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On the joining of steel and aluminium by means of a new friction melt bonding process
Available online at www.sciencedirect.com
ScienceDirect Scripta Materialia 77 (2014) 25–28 www.elsevier.com/locate/scriptamat
On the joining of steel and aluminium by means of a new friction melt bonding process Camille van der Rest, Pascal J. Jacques⇑ and Aude Simar Universite´ catholique de Louvain (UCL), Institute of Mechanics, Materials and Civil Engineering, IMAP, Place Sainte Barbe 2, B-1348 Louvain-la-Neuve, Belgium Received 26 November 2013; revised 7 January 2014; accepted 8 January 2014 Available online 15 January 2014
A new lap welding process for dissimilar materials has been developed and tested for steel-to-aluminium joints. A simple cylindrical tool is rotated and translated over the top steel plate, leading to a transient partial melting of the bottom aluminium plate and the formation of intermetallic reaction layers of FeAl3 and Fe2Al5 as thin as 2.5 lm. Lap shear tests of the welds show very good resistance with fracture mostly in the base materials. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Friction stir welding; Dissimilar welding; Aluminium alloys; Steels; Mechanical properties
Assembling dissimilar materials within integrated structures is nowadays a stringent requirement related to improved performance. It is particularly the case for the transportation industry that needs to reduce the weight of structures while still improving their strength and crashworthiness. While efforts are carried out to continuously improve the mechanical properties of new steel grades up to levels unreachable by other materials, as in the case of advanced high-strength steels, Al alloys also present very interesting levels of specific strength that easily ensure lightweight structures, as demonstrated by their ever-increasing use in transport applications. Consequently, the need for an efficient technology to join steel and Al alloys arises in the automotive industry to ensure profitability and lightness of the vehicles in addition to efficient sealing of the joints. The main difficulty in joining steel and Al alloys results from the large difference in their respective range of solid-state stability, but also from their large reactivity leading to the formation of brittle intermetallic layers at their interface [1–4]. As a consequence, conventional fusion welding, such as arc and laser welding for which both materials melt, are unsuitable for joining steel to Al alloys since the temperature that is reached causes inter-
⇑ Corresponding
author. Tel.: +32 10472432; fax: +32 10474028; e-mail: [email protected]
diffusion and growth of large intermetallic reaction layers resulting in a brittle cast structure with poor mechanical properties [5,6]. Among the various solutions that have been proposed to reduce the heat input and thus to control the interface reactivity [7], solid-state processes are particularly efficient for welding dissimilar materials since the formation of an intermetallic layer at the interface may be avoided or reduced to acceptable values [2]. In diffusion bonding, the materials to be joined are put in intimate contact for hours under pressure and at a temperature below 70% of the lowest melting point. A transient liquid may sometimes be present owing to a eutectic reaction between the materials being joined [8]. However, the native Al oxide layer of Al alloys acts as a diffusion barrier and inhibits bond formation [2]. In order to break this oxide layer and activate the surface, the friction welding process is based on the movement of one piece relative to the other under compressive axial load, thus also involving shorter reaction times and higher temperatures (near the lower melting point) [3]. However, friction welding is mostly limited to rods and tubes since the relative movement is obtained by rotation. Friction stir welding (FSW) is another solid-state joining process which has been widely investigated for welding dissimilar materials, in particular Al alloys and steel, in a butt-joint [1,5,9,10] or lap-joint [11,12] configuration. Some studies have also considered a third material
http://dx.doi.org/10.1016/j.scriptamat.2014.01.008 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
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C. van der Rest et al. / Scripta Materialia 77 (2014) 25–28
with a lower melting point, such as Zn, to act as a filler metal between Al and steel [13] or to help in reducing the thickness of the intermetallic layer at the interface [14]. Although FSW is considered a solid-state technique, Chen et al. observed that, when working with materials with very different melting points (Al 6061 and AISI 1018 steel in their case), one should locally reach the lowest melting point, leading to a method that combines the effects of both fusion and solid-state welding [1]. The present study aims at reporting a new welding technique [15] suitable for assembling dissimilar metallic materials, such as Al and steel sheets, thus circumventing the drawbacks of previous techniques. After stacking and clamping the metallic sheets, with the high-melting point material on top, a non-consumable flat cylindrical tool is rotated and pressed on the upper surface of the stack as shown on Figure 1. The rotating tool then advances along the surface at a predetermined speed. Frictional heat is generated due to the intimate contact between the pressed rotating tool and the top sheet. In addition, plastic deformation occurs below the rotating tool, also contributing to the temperature increase [16]. The generated heat is transmitted to the bottom sheet through the top sheet, and if it is sufficient, the low-melting-point bottom sheet partially melts, forming a localized pool of liquid metal in contact with the upper sheet. In contrast to conventional fusion welding processes, the melting temperature of the low-melting-point material is only exceeded slightly and very locally, resulting in a semi-solid welding process. The technique is easily implemented since it does not require any shielding gas, unlike laser welding [6], or specific preparation of the metallic sheets. In addition, in comparison with conventional FSW, the pin-less tool results in less abrasion of the tool and the absence of a keyhole at the end of the weld. Another benefit of this new technique is that the advancing speed is significantly higher than in other Fe–Al welding processes [7,11–14] and one could even foresee speeds as high as 1 m min 1 being attained. Our process could also be considered for welding other combinations of metallic materials with different melting points. In order to demonstrate the sequence of reactions while avoiding complex metallurgical issues, simple steel and Al grades were first tested. A microalloyed ultralowcarbon (ULC) steel with the composition (wt.%) 0.013C, 0.136Mn, 0.01Si, 0.005S, 0.0132P, 0.002Ti, 0.003Nb,<0.0008N was used as the high-melting-point
Figure 1. Schematic of the friction melt bonding process developed to assemble metallic sheets with different melting temperatures.
material. Cold-rolled and annealed sheets 0.8 mm thick with a mean grain size of 13 lm were used. 1 mm thick Al 1050 or 2 mm thick Al 2024 T3 sheets were used as the low-melting-point material. Sheets were 250 mm long and 80 mm wide. The surfaces were cleaned with acetone prior to stacking and welding. A tungsten carbide tool 16 mm in diameter was tilted backwards by 0.5°. The welding machine was controlled in displacement while the rotation rate was set at 2000 rpm for the whole set of experiments. The advancing speed ranged from 100 to 700 mm min 1. After welding, the joints were cross-sectioned by electrical discharge machining perpendicular to the welding direction for metallographic analyses and tensile testing. The transverse sections of the joints were observed by scanning electron microscopy (SEM) with back-scattered electrons (BSEs). In addition, the structure of the joints as well as the diffusion processes were characterixed by energy-dispersive X-ray spectroscopy (EDX) and electron backscatter diffraction (EBSD). Transverse lap shear tests were performed on rectangular specimens with a width of 10 mm, at room temperature and with a crosshead speed of 1 mm min 1. The recorded loads were normalized by the initial transverse section area of the sample at the fracture location, i.e. Al or steel sheets or lap-joints. SEM and EDX were also carried out on the fracture surfaces. The typical microstructure of the joints is illustrated in Figure 2 in the case of the Al 1050–ULC couple welded at 400 mm min 1. Three zones can be distinguished on the SEM micrographs, the Al and steel sheets being separated by a reactive layer. Close to the
Figure 2. Typical band contrast EBSD map of the Al 1050–steel interface (advancing speed = 400 mm min 1) with partial colouring as a function of the indexing phases (face-centred cubic Al in orange, monoclinic FeAl3 in blue, orthorhombic Fe2Al5 in red and bodycentred cubic Fe in green) which highlights the presence of two IMLs at the interface. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
C. van der Rest et al. / Scripta Materialia 77 (2014) 25–28
interface, the Al presents a solidification microstructure with a dendritic structure separated by an eutectic structure containing Fe-rich precipitates. Some large precipitates were indexed by EBSD (see Fig. 2) as FeAl3, in agreement with Bouche´ et al. [17] who had previously reported Fe enrichment of the liquid Al close to the solid– liquid interface. Furthermore, Bouayad et al. reported that the occurrence of these precipitates is favoured when the interaction time and the cooling rate are increased [18]. The steel sheet presents a gradient of microstructure through its thickness with a refined grain size in a layer close to the tool and coarser grains close to the Al sheet. In this thermomechanically affected zone, the grain size ranges from 1 to 35 lm. As a consequence of the melting of the Al, an intermetallic reaction layer (IML) is formed at the interface with the steel sheet due to the reactivity between the molten Al and the solid steel. For all advancing speeds investigated, two IMLs were indexed by EBSD. Figure 2 first shows that both FeAl3 on the Al side and Fe2Al5 on the steel side are formed at the interface. These conclusions were validated by BSE SEM micrographs and EDX measurements, even though the different intermetallics have similar chemical compositions requiring accurate measurements. Figure 2 also shows that nanosized grains constitute the FeAl3 layer, while grains of Fe2Al5 are tongue-shaped and larger. These results are consistent with previous studies [4,17–19] that also reported that the thickness of the IML increases when the contact time between liquid Al and steel increases, causing growth of an Fe2Al5 IML governed by interdiffusion (parabolic growth). Shahverdi et al. also highlighted that the kinetics of formation of Fe2Al5 is faster than for FeAl3, which in addition breaks and dissolves within molten Al [19]. As a consequence, Fe2Al5 is generally reported as the major part of the reaction layer. Figure 3 shows that the thickness of the IML (considering both layers as a whole) decreases similarly for the two Al alloys when increasing the tool advancing speed, which corresponds to a decrease in the heat input during the welding process. Indeed, faster advancing speeds involve locally shorter reaction times. This intuitive result was already observed for other welding techniques, such
Figure 3. Evolution of the thickness of the IML as a function of the tool advancing speed for both Al 2024 (plain symbols) and Al 1050 (open symbols)–ULC steel joints. Error bars provided when available.
27
as friction stir diffusion bonding [9] and friction stir brazing [13]. More importantly, the global level of thickness of the IML remains small, going down to 2.5 lm at 700 mm min 1. Furthermore, Fe2Al5 is the major part of the reaction layer except for very fast advancing speeds for which the Fe2Al5 and FeAl3 layers are of similar thicknesses. Indeed, the thickness of the Al-rich FeAl3 IML is almost constant (800 nm–1.5 lm) regardless of the advancing speed (and thus heat input) while the thickness of the Fe2Al5 layer is highly dependent on the heat input . In addition to the IML thickness, the tool advancing speed also influences to a large extent the Al resolidification process through the resulting cooling rates. In the case of the Al 2024 grade, porosities several microns in size appear in the resolidified Al, 15–20 lm below the IML, when the advancing speed is >400 mm min 1 (see Fig. 4c). This zone corresponds to the last liquid Al that solidifies. This phenomenon was confirmed by EDX analyses that showed larger concentrations of the Al alloying elements (Cu and Mg) close to the porosities, in agreement with the solidus on the phase diagrams [20]. In addition, SEM observations revealed that the quantity and size of these porosities increase when the advancing speed increases. Figure 4a presents the evolution as a function of the tool advancing speed of the engineering strength at the onset of necking of transverse lap shear tests. Three sets of constant levels of strength with the advancing speed can be observed that are related to the location of the necking and fracture; these levels depend on the composition of Al alloy. On the one hand, in the case of the 1 mm thick 1050 Al–ULC steel joints, fracture occurs in the Al sheet at 4.5–5 mm from the centre of the joint, namely in the heat-affected zone. Whatever the advancing speed and the resulting IML thickness, the maximum engineering strengths are similar with a mean value of 76 MPa. The Al 1050 base metal exhibits an ultimate tensile strength (UTS) of 100 MPa, while heat treatments at 620 °C for 3 and 10 min, mimicking the heat-affected zone, lead to UTSs of 90 and 65 MPa, respectively. On the other hand, in the case of the 2 mm thick Al 2024–ULC steel joints, fracture occurs mostly
Figure 4. (a) Evolution of the engineering strength of the Al alloy–steel joints as a function of the tool advancing speed. The mode of fracture is also indicated. (b) Picture showing an Al 2024–steel joint (advancing speed = 100 mm min 1) that broke at the steel plate. (c) SEM micrograph of the cross-section of typical Al 2024–steel joints at 600 mm min 1, showing the porosities appearing in the last liquid Al zone.
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C. van der Rest et al. / Scripta Materialia 77 (2014) 25–28
in the steel sheet for low advancing speeds (200– 400 mm min 1), as shown in Figure 4b. The mean strength is roughly similar to the UTS of the base ULC steel (289 MPa vs. 300 MPa). The joints processed at a faster advancing speed (500–700 mm min 1) broke at lower engineering strengths. Indeed, fracture occurs following the porosities resulting from the last solidified liquid Al reported above (see Fig. 4c). Between 400 and 500 mm min 1, both fracture locations are observed, suggesting that this range of welding speeds corresponds to a transition in the fracture mode resulting from a significant level of solidification porosities in the weld. Even though the results of Figure 4c clearly show that the IML does not initiate fracture, no conclusions can be drawn about the effect of the IML thickness on the weld transverse lap shear strength. For FSW or diffusion bonding in a butt-joint configuration, Tanaka et al. predicted that stronger joints result from thinner IML, with an optimum thickness around 0.5–1.5 lm [10]. While the same tendency is expected in the case of the present lap-joint configuration, specific mechanical tests need to be designed to better characterize the intrinsic strength of the weld interface, and to avoid fracture in the base materials. Meanwhile, it seems that the present welding process results in very fine grains in the IML (Fig. 2) that do not deteriorate the weld strength with respect to the base material, even for IMLs up to 6–8 lm thick. ULC steel has been successfully welded to 1050 or 2024 Al alloys by a new simple and efficient process based on transient liquid bonding. Two IMLs were identified at the Fe–Al interface: Fe2Al5 on the steel side and FeAl3 on the Al side. The thickness of the Fe2Al5 IML decreases with increasing tool advancing speed, whereas the FeAl3 IML thickness remains constant at 1 lm. A solidification microstructure with FeAl3 precipitates was observed in the Al alloy. For advancing speeds >400 mm min 1, porosities appear in the last liquid zone of the Al 2024 where larger Cu and Mg concentrations were highlighted. Fracture occurred either in 1050 Al or ULC steel at UTSs close to that of the base materials. However, when porosities appeared in the 2024 Al, fracture occurred in the last liquid zone of the 2024 Al, 15– 20 lm below the interface, leading to a slight decrease in the strength. Based on these results, the welding process developed here can be considered a promising technique for joining steel to Al alloys or other combinations of metallic materials with different melting points.
The authors acknowledge financial support from the Interuniversity Attraction Poles Program from the Belgian state through the Belgian Policy agency; contract IAP7/21 “INTEMATE”. C.v.d.R. is funded by a FRIA grant from the National Funds for Scientific Research (FRS-FNRS). P.J.J. acknowledges the FRS-FNRS. [1] C.M. Chen, R. Kovacevic, Int. J. Mach. Tools Manuf. 44 (2004) 1205–1214. [2] S. Elliott, E. Wallach, Met. Constr.-Brit. Weld. 13 (1981) 167–171. [3] S. Elliott, E. Wallach, Met. Constr.-Brit. Weld. 13 (1981) 221–225. [4] H. Springer, A. Kostka, E.J. Payton, D. Raabe, A. Kaysser-Pyzalla, G. Eggeler, Acta Mater. 59 (2011) 1586– 1600. [5] W.-B. Lee, M. Schmuecker, U. Mercardo, G. Biallas, S.B. Jung, Scripta Mater. 55 (2006) 355–358. [6] S. Meco, G. Pardal, S. Ganguly, R.M. Miranda, L. Quintino, S. Williams, Int. J. Adv. Manuf. Technol. 67 (2013) 647–654. [7] T. Ogura, Y. Saito, T. Nishida, H. Nishida, T. Yoshida, N. Omichi, M. Fujimoto, A. Hirose, Scripta Mater. 66 (2012) 531–534. [8] T. Araki, M. Koba, S. Nambu, J. Inoue, T. Koseki, Mater. Trans. 52 (2011) 568–571. [9] M. Girard, B. Huneau, C. Genevois, X. Sauvage, G. Racineux, Sci. Technol. Weld. Joining 15 (2010) 661–665. [10] T. Tanaka, T. Morishige, T. Hirata, Scripta Mater. 61 (2009) 756–759. [11] A. Elrefaey, M. Gouda, M. Takahashi, K. Ikeuchi, J. Mater. Eng. Perform. 14 (2005) 10–17. [12] Y. Chen, T. Komazaki, Y. Kim, T. Tsumura, K. Nakata, Mater. Chem. Phys. 111 (2008) 375–380. [13] G. Zhang, W. Su, J. Zhang, Z. Wei, Metall. Mater. Trans. A 42 (2011) 2850–2861. [14] A. Elrefaey, M. Takahashi, K. Ikeuchi, Q. J. Jpn. Weld. Soc. 23 (2005) 186–193. [15] C. van der Rest, A. Simar, P.J. Jacques, Method for welding at least two layers (patent), International Publication No. WO2013164294 (A1), International Publication Date: 7 November, 2013. [16] R. Mishra, Z. Ma, Mater. Sci. Eng., R 50 (2005) 1–78. [17] K. Bouche´, F. Berbier, A. Coulet, Mater. Sci. Eng., A 246 (1998) 167–175. [18] A. Bouayad, Ch. Gerometta, A. Belkebir, A. Ambari, Mater. Sci. Eng., A 363 (2003) 53–61. [19] H.R. Shahverdi, M.R. Ghomashchi, S. Shabestari, J. Hejazi, J. Mater. Process. Technol. 124 (2002) 345–352. [20] T.B. MassalkiBinary Alloy Phase Diagrams, vol. 1, Springer, New york, 1986.
ScienceDirect Scripta Materialia 77 (2014) 25–28 www.elsevier.com/locate/scriptamat
On the joining of steel and aluminium by means of a new friction melt bonding process Camille van der Rest, Pascal J. Jacques⇑ and Aude Simar Universite´ catholique de Louvain (UCL), Institute of Mechanics, Materials and Civil Engineering, IMAP, Place Sainte Barbe 2, B-1348 Louvain-la-Neuve, Belgium Received 26 November 2013; revised 7 January 2014; accepted 8 January 2014 Available online 15 January 2014
A new lap welding process for dissimilar materials has been developed and tested for steel-to-aluminium joints. A simple cylindrical tool is rotated and translated over the top steel plate, leading to a transient partial melting of the bottom aluminium plate and the formation of intermetallic reaction layers of FeAl3 and Fe2Al5 as thin as 2.5 lm. Lap shear tests of the welds show very good resistance with fracture mostly in the base materials. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Friction stir welding; Dissimilar welding; Aluminium alloys; Steels; Mechanical properties
Assembling dissimilar materials within integrated structures is nowadays a stringent requirement related to improved performance. It is particularly the case for the transportation industry that needs to reduce the weight of structures while still improving their strength and crashworthiness. While efforts are carried out to continuously improve the mechanical properties of new steel grades up to levels unreachable by other materials, as in the case of advanced high-strength steels, Al alloys also present very interesting levels of specific strength that easily ensure lightweight structures, as demonstrated by their ever-increasing use in transport applications. Consequently, the need for an efficient technology to join steel and Al alloys arises in the automotive industry to ensure profitability and lightness of the vehicles in addition to efficient sealing of the joints. The main difficulty in joining steel and Al alloys results from the large difference in their respective range of solid-state stability, but also from their large reactivity leading to the formation of brittle intermetallic layers at their interface [1–4]. As a consequence, conventional fusion welding, such as arc and laser welding for which both materials melt, are unsuitable for joining steel to Al alloys since the temperature that is reached causes inter-
⇑ Corresponding
author. Tel.: +32 10472432; fax: +32 10474028; e-mail: [email protected]
diffusion and growth of large intermetallic reaction layers resulting in a brittle cast structure with poor mechanical properties [5,6]. Among the various solutions that have been proposed to reduce the heat input and thus to control the interface reactivity [7], solid-state processes are particularly efficient for welding dissimilar materials since the formation of an intermetallic layer at the interface may be avoided or reduced to acceptable values [2]. In diffusion bonding, the materials to be joined are put in intimate contact for hours under pressure and at a temperature below 70% of the lowest melting point. A transient liquid may sometimes be present owing to a eutectic reaction between the materials being joined [8]. However, the native Al oxide layer of Al alloys acts as a diffusion barrier and inhibits bond formation [2]. In order to break this oxide layer and activate the surface, the friction welding process is based on the movement of one piece relative to the other under compressive axial load, thus also involving shorter reaction times and higher temperatures (near the lower melting point) [3]. However, friction welding is mostly limited to rods and tubes since the relative movement is obtained by rotation. Friction stir welding (FSW) is another solid-state joining process which has been widely investigated for welding dissimilar materials, in particular Al alloys and steel, in a butt-joint [1,5,9,10] or lap-joint [11,12] configuration. Some studies have also considered a third material
http://dx.doi.org/10.1016/j.scriptamat.2014.01.008 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
26
C. van der Rest et al. / Scripta Materialia 77 (2014) 25–28
with a lower melting point, such as Zn, to act as a filler metal between Al and steel [13] or to help in reducing the thickness of the intermetallic layer at the interface [14]. Although FSW is considered a solid-state technique, Chen et al. observed that, when working with materials with very different melting points (Al 6061 and AISI 1018 steel in their case), one should locally reach the lowest melting point, leading to a method that combines the effects of both fusion and solid-state welding [1]. The present study aims at reporting a new welding technique [15] suitable for assembling dissimilar metallic materials, such as Al and steel sheets, thus circumventing the drawbacks of previous techniques. After stacking and clamping the metallic sheets, with the high-melting point material on top, a non-consumable flat cylindrical tool is rotated and pressed on the upper surface of the stack as shown on Figure 1. The rotating tool then advances along the surface at a predetermined speed. Frictional heat is generated due to the intimate contact between the pressed rotating tool and the top sheet. In addition, plastic deformation occurs below the rotating tool, also contributing to the temperature increase [16]. The generated heat is transmitted to the bottom sheet through the top sheet, and if it is sufficient, the low-melting-point bottom sheet partially melts, forming a localized pool of liquid metal in contact with the upper sheet. In contrast to conventional fusion welding processes, the melting temperature of the low-melting-point material is only exceeded slightly and very locally, resulting in a semi-solid welding process. The technique is easily implemented since it does not require any shielding gas, unlike laser welding [6], or specific preparation of the metallic sheets. In addition, in comparison with conventional FSW, the pin-less tool results in less abrasion of the tool and the absence of a keyhole at the end of the weld. Another benefit of this new technique is that the advancing speed is significantly higher than in other Fe–Al welding processes [7,11–14] and one could even foresee speeds as high as 1 m min 1 being attained. Our process could also be considered for welding other combinations of metallic materials with different melting points. In order to demonstrate the sequence of reactions while avoiding complex metallurgical issues, simple steel and Al grades were first tested. A microalloyed ultralowcarbon (ULC) steel with the composition (wt.%) 0.013C, 0.136Mn, 0.01Si, 0.005S, 0.0132P, 0.002Ti, 0.003Nb,<0.0008N was used as the high-melting-point
Figure 1. Schematic of the friction melt bonding process developed to assemble metallic sheets with different melting temperatures.
material. Cold-rolled and annealed sheets 0.8 mm thick with a mean grain size of 13 lm were used. 1 mm thick Al 1050 or 2 mm thick Al 2024 T3 sheets were used as the low-melting-point material. Sheets were 250 mm long and 80 mm wide. The surfaces were cleaned with acetone prior to stacking and welding. A tungsten carbide tool 16 mm in diameter was tilted backwards by 0.5°. The welding machine was controlled in displacement while the rotation rate was set at 2000 rpm for the whole set of experiments. The advancing speed ranged from 100 to 700 mm min 1. After welding, the joints were cross-sectioned by electrical discharge machining perpendicular to the welding direction for metallographic analyses and tensile testing. The transverse sections of the joints were observed by scanning electron microscopy (SEM) with back-scattered electrons (BSEs). In addition, the structure of the joints as well as the diffusion processes were characterixed by energy-dispersive X-ray spectroscopy (EDX) and electron backscatter diffraction (EBSD). Transverse lap shear tests were performed on rectangular specimens with a width of 10 mm, at room temperature and with a crosshead speed of 1 mm min 1. The recorded loads were normalized by the initial transverse section area of the sample at the fracture location, i.e. Al or steel sheets or lap-joints. SEM and EDX were also carried out on the fracture surfaces. The typical microstructure of the joints is illustrated in Figure 2 in the case of the Al 1050–ULC couple welded at 400 mm min 1. Three zones can be distinguished on the SEM micrographs, the Al and steel sheets being separated by a reactive layer. Close to the
Figure 2. Typical band contrast EBSD map of the Al 1050–steel interface (advancing speed = 400 mm min 1) with partial colouring as a function of the indexing phases (face-centred cubic Al in orange, monoclinic FeAl3 in blue, orthorhombic Fe2Al5 in red and bodycentred cubic Fe in green) which highlights the presence of two IMLs at the interface. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
C. van der Rest et al. / Scripta Materialia 77 (2014) 25–28
interface, the Al presents a solidification microstructure with a dendritic structure separated by an eutectic structure containing Fe-rich precipitates. Some large precipitates were indexed by EBSD (see Fig. 2) as FeAl3, in agreement with Bouche´ et al. [17] who had previously reported Fe enrichment of the liquid Al close to the solid– liquid interface. Furthermore, Bouayad et al. reported that the occurrence of these precipitates is favoured when the interaction time and the cooling rate are increased [18]. The steel sheet presents a gradient of microstructure through its thickness with a refined grain size in a layer close to the tool and coarser grains close to the Al sheet. In this thermomechanically affected zone, the grain size ranges from 1 to 35 lm. As a consequence of the melting of the Al, an intermetallic reaction layer (IML) is formed at the interface with the steel sheet due to the reactivity between the molten Al and the solid steel. For all advancing speeds investigated, two IMLs were indexed by EBSD. Figure 2 first shows that both FeAl3 on the Al side and Fe2Al5 on the steel side are formed at the interface. These conclusions were validated by BSE SEM micrographs and EDX measurements, even though the different intermetallics have similar chemical compositions requiring accurate measurements. Figure 2 also shows that nanosized grains constitute the FeAl3 layer, while grains of Fe2Al5 are tongue-shaped and larger. These results are consistent with previous studies [4,17–19] that also reported that the thickness of the IML increases when the contact time between liquid Al and steel increases, causing growth of an Fe2Al5 IML governed by interdiffusion (parabolic growth). Shahverdi et al. also highlighted that the kinetics of formation of Fe2Al5 is faster than for FeAl3, which in addition breaks and dissolves within molten Al [19]. As a consequence, Fe2Al5 is generally reported as the major part of the reaction layer. Figure 3 shows that the thickness of the IML (considering both layers as a whole) decreases similarly for the two Al alloys when increasing the tool advancing speed, which corresponds to a decrease in the heat input during the welding process. Indeed, faster advancing speeds involve locally shorter reaction times. This intuitive result was already observed for other welding techniques, such
Figure 3. Evolution of the thickness of the IML as a function of the tool advancing speed for both Al 2024 (plain symbols) and Al 1050 (open symbols)–ULC steel joints. Error bars provided when available.
27
as friction stir diffusion bonding [9] and friction stir brazing [13]. More importantly, the global level of thickness of the IML remains small, going down to 2.5 lm at 700 mm min 1. Furthermore, Fe2Al5 is the major part of the reaction layer except for very fast advancing speeds for which the Fe2Al5 and FeAl3 layers are of similar thicknesses. Indeed, the thickness of the Al-rich FeAl3 IML is almost constant (800 nm–1.5 lm) regardless of the advancing speed (and thus heat input) while the thickness of the Fe2Al5 layer is highly dependent on the heat input . In addition to the IML thickness, the tool advancing speed also influences to a large extent the Al resolidification process through the resulting cooling rates. In the case of the Al 2024 grade, porosities several microns in size appear in the resolidified Al, 15–20 lm below the IML, when the advancing speed is >400 mm min 1 (see Fig. 4c). This zone corresponds to the last liquid Al that solidifies. This phenomenon was confirmed by EDX analyses that showed larger concentrations of the Al alloying elements (Cu and Mg) close to the porosities, in agreement with the solidus on the phase diagrams [20]. In addition, SEM observations revealed that the quantity and size of these porosities increase when the advancing speed increases. Figure 4a presents the evolution as a function of the tool advancing speed of the engineering strength at the onset of necking of transverse lap shear tests. Three sets of constant levels of strength with the advancing speed can be observed that are related to the location of the necking and fracture; these levels depend on the composition of Al alloy. On the one hand, in the case of the 1 mm thick 1050 Al–ULC steel joints, fracture occurs in the Al sheet at 4.5–5 mm from the centre of the joint, namely in the heat-affected zone. Whatever the advancing speed and the resulting IML thickness, the maximum engineering strengths are similar with a mean value of 76 MPa. The Al 1050 base metal exhibits an ultimate tensile strength (UTS) of 100 MPa, while heat treatments at 620 °C for 3 and 10 min, mimicking the heat-affected zone, lead to UTSs of 90 and 65 MPa, respectively. On the other hand, in the case of the 2 mm thick Al 2024–ULC steel joints, fracture occurs mostly
Figure 4. (a) Evolution of the engineering strength of the Al alloy–steel joints as a function of the tool advancing speed. The mode of fracture is also indicated. (b) Picture showing an Al 2024–steel joint (advancing speed = 100 mm min 1) that broke at the steel plate. (c) SEM micrograph of the cross-section of typical Al 2024–steel joints at 600 mm min 1, showing the porosities appearing in the last liquid Al zone.
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in the steel sheet for low advancing speeds (200– 400 mm min 1), as shown in Figure 4b. The mean strength is roughly similar to the UTS of the base ULC steel (289 MPa vs. 300 MPa). The joints processed at a faster advancing speed (500–700 mm min 1) broke at lower engineering strengths. Indeed, fracture occurs following the porosities resulting from the last solidified liquid Al reported above (see Fig. 4c). Between 400 and 500 mm min 1, both fracture locations are observed, suggesting that this range of welding speeds corresponds to a transition in the fracture mode resulting from a significant level of solidification porosities in the weld. Even though the results of Figure 4c clearly show that the IML does not initiate fracture, no conclusions can be drawn about the effect of the IML thickness on the weld transverse lap shear strength. For FSW or diffusion bonding in a butt-joint configuration, Tanaka et al. predicted that stronger joints result from thinner IML, with an optimum thickness around 0.5–1.5 lm [10]. While the same tendency is expected in the case of the present lap-joint configuration, specific mechanical tests need to be designed to better characterize the intrinsic strength of the weld interface, and to avoid fracture in the base materials. Meanwhile, it seems that the present welding process results in very fine grains in the IML (Fig. 2) that do not deteriorate the weld strength with respect to the base material, even for IMLs up to 6–8 lm thick. ULC steel has been successfully welded to 1050 or 2024 Al alloys by a new simple and efficient process based on transient liquid bonding. Two IMLs were identified at the Fe–Al interface: Fe2Al5 on the steel side and FeAl3 on the Al side. The thickness of the Fe2Al5 IML decreases with increasing tool advancing speed, whereas the FeAl3 IML thickness remains constant at 1 lm. A solidification microstructure with FeAl3 precipitates was observed in the Al alloy. For advancing speeds >400 mm min 1, porosities appear in the last liquid zone of the Al 2024 where larger Cu and Mg concentrations were highlighted. Fracture occurred either in 1050 Al or ULC steel at UTSs close to that of the base materials. However, when porosities appeared in the 2024 Al, fracture occurred in the last liquid zone of the 2024 Al, 15– 20 lm below the interface, leading to a slight decrease in the strength. Based on these results, the welding process developed here can be considered a promising technique for joining steel to Al alloys or other combinations of metallic materials with different melting points.
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