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FRACTURE MECHANICS AND MECHANICAL BEHAVIOUR IN
AA5083-H111 FRICTION STIR WELDS (FSW)
Oluwaseun John Dada
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DOI:
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S2468-2276(20)30003-X
https://doi.org/10.1016/j.sciaf.2020.e00265
SCIAF 265
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Scientific African
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23 January 2019
28 November 2019
13 January 2020
Please cite this article as: Oluwaseun John Dada , FRACTURE MECHANICS AND MECHANICAL BEHAVIOUR IN AA5083-H111 FRICTION STIR WELDS (FSW), Scientific African (2020), doi:
https://doi.org/10.1016/j.sciaf.2020.e00265
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Prof. Oluwaseun John Dada
FRACTURE MECHANICS AND MECHANICAL BEHAVIOUR IN
AA5083-H111 FRICTION STIR WELDS (FSW)
1,2
1
Oluwaseun John Dada
Signature EcoSystems Technologies,
DLF Center, 12A Pok Man Street, Tai Kok Tsui, Hong Kong.
2
Department of Metallurgical and Materials Engineering,
Federal University of Technology, Akure, Ondo State, Nigeria.
oluwaseundd@yahoo.com; john.dada@signecotech.org; ojdada@futa.edu.ng
ABSTRACT: Friction Stir Welding (FSW) is one of the fastest developing welding
technologies being implemented in the transport industry today.
Fracture surfaces of
AA5083-H111 specimens, which were friction-stir-welded using different parameters, were
studied with a scanning electron microscope (SEM). The changes in the microstructure of the
weld, the location and type of defects that accompanied each stirring process determined the
magnitude of reduction in strength and soundness of joints for flawed samples. The paths of
crack propagation and modes of fast fracture under static loads were also determined. The
sharp edges of the defects were stress concentrators which aided crack initiation.
Understanding of the effect of microstructure and defect on the failure of FSW joints will aid
optimization of the process variables, tool design, weld quality assurance and decision
making. On the long run, it will boost the confidence of medium-scale manufacturers in the
adoption and implementation of this fairly new technology.
KEYWORDS: Friction Stir Welds (FSW); Aluminium; Microstructure; Failure mechanism;
Defects.
Prof. Oluwaseun John Dada
Bio: Dr Oluwaseun John Dada is Professorial Chair of Signature EcoSystems
Technologies. He has worked over over the past 8 years on NanoScale materials, advanced
materials manufacturing and applications, and over 10 years in Metallurgy and Aluminium
Manufacturing.
I have achieved Electrolytic syntheses, characterizations, and applications of novel
graphene/nano-Carbons, resulting in Superior Electrical, Electronic, Thermal and Dielectric
characteristics; also resulting in the Highest Capacities when applied to Lead Acid and
Lithium-ion Batteries, resulting from higher Utilization of Electrode Materials, 1st to
demonstrate electrode precipitation and dissolution mechanisms in the gel-crystal
electroactive zone of the lead-acid battery; synthesized graphene also resulted in higher
viscoelastic mechanical behaviour in graphene-enhanced epoxy nanocomposites prepared by
in-situ reduction of GO by solvent evaporation and hot stirring. Synthesized highly scalable
and diversely applicable crystalline pristine graphene directly from graphite, with high basal
planarity using solid-state processes, or alternative colloidal sonication process. This has
resulted in the high-performance graphene anode for lithium-ion battery.
On my project on FSW, I have achieved parametric analysis, 1st SEM demonstration of
material flow pattern, and failure characteristics of Friction Stir Welds of industrial and
aeronautical alloys. Experience dates back to Alussuisse owned Nigerian Aluminium
Extrusion, a Laboratory Extrusion project, and AIRBUS FSW project at Wits University.
Dr Dada received PhD in Mechanical Engineering Nanotechnology Concentration at Hong
Kong University of Science and Technology, Hong Kong. Dr Oluwaseun John Dada is the
CEO/ Director of Research and Engineering at Signature EcoSystems Technologies, Hong
Kong. Dr. Dada was ORTg, Hong Kong PhD Fellow (Mechanical Engineering,
Nanotechnology Concentration), Senate Representative, Entrepreneurship Centre (EC)
Venture Entrepreneurship (VEN) Program Manager, and Engineering School Board member
at The Hong Kong University of Science and Technology (HKUST). He was on AMSEN
(Carnegie-RISE network) MSc scholarship at University of Witwatersrand, and won the
AMSEN conference 3rd presentation prize in Nairobi. Professor Dada was Lecturer, 1st
Class and Best Engineering School Graduate at the Federal University of Technology FUT,
Akure. I am currently an Chief\Editor for 3 materials science journal papers and published 3
books including: 1. Interfacial Performance of Applied Lead/Graphene; 2. Innovative
Friction Stir Welds of Aluminium Alloys Characterizations), and contributed to more than 60
scholarly journal publications.
Prof. Oluwaseun John Dada
1. INTRODUCTION
Friction Stir Welding (FSW) is one of the fastest developing welding technologies being
implemented in the transport industry today. It utilizes a robust rotating and non-consumable
tool comprising a shoulder and a probe, which is pushed into the adjoining surfaces of the
plates to stir, plasticise, forge and consolidate the material as it moves along the joint-line
(Figure 1) [1–5].
Figure 1. Schematic diagram of the FSW process and forces on the tool [6].
The application of FSW to aluminium structures, especially those made of 5083-H111
aluminium alloy, is particularly significant for body structures of road vehicles and marine
automobiles such as fast ferries and catamarans where higher speed, lower weight and energy
savings is essential. AA5083-H111 has been notable for these applications because of its high
strength,
formability,
weldability
and
corrosion
resistance.
Cost-effectiveness,
Prof. Oluwaseun John Dada
environmentally friendliness, enhanced superplasticity and superior mechanical properties
have made FSW preferred to other techniques like metal inert gas (MIG) and tungsten inert
gas welding (TIG). Also, FSW’s application scope has been extended beyond aluminium, to
other materials such as steels, titanium, and polymers.
FSW results in the creation of the following microstructural zones:
Weld nugget, which has undergone mechanical stirring, resulting in severe
plastic deformation and recrystallization caused by frictional heat generated by
the tool in the material workpiece [7, 8]. It has fine equiaxed grains in a banded
structure. The flow arm zone is at the top of the weld nugget, where the grains
are finer as a result of more intense stirring from the tool shoulder.
Thermo-mechanically affected zone (TMAZ) is located between the weld
nugget and HAZ. It has undergone some plastic strain without recrystallisation.
Visible elongation and rotation of the parent material grain structure occur in
this zone during welding [9].
Heat affected zone (HAZ), which is typical in all welds; the material has
undergone some thermal cycling, but no deformation and has a different
microstructure from the base material.
The difference in local properties of these microstructural zones and the presence of defects
makes mechanical behaviour and failure of FSWs complex. Rotational and transverse speeds
of the tool have been identified as key welding parameters that influence microstructure,
mechanical behaviour and defect formation. The industrial viability of FSWs depends on the
mechanical properties and the effect of defects. FSW AA5083-H111 of less than 5mm
thickness, which is less studied compared to other alloys in the aluminium series such as
2XXX, 6XXX and 7XXX, was used in this investigation. The mechanical and microstructural
Prof. Oluwaseun John Dada
characteristics of FSWs of 2.8 mm thick AA5083-H111 and their effect on tensile fracture
and failure mechanisms were ascertained.
2. EXPERIMENTAL PROCEDURE
AA5083-H111 plates of 360mm × 120mm × 2.8mm, supplied by Hulamin Ltd, were welded
using a cylindrical shoulder of 20mm diameter and tapered probe of 6mm diameter on a
converted CNC milling machine along the direction of rolling in the plates. The FSW tool
was made from H13 tool steel. The chemical composition of AA5083-H111 is given in Table
1. The surfaces of plates were cleaned with acetone to degrease them prior to welding.
Combinations of rotational speeds between 400, 500 and 600 rpm and welding speeds
between 50, 60 and 70 mm/min were used to produce six samples in Table 2.
Table 1. Chemical composition of AA5083-H111 (wt%).
Elements
Si
Fe
Cu
Mn
Mg
Cr
Zn
Ti
Al
wt %
0.12
0.33
0.03
0.51
4.39
0.08
0.01
0.02
Balance
Table 2. Specimens and their welding parameters.
FSW Samples
A
B
C
D
E
F
Rotational speeds (rpm)
400
400
500
630
630
630
Welding speeds (mm/min)
50
60
50
50
70
60
Pin loaded dog-bone specimens were prepared according to ASTM standard E8/E8M.
Transverse static tests were performed on an ESH hydraulic testing machine at a strain rate of
0.007mm/mm/s. Fracture surfaces were analysed using the Zeiss LEO 1525 Field Emission
Gun Scanning Electron Microscope (FEGSEM) in mainly secondary electron mode.
Prof. Oluwaseun John Dada
Metallographic preparation FEG-SEM according to Struer’s guide [10] and etching was done
using Keller’s reagent. A Leica microscope was used to observe the microstructure. The sizes
of defects in the flawed samples were measured across the weld. Microhardness profiles
across the welds were obtained using FM-700 Future-Tech Vickers hardness tester with a
300kgf load and 15s dwell time.
3. RESULTS AND DISCUSSION
3.1. Microstructural characterisation
All expected microstructural zones such as the weld nugget (WN), thermo-mechanically
affected zone (TMAZ) and heat-affected zone (HAZ) were present. There were no clear
boundaries between the zones at the retreating side of the weld tool (LHS of Figure 2),
although the size of the grains increased from the weld nugget outwards. On the advancing
side of the weld tool, there was a clear boundary between the nugget and the HAZ (RHS of
Figure 2). A higher shear rate in a narrower band, which resulted in a sharp change in
microstructure was reported to be responsible for the advancing weld nugget/TMAZ interface
[11]. A joint-line remnant, kissing bond, or “lazy-S”, which appeared as a thin dark line more
towards the advancing side, was found in the weld nugget. It has been identified to comprise
amorphous clusters of alumina both within grains and at grain boundaries [12], because it
comprises the surface of aluminium which is nature of surface alumina, which are grains
attacked by atmospheric oxygen. The finest grains were found at the crown of the weld in the
flow arm zone due to additional torque and frictional heating from the tool shoulder. Tunnel
defects were found at the base of the weld nugget/TMAZ interface on the advancing side
(Figure 3) in samples B-F and their sizes are shown in Table 3. Pinholes are in samples D-F
(Figure 4). Root flaws were also found at the root of all the welds.
Prof. Oluwaseun John Dada
Figure 2. Light microscope image through the FSW AA5083-H111, showing the
microstructural zones in Sample A.
Figure 3. Light microscope image through the FSW AA5083-H111, showing tunnel defect at
the base of the weld nugget/TMAZ interface (Sample A).
Prof. Oluwaseun John Dada
Figure 4. The microstructure of Samples B-F showing positions of tunnel defects and
pinholes.
Table 3. Average sizes of tunnel defects in Sample B-F.
Specimen
Tunnel Size (µm2)
B
C
D
E
F
17.8
29.3
63.3
35.8
37.7
3.2. Micro-hardness profiles
There was an increase in hardness in the weld zone, more towards the retreating side, due to
grain refinement and better weld consolidation (Figure 5). There was no discernible
Prof. Oluwaseun John Dada
relationship between weld parameters and hardness profiles. The hardnesses towards the
retreating side are higher, while hardnesses are lower towards the advancing side due to
greater frictional heat and grain enlargement, while there is reduced frictional heat in the
retreating side. The grain size differences follow hall-perch relationship resulting in the slope
in hardness profile inverse to the direction of tool flow.
Figure 5. Microhardness profiles of Samples A-F across the weld nugget.
3.3. Static testing and failure
The average static properties of the welds are given in Table 4. The fracture paths in Samples
A-E were similar. The welds were expected to fail in the base material since the hardnesses
there were lower than in the weld. The presence of tunnel defects changed the fracture
location under tensile loading. In Samples A-E, fast fracture initiated due to the stress
concentration from the sharp corners of the triangular weld tunnel. Fast fracture propagated in
Prof. Oluwaseun John Dada
two directions: first, upwards at 45°, slightly bending through the TMAZ and HAZ, and
finally vertically upward in the flow arm zone. Secondly, fracture initiated from the tool
marks and propagated vertically downwards from the right angle of the tunnel through the
HAZ. Sample F fractured almost along the joint line, despite the tunnel defect at the base of
the weld nugget/TMAZ being larger than the pin-hole at the centre of the weld nugget
(Figures 5-6). This showed that the pinholes in the nugget determined the fracture location,
and were more deleterious to the ductility and UTS (Tables 3-4, Figures 4,6-7). Sharp corners
of the pinholes were stress concentrators for fast fracture within the weld nugget. Liu et al.
[13] deduced that the tensile failure of FSWs near the weld nugget/TMAZ interface was due
to significant differences in their internal structures which made the interface a weaker region.
This was not found in the current study. The fracture path was bent away from this interface,
despite the presence of tunnels defect with sharp corners in Specimen A, as shown in Figure
3. There was no evidence of weakness at the weld nugget/TMAZ interface, despite the
microstructural differences. Fracture of welds at locations where microhardness was higher,
as in this case, was an indication of the presence of defects. In this case, the fracture path
approached the areas of lowest hardnesses in the weld, in agreement with Liu et al. [14],
Biallas et al. [15] and Okamura et al. [16].
Prof. Oluwaseun John Dada
Table 4. Static testing results of the welds and parent material.
SPECIMEN
PARENT
A
B
C
D
E
F
UTS (MPA)
295
265
246
230
245
172
300
0.2% PROOF STRESS
140
135
139
138
139
135
145
STRAIN-TO-FAILURE
0.3
0.12
0.08
0.08
0.07
0.02
0.23
WELD EFFICIENCY (%)
98.3
88.3
82.0
76.7
81.0
57.0
100
MATERIAL
PROPERTIES
(MPA)
Figure 6. Pictures of fractured tensile specimens A-F of Al Friction Stir Welded (FSWs)
with different parameters.
Prof. Oluwaseun John Dada
(a)
(b)
Figure 7. Schematic diagrams of the fracture path for (a) Samples A, and (b) Sample F.
3.4. Fracture surface analysis
Weld tunnels at the roots in Samples A-E
Prof. Oluwaseun John Dada
The overall fracture surface (Sample A, inverted) is shown in Figure 8. Features such as the
root of the weld (root is at the top, Figure 9), tool marks (e.g. 1, Figure 9), tunnel defect
(indicated by the arrow), part of the onion ring structure (e.g. 4, Figure 9) and the region of
fast fracture (e.g. 5, Figure 9) were observed on the fracture surface towards the lower part of
the weld (which is inverted). Tool marks and bands of unconsolidated material (lap pattern)
that made up the weld tunnels resulted from the reduced bonding between the onion structure
and the weld root. The tool mark was an incomplete semi-circular feature created by the
motion of the tool on the material at the root of the weld, while the onion ring was created by
successive deposition of material, rotated and pushed down around the tool pin, but which
failed to be consolidated with the root of the weld at the advancing side of the tool. Similar
features were found by Gratecap et al. [17] and Chen et al. [18]. Poor consolidation that
produced the weld tunnel resulted from the reduction in fluidity and flow of material to the
end of the pin on the advancing side [18]. Also, complete weld consolidation did not take
place, because the vertical forging pressure was too low. Tilting the tool, which could have
compensated for the low vertical pressure and improved downward flow of material [18], as
the tool translated along the joint, could not be done on the current CNC milling machine. The
unevenness in the depth of the unconsolidated material in Specimen 1 indicated that the tool
lost vertical pressure occasionally while rotating and translating. This showed that position
controlled welding was not optimal, if defects were to be avoided. Greater vertical pressure in
force controlled welding mode should be necessary to avoid tunnel defects.
Prof. Oluwaseun John Dada
(a)
(b)
Figure 8.(a) SEM-SE image of the overall fracture surface of Specimen 1 along the weld, (b)
SEM-SE image of the lower part of the fracture surface taken along the weld in Sample A.
Prof. Oluwaseun John Dada
The river marks diverging downwards on the surface of the tool marks (Figure 10a) and
upwards at the ending of the unconsolidated onion rings (Figure 10b), indicated that they
were the primary regions of crack initiation in the weld tunnel.
(a)
(b)
Figure 9. SEM-SE image (Sample A) of (a) one of the tool marks with arrows showing lines
indicating the initiation of fast fracture at the root of the weld, and (b) the unconsolidated
materials that formed the hypotenuse of the triangle weld tunnel with arrows showing the
direction of fracture in the TMAZ.
Failure mechanisms and stress states in Sample A
The weld root (below the tunnel) was characterized by a large number of small parabolic
dimples with large lips on the cleavages, showing a failure by mixed micro-void coalescence
and quasi-cleavage (e.g. 2-3, Figure 8). Closer to the base of the weld, there were areas both
with and without dimples, indicating the quasi-cleavage mechanism. This is consistent with
the results obtained by Elangovan and Balasubramanian [19], who identified poor weld
consolidation in the root was being responsible. The grains in this region were fine, and high
stresses were needed for failure, due to the increased hardness. Quasi-cleavage fractures are
brought about by crack initiation by microvoid coalescence and propagation by tearing of
cleavage facets ahead of the moving crack front [14].
Prof. Oluwaseun John Dada
Above the unconsolidated material region (in the TMAZ/HAZ, Figure 10a), there were more
equiaxed and small parabolic dimples. In the middle, there were larger parabolic dimples than
those near the root, which indicated larger grains, larger precipitates and higher shear
ductility. A large number of larger dimples, which sometimes had precipitated, indicated that
the precipitates were active originators of microvoids under applied stresses.
In the flow arm zone at the crown (Figure 10b), cleavages carrying a few tiny dimples were
found, which indicated failure by quasi-cleavage mode. The precipitates that caused dimples
were regions of local crack initiation, but the matrix failed in a brittle manner. Smaller
dimples were from the smaller precipitates which were reduced by the extra torque from the
tool shoulder. The lower ductility was due to the reduction of grain size by work-hardening
induced by tool shoulder. There was a transition from the ductile tear to quasi-cleavage from
the middle to the top of the weld fracture surface [20]. This is consistent with Zadpoor et al.
[20] where brittle regions were formed in the top of the sheet which underwent severe plastic
deformation when in contact with the rotating tool shoulder. Cleavage-like fracture is not
common in FCC metals where there are many active slip systems [21]. Zadpoor et al. [20]
suggested that high rotational speed could have resulted in the non-optimal grain structure and
strain rate that weakened the recrystallized microstructure, and the mixture of ductile and
brittle fracture could be the result of the blocking of dislocation movement by the precipitates.
Quasi-cleavage fracture is different from cleavage, which is a low energy fracture mechanism
that moves through well-defined low index crystallographic planes [21]. Quasi-cleavage is
used to describe cleavage that has signs of plastic deformation, although limited [20, 21].
Prof. Oluwaseun John Dada
(a)
(b)
Figure 10. SEM-SE image (Sample A) of (a) the middle of fracture surface along the weld
HAZ showing dimples and precipitates, and (b) the crown (flow arm zone) of the weld
(inverted) showing the fine dimples on cleavages.
The shape of dimples on fracture surfaces indicates the responsible stress states [22].
Equiaxed dimples are associated with tensile stresses (Mode I), while shear dimples are from
shear or tear stress systems (Mode II and III) and point in the direction of shear [14]. The lips
of the shear and elongated dimples at the middle and near the top of the fracture surface
pointed upwards (in the direction of the crown of the weld), while those at the root pointed
downwards, and this was not associated with normal tensile loading. The presence of tunnels
created a response of both tensile and tearing stresses when the tensile loads were applied.
The sizes and shapes of the dimples were directly proportional to the sizes of the grains and
precipitates in the different zones and ductility [23]. Zones, where the precipitates were large,
had larger and deeper dimples (e.g. at the middle of the fracture surface). Where the grains
were larger, the lips of the dimples were wider, indicating that the material was more ductile.
Fine dimples, which are associated with work-hardening [23], were found in the flow arm
zone and in the root, which indicates that more plastic deformation took place in these two
Prof. Oluwaseun John Dada
zones. The different fracture mechanisms observed in the SEM showed that the cracks
propagated through different microstructural zones in Samples A-E.
(a)
(b)
Figure 11. Schematic diagrams fracture path with arrows indicating dominant stress systems
responsible for fracture in (a) Specimen A, and (b) Specimen F.
Prof. Oluwaseun John Dada
Failure mechanisms and stress states in Sample F
The overall fracture surface is shown in Figure 12a. The lips of the dimples at the upper part
of the fracture surface in the flow arm zone were elongated downwards (Figure 12b), opposite
to the direction of failure. Where the direction of dimple elongation is opposite to that of fast
fracture, as in this case, it is an indication of ductile shear fracture [22]. Towards the middle to
the root of the fracture surface, the dimples were elongated downwards in the direction of the
root, and in the direction of fast fracture in the root of the weld. This indicates that the
mechanism of fast fracture here was a ductile tear. At the top in the flow arm zone and at the
root, there was a gradual transition from ductile shear and tear to quasi-static fracture. The
middle of the fracture surface contained shallow dimples which showed that the precipitates
involved in dimple rupture were finer and the linear pattern on the surface was from the weld
nugget bands, indicating that fracture occurred across the onion ring structure (Figure 12c).
The failure mechanism here was ductile. Smaller and shallower dimples showed that failure in
the weld nugget was less ductile than in Specimens A-E, which failed in the TMAZ/HAZ.
When the tensile load was applied on Specimen 6, the presence of defects created a system of
shear stresses at the upper part of the weld and a tear system at the lower part, which resulted
in the fast fracture through the weld nugget (Figure 12b).
Prof. Oluwaseun John Dada
(a)
(b)
(c)
Figure 12. SEM-SE image (Sample F) showing (a) Overall fracture surface, (b) weld crown
(flow-arm zone) showing shear dimples on cleavages, and (c) Middle of the weld (in the weld
nugget) showing a linear pattern from onion rings.
3.5. Effect of welding parameters on defect formation
The defect size had no relationship with the advance of the tool per revolution. Increased
welding speed led to decreased vertical force, thereby reducing weld consolidation and
causing widening of the tunnel defects. Leonard and Lockyer [24] stated that the material is
Prof. Oluwaseun John Dada
given less work per unit of the weld length when welding speed is increased at a constant
rotational speed. Thus, materials are cooler and more difficult to consolidate at higher welding
speeds. Also, increased rotational speed at constant welding power of the CNC milling
machine led to the formation of a pinhole at the advancing side and another pinhole in the
nugget due to poor stirring (caused by sliding) in Specimens E and F. Increased rotational and
welding speed has been linked to higher heat input and more intense stirring [25]. However,
this is only obtainable when the welding power of the FSW machine is high enough to
accommodate the increases, by ensuring other parameters such as the downward and
transverse forces are not significantly reduced. Also, increased rotational speed to an
optimum, without a constant or reducing downward force (pressure) leads to a reduction in
the frictional forces that provides intense heating and stirring [26]. Depending on tool
geometry, increased rotational speed causes sliding of the material on the tool. In this case,
the tool used was a disadvantage, because it was unthreaded. The absence of a tilt angle on the
tool was a major disadvantage. Tilting the tool (2-3°) causes the rear of the shoulder and pin
to be lower than the front, which aids the downward flow of the material around the rotating
pin. Also, the rear of the shoulder forges the material as the tool moves along the joint line
[27]. A deeper plunge depth in position control would have ensured greater tool penetration,
reducing or avoiding the root flaw and also achieving forging within the weld. Root flaws
could have been avoided or reduced with increased rotational speed, which causes increased
stirring at the root and improved weld penetration at constant downward pressure [25].
3.6. Effect of defects on weld integrity
Hardness results showed that AA5083-H111 FS welds were stronger than the base material,
but tunnel defects reduced tensile properties. Jamshidi-Aval et al. [28] and Squillace et al.
[29] found tunnel defects in joints produced using conical tools; reduced static mechanical
Prof. Oluwaseun John Dada
properties [28], or not [29]. In both cases, the sizes of the tunnels across the welds were not
measured. In this study, the size of the tunnel defects affected the static properties. There is an
optimum size of tunnel defect below which the overall strength of the weld was unaffected
[29, 30]. In Sample A where the pin-hole (discontinuous tunnel) was 2µm2 across, the tensile
properties were similar to that of the base material (Table 4). UTS, strain-to-failure and joint
efficiency reduced with increasing tunnel size, while yield strength was almost unaffected
(Figure 13), in agreement with Chen et al. [18]. The closeness of defects to the centre of the
weld (in Sample F, FSW at RS 630rpm; WS 60mm/min) was more deleterious to UTS and
ductility. The presence of root flaws, lazy-S’s (A-F) and pinhole at the upper part of the
advancing side (E-F) had no noticeable effect on fracture or overall strength of FSW
AA5083-H111.
(a)
(b)
(c)
Figure 13. Effect of tunnel defect size on UTS, proof stress, strain-to-failure and weld
efficiency for FSWs AA5083-H111 (Samples A-E).
Prof. Oluwaseun John Dada
4. CONCLUSIONS
The difference in failure mechanisms in the welds were due to different microstructural zones
which differed in grain structure. Tunnel defects were caused by a lack of consolidation of the
onion ring structure at the lower part of the advancing side of the tool. This resulted from
differences in the rate of material flow away and from the bottom of the advancing side. This
study confirmed that fact that low vertical pressure and non-tilting of the tool are the overall
causes of tunnels. Tool penetration was inadequate as the height of the root flaw was
significantly high. The degree of severity of weld defects: Tunnels, pinholes and lazy-S’s
determined the strength. The presence of a pinhole at the weld nugget significantly reduced
tensile properties and caused a fracture, even when the tunnel at the weld nugget/TMAZ
interface was larger. Thus, the tunnels were more severe than the pinholes, while the lazy-S’s
had no effect on quasi-static fracture. Tensile fracture in regions of the weld where hardness
values were higher indicated the likely presence of weld defects such as tunnels and pinholes.
The presence of sharp transitions as a result of differences in microstructure between the weld
nugget and TMAZ at the advancing side had no effect on hardness or fracture path. The weld
nugget/TMAZ interface was not a weak region in FSW AA5083-H111, as previously
reported, because cracks propagated away from tunnels located along which initiated the
failure. At constant welding power, increased rotational and welding speeds resulted in a
reduction in vertical pressure that caused increased size, number and severity of weld defects.
In all the specimens, the presence of defects transformed the applied tensile loads into tear
and shear stresses. Also, there was a transition from ductile shear and tear in the center of the
fracture surface to mixed modes and quasi-cleavage at the top and root of the weld.
Prof. Oluwaseun John Dada
ACKNOWLEDGEMENT
The author wishes to acknowledge Hulamin South Africa for supplying the materials.
Professors Lesley Cornish, Claudia Polese, Tony Paterson of the University of the
Witwatersrand, South Africa, and George F. Van der Voort for their contributions.
REFERENCES
1.
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