Materials Characterization 157 (2019) 109879

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Materials Characterization
journal homepage: www.elsevier.com/locate/matchar

Interfacial reaction behavior and mechanical properties of pure aluminumand magnesium alloy dissimilar materials fabricated by hot press and heat
treatmentT

Junko Umedaa, , Katsuyoshi Kondoha, Hiroyuki Sannomiyaa, Tachai Luangvaranuntb,
Makoto Takahashia, Hiroshi Nishikawaa⁎abJoining and Welding Research Institute, Osaka University, Japan
Faculty of Engeering, Chulalongkorn University, Thailand

1. Introduction
A layered dissimilar material is a composite that consists of morethan two different materials and exhibits a unique performance by the
combination of different material properties. In particular, to achieve
the balance between a weight reduction and high strength, the development of the advanced dissimilar materials is important in the industrial applications. A previous study on the impact energy absorption
behavior of aluminum (Al) alloy and silicon carbide dissimilar materials
with three-layer structures indicated that these dissimilar material
plates have a higher energy absorption ability compared with themonolithic Al alloy [1]. The Charpy impact properties of ultrahighcarbon steel and CueZn (brass) alloy layered dissimilar materials havealso been investigated at ambient temperatures [2]. The results clarified
that the impact toughness of the dissimilar materials was approximately
9 times higher than those of the monolithic ultrahigh carbon steel andbrass alloy. The interfacial energy absorption mechanism due to delamination and separation behaviors was previously discussed in detail[3]. The properties of dissimilar materials strongly depend on theirinterface, and control of the interfacial microstructure is critical toproducing layered materials with the excellent mechanical properties.
The preparation conditions of dissimilar materials, particularly bonding
temperature and time, are also important because the intermetallic
compounds (IMCs) that form at interfaces during bonding are a dominant metallurgical factor in performance. For example, for a three-layer
dissimilar material consisting of 5052 Al alloy/AZ31B Mg alloy/5052
Al alloy fabricated by a hot rolling process, there was a strong dependence of normal and shear bonding strengths on IMC layer thickness[4]. A suitable IMC thickness must be determined to obtain superiormechanical properties in dissimilar materials. Alternatively, from theviewpoint of solid-state dissimilar metal material preparation, thefriction stir welding (FSW) technique has often been employed in previous studies [5,6]. Al alloy, AA2024 and Mg alloy, and ZE41 dissimilar
jointed materials have been fabricated by FSW, and detailed discussions⁎

have been reported regarding the formation mechanisms of various
IMCs in the joint stir zone and heat-affected zone. The solid-state reaction between the starting materials occurs easily at the interface dueto the high pressure and high temperature fields during FSW. The
plastic deformation caused by FSW, however, also leads to the mechanical breakage of in-situ formed IMC layers, which results in a
uniform dispersion of fine IMC particles in the stir zone and heat-affected zone. In short, the strengthening mechanism of dissimilar jointed
materials by both IMC dispersion and grain refinement have been thedominant focus in these previous studies. However, there has been no
discussion regarding the effect of IMC layers on the mechanical strengthof layered dissimilar materials.
Toward component weight reduction, this study employed a combination of pure Al and Mg alloys, where the commercial AZ31B (Mg3%-1%Zn/wt%) was selected, as starting materials to fabricate dissimilar materials by hot pressing (HP) [7] and heat treatment (HT). The
effects of HT temperature and time on the interfacial microstructures of
AZ31B/pure Al/AZ31B dissimilar materials were discussed. In particular, the IMCs were identified by X-ray diffraction (XRD), and theformation mechanism and growth kinetics of IMCs layers were also
discussed by using the activation energy of each compound. The microstructural analysis helped to clarify the effect of IMC layer thickness
on the dynamic Charpy impact and static tensile properties of the dissimilar materials. In-situ scanning electron microscopy (SEM) observations of the fracture behavior during tensile tests was a powerfultool for understanding crack initiation and propagation phenomena
near the interfaces between the IMC layers and each monolithic material.
2. Experimental procedure
2.1. Specimens fabrication
Pure Al (A1050, Fe: 0.26, Si: 0.1, Cu: 0.01, Al: Bal./mass%) plates

Corresponding author.
E-mail address: umedaj@jwri.osaka-u.ac.jp (J. Umeda).

https://doi.org/10.1016/j.matchar.2019.109879
Received 15 April 2019; Received in revised form 15 August 2019; Accepted 19 August 2019
Available online 21 August 20191044-5803/ © 2019 Elsevier Inc. All rights reserved.

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dispersive X-ray spectroscopic (EDS) analysis was also carried out toinvestigate interfacial element mapping. Transmission electron microscopy (TEM) and XRD analyses were conducted for phase identificationof the IMC layers. To evaluate the mechanical properties of the bonded
dissimilar materials, Charpy U-notch impact tests (maximum impact
energy: 27.4 mJ) and tensile tests (strain rate: 5 × 10−4 s−1) werecarried out at room temperature. Fig. 2 illustrates the specimen geometry used in the Charpy impact tests and the test procedure. In thisstudy, the total energy schematically shown in Fig. 2 was used toevaluate the impact toughness of the specimens. After both tests, theinterfacial fracture profiles of the specimens were also investigated by
SEM.

Punch
Stroke end blockContainer
Die
HeaterLiner (upper)
AZ31B
Al
AZ31B

3. Results and discussion
3.1. Microstructural analysis of IMC layers at interfaces of dissimilarmaterials

Liner (lower)

Fig. 3 shows the SEM images of the microstructures at the interface
between the pure Al and AZ31B plates of the as-HPed specimen andthose after HT with various conditions. It is obvious that the as-HPed(a) and HTed at 598 K for 3.6 ks (b) materials showed no formation ofinterfacial IMCs. As shown in (c-1), elliptical or spherical IMCs discontinuously formed at the interface after HT at 598 K for 7.2 ks. Thissuggests the possibility of synthesizing IMCs via solid-state reaction
during a long HT period at 598 K. According to the AleMg binary phase
diagram [9], the solubility of Mg in Al and that of Al in Mg are approximately 9.0 at.% and 6.0 at.%, respectively, at 598 K. With an increase in HT time, the interdiffusion between pure Al and AZ31B proceeded gradually through the interface, and each solid solution
concentration over the above mentioned solubility resulted in the formation of IMC particles, as schematically illustrated in Fig. 3 (c-2).
Finally, layered IMCs with a uniform thickness formed at the interface
upon HT at 598 K for 14.4 ks, as shown in Fig. 3 (d-1) and (d-2), because the diffusion of Al and Mg atoms occurred along the interface [4].
When the HT temperature was increased from 623 K to 673 K, thethickness of the IMC layers increased and two different layers clearly
formed. In the cases of HT at 648 K for 14.4 ks and at 673 K for 7.2 ks/
14.4 ks, however, large cracks propagated inside the layers along theinterfaces because of thermal stress induced by thermal expansion
differences between the two brittle IMC layers during long HT times
and subsequent cooling to room temperature.
XRD and electron-probe X-ray microanalysis (EPMA) were carriedout to identify the form of each interfacial IMC layer in the Al/AZ31B
dissimilar materials after HT at 673 K for 3.6 ks. As shown in Fig. 4 (a),in addition to Al and Mg diffraction peaks for the starting materials,
peaks corresponding to Al3Mg2 (β-phase) and Al12Mg17 (γ-phase) wereclearly detected at the interface. Fig. 4 (b) shows the EPMA line analysis
results of Al and Mg elements across the interface. The IMC layer next to
the AZ31B plate with an approximate thickness of 15 μm (“layer-1”)

Fig. 1. Schematic illustration of direct bonding of Mg alloy (AZ31B) plates topure Al plate by hot pressing process.

with thicknesses of 0.5 mm and commercial Mg alloy (AZ31B, Al: 2.95,
Zn: 1.02, Mn: 0.3, Si: 0.015, Fe: 0.001, Cu < 0.01, Mg: Bal.) plates withthicknesses of 1.0 mm were used as the starting materials. Surface
treatments by mirror polishing to prepare a smooth surface with nooxide film and by ultrasonic cleaning to completely remove oil and fats
were conducted for each starting material. The three-ply plate specimens (AZ31B/Al/AZ31B) were pre-heated at 673 K for 180 s in an
argon gas atmosphere as illustrated in Fig. 1. Then, they were immediately inserted into a container mold and consolidated by conventional HP, where the temperature of the mold and liners was 673 K and
600 MPa of pressure was applied to the specimens for 10 s by the upper
punch. A previous study suggested that no new formation of IMCs at theinterface of Al/Mg hot-rolled materials occurs after HT at 473 K [8].
Therefore, pre-heating and an HT temperature of 673 K were used inthis study. Controlling the extent of plastic deformation induced by HP
in the specimens was important for obtaining completely bonded materials at the interface. In this study, this was accurately controlled byusing a stroke end block as shown in Fig. 1. HTs with various conditions(temperature: 473 K ~ 673 K, holding time at each temperature: 3.6 ks~ 28.8 ks) under vacuum (~200 Pa) were applied to the bonded specimens to control the phases and thicknesses of the synthesized IMC
layers.
2.2. Characteristics evaluation
For microstructural analysis, compositional images in backscatteredelectron mode were captured by field emission SEM, and energy

Fig. 2. Schematic illustration of micro Charpy impact test and its specimen geometry of AZ31B/Al/AZ31B layer-structure dissimilar material.

2

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Fig. 3. SEM observation on AZ31B/Al bonding interfaces of as-bonded dissimilar material (a) and those via various heat treatment conditions (b)~(h). Illustrations ofintermetallic compounds formation at interface is also attached as (c-2) and (d-2).

was rich in Mg, whereas that close to pure Al with a thickness of 50 μm(“layer-2”) was poor in Mg. The point analysis results for No. 1–7 in
Fig. 4 (b-1) are summarized in the table in Fig. 4. At points 2 and 3
located in layer-1, the Al and Mg contents were 40–44 at.% and
56–60 at.%, respectively. These atomic ratios are close to the chemicalcomposition of Al12Mg17 (Al: Mg = 41.4: 58.6 at.%). Similarly, the ratioof Al to Mg at points 4 and 5 located in layer-2 was 60 at.% to 40 at.%,which is in very good agreement with that of Al3Mg2 (Al: Mg = 60:
40 at.%). For further microstructural investigation, TEM-EDS analysis
was conducted, which successfully identified layer-1 and layer-2 as
Al12Mg17 (γ-phase) and Al3Mg2 (β-phase), respectively, by calculationof the interplanar d-spacings using the diffraction patterns from points 1
and 2 shown in Fig. 5 (a).

3.2. Growth mechanism of interfacial IMC layers by HT
According to the microstructural observation (Fig. 3) and phase
identification (Figs. 4 and 5) results, Fig. 6 shows the relationship between IMC layer thickness (X) and the square root of the holding time( t ) of the (a) Al3Mg2 (β-phase) and (b) Al12Mg17 (γ-phase) layers resulting from different HT conditions. Both IMC layers clearly increasedin thickness with increasing holding time and HT temperature. In addition, the increase in thickness of the Al3Mg2 layer was larger than thatof the Al12Mg17 layer. According to the relationships of both IMC layers
shown in Fig. 6, the below equation was obtained [10]:

X=

Dt ,

(1)

Fig. 4. XRD profile (a), EPMA line scanning analysis result (b) and element point analysis of AZ31B/Al interface of dissimilar material after heat treatment in vacuumat 673 K for 3.6 ks.

3

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Fig. 5. TEM-EDS line scanning analysis and TEM diffraction analysis results of interface of dissimilar material after annealing treatment at 598 K for 14.4 ks. Table
shows calculation results of interplanars (d-spacing) using diffraction patters at point 1 and 2 shown in Fig. 5 (a).

Fig. 6. Relationship between each IMC layer thickness and square root of holding time at various HT temperatures of dissimilar materials: β-phase of Al3Mg2 (a) and
γ-phase of Al12Mg17 (b).

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where X (m) is the layer thickness at time t and D (m2/s) is the reaction
rate constant.
To calculate the activation energy necessary for the growth of each
IMC layer during HT [11], the Arrhenius plots [12] shown in Fig. 7
were composed using the gradient of the straight data line under eachtemperature obtained in Fig. 6. The activation energy can be calculated
by expressing the Arrhenius equation, Eq. (2), as a natural logarithm,
Eq. (3), with which the growth rate of each IMC layer at certain temperatures can be estimated [12].

D = D0 exp

Q,RT

ln D = ln D0

Q
,
RT

(2)
(3)
2

where D is the diffusion coefficient (m /s), D0 is the diffusion constant(m2/s), Q is the activation energy (J/mol), R is the gas constant(=8.314 J/mol∙K), and T is temperature (K).
In Fig. 7, the y-axis, ln D, is derived from the slope of each straightline obtained by the least squares method from the relationship between the thickness of each IMC layer (Fig. 6) and square root of theretention time t , and the x-axis is the reciprocal T−1 of each HTtemperature. The apparent activation energies for the growth of βphase and γ-phase layers calculated from the Arrhenius plot were39.0 kJ/mol and 65.8 kJ/mol, respectively. This implies that the β-

Fig. 7. Arrhenius plots for growth rate of β-phase (Al3Mg2) and γ-phase(Al12Mg17) layers at interface of AZ31B/Al/AZ31B dissimilar materials afterHT.

Fig. 8. Dependence of each intermetallic compound layer thickness ratio on HT temperature.

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Fig. 9. SEM micrographs of AZ31B/Al interface via various HT conditions: as-bonded dissimilar material (a), heat treated at 473 K for 3.6 ks (b), 523 K for 3.6 ks (c),523 K for 5.4 ks (d), 523 K for 7.2 ks (e) and 573 K for 3.6 ks (f).

r ( ) = d ( )/(d ( ) + d ( )), r ( ) = d ( )/(d ( ) + d ( )).

(4)Fig. 8 shows the relationship between each layer thickness ratio andHT temperature, where the subscript of each symbol indicates theholding time at the HT temperature and the subscript “mean” indicatesan average layer thickness ratio. With an increase in temperature,rmean(β) gradually decreased, whereas rmean(γ) increased. According
to the relationship between ln D and T−1 shown in Fig. 7, the difference
between ln D (= −14.22) in the β-phase and ln D (= −14.95) in the γphase at T = 673 K was 0.73 m2/s. On the other hand, the difference
between ln D (= −14.77) in the β-phase and ln D (= −15.88) in the γphase at T = 623 K was 1.12 m2/s. As shown in Fig. 6, the present data
have a linear relationship in the Arrhenius plot. Therefore, the formation and growth of each IMC phase is based on a thermally activateddiffusion mechanism, and with an increase in HT temperature, thediffusion coefficient of each phase increases rapidly. Because the rate ofincrease with temperature is different, however, the relative difference
between the phases decreases. Accordingly, as the HT temperature increases, rmean(β) decreases and rmean(γ) tends to increase in theformation of interfacial IMC layers.
In the range of HT experimental conditions used in this study, theratio of β-phase thickness to the total thickness of both IMC layers wasmuch larger than that of the γ-phase layer. As a further consideration,
the temperature at which the thickness ratios, r(β) and r(γ), intersect
was determined as T = 873 K, as shown in Fig. 8. The AleMg binary
phase diagram suggests that the melting point, Tm, of the β-phase(Al3Mg2) is 723 K and that of the γ-phase (Al12Mg17) is approximately
733 K. That is, it is not possible for r(β) and r(γ) to be the same in solidstate conditions. In addition, the r(γ) limit of the β-phase (Al3Mg2) at
Tm (= 723 K) is r(γ) = 0.335. In the solid-state synthesis range of thesecompounds, it is found that the proportion of β-phase to the total IMC

Fig. 10. Micro Charpy impact test results of monolithic AZ31B, pure Al and
AZ31B/Al/AZ31B layer-structure dissimilar materials after HT.

phase grows more easily at the interface of the dissimilar material
compared with the γ-phase. The thicknesses of the β- and γ-phases
synthesized at each HT temperature were denoted as d(β) and d(γ), andthe ratio of each layer thickness to the total thickness of the IMC layers
was determined as follows:6Materials Characterization 157 (2019) 109879

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Fig. 12. Dependence of shear bond strength of AZ31B/Al bonded materials on
IMC layer thickness. Details of test specimen and evaluation are also inserted in
test results.

Fig. 11. Dependence of Charpy impact toughness on total thickness of IMC
layers formed at interface.

layer thickness is always larger than that of the γ-phase.
3.3. Mechanical property evaluation of AZ31B/Al/AZ31B layered
dissimilar materials
Charpy impact tests and static tensile tests were conducted to investigate the effect of IMC layer thickness on the mechanical properties
of AZ31B/Al/AZ31B layered dissimilar materials induced by different
HT conditions after the HP process. After fabrication of the dissimilarmaterials by HP, the HT conditions of 473 K, 523 K, and 573 K for 3.6 ks
under vacuum and 523 K for 5.4 ks and 7.2 ks under vacuum were
employed. Fig. 9 shows the SEM images of the AZ31B/Al interface after
various HT conditions and the mean IMC layer thickness of eachbonded material, which is the total thickness of the β- (Al3Mg2) and γphase (Al12Mg17) layers. In the as-HPed material with no HT (a), an
IMC layer with a thickness of approximately 1 μm formed discontinuously at the interface. The other dissimilar materials contained
continuously synthesized layers with uniform thicknesses. Fig. 10shows the Charpy impact test results of the specimens, where A31B andpure Al plates were also evaluated as reference materials. First, the pureAl single specimen had the highest impact toughness, which was approximately 32% higher than that of the AZ31B single plate, although
its variation was large. All the dissimilar materials, except for the
specimen subjected to HT at 573 K for 3.6 ks, had much higher impact
toughnesses than the AZ31B single specimen. Fig. 11 shows the relationship between the Charpy impact toughnesses of the AZ31B/Al/
AZ31B layered bonded materials and total IMC layer thickness determined according to these results. When the IMC layer thickness was
between approximately 1 μm and 2 μm, the Charpy impact toughness ofthe dissimilar material increased as layer thickness increased. However,
the impact toughness decreased for bonded materials with relatively
thick layers from approximately 9 μm to 16 μm. Similarly, a previous

Fig. 13. Dependence of normal and shear bond strength on IMC layer thicknessof AZ31/A5052 dissimilar materials fabricated by hot rolling process [4].

study [3] reported that the thickness of the IMC layer formed at thebonding interface greatly affects the interfacial bonding strength between different phases of dissimilar materials due to the brittleness of
the IMCs.
Next, shear test specimens were machined from the as-HPed
AZ31B/Al/AZ31B dissimilar material and those subjected to HT at
473 K, 523 K, and 573 K to quantitatively evaluate the bonding strengthof the AZ31B/Al joint parts. As shown in Fig. 12, the shear bond stressof the AZ31B/Al joint increased for IMC layer thicknesses less thanapproximately 2 μm. The stress, however, was the smallest for a7

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Fig. 14. Fractured profiles of AZ31B/Al/AZ31B layer-structure dissimilar materials after micro Charpy impact test: as-bonded dissimilar material (a), heat treated at
473 K for 3.6 ks (b), 523 K for 3.6 ks (c), 523 K for 5.4 ks (d), 523 K for 7.2 ks (e) and 573 K for 3.6 ks (f)Fig. 15. Tensile properties of AZ31B/Al/AZ31B dissimilar materials after HT with various conditions compared to AZ31B and pure Al plates (a) and elongation
dependence on IMC layer thickness of dissimilar materials (b).

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AZ31B
AZ31B

Longitudinal
IMC layer
cracksAl
Transverse crack at
IMC/Al interface

AZ31B

Al

200 µm

10 µm

IMC layer (γ-Mg17Al12)
Longitudinal Crack
cracks

AZ31B

AZ31B
Al
Crack inside IMC layer
AZ31B

Al

200 µm

IMC layer (β-Mg2Al3)
10 µm

Fig. 16. Fractured profiles of AZ31B/Al/AZ31B dissimilar materials after tensile test: 523 K for 3.6 ks (a) and 573 K for 3.6 ks (b).

relatively large layer thickness of approximately 16 μm, which agreed
with the relationship between Charpy impact toughness and IMC layer
thickness shown in Fig. 11. Furthermore, a previous study [4] reported
the relationship between shear bond strength (closed circles, ●) and
IMC layer thickness at the bonding interface of AZ31/A5052 dissimilarclad materials, as shown in Fig. 13. The results indicated that shear
bonding strength gradually increased when the thickness was 4 μm orless and decreased drastically when the thickness was over 4 μm. This isin good agreement with the shear test results from this study shown in
Fig. 12.
The crack propagation behavior was investigated by SEM observations of laminated sections of the specimens after Charpy impact testing
to clarify why the AZ31B/Al/AZ31B dissimilar material subjected to HT
at 573 K with an IMC layer thickness of 16.3 μm had the lowest impact
fracture property. As shown in Fig. 14 (a) ~ (d), the bonded specimens
with IMC layer thicknesses of 1 μm ~ 6 μm clearly showed interfacialdelamination behavior. For the dissimilar materials with relatively
large layer thicknesses of 9 μm and 16 μm shown in (e) and (f), however, crack propagation occurred inside the brittle IMC layers, but nointerfacial delamination was observed because the thick IMC layers
easily became a path for crack propagation. In addition, since the IMC
layer has a high elastic modulus compared with those of monolithic
AZ31B and pure Al plates, a larger tensile stress is applied to the insideof the layer with an increase in layer thickness when bending deformation is applied. As a result, the dissimilar material with a 16.3 μm

IMC layer had the lowest impact toughness of all the specimens. In aprevious study [3], the toughening mechanism of laminate materials
was investigated by applying impact energy and introducing 4 types offracturing behavior: crack deflection, crack blunting, crack bridging,and stress redistribution. According to the above discussions of crack
propagation behavior at the interface of AZ31B/Al/AZ31B dissimilarmaterials with relatively thick IMC layers, the crack initiated from thebrittle IMC layer and propagated inside the layer, not at the interface.
This is mainly because of the residual stress caused by a thermal expansion difference between the IMC layers and the metal (AZ31B andAl) plates. As a result, crack-deflection-like fracturing occurred at theinterface of this bonded dissimilar material with IMC layers.
Fig. 15 (a) shows the tensile test results of the AZ31B/Al/AZ31B
dissimilar materials compared with those of the monolithic AZ31B andpure Al plates, where the mean values of 2 test specimens were used.
The results clarified that the ultimate tensile strength and yield stress(0.2%YS) were almost constant regardless of IMC layer thickness induced by different HT conditions. As shown in Fig. 15 (b), however, theelongation to failure did depend on layer thickness. In the thickness
range of 1 μm ~ 2 μm, the elongation gradually increased with increasing thickness because the interfacial shear bonding strength also
increased, as shown in Fig. 12. In order to analyze the factors reducing
the elongation of the bonded specimens with IMC layer thicknesses over
5 μm, the fractured profiles, in particular the micro-crack initiation
behaviors, at the interfaces of the tensile specimens HTed at 523 K for

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with relatively thick IMC layers was mainly due to large transverse
cracks inside the brittle IMC layer.
To further investigate the crack initiation and propagation behaviorof the bonded specimen with a relatively thick IMC layer of 16.3 μm
shown in Fig. 16 (b), in-situ SEM observations of the fracturing phenomenon during tensile testing were carried out. As shown in Fig. 17(a), a small tensile test system was installed in the SEM equipment toevaluate the fracturing behavior in this study. An example of the loaddisplacement profile from this tensile test of the AZ31B/Al/AZ31B
dissimilar bonded material subjected to HT at 573 K for 3.6 ks is shownin Fig. 17 (b). During the tensile test, the cross-head was temporarily
stopped to observe interfacial fractures at each point (a)~(f). Fig. 18 (a)shows the well-bonded AZ31B/Al interface, where no cracks were
present and an IMC layer with a thickness of approximately 15 μm
existed between both monolithic plates, which was close to the layer
thickness shown in Fig. 9 (f). When an approximately 50 N load wasapplied at point (b) in Fig. 17 (b), some longitudinal cracks initiatedinside the IMC layer, as shown in Fig. 18 (b). With an increase in theapplied load from 65 N to 70 N (corresponding to points (c) and (d),respectively), the transverse crack mainly propagated inside the IMC
layer, and the crack opening displacement became increasingly larger,
as shown in Fig. 18 (c) and (d). The results also clarified that the small
longitudinal cracks and main transverse crack never propagated into
the ductile monolithic materials of the AZ31B and pure Al plates withincreasing displacement, as shown in Fig. 18 (e) and (f). This fracturing
behavior is in good agreement with the SEM observations shown in
Fig. 16 (b). The results indicated that the plastic deformation ability,
particularly the elongation, of the dissimilar material with a relatively
thick IMC layer was strongly limited by the brittleness of the β- and γphases.
4. Conclusions
The formation mechanism and growth kinetics of IMCs at the interface of AZ31B/pure Al/AZ31B layered dissimilar materials after HTwith various conditions were quantitatively discussed by using the activation energy for IMC layer growth calculated from the Arrhenius plot
and XRD, SEM, and TEM analyses. The apparent activation energies forthe growth of β-phase (Al3Mg2) and γ-phase (Al12Mg17) layers were39.0 kJ/mol and 65.8 kJ/mol, respectively; that is, the β-phase grew
more easily at the interface of the dissimilar material compared withthe γ-phase. Regarding the dynamic impact properties of the dissimilarmaterials and their fracturing behavior, a maximum Charpy impact
toughness was obtained for an IMC layer thickness of approximately
2 μm, and the toughness drastically decreased with increasing IMC
layer thickness. In-situ SEM observations during tensile testing indicated that cracks propagated inside the brittle IMC layers, but nointerfacial delamination was observed for dissimilar materials withrelatively thick IMC layers. It was concluded that these MgeAl dissimilar materials showed the important advantages to reduce thecomponent weight and improve both strength, ductility and toughness
compared to the conventional Mg or Al alloys.

Fig. 17. Small tensile test system installed in SEM equipment (a) and example
of load-displacement profile in tensile test of AZ31B/Al/AZ31B dissimilar
bonded material with HT at 573 K for 3.6 ks (b).

3.6 ks (a) and 573 K for 3.6 ks (b) were investigated by SEM observation, as shown in Fig. 16. In the specimen with an IMC layer thickness of
2.1 μm (a), numerous micro-cracks perpendicular to the IMC layer(“longitudinal cracks”) were detected inside the layer, and the main
fracture was due to a transverse (parallel) crack at the interface between the IMC layer and pure Al plate. On the other hand, the specimen
with an IMC layer thickness of 16.3 μm (b) showed not only longitudinal cracks but also large transverse cracks inside the β-phase(Al3Mg2) layer. Longitudinal crack propagation, however, was completely halted at both the monolithic AZ31B and pure Al plates. According to these results, the poor elongation of the bonded specimens

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Fig. 18. Crack initiation and propagation inside IMC layers formed at interface during in-situ tensile test at (a) ~ (f) points in Fig. 17 (b).

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This study was financially supported by the Project to Create
Research and Educational Hubs for Innovative Manufacturing in Asiaand the research program on establishment of a virtual joint laboratory
between RMIT and JWRI, Osaka University. The authors showed great
thanks for Mr. Hiroyuki Sannomiya for his supports in this experiment.
Declaration of competing interest
None.
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