Laser-tungsten inert gas hybrid welding of dissimilar metals AZ31B Mg alloys to Zn coated steel

Laser-tungsten inert gas hybrid welding of dissimilar metals AZ31B Mg alloys to Zn coated steel

Materials and Design 49 (2013) 766–773 Contents lists available at SciVerse ScienceDirect Materials and Design journal homepage:

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Materials and Design 49 (2013) 766–773

Contents lists available at SciVerse ScienceDirect

Materials and Design journal homepage:

Laser-tungsten inert gas hybrid welding of dissimilar metals AZ31B Mg alloys to Zn coated steel Caiwang Tan a,b, Liqun Li a,⇑, Yanbin Chen a, Wei Guo a a b

State Key Laboratory of Advanced Welding and Joining, School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China Center for Advanced Materials Joining, Department of Mechanical and Mechatronics Engineering, University of Waterloo, Waterloo, Canada N2L 3G1

a r t i c l e

i n f o

Article history: Received 4 January 2013 Accepted 16 February 2013 Available online 26 February 2013 Keywords: C. Lasers D. Welding F. Microstructure

a b s t r a c t Laser-tungsten inert gas (TIG) hybrid welding has been developed for joining Mg alloys to Zn coated steel in a lap joint configuration. The joint could not be produced in laser or arc welding only, while acceptable joints without obvious defects were obtained with a relatively wide processing window in the hybrid process. Two reaction layers were observed to form at the interface and were identified as Mg–Zn eutectic structure (a-Mg + MgZn) and Fe3Al phase by TEM analysis. In some cases, Al6Mn phase also formed adjacent to the Fe–Al reaction layer. The tensile-shear strength attained the maximum value of 68 MPa, representing 52.3% joint efficiency relative to Mg base metal. The element Al from AZ31B Mg alloys diffused to the liquid/solid interface and then reacted with the elements from steel, such as Fe and Mn, contributing to the metallurgical bonding at the interface. The weak bonding between Mg–Zn reaction layer and newly formed Fe–Al layer resulted in the interfacial failure. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction Recently, the demand for reducing fuel consumption and greenhouse gas emissions has been continuously increasing in transportation industry. These requirements can be satisfied to a great extent by reducing the total weight of vehicles since weight reduction has been confirmed as one of the important factors for improving fuel efficiency [1]. Therefore, the hybrid structural components which can integrate the advantages of dissimilar metals in one part are highly needed, thereby facilitating the research on dissimilar joining such as Al to steel [2,3], Al to Ti [4,5], Mg to steel [6–8] and Mg to Ti [9,10]. Among these metals, magnesium alloys have been considered as promising materials for automotive applications due to their low density and high specific strength, which draws considerable attentions in the last few years [11,12]. Steel, on the other hand, is currently the most common metal used in the automotive industry. Therefore, reliable joining of magnesium alloys to steel becomes a feasible method to fabricate the lightweight automotive components. Nowadays, Zn coated steel has been widely using in the car body panels due to its excellent corrosion resistance. Thus, direct joining Mg alloys to Zn coated steel is inevitable in some cases. Friction stir welding (FSW) and resistance spot welding (RSW) have been considered as the effective joining techniques in joining dissimilar AZ31 Mg alloys and Zn coated steel. This is because both processes have the ability to accelerate the diffusion of Al from ⇑ Corresponding author. Tel./fax: +86 451 86415506. E-mail address: [email protected] (L. Li). 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved.

base metal Mg alloys to steel producing Fe–Al phase through the external force and strong stirring action [13,14]. As a result, metallurgical bonding was achieved between immiscible Mg and Fe. Meanwhile, the effective joining of Mg to steel was found to fail without Zn coating, indicating Zn was an essential part in the whole process [13,15]. However, both processes have some disadvantages in joining efficiency and flexibility, which limited their widespread application in industry. Laser welding process, as an advanced welding technique, has shown great advantages in high flexibility and adaptability for practical application. The characteristics of laser welding AZ31B magnesium alloy to Zn coated steel were investigated by Wahba and Katayama [16]. An acceptable joint could be obtained in the laser conduction mode. The keyhole welding mode easily caused the unstable process with obvious defects such as underfill and spatter. Therefore, the heat input must be controlled precisely. To expand the processing window while preventing the severe evaporation of Mg, the method of laser-TIG hybrid welding is proposed. Liu and Qi [17] investigated the feasibility of joining Mg to uncoated steel by laser-TIG hybrid welding and reported the combined action of laser and TIG could not only stabilize the arc causing good joint appearance, but also increase the penetration in the lower steel. However, some metallic oxides were easily produced at the Mg/Fe interface, deteriorating the joint strength sharply [18]. Nevertheless, limited literature has focused on the joining of Mg to Zn coated steel by laser-TIG hybrid welding. The presence of Zn coating was found to have a significant influence on interfacial reaction and mechanical properties during laser welding Mg to steel [8]. The aim of the present investigation is, therefore, to investigate


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the characteristics of joining Mg to Zn coated steel using the laserTIG hybrid welding process. The microstructure at the interface was characterized and mechanical properties were evaluated. Additionally, the role of Zn coating on laser-TIG hybrid welding of Mg to Zn coated steel and joining mechanism were elucidated. 2. Experimental details Commercially available AZ31B magnesium alloys sheet and Zn coated DP980 steel sheet were used in the present study. The dimension of these sheets was 100  30  1.5 mm. The chemical compositions of the base metals are listed in Table 1. The thickness of Zn coating was approximately 10–15 lm. The pre-existing Fe–Al phase was observed at the interface and its thickness was about 0.9 lm. Detailed description about the microstructure has been reported in our previous investigation [19]. The experiments were performed with a 10 kW fiber laser (IPG YLR-10000) and TIG welding machine (Magic Wave 4000) with a maximum output of 400 A. The laser beam had a wavelength of 1070 nm and a beam parameter product (BPP) of 7.2 mm mrad. It was transmitted by a 200-lm core diameter fiber and focused by a 200-mm lens to obtain a spot size of 0.2 mm. TIG welding machine was used in continuous mode. The schematic diagram of LWB process is illustrated in Fig. 1. The assembly was fixed in a lap configuration with the Mg sheet placed on top of steel sheet. Prior to welding, both sheets were ground and degreased. Laser beam was irradiated on the surface of workpiece vertically. Argon shielding gas was provided to prevent oxidation of the molten pool. The process parameters used in the experiment are listed in Table 2. After laser-TIG hybrid welding process, typical transverse crosssections of the welded specimens were cut and mounted in epoxy resin. Standard metallographic preparation procedures were then utilized. Macrostructures of the joints were observed using an optical metallographic microscope. The reaction layer at the interface between weld seam and steel was observed using scanning electron microscopy (SEM) in back-scattered electron (BSE) mode. A transmission electron microscopy (TEM) foil of the bonded region was prepared using the focused ion beam (FIB) technique. An insitu lift out method was applied for the preparation of FIB-TEM specimen. TEM with a Tecnai-G2 F30 operating at a nominal voltage of 300 kV was used to characterize the microstructure in detail. Phase identification was investigated by selected-area electron diffraction (SAED) combined with energy dispersive spectroscopy (EDS). Z-contrast images were acquired using a high angle annular dark field (HAADF) detector in scanning transmission electron microscopy (STEM) mode. According to GB/T 2651-2008 standard (equivalent to ISO 9016: 2001) [20], the tensile-shear tests were operated at room temperature using INSTRON 5569 at a crosshead speed of 1 mm/min. Joint strength was calculated via the tensile testing of at least three specimens. The phase structure at fracture surface was identified using X-ray diffraction (XRD). 3. Results and discussion 3.1. Joint appearances and cross section Fig. 2 shows the representative appearances of the joints produced by laser-TIG hybrid welding, laser welding and TIG welding, respectively. In the hybrid welding process, the arc current was

Fig. 1. Schematic of laser-TIG hybrid welding process.

Table 2 Experimental parameters used in the process. Welding parameters


Laser power (W) Arc current (A) Defocus distance from steel surface (mm) Welding speed (m/min) Distance between laser and arc DLA (mm) The angle of the arc and the workpiece (h) Flow rate of shielding gas Ar (L/min)

1200–2100 30–120 +20 0.5 5 45 22

kept constant at 30 A after preliminary trials, while the laser power employed was varied from 1200 W to 2100 W. Smooth and uniform weld surfaces without obvious defects were observed from Fig. 2a–c, suggesting stable processes were obtained at the laser power ranging from 1200 W to 1800 W. However, when the laser power increased further to 2100 W a complete penetration occurred resulting in Mg surface collapse and crater defect, as shown in Fig. 2d and e. Meanwhile, for the purpose of comparison, the laser welding and arc welding only were performed. It was obvious that joining Mg to Zn coated steel was difficult when using laser or arc welding only, as indicated in Fig. 2f–j. More specifically, in the case of laser welding only (Fig. 2f and g), the expulsion of molten pool in the Mg weld surface was observed regardless of the applied laser power. With regard to the arc welding only, the unstable process was noticed from the weld surface in Fig. 2h mainly due to the low arc current (60 A) and relatively high welding speed (0.5 m/min). The upper Mg sheet was found to be totally melted by the selected arc current, the faying surface was however observed without any penetration in the lower steel sheet. When the arc current increased further to 120 A the severe vaporization and burning loss of Mg occurred, without any improvement in interfacial reaction between Mg and steel. In sum, acceptable joint was successfully achieved using the laser-TIG hybrid welding. This phenomenon was closely associated with the characteristics of laser and arc. It was known that arc welding was limited to low penetration depth. The melting of the lower steel sheet, even slight, could not take place using the arc welding only. Severe

Table 1 Chemical compositions of base metals (wt.%).

DP980 steel AZ31B










Bal. 0.005

– Bal.

0.05 2.5–3.5

– 0.5–1.5

2.1 0.2–0.5

0.05 0.10

0.135 –

0.35 –

0.15 –


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Fig. 3. Typical cross section of Mg/Zn coated steel joints by laser-TIG hybrid welding.

Fig. 2. Typical joint appearances: (a–e) laser-TIG hybrid welding, (f and g) laser only and (h–j) arc only.

evaporation of Mg (boiling point 1091 °C) would occur if Fe gets melted due to the huge difference in their melting point (Mg  630 °C, Fe  1538 °C), as shown in Fig. 2j. While in the case of the laser welding only, the peak temperatures on the weld pool surface increased to a high level instantly, and often exceeded the boiling points of materials. In such situations, the equilibrium pressure on the weld pool surface was higher than the atmospheric pressure and the escaping vapor exerted a large recoil force on the weld pool surface [21]. Consequently, the molten metal was expelled from the weld pool surface, as shown in Fig. 2f and g. During the hybrid welding, the arc played a significant role in preheating the lower steel sheet and increasing the absorption of laser beam by the steel sheet. The upper Mg sheet was first melted by the low arc current. The molten pool reduced the reflectivity of Mg to laser beam and thus increased the absorption of laser beam [17]. The accumulated heat was then transferred to melt steel sheet. By this laser–arc interaction, the molten pool would not be expelled causing good joint appearance, corresponding to the observation in Fig. 2a–c. Because of this characteristic, the hybrid welding process was also applied to welding Mg alloys. Gao et al. [22,23] realized high power laser hybrid welding of Mg alloys and obtained acceptable joints with complete penetration and few weld defects. These results confirmed hybrid welding had the great potential for suppressing weld defects and improving the joint quality, whether in similar welding or in dissimilar welding. Fig. 3 shows the typical cross section of dissimilar joints made by the laser-TIG hybrid welding (laser power 1800 W, arc current 30 A). Obviously, the seam was defect-free without any porosity or cracking, indicating it was a feasible way to join Mg to Zn coated steel using the hybrid process. Furthermore, it could be observed from the seam width that the upper Mg sheet was melted largely by the arc, while the lower steel sheet at the center position of heat source was slightly melted by the defocused laser. On this occasion, alloying elements in base metals may get involved in the diffusion, which further induced the interfacial reaction between Mg and Zn coated steel. Therefore, it was necessary to investigate the interfacial microstructure in detail.

3.2. Morphology of interfacial microstructure and phase identification Fig. 4 presents the SEM morphologies of reaction layers along the joint interface at different laser powers. The location for the

observation of the interface was indicated by rectangle in Fig. 3. It was clear that the newly formed reaction layers were in intimate contact with each other. Two layers with different morphologies could be identified at the interface. The first layer was found to exhibit a continuous and irregular morphology. High magnification image revealed the phase was characterized by lamellar morphology indicated in the inset of Fig. 3b. According to the EDS analysis, this gray phase primarily consisted of 70.9 at.% Mg, 20.46 at.% Zn and 6.88 at.% Al. Based on the Mg–Zn binary diagram, it was therefore identified as the eutectic structure (a-Mg + MgZn), which was reported by Ma et al. in previous studies [24,25]. They designed a Zn based filler metal to braze similar AZ31B magnesium alloy using high-frequency induction brazing technique. The reaction occurred between Zn and Mg, resulting in this eutectic structure. The results obtained in our work differed from the aforementioned studies in thickness of eutectic structure and morphology. This could be attributed to the employed heat source and thickness of Zn. In the study, thin Zn coating would melt first and then react with molten Mg at the combined action of laser-TIG welding. The newly formed structure did not have the sufficient time to grow up, giving rise to the compact and thin characteristics. A detailed description about reaction mechanism has been reported in our previous work [8]. The second layer adjacent to steel substrate was characterized by homogeneous morphology, suggesting it was an intermetallic compounds (IMCs). The morphology was totally different from that of reaction phase achieved at the action of laser welding only [16]. Dispersion of fine particles was observed at the transition zone which was in between Mg–Zn eutectic and steel matrix. The formed layer was so thin that the EDS result was not accurate to identify its phase. Therefore, further investigation via TEM analysis was required. In addition, it could be seen that the thickness of the eutectic structure, by and large, reduced gradually with the increase of laser power from 1200 W to 2100 W. To reveal the concentration fluctuation of main elements across the interface between the seam and the steel, the EDS line scan analysis was performed. Fig. 5 shows the EDS line scan results corresponding to the interfaces presented in Fig. 4. As shown in Fig. 5, the Mg content reduced first from the seam to the Mg–Zn eutectic structure, while the Zn content increased at this stage. After reaching the eutectic structure, both elements went to the steady stage suggesting constant proportion between Mg and Zn [8]. At last, the Mg and Zn content both went down sharply as approaching the steel side. The Fe content was varied in the opposite direction to the Mg. It was worth noticing that the Al content enriched at the interface adjacent to the steel side, suggesting atom diffusion occurred between Al and Fe. Additionally, this Fe–Al phase was found to grow from 1.9 lm to 2.9 lm with the increasing laser power ranging from 1200 W to 2100 W.

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Fig. 4. Interfacial microstructure morphologies at different laser powers: (a) 1200 W, (b) 1500 W, (c) 1800 W and (d) 2100 W.

Fig. 5. EDS line scan results corresponding to the interfaces presented in Fig. 4.

The representative concentration profiles of Mg, Zn, Al, Mn and Fe across the interface between the seam and the steel are shown in Fig. 6. It was evident that the Al content enriched at the interface corresponding to the observation in Fig. 5. Note that the Mn content also peaked at the interface, which was not detected in EDS

line scan. It suggested that Al–Mn phase possibly formed at the interface in some situations. TEM analysis was conducted to identify the composition and structure of the reaction layers formed between the seam and the steel. Fig. 7 shows the TEM bright field micrographs and HAADF


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images with corresponding SAED pattern at the interface of Mg/ steel. As shown in Fig. 7a, an ultra-thin reaction layer consisting of a continuous reaction phase and a particle was clearly observed from the bright field TEM images. It was evident in Fig. 7b the feature of the phases in the Mg seam was lamellar. The rod-shaped gray phases corresponded to MgZn phases and the dark part was confirmed as a-Mg solid solution. These observations were in good agreement with aforementioned EDS analysis in Figs. 5 and 6. In the HAADF STEM micrograph, Z-contrast imaging of the newly formed reaction phases differing from steel substrate were observed, which further confirmed the observation in Fig. 7a. Similarly, the Mg–Zn eutectic structure was observed using HAADF in Fig. 7d. According to the SAED pattern calibration result, the reaction phases at the interfacial layer were indexed as Fe3Al and Al6Mn, respectively. The presence of Al–Fe and Al–Mn phases at the Mg/steel interface indicated metallurgical bonding was produced by Al element diffusing from AZ31B Mg alloys to react with Fe and Mn. That was to say, the interfacial bonding of immiscible Mg and Fe was realized by diffusion and combination of alloying elements from base metals despite fast heating and cooling rate of laser-TIG hybrid fusion welding [16].

3.3. The role of Zn coating in joining Mg to steel In the case of joining Mg to uncoated steel, the reaction phase was not observed using the same process parameters. Severe oxidation was found along the interface. These findings indicated that the Zn coating was an indispensable part for producing interfacial reaction and metallurgical bonding of Mg and Fe. The feature of Zn coating in protecting joining area was also reported when using laser conduction welding [16]. Similar phenomenon was observed during FSW and RSW, but in different protection mechanisms [13–15]. In both processes, although Zn coating did not contribute to the interfacial reaction product, it could protect the joining interface by reacting with Mg and dissolving other oxide films and finally was squeezed out of the bonded region. On the other hand, local embrittlement at the Mg base metal may be induced by penetration of a small amount of Zn coating, which became the crack initiation in the Mg/steel joint [26]. In this study, the Zn coating melted and dissolved some Mg producing liquid Mg–Zn product and flowed to the peripheral region of bonded interface due to its good wettability [8]. The liquid product impeded the oxygen to enter into the joining interface for a long

time since it was the last solidified structure (solidification point  325 °C) in the molten pool. As a result, an ideal condition with intimate contact between the Mg and the steel was produced, which was beneficial to the metallurgical bonding.

3.4. Joint strength and fracture behavior Fig. 8 shows the tensile-shear strength of hybrid welded Mg/Zn coated steel joints as a function of laser power. As shown in the figure, the tensile-shear strength increased first as the laser power increased from 1000 W to 1800 W. It attained a maximum value of 68 MPa at the laser power of 1800 W, representing 52.3% joint efficiency relative to Mg base metal. The calculation of joint efficiency was reported elsewhere [27]. As shown in the inset of Fig. 8, the serious deformation of the lower steel sheet was noticed after the tensile-shear test, suggesting relatively high strength was achieved at the laser power of 1800 W. The joint strength dropped drastically with further increase in the laser power. The sharp decrease could be attributed to the cracks formed at the periphery of joining zone caused by the welding distortion. At high heat input, the two ends of Mg sheet would tilt after the welding process. During the test, the cracks preferentially moved through the top Mg sheet, which deteriorated the joint strength. This failure behavior was similar to the fracture observation in Mg-to-steel joint by RSW [26]. All joints fractured along the interface, indicating the interface was the weakest region at the joint. Fig. 9 shows the SEM morphologies of fracture surfaces of dissimilar joints made at the laser power of 1200 W and 1800 W, respectively. It could be observed that the feature of fracture surface at the Mg side exhibited tear ridges in Fig. 9a. Distinct lamellar morphology corresponding to Mg–Zn eutectic structure was observed without severe deformation at higher magnification, implying it was brittle fracture [8]. At the steel side, the fracture morphologies were characterized by smooth surface attached with some dispersed particles, as shown in the inset of Fig. 9b. According to EDS result, the fracture surface was mainly composed of 2.34 at.% Mg, 32.53 at.% Al, 4.68 at.% Mn and 59.53 at.% Fe, indicating crack propagated along the interface of Mg–Zn eutectic structure and Fe–Al layer. In the case of joints produced at the laser power of 1800 W, some tearing of Mg–Zn eutectic structure indicated by arrows was noticed at the fracture surface of the corresponding Mg side, which was quite different from Fig. 9a. Similarly, the presence of more tearing

Fig. 6. Elemental distribution of Mg/steel interface: (a) SEM micrograph of interfacial microstructure, (b–f) Mg, Zn, Al, Mn and Fe mapping.

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Fig. 7. TEM investigation of the interface region in Mg/Zn coated steel joint: (a and b) bright field micrograph taken at the interface and the seam, (c and d) HAADF micrographs corresponding to (a) and (b), (e and f) SAED patterns of interfacial reaction phases.

was revealed on the steel side, suggesting high resistance to crack propagation. In addition, the fracture surface consisted of 4.73 at.% Mg, 44.4 at.% Al, 12.84 at.% Mn and 37.34 at.% Fe based on the EDS analysis. Compared with the findings in Fig. 9b, Al and Mn elements with higher content were detected suggesting the occurrence of more interfacial reaction. The formation of reaction phases would improve the joint strength to a great extent, which was in agreement with the result shown in Fig. 8. XRD analysis was performed on the fracture surfaces of both Mg and steel sides to confirm the observation in Fig. 9c and d. Fig. 10 shows the XRD patterns of fracture surfaces. Reaction phase Fe3Al was detected in the Mg side with low diffraction peak, suggesting some residual Fe–Al phase was attached to the Mg–Zn structure. Accordingly, the remaining eutectic structure was found on the steel side besides the newly formed phases Al6Mn and Fe3Al.

3.5. Joining mechanism of Mg to Zn coated steel Based on the analyses above, the joining mechanism of Mg to Zn coated steel by laser-hybrid welding was elucidated with the assistance of schematic diagram shown in Fig. 11. First, melting of the top AZ31B Mg sheet and Zn coating occurred. The liquid Zn atoms and liquid Mg atoms dissolved into each other. The presence of molten Mg–Zn reaction products at the periphery area protected the joining region from oxidation by dissolving oxide film and impurity. This was the main reason for unsuccessful joining Mg to uncoated steel. At the same time, the pre-existing Fe–Al phase also melted after suffering high laser power density. The atoms in the steel such as Fe and Mn were activated and dissolved into the molten pool. Then Al diffused from molten AZ31B to the interface and segregated at the front of liquid/solid interface since atoms tended to diffuse from high


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Fig. 8. Tensile-shear strength versus laser power.

chemical potential to low chemical potential, as shown in Fig. 11a and b [6]. The maximum content of Al element at the interface could

reached (20–30) at.% at fast thermal cycle [28]. As the temperature decreased to about 1000 °C, the solid solubility of Al in Fe at the solid/liquid interface was saturated inducing the crystallization of FeAl and subsequently transformed to Fe3Al, indicated in Fig. 11c. At some local region, the content of Al and Mn reached the stoichiometric composition of Al6Mn, which was precipitated when the temperature decreased to 705 °C [28]. When the temperature further decreased to 650 °C below but higher than 325 °C, liquid AZ31B was solidified as the Mg seam, as shown in Fig. 11d. Finally, as the temperature decreased to 325 °C or below, a eutectic reaction took place in the remaining liquid phase. Lamellar phases formed in between the Mg seam and Fe–Al or Al–Mn phase, as shown in Fig. 11e. Therefore, the bonding status of the Fe–Al phase and Mg– Zn eutectic structure was similar to brazing since the former was the first precipitated phase with high melting point and the latter was the last formed structure with low melting point. The weak metallurgical bonding between the two layers was formed, owing to the big mismatch in their lattice structure. This was why the weakest region was at the Mg–Zn eutectic structure/Fe3Al interface, giving rise to the interfacial failure indicated in Fig. 11e.

Fig. 9. Fracture surface morphologies of joints produced at 1500 W and 1800 W: (a) Mg side of 1500 W, (b) steel side of 1500 W, (c) Mg side of 1800 W and (d) steel side of 1800 W.

Fig. 10. XRD patterns of fracture surfaces of Mg/Zn coated steel joint produced at laser power of 1800 W: (a) Mg side and (b) steel side.

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the authors (Caiwang Tan) is grateful for the financial support by the China Scholarship Council for one year study at the University of Waterloo. The assistance of Dr. Jiecai Feng in performing the experiments is gratefully acknowledged. The authors also express their gratitude to Professor Y. Zhou for his helpful discussion. References

Fig. 11. Schematic of joining mechanism: (a) preparation for laser-TIG hybrid welding, (b) melting of AZ31B Mg alloys, Zn coating and pre-existing Fe–Al phase, (c–e) solidification at different temperature ranges.

4. Conclusions Laser-TIG hybrid welding has been developed for joining Mg alloys to Zn coated steel. The appearance and cross-section were observed. The morphologies of interfacial microstructure were characterized and identified. The tensile-shear strength was evaluated. The role of Zn coating on joining Mg to steel and joining mechanism were elucidated. The following conclusions could be drawn: (1) Successful joining of Mg alloys to Zn coated steel could be realized by laser-TIG hybrid welding process. The process stability and weld quality was highly improved in comparison with the laser or arc welding only. (2) The microstructure formed at the interface was observed to consist mainly of two layers with different morphologies. According to the SEM observation and TEM analysis, they were identified as Mg–Zn eutectic structure (a-Mg + MgZn) and Fe3Al phase. In some cases, the Al6Mn phase also formed adjacent to the Fe–Al reaction layer. The thickness of the eutectic structure, by and large, reduced gradually with the increase of laser power. The Fe–Al reaction layer, however, was varied in the opposite way. (3) The tensile-shear strength increased first with the increasing laser power. A maximum value of 68 MPa was achieved at the laser power of 1800 W, representing 52.3% joint efficiency relative to Mg base metal. The severe deformation of the lower steel sheet was observed after the tensile-shear test, indicating relatively high strength. The decrease of joint strength with further increase in laser power was attributed to the cracks formed at the periphery of joining zone. (4) The presence of Zn coating protected the interfacial reaction from oxidation. The element Al from AZ31B Mg alloys diffused to the solid/liquid interface and reacted with the elements from steel, such as Al and Mn, contributing to the metallurgical bonding at the interface. The weak bonding between Mg–Zn reaction layer and Fe–Al layer resulted in the interfacial failure.

Acknowledgements This work is supported by special foundation for scientific and technical innovation, Harbin (Grant No. 2012RFLXG028). One of

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