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Materials Science and Engineering A 527 (2010) 29512961
Contents lists available at ScienceDirect
Materials Science and Engineering A
journa l homepage: www.e lsev ier .co
Tensile iorand fric oy
S.M. Cho idaa Department o to, Onb Aerospace Ma uncil Cc
Department o enue W
a r t i c l
Article history:Received 17 NReceived in reAccepted 11 Ja
Keywords:AZ31 magnesium alloyDouble-sided arc weldingFriction
stir weldingMicrostructureTensile propertiesStrain hardeni
d wo31B-Hgher,ed s
ng exation
twice that of the base metal. The DSAWed samples exhibited
stronger strain-hardening capacity due tothe larger grain size
coupled with the divorced eutectic structure containing -Mg17Al12
particles in thefusion zone, compared to the FSWed samples and base
metal. KocksMecking type plots were used toshow strain-hardening
stages. Stage III hardening occurred after yielding in both the
base metal and thewelded samples. At lower strains a higher
strain-hardening rate was observed in the base metal, but it
1. Introdu
Weightimportantmronment [1alloys, haveandaerospato-weight
rcastability [to further exwelding (FSWelding Intial for joinweld
defec[1,5]. FSWplastic defoworkabilityOn the othe
CorresponE-mail add
0921-5093/$ doi:10.1016/j.ng decreased rapidlywith
increasingnetowstress. Athigher strains the strain-hardening rateof
theweldedsamples became higher, because the recrystallized grains
in the FSWed and the larger re-solidied grainscoupledwith particles
in the DSAWed providedmore space to accommodate
dislocationmultiplicationduring plastic deformation. The
strain-rate sensitivity evaluated via Lindholms approach was
observedto be higher in the base metal than in the welded
samples.
2010 Elsevier B.V. All rights reserved.
ction
reduction in ground vehicles and aircraft is one of theeasures
to improve fuel economyandprotect the envi-
]. Magnesium alloys, as the lightest metallic structuralbeen and
will be increasingly used in the automotivece industries due to
their lowdensity [1], high strength-atio [13], environmental
friendliness, recyclability and2,3]. However, effective joining
techniques are requiredpand the applications
ofmagnesiumalloys.Friction stirW), a solid-state joining technique
developed by Thestitute of Cambridge, UK, in 1991 [4], has great
poten-ing magnesium alloys, since it can signicantly reducets
normally associated with fusion welding processeshas also been used
to rene the grain size via severermation and recrystallization [69]
so as to improve theof Mg alloys and increase the strength of
welded joints.r hand, the weldability of magnesium alloys by
some
ding author. Tel.: +1 416 979 5000x6487; fax: +1 416 979
5265.ress: [email protected] (D.L. Chen).
arc welding processes such as gas tungsten arc welding (GTAW)
isconsidered to be excellent as well [10]. In 1999, Zhang and
Zhang[11] developed and patented a novel arc welding process
referredto as double-sided arc welding (DSAW). The DSAW process
usesone welding power supply and two torches; frequently a
plasmaarc welding (PAW) and GTAW torch each connected directly to
oneof the power supply terminals. The torches are positioned on
oppo-site sides of a work-piece such that the welding current ows
fromone torch through the work-piece to the opposite torch. Zhang
etal. [1215] have examined the feasibility of using the DSAW
pro-cess to make vertical-up, keyhole-mode welds in 612mm
thickplain carbon steel, stainless steel or aluminum alloy plates.
Morerecently,Weckman and co-workers [16,17] have examined the
fea-sibility of using the DSAW process for conduction-mode
weldingof 1.2mm thick AA5182-O aluminum sheet for tailor welded
blankapplications. It was noted that the opposing welding torches
andsquare-wave AC welding current successfully cleaned the
oxidefromboth sides of the joint and produced visually
acceptableweldsat speeds up to 3.6m/min. Through-thickness heating
was moreuniform with DSAW than with other single-sided welding
pro-cesses allowing symmetric welds to be produced with
minimalangular distortion of the sheets [16,17].
see front matter 2010 Elsevier B.V. All rights
reserved.msea.2010.01.031properties and strain-hardening behavtion
stir welded AZ31B magnesium all
wdhurya, D.L. Chena,, S.D. Bholea, X. Caob, E. Powf Mechanical
and Industrial Engineering, Ryerson University, 350 Victoria
Street, Toronnufacturing Technology Centre, Institute for Aerospace
Research, National Research Cof Mechanical and Mechatronics
Engineering, University of Waterloo, 200 University Av
e i n f o
ovember 2009vised form 8 January 2010nuary 2010
a b s t r a c t
Microstructures, tensile properties anand friction stir welded
(FSWed) AZrates. While the yield strength was hithe FSWed samples
than in the DSAWin the FSWed samples. Strain-hardeniLudwik equation
and a modied equm/locate /msea
of double-sided arc welded
jkoc, D.C. Weckmanc, Y. Zhouc
tario M5B 2K3, Canadaanada, 5145 Decelles Avenue, Montreal,
Quebec H3T 2B2, Canadaest, Waterloo, Ontario N2L 3G1, Canada
rk hardening behavior of double-sided arc welded (DSAWed)24
magnesium alloy sheet were studied at different strain
both the ultimate tensile strength and ductility were lower
inamples due to welding defects present at the bottom
surfaceponents were evaluated using the Hollomon relationship, the.
After welding, the strain-hardening exponents were nearly
-
2952 S.M. Chowdhury et al. / Materials Science and Engineering A
527 (2010) 29512961
Mechanical properties such as strength, ductility,
strain-hardening behavior, strain-hardening sensitivity, etc., of
all weldsmade inmagnesiumalloypartsused in structural
applicationsmustbeevaluated toensure the integrityandsafetyof the
joint andstruc-ture. Whilemagnesiumerties of we[18] studiedsile
propertreported therties innonalloy. Zhu eon the welddiode
laserbeendoneoAZ61 [25] alier results oFSWed magumented (etests on
a Mthe values oing strain raAZ31B-H24FSWed thixhardening[32]
studiesium alloy sSome authonesium allograin size smaterial
[2formation owrought mpasses (i.e.,side and thside of
thedouble-sidebeen examicould be usable mechatherefore,
wproperties,and FSWed
2. Materia
In the pwas used.2.53.5wt%[37]. Twoemployed tthe butt
joimadewith ttion of thetorch wereA detailed dparameterspieces
wereon the surfaarea of thewwere madeof 1.4 kW.10mm/s
an1.65mmroface oxideswas cleaned
The welded joints perpendicular to the welding direction werecut
and cold mounted in order to examine the microstructure ofthe
fusion zone (FZ), heat-affected zone (HAZ) and basemetal (BM).The
mounted samples were manually ground, polished, and etched
cetic picral (10mL acetic acid (99%), 4.2 g picric acid, 10mL0mL
ethanol (95%)) [38]. The microstructure was
observedopticalmicroscope equippedwith quantitative image anal-
ftware. Vickers microhardness tests were conducted withuterized
Buehler machine across the sectioned weld with
ing of 0.5mm. A load of 100g and dwell time of 15 s wereduring
the hardness tests. Sub-sized tensile specimens in
ancewithASTME8M-08 standard [39]weremachined alongling (or
longitudinal) direction for both the base metal andjoints, where
the weld was positioned at the center of the
area. Tensile testswereperformedusing a computerized ten-ting
machine at constant strain rates of 1102, 1103,4 and 1105 s1 at
room temperature. At least two sam-ere tested at each strain rate.
The fracture surfaces wereed using a scanning electron microscope
(SEM) equipped
n energy dispersive X-ray spectroscopy (EDS) system and 3Draphic
analysis capacity.
ults and discussion
icros
mic1, wes wbaseby ronealimAZ3welof t
d zoed anparised inken poy wHAZargeallizesent
Typicathere are numerous investigations on the properties
ofalloys, only a limited number of studies of the prop-lded
magnesium alloy joints are reported. Quan et al.the effects of heat
input on microstructure and ten-
ies of laser welded AZ31 Mg alloy. Liu and Dong [19]e effect of
microstructural changes on the tensile prop--autogenousgas
tungstenarcweldedAZ31magnesiumt al. [20] presented the effect of
welding parametersing defects and change of microstructure in CO2
and
welded AZ31 magnesium alloy. Tensile testing has alson friction
stirweldedAZ31 [1,5,2124], FSWedwroughtnd a ne-grained laser welded
Mg alloy [26]. Some ear-n the microstructural changes and strengths
of variousnesium alloys and other alloys have also been well
doc-.g., in refs. [2729]). Takuda et al. [30] performed
tensileg9Li1Y alloy at room temperature and observed thatf
strain-hardening exponents increased with increas-te. Afrin et al.
[24] obtained similar results for a FSWedMg alloy. Yu et al. [31]
evaluated the tensile strength ofomolded AE42 Mg alloy, but no
information on strain-and strain-rate sensitivity was given, while
Lee et al.d the formability of friction stir welded AZ31 magne-heet
and other alloys experimentally and numerically.rs have studied the
strain-hardening behavior of mag-ys with emphasis on the
relationships between thetrengthening and dislocation strain
hardening of the,24,3335]. While Shen et al. [36] have evaluated
thef macropores in double-sided gas tungsten arc weldedagnesium
AZ91D alloy plates made with two separatewelded with one partial
penetration weld on the topen a separate partial penetration weld
on the backplate), the mechanical properties of conduction-moded
arc welds made in magnesium alloy sheet have notned. It is unknown
if this novel arc welding techniqueed to produce welds in magnesium
alloys with accept-nical properties. The aim of the present
investigation,as to evaluate and compare the microstructure,
tensile
strain-hardening and strain-rate sensitivity of DSAWedAZ31B-H24
Mg alloy sheet.
ls and experimental procedure
resent study, 2mm thick AZ31B-H24 Mg alloy sheetThe nominal
chemical composition of this alloy wasAl, 0.71.3wt% Zn, 0.21.0wt%
Mn and balance Mg
different welding methods, DSAW and FSW, wereo make autogenous
welds between the work-pieces innt conguration. Both DSAWed and
FSWed joints werehewelding direction perpendicular to the rolling
direc-sheet. In the DSAW process, a PAW torch and a GTAWused with a
square-wave AC welding power supply.escription of the welding
apparatus and other processusedmay be found in [16,17]. Prior to
DSAW, thework-degreased using acetone and then alcohol. The oxidece
of the sheets was then mechanically removed in theeld using a
stainless steelwire brush. TheDSAWweldsusing a welding speed of
25mm/s and welding powerThe FSW welds were made using a welding
speed ofd a right-hand threaded pin tool having a pin length
oftating clockwise at a rate of 2000 rpm. Prior to FSW, sur-were
removed with a steel brush and then the surfaceusing ethanol as
well.
using aH2O, 7with anysis soa compa spacappliedaccordthe
rolweldedgaugesile tes110ples wexaminwith afractog
3. Res
3.1. M
Thein Fig.ing sizof thesheettial anabout 5FSWedthe
topsectionaffecteequiaxin comappearalso tathe allof thesome
lrecrystthe pr
Fig. 1.tructure
rostructure of the AZ31B-H24 Mg base metal is shownhere
elongated and pancake-shaped grains with vary-ere observed. The
heterogeneity in the grain structuremetal was due to both
deformation of the 2mm thicklling and incomplete dynamic
recrystallization (par-ng) [1]. The average grain size of the base
metal was. The typical macroscopic and microscopic structures
of1B-H24 Mg alloys are shown in Fig. 2. Fig. 2(a) showsd bead after
FSW and Fig. 2(b) presents a typical cross-he FSWed sample
including HAZ, thermomechanicallyne (TMAZ) and stir zone (SZ). As
seen in Fig. 2(c), bothd elongated grains were present in the HAZ.
However,on to the base metal (Fig. 1), far more equiaxed grainsthe
HAZ, indicating that partial recrystallization hadlace during FSW.
The recrystallization temperature of
as approximately 205 C. Thus, the temperature in partmay have
been above this value. This is conrmed bygrains observed in the HAZ
due to grain growth afteration [1]. The grain structure in the TMAZ
(Fig. 2(d)) instudy is basically equiaxed and recrystallized,
which
l microstructures of the base metal (BM) of the AZ31-H24 Mg
alloy.
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S.M. Chowdhury et al. / Materials Science and Engineering A 527
(2010) 29512961 2953
Fig. 2. TypicaAZ31-H24 allo(c) heat-affectstir zone (SZ).
was similarand Afrin etions [41] welongated gbecame notThese
chanl macroscopic and microscopic structures of a friction stir
weldedy. (a) Top weld bead surface, (b) cross-section of the welded
joint,ed zone (HAZ), (d) thermomechanically affected zone (TMAZ),
and (e)
to the recent results reported by Cao and Jahazi [1]t al.
[5,40], while it was different from earlier observa-here the TMAZ
was still characterized by deformed andrains. The grains in the SZ
were equiaxed (Fig. 2(e)) andiceably bigger in the center of the
stir zone (8m).ges were caused by dynamic recrystallization
during
Fig. 3. TypicalAZ31-H24 alloheat-affected
FSW [42]. Aand Jahazi [and Lim et
Fig. 3(a)section ofweld beadthe oxide. Tshown in Fito the
effectshows an eqple that waof DSAWedmacroscopic andmicroscopic
structures of a double-sided arcweldedy. (a) Top weld bead surface,
(b) cross-section of the welded joint, (c)zone (HAZ), and (d)
fusion zone (FZ).
larger grain size in the SZ was also reported by Cao1], Afrin et
al. [5], Fairman et al. [40], Pareek et al. [43],al. [44].and (b)
shows the top weld bead surface and a cross-
the DSAW weld, respectively. It is seen that the topquality
appeared excellent with complete cleaning ofhe bottom weld bead
surface was similar to the top. Asg. 3(b), there was a slight
sagging of the weld pool dues of gravity on the molten pool during
welding. Fig. 3(c)uiaxed microstructure of the HAZ in the DSAWed
sam-s completely recrystallized. The grain size in the HAZsample
(Fig. 3(c)) was larger than that in the HAZ of
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2954 S.M. Chowdhury et al. / Materials Science and Engineering A
527 (2010) 29512961
Fig. 4. EDS lin
the FSWedtemperaturfusion boungrain size iwith (Mge scan showing
the compositional variation across the particles in (a) base metal,
(b) hea
sample (Fig. 2(c)). This was attributed to the highere
experienced in the HAZ of the DSAWed sample. Thedary had a slight
hour glass shape (Fig. 3(b)) and then the FZ (Fig. 3(d)) became
further larger (15m)17Al12) phase particles arising from a divorced
eutec-
tic that formthe solidic
EDS anacles/inclusiFig. 4(a). St-affected zone (HAZ), (c) and
(d) fusion zone (FZ) of a DSAWed joint.
ed in the interdendritic and intergranular regions ofation
microstructure in the FZ.lysis indicated that some MnAl containing
parti-ons were present in the base material, as shown inimilar
results were observed by Lin and Chen [45].
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S.M. Chowdhury et al. / Materials Science and Engineering A 527
(2010) 29512961 2955
F
Fig. 4(b) showithEDSanMnAl contspite of themicrostructdivorced
euboth phaMnAl contboundary, aimage. Thetic -Mg anbe seen in Fdue
to the faing, since nlayers of taining onlythe coolingcooling).
Asmany -Mglaser/arc hymany spotto be the intermetalljoints. Ben-and
reportecored -Mgregions whformed to welding andAZ91 and Amanaban
eintergranul
3.2. Microh
A typicaalloy is showgradually frto approximjoints. Thistures in
thein the FZ inBM (Fig. 1).from the dethe
half-harecrystallizfactors led
ypicaly, FSW
bsered at
nsile
6 shoof that ae strhadultimof abboutratecreasr thel becin raM30h
theasesencrfacen inage
rate od Fiserventsubsof t
hat sleadig. 5. Microhardness prole across a DSAWed sample.
ws a typical particle in the heat-affected zone. Line scanalysis
signied that theseparticleswere similar to thoseaining particle
observed in the base metal (Fig. 4(a)), inremainder of the material
experiencing a transition ofure in the HAZ (Fig. 3(c)). Fig. 4(c)
revealed clearly atectic structure in the FZ of the DSAWed sample,
withse particles on the left-hand side and the unmeltedaining
particle on the right-hand side along the grainlthough both
particles had the same white color on theeutectic-like structure
consisting of alternating eutec-d eutectic -Mg17Al12 along the
grain boundary couldig. 4(d). The presence of the eutectic-like
structure wasst non-equilibrium cooling of theweld pool afterweld-o
normal eutectic structure in the form of alternatingand would be
possible in the AZ31 Mg alloy con-3wt% Al based on the equilibrium
phase diagram if
rate would be innitely slow (i.e., under equilibriumseen in Fig.
4(c) and (d) and Fig. 3(d) there were also
17Al12 particles within the grains. Liu et al. [46] studiedbrid
welding behavior on Mg alloy and reported that precipitates formed
within the Mg grain were likely-Mg17Al12 phase. Liu et al. [47]
also observed MgAlic brittle phase in the FZ of TIG welded Mg/Al
dissimilarHamu et al. [48] studied GTA welded AZ31B Mg alloyd that
the fusion zone microstructure consisted of amatrix and a divorced
eutectic in the interdendritic
ich were originally Mg17Al12 that subsequently trans-phase
(Mg32(Al,Zn)49) intermetallics. Electron beamgas tungsten arc
welding had been employed to weld
Z31B Mg alloys, respectively, by Su et al. [49] and Pad-t al.
[50], and they observed ne equiaxed grains withar -Mg17Al12
precipitates as well.
Fig. 6. Tbase allo
were ooccurr
3.3. Te
Fig.curvestestedboth thsamplehigherciencyonly astrainUTS inrate
foand %Eof stra[30], A
Botof the bthe pretomsuas showSEM imstraintion analso
obmovemtion ofsurfaceclear twouldardness
l hardness prole across the DSAWed AZ31B-H24 Mgn in Fig. 5. It
is seen that the hardness value decreased
om about HV 70 in the half-hardened H24 temper BMately HV 50 at
the center of the FZ of the welded
is due to the formation of non-equilibrium cast struc-FZ (Fig.
3(d)), in conjunction with a larger grain sizecomparison with that
in the HAZ (Fig. 3(c)) and in theFurthermore, the grain shape had a
signicant changeformed and elongated (or pancake-shaped) grains
in
rdened H24 condition (Fig. 1) to the fully annealed ored
equiaxed grains in the HAZ (Fig. 3(c)). All of theseto the hardness
change shown in Fig. 5. Similar results
a signicanmature failand ductilitand (c)), inwith those
It has re(downwardjoint in anunderstandferent toolbutt joint
otools werestir regionshand, additdefect-free,engineering stress
versus engineering strain curves of theAZ31B-H24ed and DSAWed
samples tested at a strain rate of 1105 s1.
ved for the FSWed joints, where the lowest hardnessthe center of
stir zone as well.
properties
ws typical engineering stress versus engineering straine
basemetal, FSWed andDSAWedAZ31Mg alloy sheetsstrain rate of 1105
s1. It is seen that after welding,ength and elongation were
reduced. While the FSWedhigher yield strength (YS), the DSAWed
sample had aate tensile strength (UTS) and elongation. A joint
ef-
out 83% was achieved for the DSAWed joints, but it was72% for
the FSWed joint. Fig. 7 presents the effect ofon the tensile
properties. It was clear that the YS anded and ductility (%El)
decreased with increasing strainbase metal, but the effect of
strain rate on the YS, UTSame smaller after both types of welding.
Similar effectte on the YS and UTS was also reported in
Mg9Li1Y[51], cryo-rolled Cu [52] and AZ31B alloys [2,24,53].UTS
and%El of theDSAWed joints lay in-between thosemetal and the FSWed
joints. The reason behind this wase of a signicant welding defect
observed near the bot-of FSWed samples using a right-hand
threadedpin tool,Fig. 8. The welding defect could be better seen
from the
s taken from a fracture surface after tensile testing at af 1103
s1 shown in Fig. 9(a) at a lower magnica-
g. 9(b) at a higher magnication. Similar defects wereed by Cao
and Jahazi [1] who noted that the upwardof the material in the stir
zone may cause the forma-urface porosity or even root notches near
the bottomhe work-piece when the right-hand pin was used. It isuch
defects at the bottom surface in the FSWed jointsto a strong notch
effect or stress concentration andhavet inuence on the mechanical
properties, causing pre-
ure as shown in Fig. 6. As a consequence, both the UTSy of the
FSWed samples were reduced notably (Fig. 7(b)spite of the slightly
higher YS (Fig. 7(a)), in comparisonof DSAWed joints.cently been
reported that tool proles and axial
forceforce)haveasignicanteffecton thedefect-freeFSWed
Al alloy and the subsequent tensile properties [54]. Tothe
effect of axial force and tool pin prole, ve dif-
pin proles and three different axial force levels to thef Al6061
Al alloy were examined. Square pin proledobserved to produce
defect-free, good quality friction, regardless of the applied axial
force levels. On the otherional axial force (7 kN) increased heat
input and led togood quality friction stir regions as well,
irrespective of
-
2956 S.M. Chowdhury et al. / Materials Science and Engineering A
527 (2010) 29512961
Fig. 7 , and (
tool pin proright-handeoperation ipin ends dutool. The
sufacewerenin the samewelding deFigs. 8 androtated cloc
Fig. 8. An OMspeed of 10min the clockwi
tenth w. Effect of strain rate on (a) yield strength (YS), (b)
ultimate tensile strength (UTS)
les [54]. In the present study, a cylindrical tool withd threads
was used for the FSWed joints. During FSW
Thestrengt was observed that metal owed up and stuck to thee to
the use of right-handed threads in the rotatingbsurface porosity
and notch defect at the bottom sur-ot observed in the FSWwhenusing
left-handedpin toolsclockwise rotation [5]. This suggests that the
observed
fects at the bottom surface of FSWed joints as shown in9 were
due to the use of a right-hand threaded pin toolkwise.
micrograph of a FSWed sample near the bottom surface at a
weldingm/s and rotational rate of 2000 rpm using a right-hand
threaded pinse rotation.
speed (v) twelding dewas still eqwelding spein the
preseliterature [effect on threduced dution aroundof these destrain
of thinternal toin the gaugsignicantenhanced tmoted the fit is
necessa
3.4. Strain-
The harratio of they [24,57].hardening c
Hc =UTS
y
The obtwelded samcapacity wac) ductility of the base metal,
FSWed and DSAWed samples.
sile test data are summarized in Fig. 10, where theas presented
as a function of the ratio of the welding
o the rotational rate (). Despite the presence of thefects at
the bottom surface of the FSWed joints, the UTSuivalent to those
reported in the literature at variouseds and tool rotational rates
[5,43,44]. The YS obtainednt study was indeed higher than those
reported in the
5,43,44]. The fact that the pores did not have a stronge YS but
the UTS could be reducedwas simply due to thectility. That is, the
sample failed due to strain localiza-the defects before the UTS was
reached. The number
fects appeared to have a dominant effect on the truee welded
joint at localization. These voids, which werethe material,
resulted in a reduced cross-sectional areae area of the material
[55]. Some authors also observedinteractions between the surface
defect and pores thathe formation of macroscopic shear bands, which
pro-ailure [56]. Therefore, to increase the UTS and ductilityry to
minimize or avoid the weld defects during FSW.
hardening behavior
dening capacity of a material may be considered as aultimate
tensile strength UTS, to the yield strength
Afrin et al. [24] re-dened a normalized parameter ofapacity, Hc,
as follows:
y = UTSy
1 (1)
ained hardening capacity of the base metal and theples is listed
in Table 1. It is seen that the hardenings enhanced after FSW,
similar to the results reported
-
S.M. Chowdhury et al. / Materials Science and Engineering A 527
(2010) 29512961 2957
Fig. 9. SEM imof 1103 s1defects at a hig
by Afrin et awas furthereffect of thethehardenithat was fuof the
mateYS accordindislocationity. A decreresistance b
Fig. 10. A comspeed (v) to thein the literatur
Table 1Hardening capacityof thebase alloy, FSWedandDSAWedsamples
testedatdifferentstrain rates.
Specimen Strain rate (s1) Hardening capacity
Base metal
Friction stir
Double-side
reduced thDSAWed sain Figs. 3(dcapacity init should bstrong
textverse direc
er, akeneof t
tionthe (d shld stengted asentages of fracture surfaces of a
FSWed sample tested at a strain rate. (a) Overall view at a lower
magnication, and (b) details of welding
Howevbe weachangediffracalongresolvethe yiesile strincreasthe
preher magnication.
l. [24]. The hardening capacity in the DSAWed sampleshigher than
that in the FSWed samples. There was littlestrain rate on the
hardening capacity. Based on Eq. (1)
ng capacity of amaterialwas related to its yield strengthrther
associated with the microstructure and texturerial. An increase in
the grain size would decrease theg to the HallPetch relationship
[5860] and increasestorage capacity, leading to thehigher hardening
capac-ase in the grain size reduced the difference of the owetween
the grain boundary and interior, which in turn
parison of the YS and UTS, as a function of the ratio of the
weldingrotational rate (), obtained in thepresent studywith those
reportede.
To undethe strain-hhardening eharden; theing for a gihave
propoexponent. H
= Kn,wheren is tcoefcient,To better qutted their
= y + K1where n1 iscoefcientwhardeningtion using nyielding:
= y + K
where n*,true stress,rial, respecincrement( y) = 1.Fig. 11.
Eq.tic deformaonly the datUTS are usein Fig. 11 w1102 0.371103
0.361104 0.361105 0.41
welded sample 1102 0.491103 0.671104 0.661105 0.57
d arc welded sample 1102 0.841103 0.861104 0.841105 0.84
e hardening capacity [57]. Since the grain size of themples was
larger than that of FSWed sample, as shown) and 2(e), respectively,
the higher value of hardeningthe DSAWed samples would be
anticipated. In addition,e noted that for rolled Mg plates there
was usually aure with the (0002) plane perpendicular to the
trans-tion (TD) and normal direction (ND) of the plate [61].fter
FSW and DSAW the texture would be expected tod. Yang et al. [61]
reported that after FSW a signicantexture was observed with a high
intensity of (0 002)around the TMAZ, and the FSWed joint prone to
slip0002) with an orientation of 45 at a lower criticalear stress
exhibited an approximately one third drop inrength but only a
modest decrease in the ultimate ten-h. This implied that the
hardening capacity after FSWs well, in good agreement with the
results observed instudy.
rstand strain-hardening behavior it is better to examineardening
exponent of different materials. The strain-xponent is a measure of
the ability of a metal to strainlarger its magnitude, the greater
the strain harden-
ven amount of plastic strain [58]. Several researcherssed
different equations to evaluate the strain-hardeningollomon [62]
gave the following expression:
(2)
he strain (orwork)hardeningexponent,K is the strength is the
true stress and is the true strain [60,6264].antify the
strain-hardening response, Chen and Lu [65]tensile curves using the
Ludwik equation [64,66]:
n1 (3)
the strain-hardening exponent and K1 is the strengthhich
represents the increment in strengthdue to strain
at =1. Afrin et al. [24] proposed the following equa-et ow
stress and net plastic strain of materials after( y)n
(4)
, , y and y are the strain-hardening exponent,true strain, yield
strength and yield strain of a mate-
tively. K* is the strength coefcient which reects thein strength
due to strain hardening corresponding toThe above three equations
could be better illustrated in(2) has the origin positioned at O
(including the elas-tion stage where the Hookes law holds true),
althougha in the uniform deformation stage between the YS andd. Eq.
(3) corresponds to a shift of the origin fromO toO1here the yield
stress is excluded, while Eq. (4) proposed
-
2958 S.M. Chowdhury et al. / Materials Science and Engineering A
527 (2010) 29512961
Fig. 11. Schematic illustration of a true stress versus true
strain curve.
recently by Afrin et al. [24] represents a further shift of the
originfromO1 toO*. It implies thatboth theyield stress andyield
strainareprecluded. Fig. 12 presents the evaluated strain-hardening
expo-nents (n, n1, n*) as a function of strain rate for the base
metal,FSWed and DSAWed samples. Almost no effect of strain rate
onthe strain-hardening exponents was seen in the base metal. But
forthe DSAWed and FSWed samples, the strain-hardening
exponentsincreased with increasing strain rate. The
strain-hardening expo-nents evaluated according to all the above
three equations wereobviously higher after welding, and the DSAW
resulted in some-what higher strain-hardening exponents than the
FSW, as shownin Fig. 12. It is also seen that the n values were the
smallest and n1values were the highest with n* lying in-between the
two. Similarresults were also reported by Afrin et al. [24].
One of themost important contributions to the strain hardeningis
related to the formation and multiplication of dislocations. In
theplastic deformation stage, the net ow stress due to
dislocationdensity could be expressed as [24,33,34]:
y (5)where is the dislocation density. The net ow stress
necessaryto continue deformation of a material is proportional to
the squareroot of the dislocation density. The dislocation density
in a metalincreases with deformation or cold work due to
dislocation multi-plication or the formation of new dislocations
which decreases thespacing among dislocations and their
interactions become repul-sive. The net result would be that the
motion of a dislocation isimpeded by other dislocations. As the
dislocation density increases,the resistance to dislocation motion
by other dislocations becomesmore pronounced. Thus, a higher stress
is necessary to deform ametal [58].
Fig. 13 shows a typical KocksMecking plot of
strain-hardeningrate ( =d/d) versus net ow stress ( y) at different
strainrates from 1102 s1 to 1104 s1 for the base metal. It isseen
that no stage I hardening (or easy glide) which dependsstrongly on
the orientation of the crystal and stage II linear harden-ing where
the strain-hardening rate should be constant occurred.Srinivasan
and Stoebe [67] reported that the presence of stage IIstrain
hardening could be due to the interactions of the disloca-tions in
the primary slip system with those in an intersecting slipsystem.
Stage III hardening is characterized by a hardening ratethat
decreases monotonically with increasing ow stress leadingto the
much repeated term parabolic hardening on the stressstrain curve
[68]. This stage is very sensitive to the temperatureand rate of
deformation [33]. In this present study test temper-ature remained
constant but strain rates were changed. As seenfrom Fig. 13, stage
III was somewhat strain-rate sensitive, i.e., theFig. 12. Effect of
strain rate on (a) n-value, (b) n1-value and (c) n*-value in the
base metal, FSWed and DSAWed samples.
-
S.M. Chowdhury et al. / Materials Science and Engineering A 527
(2010) 29512961 2959
Fig. 13. strain-hardening rate () versus net ow stress ( y) of
the base metaltested at different strain rates.
strain-hardening rate increased with increasing strain rate.
Simi-lar result was observed by Lin and Chen [2] in AZ31B extruded
Mgalloy as well. This could be explained on the basis of the
followingequation [2
= 0 Rd1
where isq is an expand is rathetemperaturthe presentremained
chardeningstress.
As stageappeared. Dto track thewith geom[33]. Howevextended toto
the disorhardening rstage IV instrain-hard
Fig. 14 srate versussamples tesnet ow stthe base
mestrainsornecouldbeun
Fig. 14. strainFSWed and DS
dislocation strain hardening [24,34,35]:
= 0 + HP + d, (7)
where 0 is the frictional contribution, HP = kd1/2 is
theHallPetch contribution and d =MGb1/2 is the Taylor disloca-tion
contribution (where G is the shear modulus, b is the Burgersvector,
M is the Taylor factor and is constant). Sinclair et al. [70]and
Kovacs et al. [71] reported that at lower strains the grain sizehad
a strong contribution to the strain hardening and the inu-ence of
the grain size on the strain hardening vanished at higherstrains
due todislocation screening anddynamic recovery effects atgrain
boundaries. Since the basemetal had a smaller grain size cou-pled
with its deformed grain structure, the HallPetch contribution(HP)
would be stronger at lower strains, leading to a higher
strain-hardening rate in comparison with the welded joints. On the
otherhand, at higher strains or later stage of deformation strain
harden-ing was higher in the welded joints. This could be due to
the largergrain size and the presence of particles in the fusion
zone of theDSAWed samples. It was stated that the presence of
precipitatesincreased the dislocation density [72] and larger grain
size pro-
ore space to accommodate dislocations [24]. As the grainthe Das
obetalhereAZ3
ce ofiled psamthermatlm [lowin
() +
Linde ploLindbasepectich ww stonshws:
= k2
d,33]:
/q
(6)
the strain-hardening rate in stage III, 0 is a
constant,erimental stress exponent varying with temperaturesr large
namely about n=8 at higher and n=35 at loweres, Rd is also a
temperature dependent parameter. Inexperiments Rd and q were
constant as temperatureonstant. It is evident from Eq. (6) that the
strain-rate increased with increasing strain rate at a given
III approached a saturation level, hardening stage IVue to the
low hardening level of stage IV, it is difculthardening mechanisms
because they must compete
etrical and structural instabilities including damageer,
Pantleon [69] reported that stage III hardening wasstage IV by
incorporation of excess dislocations relatedientations. He further
stated that in stage IV the workate often remains constant. It is
seen from Fig. 14 thatthe base metal remained almost constant with
a lowening rate.hows a typical KocksMecking plot of
strain-hardeningnet ow stress in the base metal, FSWed and
DSAWedted at a strain rate of 1102 s1. At lower strains orress
levels after yielding the strain-hardening rate oftal was higher
than that of the welded joints. At highertowstress levels itwas
reversed. The strainhardeningderstoodon thebasis of thegrain size
strengtheningand
vided msize ofrate wbase mtions wsizes inpresenples faFSWed
Anoior ofLindhothe fol
= 0
ThemL. Thwherefor the3.5, resing, whand oa relatias follo
ln -hardening rate () versus net ow stress ( y) of the base
metal,AWed samples tested at a strain rate of 1102 s1.
Fig. 15. A typivalue of 2.5% vsamples.SAWed samples was larger,
a strong strain-hardeningserved after DSAW in comparison with the
FSW and the(Fig. 14). This is also in agreement with other
observa-the hardening rate decreasing with decreasing grain
1 alloy was presented [24,34,73]. However, due to theweld
defects at the bottom surface, the FSWed sam-rematurely, thus the
net ow stress up to failure in theple was the lowest (Fig.
14).important parameter involving thedeformationbehav-erials was
the strain-rate sensitivity (SRS), m. The74] approach was used to
evaluate the SRS based ong equation:
1() log (8)
holm SRS is 1() in the equation commonly termed ast of versus
log at 2.5% true strain is shown in Fig. 15,holm SRS is represented
by the slope, mL. The mL valuesmetal, DSAWed joins and FSWed joints
were 7, 4 and
ively. It follows that the SRS became lower after weld-ould be
related to the difference in the microstructure
ress. del Valle and Ruano [35] have recently presentedip between
stress, SRS and grain size for the AZ31 alloy
1/2(Mc 2Mcg) + Mcg, (9)
cal plot used to evaluate the strain-rate sensitivity,mL , at a
true strainia the Lindholms approach for the base metal, DSAWed and
FSWed
-
2960 S.M. Chowdhury et al. / Materials Science and Engineering A
527 (2010) 29512961
Fig. 16. Typicstrain rate of 1
where k isstress, Mcg =resents a ligrain size iwith decreasize
and higwith that ofreported by[75].
3.5. Fractog
Fig. 16base metal,1103 s1fracture apdenoted du
sions [59]. Some inclusions were present on the fracture
surfaceof AZ31B rolled magnesium alloy. Similar fracture surface
char-acteristics were reported in [2,5]. As shown in Fig. 16(b),
somecleavage-like at facets in conjunction with dimples and
river
g coby t
lograg led
thesam
he aret a
at thAWeets whe sps an
clusmarkincausedcrystalnectinwell toFSWedfrom tby Limdefectthe
DSlike facfrom tsample
4. Conal SEM images showing the fracture surfaces of samples
tested at a103 s1. (a) Basemetal, (b) FSWedsample, and
(c)DSAWedsample.
Boltzmann constant, d is the grain size, is the ow lncg/ ln and
Mc = ln c/ ln . This equation rep-near correlation between the SRS,
ow stress andn the form of d1/2. It means that the SRS
increasedsing grain size. Since the base metal had a smaller
grainher ow stress, its SRS would be higher in comparisonthe welded
joints. This is in agreement with the resultsdel Valle and Ruano
[35] and Prasad and Armstrong
raphy
shows typical images of the fracture surfaces of theFSWed and
DSAWed samples tested at a strain rate of. It is seen from Fig.
16(a) that dimple-like elongatedpeared more apparent. This type of
fracture surfacectile fracture thatwas characterized by cup-like
depres-
1. The micrsmall elowhich wFSW resuSZ and TMDSAW, fsize werture
conintergran
2. While thFSWedsaprematuthe FSWethe causUTS in
thsamplesliteraturspeed lo
3. The straobservedhardeninThe straevaluatetimes hig
4. A higherwas obsmanydiswith strograin sizwelded
stallizatiowith p
5. The YS,rate incrstrain rabase mein compa
Acknowled
The autEngineeringwork of CeninvestigatioResearch Pruld be seen
in the FSWed sample. The rivermarkingwashe crack moving through the
grain along a number ofphic planes which formed a series of
plateaus and con-ges [59]. The fractographic observations
correspondedrelatively low percentage elongation of 4% in the
ples.While tensile fracture initiation could have startedea
between the weld nugget and the TMAZ as reportedl. [44], the crack
initiation occurred from the weldinge bottom surface in the present
study (Figs. 8 and 9). Ind sample, more dimples together with fewer
cleavage-ere observed, as shown in Fig. 16(c). The crack
initiatedecimen surface or near surface defects in the DSAWedd in
the base metal as well.
ions
ostructure of the as-received AZ31B-H24 consisted ofngated
grains with some MnAl containing inclusionsere still present in
different zones after welding. Thelted in recrystallized and
relatively small grains in theAZ, and partially recrystallized
grains in theHAZ. After
ully recrystallized grains with a relatively large graine
observed in the HAZ, and the divorced eutectic struc-taining
-Mg17Al12 particles in the interdendritic andular regions appeared
in the fusion zone.e YS was higher, the UTS and ductility were
lower in themples than in theDSAWedsamples. Thiswasdue to there
fracture caused by the presence of welding defects ind samples at
the bottom surface. The defects were also
e for the small net ow stress from yielding up to thee FSWed
samples. However, the strength of the FSWedwas still similar to or
higher than those reported in thee due to the smaller grain sizes
arising from the highw temperature FSW.in-hardening capacity of the
DSAWed samples wasto be twice that of the base metal, with the
strain-
g capacity of FSWed samples lying in-between them.in-hardening
exponents after both types of welding,d via three different
approaches, were all about twoher than those of base
metal.strain-hardening rate of the base metal at lower strainserved
due to smaller and pre-deformed grains
wherelocationshadbeengenerated in thebasemetal, couplednger
HallPetch contribution stemming from smalleres. At higher strains,
the strain-hardening rate of theamples became higher due to the
occurrence of recrys-n in the FSWed samples, and the larger grains
togetherarticles in the fusion zone in the DSAWed samples.UTS,
strain-hardening exponent and strain-hardeningeased slightly, and
ductility decreased with increasingte. Stronger strain-rate
sensitivity was observed in thetal due to the smaller grain size
and higher ow stressrison with the welded joints.
gements
hors would like to thank the Natural Sciences andResearch
Council of Canada (NSERC) and AUTO21 Net-ters of Excellence for
providing nancial support. Thisn involves part of CanadaChinaUSA
Collaborativeoject on the Magnesium Front End Research and
Devel-
-
S.M. Chowdhury et al. / Materials Science and Engineering A 527
(2010) 29512961 2961
opment (MFERD). The authors also thank General Motors
Researchand Development Center for the supply of test materials.
One ofthe authors (D.L. Chen) is grateful for the nancial support
by thePremiers Research Excellence Award (PREA), Canada
Foundationfor Innovation (CFI), and Ryerson Research Chair (RRC)
program.The authors would like to thank Q. Li, A. Machin, J.
Amankrah, D.Ostrom and R. Churaman for their assistance in the
experiments.The authors also thank Dr. S. Xu, Dr. K. Sadayappan,
Dr. J. Jackman,Professor N. Atalla, Professor S. Lambert, Professor
H. Jahed, Profes-sor Y.S. Yang, Professor B. Jordon, Dr. A.A. Luo,
Mr. R. Osborne, Dr.X.M. Su, and Mr. L. Zhang for helpful
discussion.
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Tensile properties and strain-hardening behavior of double-sided
arc welded and friction stir welded AZ31B magnesium
alloyIntroductionMaterials and experimental procedureResults and
discussionMicrostructureMicrohardnessTensile
propertiesStrain-hardening behaviorFractography
ConclusionsAcknowledgementsReferences