dissimilar Al alloy joints consisting of5182/5754, 5182/6022, and 5754/6022 combinations. Mechan-ical properties of welds were evaluated under biaxial strain, bytransverse tension test, and in stretching through OSU plane

strain testing. Tailor-welded blanks were produced by FSW ofAA 5182-H111 to AA 6016-T4 aluminum alloy (Ref 24).These welds showed rupture during formability test because ofweld root defects. AA6082-T6 and AA6061-T6 aluminum

alloys were also friction stir welded in the recent past (Ref25).Microhardness values decreased both at WN and HAZ for thesame couple. Bond strength (32%) of assembly was lowerwith respect to parent alloys, and fracture occurred throughTMAZ and HAZ. It has been concluded that the loss of T6condition was responsible for deterioration in joint efficiency. A

single alloy in two different heat-treated conditions was alsofriction stir welded; for example, FSW of AA7075 Al alloy inannealed and aged conditions with parameters 750/100, 1000/

150, 1250/200, and 1500/300 rpm/mmmin (Ref26). Plates canbe satisfactorily friction stir welded with a large window ofparameters in O-temper condition, whereas the weld parameterswere stringent to obtain adequate joint strength in T6 condition.Joints produced in O-temper condition displayed a hardness

increase in weld region with a tensile failure at base metal,whereas those produced in T6-temper condition exhibited ahardness drop in weld region with tensile fracture at weld zone.FSW of 7xxx to 2xxx series Al alloy was also reported, wherecomplexity occurred because of phase transformation (Ref27).

In few attempts FSW was compared with other joiningtechniques. One of the examples is power beam and FSW of5005_H14, 2024_T351, 6061_T6 and 7020_T6 Al alloys plateswith thickness 3 mm (Ref28). Microhardness, microtensile,macrotensile and fracture toughness were evaluated for all

joints. Fracture toughness at fusion zone was found higher forall joints with respect to that of base alloy owing to reduction in

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yield strength. It has been reported that difference in strengthlevels for two types of joints has meager effect on toughness.

From above glimpses, it is evident that very negligibleefforts were made to produce welds consisting of cast andwrought alloys. Tehyo et al. (Ref29) has produced FSW joints

between SSM 356 and AA6061-T651 alloys. SZ exhibited finegrain structure produced by mixing of two alloys. It has beenfound that 6061 Al alloy moved from retreating side to upperregion of advancing side and SSM 356 alloy travelled vice

versa. TMAZ showed distorted elongated grain structure owingto mechanical working by tool. Microhardness across weld waslow; however, at WN it was comparatively higher than TMAZand HAZ. Tensile strength was in the range of180-197 MPawith elongation being 7-9.4%, depending on welding param-eters. Lee et al. (Ref30) has carried out FSW of A356 and 6061

Al alloys at 1600 rpm with different tool speed. WNmicrostructure was dominated by retreating side alloy withonion-ring pattern at the edge of weld bead. At WN, dynamicrecrystallization resulted in fine dispersion of Si-rich particles,

dislocation cell structure, and fine equiaxed grain size. Hard-ness of SZ was dropped with respect to base alloy owing todissolution of second phase. Tensile strength in transverse

direction was close to A356 alloy, and the maximum longitu-dinal tensile strength was 192 MPa.

Thus, joining of the Cast to wrought aluminum alloy stillremains a partially explored area and needs more attention tooptimize welding parameters to fabricate joints with betterefficiency. In this respect, WN plays a major role fordetermining the quality of assembly. Therefore, in the current

investigation, major thrust has been given to quantitativeaspects of heat evolution, temperature rise, extent of straingeneration, and material transport during FSW of dissimilar Al

alloys, which are responsible for final microstructural andmechanical properties at WN.

2. Experimental

In the current endeavor, Al-Si (AS) and Al-Mg (AM) alloyswith dimensions 100 (l)930 (w)93 (t) mm3 were joined byFSW. Chemical composition and tensile properties of base

materials are furnished in Table 1and2, respectively.Welding was done in indigenously designed FSW equip-ment. Tool was made of high-speed steel with concave shoulderdiameter of15 mm, pin diameter of5 mm, and cylindricalpin length of 2.6 mm. Tool tilt angle (3) and rotationalspeed (1000 rpm) were kept constant at the time of joining.During welding, data logger was used to record normal load,traverse load, and spindle torque. AS and AM alloys were fixedat retreating and advancing sides, respectively. Before welding,

substrates were machined to obtain flat surface along transversedirection and then cleaned in acetone. Joining was done undervariable tool-traversing speeds as given in Table 3. It has been

found in most reports that rotational speed of 1600 rpm waspreferable during FSW of dissimilar Al alloys (Ref29,30). Inthe current investigation, the same has been reduced to decrease

heat generation and flash formation.Microstructural investigation was done on transverse section

of welds. Both base alloys and welded specimens wereprepared by conventional metallographic technique, etched

with Kellers reagent and examined in optical microscope(LEICA DM 2500M). Microhardness was evaluated acrossweld line on transverse section nearly at mid thickness regionunder a 50-g load with a 10-s dwelling time (LEICA VHMTAuto). Subsize tensile specimens were prepared from transverse

section of weld as per ASTM E 8/8M-11 keeping weld line atthe center of gauge length. Gauge length and thickness oftensile specimen were 10 and 1.5 mm, respectively. Testwas performed at a crosshead speed of0.1 mm/min in tensiletesting machine (Hounsfield) and repeated for four samples foreach set. Cross sections of failed samples were examined inoptical microscope to identify failure location.

3. Results

Microstructure of parent alloys is shown in Fig. 1. AS alloyconsisted of bright primarya-Al matrix and eutectic network ofAl-Si (Fig.1a). Area fraction of network was small in

comparison with conventional eutectic Al-Si alloy owing toless Si content than eutectic composition. Distribution ofSi-rich particles within grain body was scanty. AM alloy

exhibited quasi-polygonal/little elongated grains containingdark spots within matrix indicating the presence of secondphase (Fig.1b). As conjectured in literature, second phase was

Table 1 Chemical compositions of substrates

Alloy

Elements in wt.%

Si Mg Ti Fe Cu Zn Mn Al

Al-Si alloy (AS) 4.3 0.420 0.17 0.20 0.01 0.06 0.01 Bal

Al-Mg alloy (AM) 1.4 1.1 0.016 0.26 0.60 0.07 0.007 Bal

Table 2 Mechanical properties of substrates at room

temperature

Tensile properties Al-Si alloy Al-Mg alloy

UTS, MPa 233.4 11 351.8 16

Elongation, % 3 0.5 12 1

Hardness, VHN 71.4 1 94.7 1

Table 3 Friction Stir welding parameters and sample

nomenclature

Combination

Tool-traversing

speed, mm/min

Tool rotating

speed, rpm

Al-Si (AS) vs. Al-Mg (AM) alloy 70 100080 1000

130 1000

190 1000

240 1000

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presumably Mg2Si (Ref 30). Grain size of AM alloy was

30-35 lm. Grain boundary of AM alloy was decorated with

broken network of Si-rich particles.WN microstructures after FSW are revealed in Fig. 2, 3,4,

5, and 6. SZ was free from weld defect and decorated withcharacteristics of both alloys (Ref 8). A curved line wasobserved within WN. Region A in macroimage of weld wasenlarged in Fig.2-6(c), presenting predominant characteristicof retreating side. This area consisted of nearly homogeneous

distribution of fragmented Si-rich particles, arose from breakingdown of parent dendritic network of Al-Si. Breaking up andredistribution of Si-rich particles at WN and TMAZ were alsoindicated by Nandan et al. (Ref2) in friction stir-welded cast

Al-Si alloys along with healing of casting defects. Region B ofmacroimage exhibited development of fine-grained structure

because of dynamic recrystallization (Fig.2-6a). The grainsizes of this region were 23 lm (Fig.2a) and 14 lm(Fig.6a) at the lowest and the highest tool-traversing speed,respectively. The same region contained fine dispersion ofSi-rich particles, which came from parent alloys owing tomaterial churning by tool. At the slowest welding speed of

70 mm/min, within SZ, onion-ring pattern was found (Fig.2d).It was a lamellar-like structure, consisting of alternate bands ofAM and AS alloys. The stacked morphology has not beenobserved for other welded specimens. Onion-ring formationand its dependence on FSW parameters have been discussed in

Fig. 1 Optical microstructures of base materials: (a) AS and (b) AM alloy

Fig. 2 Optical microstructure of joint welded at 70 mm/min tool speed showing different regions across weld

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Fig. 3 Optical micrographs of weld processed at 80 mm/min tool speed exhibiting different regions

Fig. 4 Optical micrographs of FSW joint made at 130 mm/min tool speed presenting various regions across weld

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details by Threadgill et al. (Ref3) and Krishnan (Ref5). The

onion-ring patterns for 7075 and/2014 Al alloys were theoutcome of etching response due to variations in grain size,texture, and dislocation density between the rings (Ref 3).

According to their postulation during welding, extrusion of onelayer of semi-cylinder occurred with one rotation of tool. Therewas a small time lag between the production of heat by tool

rotation and extrusion of hot metal by forward tool advance-ment (Ref 3). This cyclic process produced continuous set ofsemicircular rings. Therefore, ring formation became the

function of tool geometry, tool rotation, and tool-traversingspeeds. It has been also indicated that increment in tool speedmay lead to disappearance of onion ring as it happened beyond70 mm/min tool-traversing speed in the current study. At highspeed of tool beyond a certain value, though weld formationwas present, one of the processes became recessive owing to

too short a time gap between them resulting in the absence ofonion ring. Practical significance of the phenomenon remainednearly unexplored as mechanical properties of the nuggetbecame satisfactory and fracture path in mechanical tests was

seldom associated with this feature (Ref3).Microhardness distribution perpendicular to weld line has

been illustrated in Fig.7. Welded assembly can be divided intothree regions as per microhardness profiles, such as the regionoutside of tool shoulder in retreating side, the zone near the

weld line, and the area outside of tool shoulder in advancingside. The first region at retreating side exhibited minimummicrohardness value. A maximum in microhardness profile wasobtained in shoulder-processed region. For all joints, this

middle zone revealed (dotted regions in Fig. 7) a number of upsand downs owing to development of composite structurethrough material mixing. The oscillation in hardness at nuggethas been also inferred by Threadgill et al. (Ref3) for friction

stir-welded heat-treatable Al alloys because of contribution oftwo constituent alloys.

Figure8 shows the tensile properties of welds along with

broken tensile specimen. Ultimate tensile strength was the lowestat the slowest tool-traversing speed (170 MPa), reached themaximum at 80 mm/min tool-traversing speed (409 MPa), andreduced upon further increment in tool-traversing speed (300-330 MPa). Previously it has been reported that FSW jointconsisting of A356 and 6061 Al alloys exhibited tensile strengthto the tune of 185-208 MPa (Ref30). Welded joints of AA5052-AA6061 alloys reported maximum bond strength of225 MPawith 7% ductility, which was 73% with respect to that of AA6061

alloy (Ref18). Ultimate tensile strength of friction stir-weldedSSM 356 to AA6061-T651 alloy joint was 180-191.3 MPa(Ref 29). FSW of high strength 6061-5083 aluminum alloysrevealed reduction in bond strength in comparison with parent

alloyand was 63%withrespect to that of 6061 alloy(Ref17).Inall these illustrations, dimensions of tensile specimens withrespect to weld assembly, fracture location, and crack propaga-tion have not been indicated clearly. Compared with these

reports, the current investigation showed substantial improve-ment in bond strength of WN for all friction stir welds except forthe joint, which was made at 70 mm/min tool-traversing speed.Ductility of friction stir-welded jointswas low in the current caseand close to the breaking strain of AS alloy.

Fig. 5 Optical photographs of weld fabricated at 190 mm/min tool speed revealing various regions across weld

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Failure location during tensile testing of FSW joints waswithin A356 Al alloy with small shift according to weldparameter. At the lowest bond strength (70 mm/min toolspeed), it was away from weld center (2.8 mm, Fig.9a); atthe highest bond strength, fracture occurred close to weld centerline (Fig.9b); and for the rest, fracture moved through region a

little away from weld center line (1.0-1.5 mm, Fig.9c).Fracture path did not propagate exactly through bond line. Thisobservation might be compared with inference of Threadgill(Ref 31). According to that author, the interface, i.e., visible

after FSW of dissimilar material could be termed as joint-lineremnant, and the presence of the same did not affect tensileproperties of joint. Existence of sharp boundaries was alsoreported for dissimilar welds like 2219-T87/7075-T6 and 5083-H321/6082-T6 alloys (Ref3).

4. Discussion

Microstructures at WN of friction stir-welded joints weredependent on quantified values of thermal and physical states.

Thermal state could be described by pseudo heat index (PHI)and peak temperature at WN. Physical state could be explainedby extrusion volume, strain rate, and Zenner Hollomon

parameter. These parameters can be elucidated as follows:PHI is a signature of heat input during welding and could be

expressed as (Ref32)

PHIRPM2

IPM EPLPD cos a ; Eq 1

where RPM is the tool rotation/min, IPM the tool-traversing

speed (mm/min), EPL is the effective pin length (mm), PD isthe plunge depth (mm), and a is the tool tilt angle ().

For given tool geometry and plunge depth, peak temperature

at WN was influenced by tool rotation and travel speed.Ignoring the minor difference in temperatures at advancing andretreating sides during welding, general expression of maxi-

mum temperature is (Ref33)

T

TmK

x2

Vf104

a; Eq 2

where T is the temperature at WN (C); Tm is the meltingpoint of alloy (C); x is the rotational speed (rev/s); Vf is the

forward travel speed (mm/s); and a and Kare the constantswith values 0.05 and 0.70, respectively.

During welding, material in front of pin moved toward theback of the tool and got consolidated. In unit time, the totalamount of material that traveled around was governed byextrusion zone width. In case the material passing through the

retreating side is greater than the material travelling around theadvancing side, then volume of material passing throughextrusion zone in unit time could be given by the expression(Ref34):

Fig. 6 Optical images of FSW joint prepared at 240 mm/min tool speed illustrating different regions across weld

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V h wr Vt Vf 2Vfkd2

; Eq 3

where Vis the volume (mm3), h is the pin length (mm), wr is

the width of extrusion zone (mm) measured from transverse

section after etching, Vt is the pin tangential velocity (mm/s),

d2 is the projected curved pin area (mm2), and k is the num-ber of threads per unit length of pin. Now, Vt was relatedwith tool rpm as

Fig. 7 Microhardness profiles along OO line of Fig. 2-4for joints produced with (a) 70 mm/min, (b) 80 mm/min, and (c) 130 mm/min

Fig. 8 Tensile properties of friction stir-welded dissimilar Al alloy (a) variation in UTS-elongation with tool speed, and (b) broken tensile spec-

imen (dimensions are not to scale)

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Vt 2R:x; Eq 4

where Rp is the pin radius (mm).

Second term of Eq 1 was zero in the current welding trials

because a cylindrical tool without thread was used.Final grain size at WN obtained through continuous

dynamic recrystallization was the function of strain rate. Highstrain at WN propelled fine grain structure. Strain rate has beenfound to be the function of Zener Hollomon (ZH) parameterand related as per the expression below (Ref34):

Z _e exp Q

RT0

Eq 5

where Z is the Zenner-Hollomon parameter, e is the strainrate (s1), Q is the activation energy for process and consid-ered to be 190 kJ/mol, and T is the WN temperature in K.

Now,

_e et ; Eq 6

where e is the total strain at WN, and t is the time requiredfor deformation and can be calculated from the relation:

t2Rp

Vf.

Total strain at WN during FSW was determined byReynolds (Ref35) assuming material only passed around tool

in streamline path. In this respect, difference in flow behaviorof two Al alloys was ignored assuming minor variation atelevated temperature. The derived expression for the same isgiven by

e ln l

APR

ln

APR

l

: Eq 7

where APR is the tool advance per revolution, and

l 22Rp cos1 2Rpx2Rp

. In the above relationship, x is the

distance perpendicular to welding direction from retreating sideof tool. Strain would be maximum when l would reach itsmaxima, and in that condition, x = 2Rp.

From the above equations PHI, T,V, e, andZwere evaluated

and presented in Fig.10, 11-12. The influence of individualparameter on microstructural and mechanical properties ofjoints is discussed in the following section.

It has been revealed in Fig.10that PHI and temperature at

WN were increased with drop in the tool-traversing speed(Fig.10). This trend supports the inference of Nandan et al.(Ref36). Moreover, the effect of temperature on second phasehas been discussed in the literature, where it has been inferred

that the temperature of 402 C helped in dissolving Mg2Siand AlFeSi precipitates of 6061 Al alloy (Ref30). Disappear-ance of hardening precipitate was also observed during FSW of6XXX Al alloy by Nandan et al. (Ref2). In the current study,temperature at WN was in the range of475-500 C (Fig.10);therefore, all second phases except Si-rich particles were

dissolved in solid solution.Figure11exhibited a monotonic relation between volume of

material transport and tool-traversing speed. This findingendorsed Arbegast Model (Ref 33) on extrusion zone

flow calculation, which proposed that higher tool speedpropelled more materials sweeping through extrusion zone.

Fig. 9 Microstructures near failure locations for different tool-traversing speeds (a) 70 mm/min, (b) 80 mm/min, and (c) 240 mm/min

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This phenomenon did not influence WN microstructuredirectly; however, it was indirectly responsible for enhancing

strain rate at nugget (Ref37).

The WN microstructure consisted of fine grain structure anddistribution of tiny Si-rich particles through severe plasticdeformation, and could be directly related with strain rate andZH parameter as shown in Fig. 12. Higher tool travelling speed

promoted faster cooling. Increased cooling rate, on the otherhand, enhanced the strain rate at WN (Ref6,38). Higher strainrate propelled finer matrix grain size with increased ZHparameter.

Microhardness profile can be co-related with welded jointmicrostructure. Inherent low hardness of A356 Al alloy was

reflected at retreating side. Increment in hardness near weld linewas due to composite microstructure where both the alloyscontributed. Further increment at advancing side was thesignature of higher strength of 6061 Al with respect to A356 Al

alloy. Though 6061 Al alloy lost its strength by precipitate sdissolution as mentioned before, yet it was compensated forby matrix grain refinement and appearance of fine dispersion ofSi-rich particles.

Tensile properties of WN exhibited an excellent micro-

structural dependence. At the lowest tool-traversing speed(70 mm/min), heat input and peak temperature were thehighest, strain rate was the lowest, and ZH parameter was thesmallest. These resulted in the highest matrix grain size

(220.8 lm). Therefore, bond strength was the least for that

particular joint. Comparing with the average microhardness asshown in Fig.7(a), the same value also displayed minimumamong all welds (78 VHN). With increment in tool-traversingspeed (80 mm/min), WN temperature and heat input decreased,

strain rate increased, and ZH parameter enhanced. Theseaspects reduced the grain size (181 lm) and produced tinySi-rich particle distribution (3-6 lm). Strength of weldedjoint reached the highest level. Average microhardness of thisregion also became the largest considering all transition joints

(94 VHN). Still further increment in traversing speed of tool(130 mm/min) encouraged lowering of WN heat input andtemperature, and increments of strain rate and ZH parameter.

Matrix grain size became still smaller (150.6 lm). However,

size of Si-rich particles within WN was increased (4-12 lm)owing to reduction in interaction time between material andtool. Matrix strengthening by Si-rich particles was reduced.This phenomenon was also revealed in the marginal drop inaverage microhardness value (89 VHN) in that region(Fig.7c) with respect to the weld processed at 80 mm/min.

Moreover, this created local notch weakening through stressconcentration. Bond strength also decreased. Decrement inbond strength continued with increment in traversing speed of

the tool beyond 130 mm/min, and at the highest speed of240 mm/min, the size of Si-rich particles became 15 lm. Inthat situation, positive contribution owing to grain sizedecrement was overshadowed by matrix weakening throughstress concentration by large-sized Si-rich particles. The low

breaking strains of all welds (2-4%), might be because ofembrittlement due to heterogeneous microstructure at WNrestricting the deformation of grains.

5. Conclusion

In the current investigation, friction stir butt welding wasperformed to join A356 and 6061 Al alloys under variable tool-

traversing speeds in the range of70-240 mm/min. Temper-ature rise at WN was in the range of 475-500 C during

Fig. 10 Variation in PHI and temperature at WN during FSW of

dissimilar aluminum alloy

Fig. 11 Volume of material passing through extrusion zone for dif-

ferent tool-traversing speeds during FSW of aluminum alloy

Fig. 12 Strain rate and ZH parameter during FSW of aluminum

alloy

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welding, resulting in dissolution of precipitates of 6061 Alloy.With increment in tool-traversing speed, strain rate and ZHparameter increased gradually leading to fine grain structures of6061 alloy within SZ. However, increase in tool speed was

responsible for the reduction in heat input and temperature riseat WN. Eutectic network of Al-Si in A356 alloy and grainboundary Si-rich phases of 6061 alloy were fragmentedbecause of severe deformation and became homogeneouslydistributed within WN. Si-rich particle distribution was depen-dant on interaction time of tool with substrate; higher tool speed

led to lower transit time to produce relatively larger Si-richparticles. Tensile strength and microhardness profile within WNwere governed by microstructure. To obtain maximum jointefficiency, fine dispersion of Si-rich particles with optimum

matrix grain size was preferable. In this respect, 80 mm/mintool-traversing speed was found optimal to achieve jointefficiency of116% with respect to that of 6061 Al alloy.

Acknowledgments

The authors are indebted to Director-NML for his kind support

during the study, as well as providing permission to publish the

research study. The cooperation received from Dr. A. K. Ray

during investigation is also gratefully acknowledged. The authorsare deeply indebted for the financial support received from the

Department of Science & Technology, Govt. of India, New Delhi

through sanction letter no. SR/S3ME/028/2007 dated 08/11/2007

to carry out the investigation.

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