CO2 laser welding–brazing characteristics of dissimilar metals AZ31B Mg alloy to Zn coated dual phase steel with Mg based filler

CO2 laser welding–brazing characteristics of dissimilar metals AZ31B Mg alloy to Zn coated dual phase steel with Mg based filler

Journal of Materials Processing Technology 213 (2013) 361–375 Contents lists available at SciVerse ScienceDirect Journal of Materials Processing Tec...

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Journal of Materials Processing Technology 213 (2013) 361–375

Contents lists available at SciVerse ScienceDirect

Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec

CO2 laser welding–brazing characteristics of dissimilar metals AZ31B Mg alloy to Zn coated dual phase steel with Mg based filler Liqun Li ∗ , Caiwang Tan, Yanbin Chen, Wei Guo, Changxing Mei State Key Laboratory of Advanced Welding and Joining, School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China

a r t i c l e

i n f o

Article history: Received 13 April 2012 Received in revised form 9 October 2012 Accepted 20 October 2012 Available online 24 October 2012 Keywords: Magnesium alloy Zn coated steel Laser welding–brazing Dissimilar joining Microstructure

a b s t r a c t The influence of process variables including heating mode, flux, laser beam offset, and travel speed on the weld bead geometry and joint strength was investigated during laser welding–brazing (LWB) of AZ31B Mg alloy to Zn coated steel. The wettability of molten filler metal on steel surface was studied via a Charge-coupled Device (CCD) camera. The reaction layers along the interface were characterized and the failure mechanism was identified. Dual beam processing was found to preheat the steel substrate and promote the wettability of molten filler metal on the steel surface, thereby improving the corresponding joint strength. Utilizing a flux was found to produce a similar effect on molten filler metal. The optimized range of laser offset was found to be between 0.5 and 1.0 mm toward the steel side of the joint. These optimized parameters led to a maximum joint strength of 228 N/mm. The joint strength was however found to decrease with increasing travel speed. Cracking was identified with travel speeds greater than 1 m/min. Microstructural characterization showed that heterogeneous interfacial reaction layers were produced from the seam head to the seam root of the joint. The reaction layer thickness varied within a certain range when applying different process parameters, suggesting the growth of interfacial layer was not essentially related to the heat input. The primary failure mode of the lap specimens was interfacial fracture. Cracks propagated along the Mg–Zn reaction layer and steel interface. Original Fe–Al phase formed during the hot-dip galvanization process hindered the metallurgical bonding of Mg–Zn reaction layer and steel substrate, which was attributed to interfacial type failure. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Magnesium and its alloys have recently attained much interest due to excellent properties such as low density, high specific strength, good damping capacity, recyclability, and low cost. As promising structural materials they offer great advantages in weight reduction in the automotive and aerospace industries. More specifically, there is a driving force in the automotive industry to preserve the global environment by reducing vehicles weight and improving fuel efficiency. High strength steels such as dual phase (DP) and transformation induced plasticity (TRIP) steels are also commonly being employed to maintain structural strength as well as safety while utilizing thinner sheets. The presence of zinc coating on these advanced steel surfaces sharply improves its corrosion resistance which contributes to widespread application of Zn coated steel in the automotive industry. Therefore, fabricating hybrid component via dissimilar combination of magnesium alloys and advanced Zn coated steels would be a great asset in further reducing vehicle weight and therefore fuel efficiency. Moreover,

∗ Corresponding author. Tel.: +86 451 86415506; fax: +86 451 86415506. E-mail addresses: [email protected], [email protected] (L. Li). 0924-0136/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jmatprotec.2012.10.009

attaining reliable hybrid joint between Mg alloys and Zn coated steels will expand the application of Mg alloys to automotive structural components. The joining of dissimilar Mg alloy and steel poses large challenges due to considerable differences in their physical and metallurgical properties. For instance, the large difference in their melting points (Mg ∼ 630 ◦ C, Fe ∼ 1535 ◦ C) makes conventional fusion welding techniques ineffective. In addition, the presence of Zn coating has a significant influence on welding process stability and weld quality since Zn is easily vaporized having a low boiling point of 906 ◦ C. In order to overcome these problems new welding techniques such as friction stir welding (FSW) and resistance spot welding (RSW) were explored. Chen and Nakata (2009, 2010) studied the friction stir lap welding characteristics of magnesium alloy and steel and found that Zn coating and probe length had a crucial effect on the microstructure and mechanical properties of the lap joints. Zn coating could improve the weldability of Mg to steel due to the formation of liquid Mg–Zn product in FSW process. With regard to probe length, the defect-free joints and high failure load were obtained with either shorter probe for Zn coated steel or longer probe for wire brushed steel surface. Jana et al. (2010) found the Zn coating became molten and alloyed with Mg sheet, resulting in the formation of solidified Zn–Mg alloy layer and reported

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Table 1 Chemical compositions of base and filler metals (mass fraction %).

DP980 steel AZ31B Filler metal

Fe

Mg

Al

Zn

Mn

Si

C

Mo

Cr

Bal. 0.005 0.0018

– Bal. Bal.

0.05 2.5–3.5 3.27

– 0.5–1.5 1.26

2.1 0.2–0.5 0.24

0.05 0.10 0.0072

0.135 – –

0.35 – –

0.15 – –

the joint strength was associated with the type of steel. Liu et al. (2010) reported that molten Zn coating was squeezed out of the fusion zone to the periphery during RSW, producing an ideal condition with intimate contact between fresh magnesium and steel surfaces. Liu et al. (2011) also reported metallurgical bonding was realized and strong joints were achieved via a nano-scaled Fe2 Al5 interlayer, although Mg and Fe were immiscible with each other. However, disadvantages of the aforementioned processes such as large limitations in joint design, the need to contact the workpiece and the maintenance requirements of the tool or electrode have limited their application in industry. As an advanced welding technique laser processes have been widely adopted by the automotive industry due to unique performance offerings such as high speed, low deformation and precise control of laser beam energy. It has advantages over FSW and RSW due to its high flexibility and adaptability for practical application. Wahba and Katayama (2012) studied the characteristics of laser welding of AZ31B magnesium alloy to Zn coated steel and concluded that an acceptable joint with high strength could be achieved via laser welding in the conduction mode. The keyhole welding mode however resulted in an unstable process with severe spatter and vaporization of the magnesium alloy. To mitigate the severe vaporization of the magnesium alloy and Zn coating while further improving the adaptability in practical application, the method of laser welding–brazing (LWB) with filler wire is proposed. This method has great potential for joining dissimilar metals, especially those with large difference in melting point. Chen et al. (2009, 2011) investigated laser welding–brazing of Al/Ti dissimilar alloys and claimed the tensile strength of the joints could be as high as 290 MPa since the interfacial reaction was suppressed and only an ultra thin reaction layer formed at the interface. Dharmendra et al. (2011) studied laser welding–brazing of zinc coated steel and aluminum alloy and reported the joint strength reached about 220 MPa with failure on aluminum side after optimizing the process parameter. To the best of the current authors’ knowledge no dissimilar combination of Mg/Zn coated steel by LWB with filler has been reported. Therefore, in this study the characteristics of joining dissimilar metals Mg alloy to Zn coated steel with filler was investigated.

formed at the interface. Therefore, the microstructure adjacent to the interface was not pure Zn, but was composed of ultra-thin Al–Fe intermetallic compound due to hot-dip galvanizing process. Elemental mapping in Fig. 2(c)–(e) confirmed the Fe–Al phase was of uniform thickness and continuity. 2.2. Laser welding–brazing process The experiments were carried out with a 3 kW CO2 laser (ROFIN DC 030 slab laser) with a wavelength of 10.6 ␮m. The laser beam was focused using a 190 mm focal length lens to obtain a spot size of 0.2 mm. Fig. 3 presents the schematic diagram of LWB process equipment and monitoring system. The assembly was fixed in a lap configuration by placing Mg sheet on top of steel sheet. The surface of Mg alloy was cleaned with abrasive paper to remove surface oxides and the steel was cleaned with acetone to remove grease prior to welding. The laser beam was irradiated on the surface of the workpiece at an angle of 90◦ (i.e., vertical). The Mg based filler wire was fed in front of the laser beam. Argon shielding gas

2. Experimental procedure 2.1. Materials Commercially available AZ31B Mg alloy and DP980 steel sheets with the dimension of 100 mm × 30 mm × 1.5 mm were used in this study. A Mg based filler wire with the diameter of 2 mm was used. The chemical compositions of AZ31B Mg alloy, DP980 steel, and filler metal are listed in Table 1. The microstructure of Mg base metal is shown in Fig. 1(a). The AZ31B Mg alloy consisted of equiaxed grains with an average grain size of 4 ␮m; measured according to ASTM E112-96. The steel was hot-dipped Zn galvanized, as shown in Fig. 1(b). The thickness of the zinc layer was found to be approximately 15 ␮m. Fig. 2 shows a SEM micrograph of cross-sectioned Zn coated steel. Line scan analysis of the Zn coating indicated in Fig. 2(b) showed that an Al enriched Fe–Al phase

Fig. 1. (a) Microstructure of AZ31B Mg alloy and (b) thickness of Zn coating on the hot-dip galvanized DP980 steel.

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Fig. 2. Elemental distribution of Zn coating. (a) SEM micrograph of Zn coating, (b) EDS line scan accross Zn coating, (c)–(e) Zn, Al, Fe mapping.

was utilized to prevent oxidation. The angle between the filler wire and the workpiece was adjusted for smooth wire feed. To completely irradiate the filler metal, the laser beam was defocused to increase the spot size. In addition, a CCD camera that operated at a frame frequency of 60 Hz was used to monitor molten liquid flow behavior. An additional diode laser light source (model: FC-W-808, output power: 10–30 W, fiber core diameter: 0.4 mm) with a center wavelength of 808 nm was selected as a backing light to provide sufficient illumination. A band pass optical filter which matched the wavelength of the diode laser light (i.e., 808 nm) was placed in front of the CCD camera. The camera observation angle was about

15◦ with respect to the steel sheet normal and the laser source was at the angle of 30◦ . The process parameters which were kept constant during LWB process included a focus distance of positive 20 mm defocused from steel surface; the ratio of wire feed speed to travel speed was between 4 and 5 (unless specified otherwise); the shielding gas flow rate was 20 l/min; and the angle between the filler wire and the workpiece was 30◦ . 2.3. Analysis methods After LWB, the welded–brazed specimens were cut perpendicular to the travel direction. Standard grinding and polishing sample preparation procedures were used followed by etching with a picric and acetic acid mixture (4.2 g picric acid + 10 ml acetic acid + 10 ml H2 O + 100 ml ethanol). Cross-sectioned joints were observed by optical microscopy (OM). The microstructures and fracture surfaces were analyzed with a scanning electron microscope (SEM) equipped with an energy dispersive X-ray spectrometer (EDS). The tensile-shear strength of the specimens (8 mm width) was evaluated by a tensile tester at a cross head speed of 0.5 mm/min. Shims were used at each end of the specimens to ensure shear loads in the lap joint while minimizing bending or torque of the specimens, as shown in Fig. 4. The contact angle and spreading area measurements provided were averages of three measurements taken at each set of welding conditions. Similarly, the joint strength was obtained via the tensile testing of three specimens and the average was reported with the standard deviation provided via error bars. 3. Results and discussion 3.1. Influence of various process parameters on weld morphology and joint strength

Fig. 3. Schematic diagram of laser welding–brazing process equipment and the monitoring system.

In the LWB process, selecting appropriate process parameters is essential to control appearance and acquire the desired weld quality. The main variables usually include heating mode (i.e., single or dual beam), flux, laser beam offset and travel speed.

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Fig. 4. Schematic of the tensile shear test specimens.

3.1.1. Heating mode (single/dual beam) Heating mode, which determines the thermal gradient distribution, is a key factor in controlling the bead profile. It is especially important with dissimilar materials combinations since the thermal cycle experienced on the brazing side has a strong impact on the wettability of molten filler on the steel surface. Iqbal et al. (2007, 2010) adopted dual beam for lap welding of galvanized steel sheets and found the zinc vapor which usually is involved in the welded area was absent in the dual beam mode. Shi et al. (2010) reported that the dual beam could effectively reduce or avoid the formation of the blowholes in the welded joints by increasing the time for molten material flow in weld pool. Therefore, dual beam was expected to play a significant role in promoting the wettability of molten filler on the sheet surface. Fig. 5 shows the schematic of beam splitting system. A wedge mirror was used to split a CO2 laser beam into two equal-power beams that were arranged side by side. The inter-beam spacing was kept constant at 0.6 mm. Fig. 6 shows representative cross sections of single and dual LWB joints. The laser power employed was varied from 1000 W to 2200 W, while other variables were kept constant with a travel speed of 0.3 m/min and wire feed speed of 1.2 m/min. The offset of the laser beam was 1 mm toward steel material. In single beam mode, smooth weld surface without obvious defects was observed even with the low laser power (P = 1200 W in Fig. 6(a)), suggesting the heat input for melting Mg based filler metal was

Fig. 5. Schematic of beam splitting required for the dual beam laser process.

very low. However, from the cross-sectional macrographs an obvious defect could be observed at the intermediate portion of the seam and steel interface. This phenomenon indicated that local incomplete brazing occurred between molten filler metal and steel, although a high quality defect free surface was obtained. This occurred because the size of the laser spot at the employed defocused distance was 1.52 mm, less than the diameter of filler metal (2 mm). Most of energy from the single beam was used to melt the filler metal leading to insufficient preheating of the steel substrate. After the droplet of filler metal was transferred to the steel surface the Zn coating and filler material reacted with each other

Fig. 6. Single and double laser beam spatial profile and corresponding joint morphology and cross-sections: single beam (a)–(c) and dual beam (d)–(f).

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Fig. 7. CCD images of molten filler: single beam (a)–(d) and dual beam (e)–(h).

producing Mg–Zn phase. The wettability of molten Mg–Zn phase was then restricted and could not spread freely on the steel substrate since the temperature of steel was too low. With increasing the laser power to 1600 W the incomplete brazing problem was significantly improved and the seam width increased correspondingly. However, the pile up of filler metal still existed and incomplete spreading was visible due to rapid cooling rate when utilizing a single beam. In the dual beam welding process incomplete fusion occurred with low laser power (1000 W) at the welded seam instead of seam and steel interface. This could be primarily attributed to the nature of the dual beam. The size of the two beam spots was 1.6 mm at the positive 20 mm defocusing distance from steel surface. The filler wire was at the center of the laser irradiation points. Since the laser beam was split to irradiate a larger area (combined into a size of 2.2 mm), part of heat input was used to preheat the steel substrate leaving less laser energy for melting filler. As a result, insufficient melting occurred at the weld seam. When the laser power increased to 1600–1800 W, incomplete fusion was eliminated and the crosssections showed a smoother transition into the base material due to better wetting compared to that obtained in the single beam mode.

Fig. 7 shows CCD images of liquid filler metal on steel surface for the single and dual beam processes. When the laser power was low (<1200 W) in the single beam mode, the molten filler shrank after transferring to the steel substrate (spreading boundary indicated by arrows). The heat input was only capable of melting the filler without inducing sufficient wetting, which was consistent with the observations in Fig. 6. This condition was improved when the laser power increased to 1800 W. However, the excessive heat input led to severe evaporation, which was identified by the large plasma in Fig. 7(d). Therefore, this resulted in a limitation of the wettability of molten filler on the steel substrate in the single beam mode. From Fig. 7(e)–(h) it was clear that the wetting and spreading ability in the dual beam mode was superior to that in the single beam mode. To clearly illustrate the comparison of the wettability of liquid filler on steel in the single beam and dual beam processing conditions, the contact angle and spreading area as a function of laser power were plotted in Fig. 8 (the contact angle and the seam width are presented in Fig. 3, spreading area = the specimen width × seam width). As shown in Fig. 8(a), with a laser power of 1200 W in both processing conditions the contact angles were similar to each other at 60.9◦ . With increasing laser power the contact angle reduced to 29.4◦ in the dual beam mode compared to 46◦ in the single beam

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Fig. 9. XRD pattern of flux QJ201.

and promote wetting of molten filler was superior to that of the single beam process. With regards to the joint strength, the maximum values in double beam and single beam processes were 147 N/mm and 112 N/mm, respectively.

Fig. 8. Comparison of specimens between single and dual beam processes: (a) contact angle, (b) spreading area and (c) joint strength.

mode. It was noticed in Fig. 8(b) that the spreading area of dual beam increased from 63.6 mm2 to 73.9 mm2 as the laser power increased from 1200 W to 1600 W. With increasing power further the spreading area decreased gradually to 71.4 mm2 . Accordingly, in the case of single beam the spreading area rose from 50.3 mm2 to 68.8 mm2 and then dropped to 65.7 mm2 . This behavior was closely associated with the heat input. The appropriate heat input could promote the wettability of molten filler, while excessive heat input caused catastrophic evaporation because the feed rate of the filler material was kept constant. In addition, this comparative analysis indicated that the ability of dual beam to reduce the contact angle

3.1.2. Flux Another significant factor affecting the welding–brazing characteristics of Mg alloy and Zn coated steel was flux which could inhibit metal oxidation during the LWB process. Moreover, the flux served the purpose of cleaning any contamination left on the brazing surfaces and enhanced the wetting of molten filler onto the steel sheet. In short, the flux facilitated the joining of both similar and dissimilar metals. The addition of flux became more important in the current work due to poor affinity of the immiscible couple Mg and Fe. Flux QJ201 in powder form was prepared in this experiment with a chemical composition of KCl 50 wt.%, LiCl 32 wt.%, NaF 10 wt.% and ZnCl2 8 wt.%. The melting temperature of the flux was between 460 and 620 ◦ C. The XRD pattern is shown in Fig. 9. Ma et al. (2010, 2012) employed the same flux to braze AZ31B sheets and found it could remove the oxide film and prevent the cleaned joint from being oxidized during brazing. Fig. 10 shows the morphologies of Mg and Zn coated steel joined with flux utilizing different laser powers. Besides laser power the other variables were kept constant with a travel speed of 0.3 m/min, wire feeding speed of 1.2 m/min and laser beam offset of 1 mm. It was noticed that the surface become smoother compared to that in Fig. 6. In the cross-section macrographs there was no incomplete brazing which occurred in the single beam case when flux was not used (Fig. 6(a)). It was also observed in Fig. 11 that the flux was molten and flowed laterally in front of the molten filler material. Another advantage of adding flux was in stabilizing the process. The spatter and burning loss in Fig. 11(c) and (d) were not obvious. This was consistent with the joint appearances in Fig. 10(c) and (d). Fig. 12 presents the contact angle and joint strength for joints produced with flux. As shown in Fig. 12(a) the contact angle of joints produced with flux was larger than that of without flux at low power (P < 1500 W) because some laser energy was used to melt flux. However, when the laser power was increased greater than 1600 W, the contact angle reduced sharply since the laser power was sufficient to melt both filler metal and flux. It can be observed in Fig. 12(b) that the mean value of the spreading area without flux was 67 mm2 , which was lower than that with flux (except the case of low power 1200 W). All these results showed that the wettability of liquid filler on steel with flux was superior to that without flux. The Zn coating without protection would be easily oxidized when the temperature increased to 100–250 ◦ C, which hindered the wettability of Mg on steel. With the addition of flux the oxide

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with flux reached 190 N/mm and increased by 29.2% compared to 147 N/mm when joined without flux.

Fig. 10. Joint morphology and cross-section images of specimens produced with flux and increasing laser power: (a) 1200 W, (b) 1400 W, (c) 1600 W and (d) 1800 W.

layer was removed by chemical reaction which produced ideal conditions for increased wettability of Mg on the steel surface. In such a case, the excess deposition of molten filler metal was avoided and the process stability was improved. The maximum joint strength

3.1.3. Laser beam offset Laser beam offset is another crucial factor affecting the weld morphology and joint strength in dissimilar materials joining. The appropriate thermal distribution can be achieved by adjusting the offset to resolve the difference in melting temperatures and thermal conductivity. Miao et al. (2010) found the defect-free joining and maximum joint strength were obtained at 0.6 mm laser offset on the Mg alloy side during laser penetration brazing of Mg to steel. Gao et al. (2011a,b) noticed that the offset played a large role in the joint properties. They found that acceptable joints were only obtained when the laser beam was offset from the edge of the weld seam by 0.2 mm toward the AZ31B side of the joint during butt welding of Mg to Ti with Mg filler wire. While in laser keyhole welding, the optimal range of the offset was found to be 0.3–0.4 mm. Based on these previous studies the investigation of laser beam offset was deemed important in the current work. Fig. 13 presents typical joint morphologies with different laser beam offsets. Constant laser power of 1600 W, constant travel speed of 0.2 m/min and constant wire feeding speed of 1 m/min were also used in these experiments. Also, the specimens were joined with flux. With 0 mm offset (i.e., the laser beam was irradiated on the edge of base metal AZ31B Mg alloy) an acceptable weld profile was obtained. Similar results were also achieved in the case of offset 0.5–1.5 mm. However, when laser beam offset increased to 2 mm, joining between the molten filler and the Mg base metal did not occur. It could also be noted that laser beam offset had little effect on the appearances of the LWB lap joints which was very different from the aforementioned studies. This was because the joints in the cited studies were of butt type configuration. Adjusting the laser beam offset essentially determined the interfacial reaction between dissimilar metals. However, in our experiments the interfacial reaction between molten filler and steel occurred at the brazing side, which was insensitive to the offset of laser beam. In addition, the molten filler and Mg base metal both were Mg–Al–Zn series alloys. They had good affinity and could easily mix once they got in contact with each other (offset 0–1.5 mm). When the distance between molten filler and Mg base metal exceeded the critical threshold (i.e., 2 mm

Fig. 11. CCD images corresponding to Fig. 10.

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Fig. 12. Comparison of joints produced with and without flux: (a) contact angle and (b) spreading area and joint strength.

offset) they failed to bond with each other resulting in the wetting of molten filler only on the steel surface. The CCD images in Fig. 14 showed differences when changing the beam offset. When the offset was 0 mm the molten filler metal was only attracted to AZ31 Mg base metal due to their good affinities. However, when the offset increased to 2 mm some evaporation of the Mg filler material occurred. These results were likely associated with the thermal gradient induced by the different thermal conductivities of the steel (34.3 W m−1 k−1 ) and Mg alloys (145 W m−1 k−1 ). For example, under the same conditions the cooling rate of Mg alloy would be much faster than that of the steel. When the offset was 2 mm toward steel side the molten filler did not come in contact with the Mg base metal, leading to lower thermal dissipation. As a result, the temperature of steel substrate would be higher, suggesting that more energy was available for melting the filler material. Fig. 15 gives the spreading area and resulting joint strength as a function of laser beam offset. It was evident that the joint strength first increased dramatically with increasing laser offset followed by a gradual decrease with offsets greater than 0.5 mm. The joint strength reached a maximum value of 228 N/mm at an offset of 0.5 mm. The optimal range of laser offset was therefore determined to be 0.5–1.0 mm toward the steel side. The main reason for this trend in measured joint strength was that the laser beam offset played an important role in heat input and distribution, influencing the temperature of base metals which affected the wettability of the filler material. The molten filler could not fully wet the steel surface due to superior wetting on the

Fig. 13. Joint morphology and cross-section images of specimens produced with different offsets: (a) 0 mm, (b) 0.5 mm, (c) 1.5 mm and (d) 2 mm.

Mg base metal and insufficient temperatures to cause significant flowing of the molten filler material when the offset was on Mg side (0 mm). On the contrary, when the offset was larger than the desirable value (i.e., 0.5 mm), more energy was available on the steel side causing high temperatures leading to vaporization losses. 3.1.4. Travel speed Travel speed is directly related to heat input. An increase in travel speed gives rise to a decrease in temperature since the laser beam irradiates the work piece for a shorter period of time. Change of travel speed determines the brazing time and affects the formation of intermetallic compounds which have a significant influence on the joint strength of the lap type joint specimens. Fig. 16 shows the joint morphologies at low travel speeds with a constant laser power of 1600 W, a constant laser beam offset of 0.5 mm and a constant ratio of wire feed speed to travel speed of

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Fig. 14. CCD images corresponding to the process conditions presented in Fig. 13.

5. When the travel speed was 0.1 m/min, ablation of liquid filler was observed as shown in Fig. 16(a). The wettability however decreased with increasing travel speed. With a higher travel speed of 0.8 m/min, the molten filler was deposited on the steel surface without adequate wetting. Fig. 17 presents the morphologies of joints at high travel speeds. Note that the ratio of wire feed speed to travel speed was adjusted to 3 when using these higher travel speeds to ensure more laser energy was distributed to preheat the steel substrate. It was observed that cracking occurred following the LWB process when the travel speed was higher than 1 m/min. This indicated a threshold above which, reliable joints were not possible due to insufficient time and therefore process energy to facilitate joining. Fig. 18 shows the corresponding CCD images during LWB with varying travel speeds. As shown in Fig. 18(a) and (b), the flux was molten and the slag flowed to both sides of the seam at low travel speeds. With an increase in travel speed the laser induced plasma was observe to decrease and flux did not flow as was required. When the travel speed was highest at 2 m/min, the flux did not melt. As a result, newly formed liquid Mg–Zn products could not adequately interact with the steel substrate, resulting in a weak reaction layer on the steel substrate.

Fig. 16. Joint morphology and cross-sections of specimens joined using low travel speeds of (a) 0.1 m/min, (b) 0.6 m/min and (c) 0.8 m/min.

Fig. 15. Joint strength and spreading area versus laser beam offset.

Fig. 17. Joint morphology and cross-sections of joints produced with high travel speeds of (a) 1 m/min and (b) 2 m/min.

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Fig. 18. CCD images of the LWB process with travel speeds of (a) 0.1 m/min, (b) 0.6 m/min, (c) 1 m/min and (d) 2 m/min.

Fig. 19 shows the joint strength as a result of varying travel speeds. It was evident that the low travel speed contributed to high joint strength. For instance, the joint strength reached 217 N/mm with a travel speed of 0.1 m/min, whereas it reduced to 114 N/mm as the travel speed increased to 0.8 m/min. When the travel speed was greater than 1 m/min, the joint strength reduced to nearly zero. This trend in joint strength resulted from higher travel speeds leading to lower heat input, thereby reducing the temperature of the steel substrate. As a result the molten Mg–Zn reaction layer formed making it difficult to brazing on the steel side of the joint giving rise to the formation of a crack. To sum up the above investigations, the optimized process parameters are as follows: dual beam mode; with flux; laser power of 1400–1600 W; laser offset of 0.5–1.0 mm toward steel side; low travel speed 0.1–0.4 m/min; and the ratio of wire feeding speed to travel speed between 4 and 5. 3.2. Microstructure morphology Zinc can react with Mg according to Mg–Zn binary diagram. Diverse stoichiometric reaction products including intermetallic

Fig. 19. Joint strength versus LWB travel speed.

compounds (IMC) and Mg–Zn eutectic structure (ES) can be generated. Fig. 20 shows typical morphologies of reaction layers along the interface on the brazing side of the joint (laser power 1600 W, travel speed 0.3 m/min, filler feeding speed 1.2 m/min and beam offset toward steel 1 mm). These different reaction products along the interface could be divided into four zones: (i) seam head zone (zone A); (ii) intermediate zone (zone B); (iii) direct laser irradiation zone (zone C); and (iv) seam root zone (zone D). Heterogeneous interfacial reactions were observed to form from the seam head to the seam root zones of the joint due to differences in thermal distribution. In the seam head zone (zone A) at the toe of lap joint, interfacial reaction layers exhibited diverse morphologies with the thickness of 150–200 ␮m. At the intermediate zone (zone B) the interfacial microstructure was found to appear homogeneous and had a continuous morphology with a thickness of 15–20 ␮m. The thickness of the reaction layer was reduced by a large extent in the direct irradiation zone (zone C). In some cases, slight melting indicated in Fig. 20(c) occurred on the steel surface due to local high power density. With regard to the seam root zone (zone D) in Fig. 20(d), similar to the seam head zone (zone A), a Zn rich area could be found having a thick reaction layer (50–150 ␮m). The identification of phase components at these zones and their formation mechanism have been described in detail elsewhere (Li et al., 2012). Fig. 21 shows the morphology of four zones at the different process conditions investigated in the present work. The thickness of the corresponding reaction layers is presented in Fig. 22. Interestingly, it was found that the microstructural characteristics were similar in each zone. Thick and diverse reaction layers at the seam head (thickness ∼150–250 ␮m) and the seam root (thickness ∼100–220 ␮m) were identified along with uniform reaction layers in the intermediate zone (thickness ∼15–35 ␮m) and relatively thin layers in the direct laser irradiation zone (thickness ∼2–10 ␮m). The thickness of the reaction layer fluctuated within a certain scope when applying different process parameters. All of the above observations further confirmed that interfacial reaction was not associated with the heat input, which differed from previous studies (Dharmendra et al., 2011; Zhao et al., 2011). In previous studies it was found that the thickness of the interfacial reaction layer did depend on the heat input. Dharmendra et al.

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Fig. 20. Typical morphology of reaction layers along the joint interface: (a) seam head (zone A), (b) intermediate zone (zone B), (c) direct laser irradiation zone (zone C) and (d) seam root (zone D).

(2011) reported that the thickness of the layer was varied with the brazing speed. Zhao et al. (2011) concluded that the thickness of intermetallic reaction layer increased with increasing laser power at a constant travel speed. The main reason for the reaction layer thickness not depending on heat input in this study may be that the thickness of Zn coating dominated in the formation of reaction layers. The thin Zn coating (15–20 ␮m) used in this study could become rapidly molten once a droplet of liquid filler hits the steel surface. Moreover, good wettability of the Zn coating induced simultaneous flow with the liquid filler, leading to a Zn rich region running from the seam head to seam root. Therefore, the thickness of these reaction layers only varied slightly with changing of process conditions. 3.3. Failure mechanism Fig. 23 shows the fracture path and fracture surface morphologies of the Mg/Zn coated steel lap joint. The failure mode exhibited typical interfacial fracture shown in Fig. 23(a). Fig. 23(b) shows the high magnification image of the area indicated in Fig. 23(a). This high magnification image revealed that the crack propagated along the Mg–Zn reaction layer and steel interface, instead of the reaction layer and seam interface. This was an indication that the cohesion between the seam and the reaction layer was stronger than that between the reaction layers and steel. The fracture surfaces on both the Mg side and steel side shown in Fig. 23(c) were quite smooth suggesting brittle fracture. Therefore, it was concluded that the joint between the reaction layer and steel was the weakest region of the joint. To further understand the failure mechanism the fracture surfaces of Mg/Zn coated steel samples were also examined by SEM at high magnification. The fracture surface of Mg side (Fig. 23(d)) exhibited distinct lamellar structures corresponding to the eutectic structure indicated in Fig. 21. Examination of

the fracture morphology of the steel side revealed the presence of a significant amount of particles. EDS analysis shown in Fig. 24(a) at point 1 consisted mainly of 34.55 at% Al, 56.59 at% Fe, implying it was an Fe–Al phase. This finding was consistent with the Al rich interface between the Zn coating and the steal in the base material indicated in Fig. 2. Therefore, it could be inferred that the Fe–Al phase produced during the hot-dipped process still existed after the LWB process due to its high melting point (Fe2 Al5 –1169 ◦ C, FeAl–1160 ◦ C). These Fe–Al phase particle hindered further metallurgical bonding between the Mg–Zn reaction layer and steel substrate; greatly deteriorating the joint strength. This was confirmed by the low amount of residual Mg–Zn eutectic structure adhering on the steel side of the joint (Fig. 23(e)) and its chemical composition was primarily composed of 67.06 at% Mg and 23.73 at% Zn according to the EDS analysis (Fig. 24(b)). Note that the role of Fe–Al phase found in this study has significant differences from previous studies (Chen et al., 2009; Wahba and Katayama, 2012). In past studies the Fe–Al intermetallic compounds were absent before the joining process. During the interfacial reaction Al as an alloying element diffused from Mg–Al–Zn series alloy to react with Fe producing Fe–Al phase. The newly formed Fe–Al phase promoted metallurgical bonding at the interface, which therefore enhanced the joint strength. However, in our experiments the Fe–Al phase was already present within Zn coating before the LWB process (Fig. 2). During the interfacial reaction the Zn could melt and react with Mg forming Mg–Zn structure. Mg–Zn reaction layer did not dissolve the Fe–Al intermetallic compounds because Fe–Al was not soluble in Mg and Zn. In this case, no metallurgical bonding occurred between the Mg–Zn structure and Fe–Al phase, thereby hindering joint strength. Based on the above analysis, the mechanisms behind the formation of the interfacial layer and failure could be identified. To aid the following discussion a schematic diagram is presented in Fig. 25.

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Fig. 21. Typical morphologies of reaction layer with different processing conditions.

Fig. 22. Thickness of reaction layer measured along the interface of joints produced with various process parameters.

Firstly, the thin Zn coating melted rapidly due to laser irradiation and conduction from the liquid filler metal. Since the melting points of Zn (420 ◦ C) and Mg (650 ◦ C) were low, the required heat input was also low during LWB process. Thus, the Fe–Al phase adjacent to the steel substrate remained after LWB due to its relatively high melting point (Fe2 Al5 –1169 ◦ C, FeAl–1160 ◦ C). With the continuous deposition of the molten filler material, the liquid front flowed while carrying liquid Zn from the seam head to the seam root leading to a Zn rich region (Fig. 25(a)). Zhang et al. (2007) made similar observations with Al/Zn coated steel joint using a metal inert gas welding–brazing process. As the process proceeded part of base metal Mg alloy became molten and mixed with filler metal creating the fusion side of the joint. At the brazing side of the joint the liquid Zn and liquid Mg dissolved creating a liquid Zn–Mg solution. The dissolution and diffusion process was quick and intense since both Zn and Mg were in the liquid state. Upon cooling, solidification occurred at Mg fusion welding side followed by the brazing side. At this point the Mg–Zn reaction layer formed along the interface

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Fig. 23. Crack orientation and fracture surface morphologies: (a) overview of the interfacial failure, (b) crack orientation at high magnification, (c) fracture surface of both Mg and steel sides of joint at low magnification, (d) enlargement of Mg side and (e) enlargement of steel side.

between welded seam and the steel. However, liquid Zn could not effectively react with steel substrate due to the separation of Fe–Al phase resulting in the weak joint between the reaction layer and steel substrate. During tensile-shear testing the crack initiated at the seam head and the seam root. After initiation, the crack

propagated along the Mg–Zn reaction layer and Fe–Al phase interface, owing to inadequate metallurgical bonding between the two layers. Thus, a large amount of Fe–Al particles could be observed at the fracture surface of the steel side and only low amounts of Mg–Zn reaction layer were identified.

Fig. 24. EDS analysis results taken at points P1 and P2 in Fig. 23(e): (a) P1 and (b) P2.

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Fig. 25. Schematic of reaction and failure mechanism: (a) enrichment with Zn, (b) interfacial reaction layer formation and (c) interfacial fracture (not to scale).

4. Conclusions In this study the influence of LWB process parameters on the joining of AZ31B Mg alloy to Zn coated steel was investigated. The interfacial microstructure and failure mechanisms were identified and discussed. The main findings of this work include: (1) The optimized process parameters for maximum joint strength include: dual beam mode, use of flux, laser power of 1400–1600 W, laser offset of 0.5–1.0 mm toward steel side, low travel speed 0.1–0.4 m/min, and a ratio of wire feeding speed to travel speed of 4–5. (2) Heterogeneous interfacial reaction layers formed along the interface between weld seam and steel. Four zones with different thickness were identified according to their microstructural characterization. The original thin Zn coating (15–20 ␮m) restricted the reaction process and was found to be responsible for the variation in microstructures identified. (3) The primary failure mode of the lap specimens was interfacial fracture. Cracks propagated along the Mg–Zn reaction layer and steel interface. The pre-existing Fe–Al phase formed during the hot-dip galvanization process suppressed the metallurgical bonding of Mg–Zn reaction layer and steel substrate which was primarily responsible for the interfacial failure. Acknowledgements The authors want to thank Dr. Dejian Liu, Huazhong University of Science and Technology, P.R. China, Dr. Lei Liu, Center for Advanced Materials Joining, University of Waterloo, Canada, for their helpful suggestions and discussions. We also gratefully acknowledge Dr. Andie Pequegnat, Center for Advanced Materials

Joining, University of Waterloo, Canada, for his help in English revision of this article.

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