Improving weld strength of magnesium to aluminium dissimilar joints ...

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aluminium dissimilar joints via tin interlayer during ultrasonic spot welding. V. K. Patel, S. D. Bhole* and D. L. Chen. Welding of magnesium to aluminium alloys ...
Improving weld strength of magnesium to aluminium dissimilar joints via tin interlayer during ultrasonic spot welding V. K. Patel, S. D. Bhole* and D. L. Chen Welding of magnesium to aluminium alloys is enormously challenging due to the formation of brittle Al12Mg17 intermetallic compounds (IMCs). This study was aimed at improving the strength of dissimilar joints of AZ31B-H24 magnesium alloy to 5754-O aluminium alloy by using a tin interlayer inserted in between the faying surfaces during ultrasonic spot welding. The addition of tin interlayer was observed to successfully eliminate the brittle Al12Mg17 IMCs, which were replaced by a layer of composite-like tin and Mg2Sn structure. Failure during the tensile lap shear tests occurred through the interior of the blended interlayer as revealed by X-ray diffraction and SEM observations. As a result, the addition of a tin interlayer resulted in a significant improvement in both joint strength and failure energy of magnesium to aluminium dissimilar joints and also led to an energy saving because the optimal welding energy required to achieve the highest strength decreased from y1250 to y1000 J. Keywords: Magnesium alloy, Aluminium alloy, Welding, Intermetallic compounds, Interlayer

Introduction In order to achieve a good combination of the properties of Mg and Al alloys, owing to their lightweight nature, high strength/weight ratio, good castability and workability,1–5 the development of reliable joints between Mg and Al alloys is required. Fusion welding of Mg and Al alloys always produces coarse grains and brittle intermetallic compounds (IMCs) in the weld metal. A number of papers have recently reported on the joining of dissimilar alloys by friction stir welding, friction stir spot welding, laser welding and resistance spot welding (RSW), especially for the lightweight and Al alloys.1–6 However, very little work has been reported for joining Mg and Al using ultrasonic spot welding (USW),which is a solid state joining process that produces coalescence through the simultaneous application of localized high frequency vibratory energy and moderate clamping forces.7,8 Previous studies showed that the formation of IMCs of Al12Mg17 and Al3Mg2 is a major disadvantage in joining Mg to Al alloys9 because the mechanical properties of the welded joints are known to be closely related to the microstructure of the intermediate layer.10 These phases are brittle and lead to fracture, making it difficult to obtain a non-brittle and strong joint between Mg and Al alloys. Therefore, the aim of this research was to improve the joint strength in dissimilar USW of

Department of Mechanical and Industrial Engineering, Ryerson University, 350 Victoria Street, Toronto, Ont. M5B 2K3, Canada *Corresponding author, email [email protected]

AZ31B-H24 with Al5754-O alloys by reducing the tendency for the formation of a brittle intermetallic layer using a tin (Sn) interlayer that could interact with both Mg and Al, as seen from their respective binary diagrams.11,12 The selection of Sn in the present study was also based on the findings that Sn improved the wettability of Mg and Al alloys during the welding process13 and also refined the grain size in the Mg alloy.14,15

Experimental In the present study, a commercial 2 mm thick sheet of AZ31B-H24 (Mg–3Al–1Zn–0?6Mn–0?005Ni–0?005Fe) Mg alloy and a 1?5 mm thick sheet of Al5754-O (Al– 3?42Mg–0?63Mn–0?23Sc–0?22Zr) Al alloy provided by General Motors Company were selected for USW. The specimens, which were 80 mm long and 15 mm wide, were sheared, with the faying surfaces ground using 120 emery papers, and then washed using acetone followed by ethanol and dried before welding. During welding, a 50 mm thick pure Sn interlayer was placed in between the workpiece of the Mg/Al weld samples. The welding was done with a dual wedge reed Sonobond-MH2016 HPUSW system. The samples were welded at an energy input ranging from 500 to 3000 J at a constant power setting of 2000 W, an impedance setting of 8 on the machine and a pressure of 0?414 MPa. Four samples were welded in each welding condition. Two of them were used for microstructural examination and microhardness tests, and the other two were used for the lap shear tensile tests. Cross-sectional samples for scanning electron microscopy were polished using diamond paste

ß 2012 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 23 January 2012; accepted 18 February 2012 DOI 10.1179/1362171812Y.0000000013

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1 Microstructure of dissimilar USW Mg/Al joints made with welding energy of 1500 J in middle of weld nugget a without (Ref. 9) and b with Sn interlayer

and MasterPrep. A computerised Buehler microhardness testing machine was used for the microindentation hardness tests diagonally across the welded joints using a load of 100 g for 15 s except for the thin IMC interlayer (,10 mm), where a load of 10 g was used for 15 s. The mean value of three indentations along the IMC interlayer was taken for better accuracy with the low indentation load of 10 g. To evaluate the mechanical strength of the joints and establish the optimum welding conditions, the researchers conducted tensile shear tests of the welds to measure the lap shear failure load using a fully computerised United testing machine with a constant crosshead speed of 1 mm min21 in air at room temperature. In the tensile lap shear testing, restraining shims or spacers were used to minimise the rotation of the joints and maintain the shear loading as long as possible. X-ray diffraction was carried out on both matching surfaces of Al and Mg sides after tensile shear tests using Cu Ka radiation at 45 kV and 40 mA. The diffraction angle (2h) at which the X-rays hit the samples varied from 20 to 100u, with a step size of 0?05u and 2 s in each step.

Results and discussion Microstructural evaluation Microstructural characterisation was conducted across the weld line of the samples. Figure 1a and b shows the microstructures at the centre of the weld nugget of the USW welded Mg/Al joints characteristic of the whole weld nugget without and with a Sn interlayer respectively. A sound joint was obtained under most of the welding conditions because no large defects were present, such as crack or tunnel type defects. It can be seen

2 a microstructure of dissimilar USW Mg/Al joints made with Sn interlayer at welding energy of 1500 J in middle of weld nugget at higher magnification; b EDS line scan across welded joint as indicated by dashed line

from Fig. 1a that there was a heterogeneously distributed IMC layer between the Mg and Al alloy sheets. In our previous study9 of the USW of Mg/Al alloys without a Sn interlayer, the non-uniform IMC layer had a solidified microstructure containing the brittle phase through the eutectic reaction, liquidRAl12Mg17zMg. In the USW of Mg/Al with an Sn interlayer, the IMCs of Al12Mg17 were not observed as the Al and Mg were separated by the Sn interlayer (thus, the ternary Al–Mg– Sn diagram was not used). The IMC layer displayed a composite-like eutectic structure, as shown in Fig. 2 at the centre of the weld nugget, where the Sn containing white fine particles were distributed homogeneously or as a network in the interlayer. This was roughly estimated via both energy dispersive X-ray spectroscopy (EDS) point analysis and line scan (Fig. 2). The chemical composition (in at. wt-%) at points A and B was 62?1Mg–36?0Al–1?9Sn and 64?9Mg–19?9Al–15?2Sn respectively, which suggests that the dark area (A) had less Sn than the light area (B). Figure 2b shows that the concentration of Al in the nugget zone is less than that of Mg, owing to the higher solubility of Sn in Mg compared to Sn in Al. Therefore, these results (Fig. 2a and b), in conjunction with the Mg–Sn phase diagram,11 suggest the presence of a Mg2Sn phase, where the

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3 Matching fracture surfaces of a Al side and b Mg side of welded sample with Sn interlayer at welding energy of 1000 J

eutectic structure consisting of b-Sn (or Mg–Sn solid solution) and Mg2Sn would occur at a temperature of 203uC.11 In the USW, the simultaneous application of localised high frequency vibratory energy and moderate clamping force leads to relative motion and friction heat at the interfaces7,8 between Al–Sn and Mg–Sn, which would cause the melting and coalescence of Sn. In the presence of the Sn interlayer in the USW, Al and Sn combine to form solid solutions, whereas Mg and Sn combine to form b-Sn and Mg2Sn IMCs. The Mg2Sn phase has an antifluorite type (CaF2) AB2 crystal structure with a moderately high melting temperature of 770uC and a lattice parameter of a50?676 nm.16 It is apparent that the large Mg2Sn particles resulted from the eutectic reaction (LRb-SnzMg2Sn). The formation of the IMCs of Mg2Sn implies that the exothermic chemical reaction MgzSnRMg2SnzQ must occur in the interlayer, where Q is the mole heat release of Mg2Sn, which is 280?75 kJ mol21 (Ref. 17), compared to the fusion heat of Al and Mg of 10 (Ref. 18) and 9 kJ mol21 (Ref. 19) respectively. Thus, it is possible that the large particles of Al could be remelted by the heat formation of Mg2Sn and resolidify as smaller particles or disappear under the condition of rapid cooling during the welding process. The addition of Sn to the lap joint was observed to refine the grain size in the fusion zone and the base Mg alloy14,15 due to the presence of a eutectic structure that restricts the growth of the Mg grains. Furthermore, it also improves the wettability of Mg and Al during the welding process.13 Thus, the surface tension of the liquid is reduced, so more liquid spreads evenly over the surface of the base metal. Figure 3a and b shows the matching fracture surface morphology of the Al and Mg sides at an energy input of 1000 J. It can be seen that on both Al and Mg sides, there was less Sn in the centre compared to the edge of the nugget zone owing to the squeeze-out caused by the

high temperature and moderate pressure generated during the USW,19 hence the heterogeneity of the thickness of the Sn interlayer, as shown in Fig. 1b. This resulted in a fracture appearance on both Al and Mg surfaces consisting of protuberances and cleavagelike features typical of brittle fracture mode in regions with low Sn (3?4 at. wt-% as determined by EDS analysis), as seen in region D in Fig. 3a and in region F in Fig. 3b. On the other hand, in the high Sn regions (12?4 at. wt-% according to EDS analysis), a fairly ductile and rough fracture surface appearance was seen, as shown in region C in Fig. 3a and in region E in Fig. 3b, suggesting that the Sn and Mg2Sn eutectic structure (Fig. 2a) played a significant role in strengthening the microstructure of the Sn interlayer joint and thus improved its strength, as will be presented later. To further verify the above microstructural observations, the X-ray diffraction patterns obtained on both matching surfaces of Al and Mg sides after tensile shear tests are shown in Fig. 4a and b. It is clear that, apart from strong peaks of Al on the Al side (Fig. 4a) and strong peaks of Mg on the Mg side (Fig. 4b), both Sn and Mg2Sn appeared on both sides without the presence of Al12Mg17 anymore. Thus, the addition of the Sn interlayer led to an effective elimination of the brittle IMCs of Al12Mg17, which was replaced by the formation of solid solutions of Sn–Mg and Sn–Al, as well as the eutectic SnzMg2Sn structure. This is in agreement with the SEM observation (Fig. 2a) and EDS analysis (Fig. 2b). Furthermore, these microstructural features of interlayer led to the tensile lap shear failure occurring predominantly through the interlayer (Fig. 3) rather than the debonding between the IMCs of Al12Mg17 and Al side.9

Mechanical properties Unlike RSW, USW displays no clearly discernible fusion zone or heat affected zone that can degrade the strength

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5 Microhardness profile across the welded joints at different energy inputs

4 X-ray diffraction patterns obtained from matching fracture surfaces of a Al side and b Mg side at welding energy of 1000 J

of the metals being joined.7,8 It can be inferred that in the USW, the heat affected zone would be small because of the welding temperature being lower than the melting points of Al and Mg alloys and the lower conductivity of the welding tip (steel) compared to the thermal conductivity of the Al and Mg alloy samples. The hardness profile across the welded joint diagonally is shown in Fig. 5. The hardness of the very thin IMC layer is the mean of measurements at three different locations along the length of the IMC layer. It is seen that the hardness decreased with increasing energy input owing to increasing grain size with increasing temperature. The interlayer between Mg and Al sheets exhibited hardness values between 200 and 300 HV due to the presence of Mg2Sn. As shown in Fig. 6a, the maximum lap shear strength of Mg/Al USW joints with the addition of an Sn interlayer was improved (y41 MPa) when compared to that of direct Mg/Al USW joints (y35 MPa). In addition, the Sn interlayer also led to an energy saving because the optimal welding energy required to achieve the highest strength decreased from y1250 to y1000 J.

The failure energy shows a similar trend, as shown in Fig. 6b, where the maximum failure energy improvement due to the Sn interlayer was y500 N mm (from y1200 to y1700 N mm). The increase in the lap shear strength and failure energy was attributed to the formation of solid solutions of Sn with Mg and Al and the composite-like eutectic structure of Sn and Mg2Sn (Figs. 2 and 4), instead of the brittle IMCs of Al12Mg17 as reported in Ref. 9. In addition, it can be seen from Fig. 6a and b that in the absence of the Sn interlayer, the lap shear fracture load and failure energy increased with increasing energy inputs and reached the maximum at 1250 J energy input and then decreased. This was a consequence of the competition between the increasing diffusion bonding arising from higher temperatures at higher energy inputs and the deteriorating effect of the brittle intermetallic layer of increasing thicknesses.9 Similarly, in the presence of the Sn interlayer, at lower energy inputs, the temperature was not high enough to soften or melt the Sn interlayer. On the other hand, at higher energy inputs, the weld specimen was subjected to higher temperatures under larger vibration amplitude for a longer time, resulting in more of the Sn interlayer being squeezed out. As a result, the lap shear strength and failure energy of the Mg/Al USW joints with an Sn interlayer also increased initially with increasing welding energy, reached the maximum values, followed by a decrease with further increasing welding energy. Figure 6c shows a summary of the maximum lap shear strength for different types of joints made at a welding energy of 1000 J. It is seen that the USW Mg/Al joint without an Sn interlayer was ,25% lower than the USW Al/Al joint and 60% lower than the USW Mg/Mg joint. However, the USW Mg/Al joint with an Sn interlayer was ,5% higher than the USW Al/Al joint and only 40% lower than the USW Mg/Mg joint. A further comparison of the lap shear strength of dissimilar Mg/Al joints produced with different joining techniques available in the literature1–5 is shown in Fig. 6d. It should be noted that the measured lap shear fracture loads of USW Mg/Al joints were converted to the lap shear strengths by using a nominal welding area corresponding to the area of the USW tip (866 mm). Likewise, the shoulder diameter and the product of the shoulder diameter and length of the tool for friction stir spot

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6 Comparison of a lap shear strength with and without an Sn interlayer as a function of energy input, b failure energy with and without an Sn interlayer as a function of energy input, c lap shear strength among different welded joints at an energy input of 1000 J and d lap shear strength of the dissimilar Mg/Al welded joints made with different welding techniques

welding and friction stir welding and the nugget diameter for RSW have been used to calculate the areas of the respective nugget zones. It is seen that USW joints with the addition of Sn interlayer displayed better lap shear strengths compared with the other techniques shown. However, it must be emphasised that there were differences in the thickness of the samples and that different Mg and Al alloys were used. More detailed studies are needed to identify the optimised processing parameters in different welding techniques.

Conclusions Ultrasonic spot welding of AZ31B-H24 Mg alloy sheet to Al5754-O Al alloy sheet with an Sn interlayer was successfully achieved. The maximum lap shear strength and failure energy of Mg to Al with an Sn interlayer were achieved at a welding energy of ,1000 J, which was significantly higher than those of an Mg to Al joint without an interlayer. This improvement was attributed to the formation of solid solutions of Sn with Mg and Al as well as the composite-like Sn and Mg2Sn eutectic structure in the interlayer instead of the brittle Al12Mg17 IMC present in Mg to Al joints without an interlayer. The fact that Sn and Mg2Sn were located on both Al

and Mg sides of matching fracture surfaces indicated that the tensile shear failure occurred through the interior of the interlayer. The presence of Sn on the fracture surfaces reduced the tendency for brittle cleavage-like fracture and encouraged fairly ductile type fracture behaviour with rougher fracture surface. Compared with other welding techniques, the dissimilar USW of Mg/Al alloys with an Sn interlayer displayed better lap shear strengths. Further research is needed to improve the strength of dissimilar USW joints using other suitable interlayers.

Acknowledgements The authors would like to thank the Natural Sciences and Engineering Research Council of Canada (NSERC) and AUTO21 Network of Centers of Excellence for providing financial support. This investigation involves part of the Canada–China–USA Collaborative Research Project on the Magnesium Front End Research and Development (MFERD). The authors thank Dr A. A. Luo, General Motors Research and Development Center, for the supply of test materials. One of the authors (D. L. Chen) is grateful for the financial support given by the Premier’s Research Excellence Award

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(PREA), NSERC-Discovery Accelerator Supplement (DAS) Award, Canada Foundation for Innovation (CFI) and the Ryerson Research Chair (RRC) program. The assistance of Q. Li, A. Machin, J. Amankrah and R. Churaman in performing the experiments is gratefully acknowledged.

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