Engineering, 2013, 5, 567-576
http://dx.doi.org/10.4236/eng.2013.57069 Published Online July 2013 (http://www.scirp.org/journal/eng)
Pulse-Impact on Microstructure of
Liquid-Phase-Pulse-Impact Diffusion Welded Joints of
Particle Reinforcement Aluminum Matrix Composites at
Various Temperatur es
Kelvii Wei Guo
Department of Mechanical and Biomedical Engineering, City University of Hong Kong, Hong Kong, China
Email: kelviiguo@yahoo.com
Received February 28, 2011; revised January 1, 2013; accepted January 8, 2013
Copyright © 2013 Kelvii Wei Guo. This is an open access article distributed under the Creative Commons Attribution License,
which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.
ABSTRACT
Investigation was to study the influence of pulse-impact on microstructure of Liquid-Phase-Pulse-Impact Diffusion
Welding (LPPIDW) welded joints of aluminum matrix composite SiCp/A356, SiCp/6061Al, Al2O3p/6061Al. Results
showed that under the effect of pulse-impact: 1) the interface state between reinforcement particle (SiC, Al2O3) and ma-
trix was prominently; 2) the initial pernicious contact-state of reinforcement particles was changed from reinforcement
(SiC, Al2O3)/reinforcement (SiC, Al2O3) to reinforcement (SiC, Al2O3)/matrix/ reinforcement (SiC, Al2O3); 3) the den-
sity of dislocation in the matrix neighboring to and away from the interface in the matrix was higher than its parent
composite; and 4) the intensively mutual entwisting of dislocation was occurred. Studies illustrated that: 1) deformation
was mainly occurred in the matrix grain; and 2) under the effect of pulse-impact, the matrices around reinforcement
(SiC, Al2O3) particles engendered intensive aberration offered a high density nucleus area for matrix crystal, which was
in favor of forming nano-grains and improved the properties of the successfully welded composite joints.
Keywords: Aluminum Matrix Composite; Particle Reinforcement; Pulse-Impact; Microstructure; Diffusion Welding
1. Introduction
The high specific strength, good wear-ability and corro-
sion resistance of Aluminum Matrix Composites (AMCs)
attract substantial industrial applications. Typically, AMCs
are currently used widely in automobile and aerospace
industries, structural components, and heat resistant-wea-
rable parts in engines, etc. [1-5]. The particles of rein-
forcement elements in AMCs may be either in form of
particulates or as short fibers, whiskers and so forth [5,6].
These discontinuous natures create several problems to
their joining techniques for acquiring their high strength
and good quality weld-joints. Typical quality problems
of those welding techniques currently available for join-
ing AMCs [7-15] are as elaborated below.
1) The distribution of particulate reinforcements in the
weld
As properties of welded joints are usually influenced
directly by the distribution of particulate reinforcements
in the weld, their uniform distribution in the weld is likely
to give tensile strength higher than 70% - 80% of the
parent AMCs. Conglomeration distribution or absence
(viz. no-reinforcements-zone) of the particulate reinfor-
cements in the weld generally degrades markedly the
joint properties and subsequently resulted in the failure
of welding.
2) The interface between the particulate reinforce-
ments and aluminum matrix
High welding temperature in the fusion welding me-
thods (typically: TIG, laser welding, electron beam etc.)
is likely to yield pernicious Al4C3 phase in the interface.
Long welding time (e.g. several days in certain occasions)
in the solid-state welding methods (such as diffusion
welding) normally leads to a) low efficiency and b) for-
mation of harmful and brittle intermetallic compounds in
the interface.
To alleviate these problems incurred by the available
welding processes for welding AMCs, a liquid-phase-
pulse-impact diffusion welding (LPPIDW) technique has
been developed [16-18]. This paper aims at providing
C
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568
some specifically studies the influence of pulse-impact
on the microstructures of welded joints. Analysis by
means of scanning electron microscope (SEM), trans-
mission electron microscope (TEM) and X-Ray Diffrac-
tion (XRD) allows the micro-viewpoint of the effect of
pulse-impact on LPPIDW to be explored in more detail.
2. Experimental Material and Process
2.1. Specimens
Stir-cast SiCp/A356, P/M SiCp/6061Al and Al2O3p/
6061Al aluminum matrix composite, reinforced with
20%, 15% volume fraction SiC, Al2O3 particulate of 12
m, 5 m mean size, were illustrated in Figures 1-3.
Figure 1. Microstructure of aluminum matrix composite
SiCp/A356.
Figure 2. Microstructure of aluminum matrix composite
SiCp/6061Al.
Figure 3. Microstructure of aluminum matrix composite
Al2O3p/6061Al.
2.2. Experiment
The quench-hardened layer and oxides, as induced by
wire-cut process, on the surfaces of aluminum matrix
composite specimens were removed by careful polishing
using 400 # grinding paper. The polished specimens were
then properly cleaned by acetone and pure ethyl alcohol
so as to remove any contaminants off its surfaces. A DSI
Gleeble®-1500D thermal/mechanical simulator with a 4
× 101 Pa vacuum chamber was subsequently used to
perform the welding.
The microstructures and the interface between the re-
inforcement particle and the matrix of the welded joints
were analyzed by SEM and TEM.
2.3. Operation of LPPIDW
Figure 4 illustrated a typical temperature and welding
time cycle of a LPPIDW. It basically involved with: 1)
an initially rapid increase of weld specimens, within a
time of ta, to an optimal temperature Ta at which heat was
preserved constantly at Ta for a period of (tb ta), 2) at
time tc, a quick application of pulse impact to compress
the welding specimens so as to accomplish an anticipated
deformation δ within a glimpse of 104 - 102 s, whilst
the heat preservation was still maintained at the opera-
tional temperature Ta; and 3) a period of natural cooling
to room temperature after time tb.
3. Results and Discussion
3.1. Microstructure of Welded Joint
Figure 5 showed the microstructures of welded joints of
SiCp/A356 at various temperatures with VI = 560 mm/s, tI
= 104 - 102 s, t = 30 s, P0 = 5 MPa, δ = 1 mm, where VI
was velocity of pulse-impact, tI was the impacting time, t
was holding time for heat preservation, P0 was holding
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Figure 4. Schematic diagram of liquid-phase-pulse-impact
diffusion welding.
400 μm
(a)
400
μ
m
(b)
400 μm
(c)
Figure 5. Microstructures of welded joints of SiCp/A356 at
various temperatures by LPPIDW; (a) 563˚C; (b) 565˚C; (c)
570˚C.
pressure during the welding, δ was the horizontal defor-
mation. It elucidated that when the welding temperature
was 563˚C, under the effect of pulse-impact, the liquid
phase matrix alloy wasn’t formed enough to wet the par-
ticle reinforcements. In addition, at this temperature, the
diffusion capability of the atoms within the matrix was
relatively low. As a result, the welding interface between
two specimens could be observed obviously as shown in
Figure 5(a) and followed by the unsuitable strength
(about 118 MPa). Moreover, because of lower welding
temperature, the area of the formed solid-liquid phase
was smaller, which led to some streamlines scattered in
the matrix (Figure 5(a)) after the pulse-impact acting on
the substrates. When the temperature reached 565˚C, the
rate of the atom diffusion in the joint region within the
matrix was accelerated (Figure 5(b)). At the same time,
more liquid phase matrix alloy was formed to wet rein-
forcements (SiC). Therefore, the interface state of rein-
forcement and reinforcement was improved and the re-
inforcements were distributed uniformly to some extent.
Also, the streamlines scattered in the matrix were disap-
peared, and the tensile strength of welded joints was
about 134 MPa higher than that of 563˚C. When the
temperature is up to 570˚C, the formed liquid phase ma-
trix alloy was enough and suitable for wetting reinforce-
ments effectively, and the rate of the atom diffusion was
more active. As a result, for reinforcements the welding
mode in the joint region changed from reinforcement-
reinforcement to reinforcement-matrix-reinforcement.
Con- sequently, the joint was welded successfully (Fig-
ure 5(c)). The average strength of 179 MPa for the
welded joints produced at welding temperature of 570˚C
was about 74.6% of the 240 MPa for the strength of par-
ent aluminum matrix composite.
The relevant fractographs were shown in Figure 6. It
illustrated that when the welding temperature was 563˚C,
the initial morphology of substrate could be detected
obviously, and some sporadic welded locations appeared
together with some rather densely scattering bare rein-
forcement particles as shown in Figure 6(a). With the
welding temperature was increased to 565˚C, more liquid
phase was formed. Under the effect of pulse-impact,
some wet locations in the joint had been excellently
welded and the aggregated solid reinforcement particles
were improved. However, the bare reinforcement parti-
cles were still distributed on the fractographic surface. It
indicated that substrates did not weld ideally the pieces
together and it consequently resulted in a low strength
joint (Figure 6(b)). Figure 6(c) showed the fractograph
of welded joint at 570˚C. It illustrated that the fracture
was dimple fracture. Moreover, SEM of the fracture sur-
face showed some reinforcement particles (SiC) in the
dimple. In order to confirm the state of these reinforce-
ment particles, particles itself and matrix neighboring to
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570
(a) (b)
(c) (d)
Figure 6. Fractographs of SiCp/A356 at various temperatures; (a) 563˚C; (b) 565˚C; (c) 570˚C; (d) 575˚C.
these particles were analyzed by energy dispersive X-ray
analysis (EDX) respectively. Result was shown in Fig-
ure 7. It indicated that reinforcement particles (SiC) were
wet by matrix alloy successfully suggesting that the re-
inforcement particles had been perfectly wet and the
composite structure of reinforcement/reinforcement had
been changed to the state of reinforcement/matrix/rein-
forcement.
As welding temperature increasing to 575˚C, it led to
more and more liquid phase matrix alloy distributing in
the welded interface, meanwhile, more liquid phase ma-
trix alloy reduced the effect of impact on the interface of
the welded joints, subsequently the application of tran-
sient pulse-impacting would cause the relative sliding of
the weldpieces that jeopardized ultimately the formation
of proper joint as shown in Figure 6(d). It demonstrated
that results of fractographs were agreed with the corre-
sponding microstructures well.
Figure 7. Energy dispersive X-ray analysis of the fracture
surface of SiCp/A356.
6061Al at various welding temperatures were shown in
Figures 8-11.
It showed that the microstructure evolutions and its
corresponding fracture surfaces under the effect of pulse-
impact are similar to that of SiCp/A356.
Figures 8(a) and 10(a) showed when the welding tem-
perature was too low to form enough liquid phase matrix
The relevant results of SiCp/6061Al and Al2O3p/
K. W. GUO 571
(a) (b) (c)
Figure 8. SEM micrographs of SiCp/6061Al we ld ed joints at various welding temperatures; (a) 620˚C; (b) 623˚C; (c) 625˚C.
(a) (b) (c)
Figure 9. Fractographs of SiCp/6061Al at various temperatures; (a) 620˚C; (b) 623˚C; (c) 625˚C.
(a) (b) (c)
Figure 10. SEM micrographs of Al2O3p/6061Al welded joints at various welding temperatures; (a) 641˚C; (b) 644˚C; (c)
647˚C.
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(a) (b) (c)
Figure 11. Fractographs of Al2O3p/6061Al at various temperatures; (a) 641˚C; (b) 644˚C; (c) 647˚C.
alloy to wet the reinforcement particles, and the diffusion
capability of the atoms within the matrix was relatively
low. Therefore, the indistinct welding interface between
two specimens could be observed resulted in low tensile
strength (about 240 MPa for SiCp/6061Al and 270 MPa
for Al2O3p/6061Al). When the temperature is higher
(623˚C for SiCp/6061Al and 644˚C for Al2O3p/6061Al),
the liquid phase matrix alloy was formed enough to wet
the reinforcement particles (SiC), together with higher
rate of the atom diffusion in the joint region (Figures 8(b)
and 10(b)). Consequently, the joints could be welded
successfully with the average strength of 260 MPa for
SiCp/6061Al (about 72.2% of the 360 MPa for the
strength of parent aluminum matrix composite) and 282
MPa for Al2O3p/6061Al (about 70.5% of the 400 MPa for
the strength of parent aluminum matrix composite). As
welding temperature increasing further (such as 625˚C
for SiCp/6061Al and 647˚C for Al2O3p/6061Al), more
and more liquid phase matrix alloy would be distributed
in the welded interface, at the same time, more liquid
phase matrix alloy reduced the effect of impact on the
interface of the welded joints, subsequently prompted for
the descending of the joint strength (Figures 8(c) and
10(c)).
Moreover, according to the fractures of welded joints
at various temperatures shown in Figures 9 and 11, it
showed that it agreed with Figures 8 and 10 very well,
and the fractures were all dimple fractures with some
reinforcement particles (SiC, Al2O3) in the dimple. Also,
the results of SiCp/6061Al were better than that of
Al2O3p/6061Al due to a mild interfacial reaction between
the reinforcement and matrix, which released the thermal
mismatch stress to an acceptable extent between the re-
inforcement and matrix to allow load transfer from the
matrix to reinforcement successfully. As a result, it had
advantageous effect of improving the strength of welded
joints further [18].
Based on microstructures of the welded joints with the
optimal parameters (i.e., p
SiC/ A356 = 570˚C, p
SiC/6061Al
= 623˚C, 23p
Al O = 644˚C, VI = 560 mm/s, tI = 102
104 s, δ = 1 mm, t = 30 s, P0 = 5 MPa) and its corre-
sponding fracture surfaces as shown in Figures 5, 6, 8-11,
the welded joint displayed with uniformly distributing
reinforcement particles and microstructure almost similar
to that of its parent composite (Figures 1-3). SEM of the
fracture surface showed that the reinforcement particles
had been perfectly wet and the composite structure of
reinforcement/reinforcement had been changed to the
state of reinforcement/matrix/reinforcement. XRD pat-
tern of the fracture surfaces (Figure 12) did not illustrate
the existence of any harmful phase or brittle phase of
Al4C3. This suggested the effective interface transfers
between reinforcement particles and matrix in the welded
joint that subsequently provided favorable welding
strength [16-18].
T T
T
3.2. Distribution of Dislocation in the
Welded Joint
The distribution of dislocation in the matrix neighboring
to the interface of the welded joint by LPPIDW in com-
parison with its parent composite was shown in Figure.
13. The clearly distinctive interface between reinforce-
ment particle and matrix indicated that the integration
between the reinforcement particle and matrix was
prominent. The effect of pulse-impact subsequently led
to dislocation in the matrix lattices and showed sign of
mutually entwisting to give higher welded strength.
Comparatively, its dislocation distribution in the matrix
neighboring to the interface was relatively denser than
that in its parent composite (cf. Figures 13(a) and (b)).
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(a)
(b)
(c)
Figure 12. XRD pattern of the fracture surfaces; (a) SiCp/A356; (b) SiCp/6061Al; (c) Al2O3p/6061Al.
Similarly, the density of dislocation and dislocation en-
twisting in the matrix away from the welded interface
was also higher than that of its parent composite (cf.
Figures 14(a) and (b)). Such favorable characteristics
ultimately gave relatively superior strength of the welded
joint to that of conventional diffusion welding [16-18].
3.3. Formation of Nano-Grains in the Weld
SEM micrograph (Figure 15) of a weld by LPPIDW
displayed some newly-formed nano-grains in the lattices
of the joint. These nano-grains would seat in the intersti-
ces of crystal lattices and create new grain boundary in
hindering the movement of neighbouring grains and
subsequently improved obviously the properties of the
welded joints. The formation of new nano-grains was the
advantageous effect of pulse-impact in LPPIDW. In ad-
dition, XRD pattern of the fracture surface (Figure 12)
did not illustrate the existence of any harmful phase or
brittle phase of Al4C3. This suggested the effective inter-
face transfers between reinforcement particles and matrix
in the welded joint that subsequently provided favorable
welding strength [16-18].
4. Conclusions
Results of this study on the microstructures of welded
joints of particle reinforcement aluminum matrix com-
posites (SiCp/A356, SiCp/6061Al, Al2O3p/6061Al) using
liquid-phase-pulse-impact diffusion welding process show
that:
1) Pulse-impact in liquid-phase-pulse-impact diffusion
welding in joining particle reinforcement aluminum ma-
trix composites (SiCp/A356, SiCp/6061Al, Al2O3p/6061Al)
resulted in higher density of dislocation in the matrix
neighboring to and away from the interface than their
parent composite. Simultaneously, the dislocation en-
twisted mutually and intensively in the welded joint pro-
pitious to improve the strength of welded joints.
2) There was distinctly clear interface between reinfor-
cement particle and matrix. It overcame some diffusion
problems normally encountered in conventional diffusion
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574
SiC
Matrix
0.25μm
SiC
Matrix
0.25μm
SiC
Matrix
0.25μm
SiC
Matrix
Matrix
SiC
(a) welded joint (b) parent composite (a) welded joint (b) parent composite
(A) (B)
0.25μm
Matrix
Al
2
O
3
0.25μm
Matrix
Al
2
O
3
Al
O
(a) welded joint (b) parent composite
(C)
Figure 13. Distribution of dislocation in the matrix neighbori ng to the interface of the welded joint and parent composite re-
spectively; (A) SiCp/A356; (B) SiCp/6061Al; (C) Al2O3p/6061Al.
Matrix Matrix
0.25μm 0.25μm
0.25μm
Matrix
0.25μm
Matrix
(a) welded joint (b) parent composite (a) welded joint (b) parent composite
(A) (B)
0.25μm 0.25μm
Matrix Matrix
(a) welded joint (b) parent composite
(C)
Figure 14. Distribution of dislocation in the matrix away from the interface of the welded joint and parent composite respec-
tively; (A) SiCp/A356; (B) SiCp/6061Al; (C) Al2O3p/6061Al.
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K. W. GUO 575
0.25μ
0.25μm
0.25μm
(a) (b) (c)
Figure 15. Nano-grains formed in the weld of particle reinforcement aluminum matrix composites during the LPPIDW; (a)
SiCp/A356; (b) SiCp/6061Al; (c) Al2O3p/6061Al.
welding, and prevented the formation of harmful micro-
structure or brittle phase in the welded joint.
3) The joint by LPPIDW process would form nano-
grains. The newly-formed nano-grains would improve
the properties of welded joints resulted in higher tensile
strength.
5. Acknowledgements
This work is supported by City University of Hong Kong
Strategic Research Grant (SRG) No. 7002582.
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