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Generation of aluminium–steel joints with laser-induced reactive wetting
Materials Science and Engineering A 444 (2007) 327–338
Generation of aluminium–steel joints with laser-induced reactive wetting P. Peyre a,∗ , G. Sierra b,c , F. Deschaux-Beaume c , D. Stuart a , G. Fras c a
GERAILP-LALP, UPR 1578 CNRS, 94114 Arcueil, France b CEA/DRT/DTEN/GERAILP, 94114 Arcueil, France c LMGC, UMR 5508 CNRS, Universit´ e de Montpellier II, 30907 Nˆımes, France Received 1 December 2005; received in revised form 1 July 2006; accepted 1 September 2006
Abstract A new mean of assembling steel to aluminium was developed, following previous work by German workers [1]. In this new method, a laser-induced aluminium melt pool spreads and wets a solid steel, to generate, after solidification a sound and resistant interface layer. Joint properties were investigated, in terms of surface aspects, interface microstructures and mechanical resistances under tensile testing, for non-galvanized and galvanized DC04 steels. Thermal and diffusional finite element (FE) simulations were also carried out to calculate temperature history at interfaces, and reaction layer thickness. The 2–20 m thick reaction layers formed all along the interface were found to be mostly composed of Fe2 Al5 intermetallic compound with a high hardness (1200 HV) and rather low ductility (presence of solidification cracks). The presence of a 10 m thick Zn layer on the steel was shown to have a beneficial influence on the wetting characteristics of the joint, despite the formation of occluded pores in the melt pool due to Zn vaporisation. FE thermal modelling evidenced 760–1020 ◦ C wetting temperatures at the interface between DC04 low carbon steel and 6016 aluminium sheets, with time maintains of the melt pool in the 0.2–0.5 s range, resulting in high-speed reaction kinetics. Using these temperature data, diffusion calculations were shown to provide a rather good prediction of intermetallic thicknesses. Tensile tests were considered on aluminium–steel lap joints and evidenced higher mechanical resistances (220 N/mm linear tensile strength) on galvanized steels, provided that fluxing of the steel surface was carried out prior to welding to avoid zinc vaporisation. Comparatively, non-galvanized assemblies exhibited much lower mechanical resistances (170 N/mm resulting in a 90 MPa interfacial shear strength). It was concluded that the laser-induced wetting technique is a rather effective way for generating Al-steel joints without filler material, and that it should be considered as a competitive technique versus solid assembly modes (friction stir welding . . .). © 2006 Elsevier B.V. All rights reserved. Keywords: Aluminium–steel; Laser; Wetting; Intermetallic compounds
1. Introduction Many processes like hot dip coating against corrosion [2,3], generation of metal matrix composites, or brazing and welding of dissimilar materials [1], are based on the interaction of a solid metal or alloy with a liquid metal [4]. During the interaction time, the solid surface dissolves in the melt, and intermetallic layers are subsequently formed at the interface. Usually, the formation of brittle intermetallic layers limits the mechanical resistance of the transition zone between dissimilar metals [5], when their thickness exceeds a permissible value. This is particularly true for aluminium–steel
∗
Corresponding author. E-mail address: [email protected] (P. Peyre).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.09.111
interfaces which have been investigated already [5–7] because of a huge industrial interest, and which are known to be unfavourable with respect to the mechanical resistance of the joints. The wide diversity of interfacial processes involved in materials joining (diffusion, chemical reaction and fluxing) suggests complicated phenomena during wetting, and wetting kinetics which can be either diffusion controlled or reaction controlled, depending on the systems considered. Several authors have focused on the interface reactions between aluminium and steel in isothermal conditions, to estimate the reaction kinetics and diffusion constants at various temperature, by dipping iron or steel samples in molten aluminium baths at temperatures ranging between 700 and 900 ◦ C [3,4,8]. Microstructure analysis have been carried out on the interface layers, and most of the reported investigations exclu-
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sively identified two distinct intermetallic phases: Fe2 Al5 as the major constituent close to the steel substrate, growing according to a parabolic law, and FeAl3 as the minor phase close to the solidified aluminium [8,9]. After solidification, these intermetallic phases provide bonding of aluminium with steel. Recently, authors [10] have also considered theoretically the interfacial reactions between solid iron and liquid zinc–aluminium alloys. Since brittle intermetallic phases are considered as a critical issue, other techniques, like friction welding [11–13] have been applied successfully to produce butt joints between aluminium and steel without generation of a melt pool, but through the friction of a rotating pin which puts the aluminium in a fluid-like plastic state adhering to the steel surface. In spite of intermetallic issues, several attempts have been made recently (2001–2004) to realize aluminium–steel assemblies with Nd:YAG lasers. Different approaches have been used which can be classified into: (1) a laser brazing technique with filler wire [7], (2) a classical key-hole (vapour capillary) welding mode [14], and (3) a laser roll welding mode where steel remains solid and laser heating provides aluminium melting and subsequent dissolution–diffusion–reaction of steel [6]. All these joining techniques have succeeded in providing acceptable mechanical strengths (in tensile or shear modes), when limiting the interface layer thickness (<10 m) for brazing or reactive wetting, or limiting the aluminium–steel dilution, for the key-hole technique where both metals are melted. Analytical work about diffusion aspects in laser-induced reactive wetting of aluminium with steels have never been published, and seem rather difficult to perform because of: (1) the non-homogeneous spatial profile of lasers, (2) the out-ofequilibrium phenomena involved at high heating and cooling rates, that limit the use of phase diagrams and (3) the strongly non-isothermal nature of laser-heating, and the shorter reaction times than those previously used during isothermal solid steel-liquid aluminium investigations (usually superior to 100 s [3,4,8]). More, the complex geometries of laser assemblies require thermal simulations (with finite element codes), to evaluate the time–temperature histories T = f(t) of interfaces before investigating growth kinetics aspects. In this paper, the generation of aluminium–steel assemblies with a highly non-isothermal laser-induced reactive wetting was addressed. The different aims of the present work were: (1) First, to optimise experimental conditions (welding speeds, overlaps) for the realisation of low carbon steel—6000 aluminium assemblies, using a dedicated experimental set-up, with a specific attention paid to the fast camera analysis of melt pools and to temperature measurements. (2) To make metallurgical and mechanical analysis of assemblies, focusing on interfacial layers. (3) To assess temperature evolution at the aluminium–steel interface with 3D FEM simulations, and to compare them with interfacial thicknesses, in order to envisage further diffusion calculations.
Table 1 chemical composition of DC04 low carbon steel and 6016-T4 aluminium alloy Element (wt.%)
Al
Fe
Mg
Si
Mn
P+S
C
DC04 6016
– Bal.
Bal. 0.5
– 0.3–0.5
– 1–1.3
0.4 0.2
0.06 –
0.08 –
(4) To compare the reactive wetting of galvanized and nongalvanized steels and the mechanical resistance of joints. 2. Materials The materials used for these investigations were a low carbon steel DC04 mostly used for automotive applications, defined by the european normalisation NF EN 10130 and a 6016-T4 aluminium alloy for both aeronautical and automotive issues. Their chemical composition, temperature-dependent physical and mechanical properties are reported in Tables 1 and 2, respectively. The 1.2 mm thick DC04 sheets were provided with a non-galvanized state or a 20 m zinc-coated layer containing nearly 0.2% Al to avoid formation of brittle Fe(Zn)x compounds. On the other hand, 1 mm thick 6016 aluminium sheets were used. Mechanical properties reported in Table 2 were evaluated by quasi-static tensile tests. 3. Experimental procedure 3.1. Reactive wetting laser experiments Respectively, 1.2 and 1 mm thick sheet specimens of DC04 steel and 6016 aluminium alloy were used. Samples were prepared by grade 800 SiC paper polishing, degreasing and (for the non-galvanized steel surface) fluxing with a fluor-base anticorrosive brazing flux. The flux suspension (flux powder dissolved in ethanol), was sprayed rather homogeneously on the surface (10 m < estimated thickness < 50 m), and dried before welding. On galvanized steel sheets, as already shown in Ref. [7], no fluxing was shown to be required to ensure wetting of the zinccoated steel surface, due to the good metallurgical compatibility between the Fe, Al and Zn. However, a few comparative tests have also been carried out with fluxing. Lap wetting tests (Fig. 1) were carried out with a Nd:YAG continuous wave (cw) laser operating in the 2.5–3 kW power range, with scanning speeds between 0.6 and 1.4 m/min. Operating with a 200 mm focal lens, laser was used in defocused condition, with a 30◦ irradiation angle, resulting in a elliptic spot of 2.3 mm × 1.8 mm, and laser intensities in the 70–90 kW/cm2 . All wetting tests were carried out in flat position (using the gravity contribution to the aluminium spreading) on a 100 mm testing length, with the laser beam located on the upper aluminium sheet (Fig. 1a and b), and with an argon shielding gas nozzle (at nearly 10 l/min) to protect heated surfaces from oxidation. For each assembly, a fast camera (Cmos 2000 Hz) recording of the aluminium melt pool (Fig. 2) was done, and temperature
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Table 2 Thermo-physical and mechanical data of DC04 steel and 6016-T4 aluminium alloy [15] Physical properties
ρ (kg/m3 )
λ (W/m/K)
Cp (J/Kg/K)
(20 ◦ C)
(20 ◦ C)
Latent heat (J/kg)
σ Y (MPa)
σ UTS (MPa)
160
280
140
230
DC04 steel
7810
65 45 (400 ◦ C) 28 (800 ◦ C)
470 595 (400 ◦ C) 875 (800 ◦ C)
3.105
6016-T4 alloy
2700
172 (20 ◦ C) 190 (560 ◦ C)
887 (20 ◦ C)
3.8.105 (560 ◦ C–640 ◦ C)
(1450–1510 ◦ C)
Fig. 1. Schematic description of the experimental set-up: (a) global overview and (b) detailed view of the reactive aluminium–steel wetting experiments (α = wetting angle).
measurements near the interaction zone were carried out with 0.2 mm diameter chromel–alumel thermocouples. This allowed us to collect input data for FEM simulations of the process. 3.2. Metallurgical characterizations From the 100 mm-long wetting joints, specimens were cut across the fillet joint seam for macrostructure and microstructure observations. Cross sections of the joints were polished by 220, 500 and 1000 SiC grades, followed by a 3 and 1 m diamond suspension, to a mirror-like surface aspect. Etching with Keller’s reagent (HNO3 + HCL + HF + H2 O) was used to reveal the aluminium microstructure and Nital 4% was used to attack the steel structure. Then, microstructures of the interfaces were observed using: (1) optical microscopy, at magnifications ranging between ×50 and ×200, and (2) scanning electron microscopy (SEM) with an energy dispersive X-ray (EDX) analysis facility to determine the chemical composition of intermetallic phases formed during steel wetting by aluminium melt pool.
Microhardness (10–25 g loading force) and nanoindentation1 (0.1–2 g) tests were also carried out to estimate local mechanical properties of affected zones and interfaces. A more detailed presentation of the nano-indentation technique is available in Ref. [16]. 4. Thermal simulations with the Finite Element Method 4.1. FEM conditions A three dimensional (3D) finite element model was implemented in the finite element software ABAQUSTM Standard [17] with a *Dflux Fortran subroutine to simulate a heat moving source, and a *Heat Transfer card to solve the heat flow equation. The heat was considered to be deposited onto the metal surface with a elliptic intensity distribution, corresponding to a defocused laser irradiation, and a 32% beam absorption
1 The authors wish to thank MC.Ste Catherine and L. Quiniou from CEPArcueil (DGA) for nano-indentation tests.
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used: a uniform 10 W/m2 K value on all the external sides of the body, and h = λ/e = 1500 and 700 W/m2 K values on the lateral vertical edges of, respectively, aluminium and steel sheets. The influence of a 20 m Zn layer was also considered in our FEM description of the thermal history of interfaces. The only effect evidenced was a small temperature decrease (−40 ◦ C maximum value) at the steel–Al interface. 4.2. Process modelling and experimental validation
Fig. 2. Fast camera (1000 Hz) coaxial analysis of the melt pool dimensions during a reactive wetting aluminium–steel assembly (V = 1.2 m/min and P = 3 kW): melt pool length is between 5.3 mm (a) and 6 mm (b).
on aluminium measured by integrated sphere in solid state. Up to 12,000, eight nodes diffusive-conductive hexaedral elements (DC3D8) were used to make a 3D meshing of the overlap aluminium–steel joint. The mesh sizes were chosen in such a way that smaller elements corresponded to the irradiated zone, where flux gradients are maximal. The computation domain was made out of two sheets (20 mm long × 11 mm width × 1 mm thick) with a 5 mm overlap on X-axis (Fig. 3a). Temperature dependent (non-linear) physical properties, reported in Table 2, were taken from Ref. [15]. In the numerical simulation, the latent heat of fusion of each material was taken into consideration, but not the hydrodynamic flow contribution to heating (due to fluid flow in the molten pool, which velocity was estimated to 0.2–0.5 m/s by fast camera analysis). It should be mentioned that in previous work, the thermal conductivity of the liquid state was artificially increased [7] to account for the enhanced heat transport due to surface-tensions-driven convections. The aim of these investigations was to evaluate the thermal history T = f(t) at the aluminium–steel interface, fitting our simulations (and especially the absorption coefficient of the Nd:YAG beam by aluminium or steel) with experimental data from thermocouple measurements near the interaction zone. Data from real-time fast camera analysis of melts (fusion zone dimensions) were also used to optimise our simulations. Concerning the boundary conditions, heat loss coefficients h were used, which traduced the cooling flux due to a laminar flow of protective gas (Qgas = −hgas (T − T0 )), but also the heat dissipation through the clamping device and the lateral sides of the sheets (Qsolid = −hsolid (T − T0 ) with hsolid = λ/e, with λ = conductivity (W/m/K) and e (m) = thickness/length of the part). For our simulations, different h values were
All calculations were carried out with a standard resolution algorithm. The energy supplied by the moving laser head was modelled by a Fortran routine in form of a distributed heat flux (W/m2 ) applied on the external edge of surface elements. It enabled us to define the position of the elliptic laser deposit all along the time range (less than 1 s total time), and for different scanning speeds (between 0.7 and 1.2 m/min). Results of the simulations were obtained in the form of timedependent temperature distributions (local thermal history), and 3D thermal fields. This enabled us to compare T = f(t) simulations with thermocouple data recorded on the steel surface at a “d” distance (between 3 and 6 mm) from the interface. Also, the dimensions of the fusion zone could be compared with fast camera recordings (Fig. 2). Results of calculations, carried out with a 32% absorption of the Nd:YAG light, are shown in Fig. 3 for a 1.2 m/min scanning speed. A 1600 ◦ C peak temperature was obtained in the energy deposit zone of the melt pool, that seemed to be rather high. On the other hand, dimensions of the melt pool (nearly 6.9 mm length and 3.2 mm width) and most of all, thermal history taken
Fig. 3. FEM simulation of a 1.2 m/min–3 kW aluminium–steel laser-induced reactive wetting: (a) thermal fields in the fusion zone in quasi-steady condition (aluminium up, steel down) and (b) upper view of the simulated fusion zone: the bead is 6.9 mm length can be compared with the 5.6 mm experimental value from Fig. 2.
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tion could be mostly controlled by reaction phenomena instead of diffusion mechanisms. 5. Macrostructure and microstructure of joints The results presented hereafter correspond to morphological and metallurgical analysis carried out on Al–steel interfaces produced by the reactive wetting assembly mode. 5.1. Bead aspects
Fig. 4. (a) Comparison between experimental thermocouple values and simulated T = f(t) profiles at 4 mm distance from the interface; (b) thermal histories T = f(t) at the aluminium–steel interface during a reactive wetting test (3 kW, 0.7–1.2 m/min).
at 4 mm distance from the interface, were in correct agreement with experimental data. Following the same procedure, simulations of 0.7–1.2 m/min processes could be carried out and successfully compared with experience (Fig. 4a). The most attractive information from these simulations, is the interfacial temperature history T = f(t), corresponding to the wetting temperature at the molten aluminium–solid steel interface. For the scanning speed range considered, 750–1000 ◦ C peak temperature were obtained (Fig. 4b), for 0.2–1 s time maintain of the wetting. This corresponds to very sharp and short thermal loadings compared to classical steel immersion into aluminium baths (usually 700–900 ◦ C constant temperature during several minutes immersion times) [4,8,9]. Following previous work by Refs. [4,8], it should be expected that, for such short interaction times, intermetallic layer forma-
The effect of travel speed on the appearance of laser-wetted joints is shown in Figs. 5 and 6 on a non-galvanized steel surface (Ra = 0.1 m). In spite of different thermal cycles, the global morphology of beads is kept rather constant, even though the bead widths are reduced with increasing speeds (nearly 4 mm at 0.8 m/min and 3 mm at 1.2 m/min). Also, different wetting (contact) angles are obtained (between 30◦ and 80◦ ), revealing more spreading of aluminium on steel for lower velocities (Fig. 6 and Fig. 7). However, for the same travel speed, contact angle may vary drastically (±20◦ ), due to local changes in heat dissipation (local gaps between upper and lower sheets). When comparing wetting angles obtained on galvanized and non-galvanized steels, one thing is obvious: the zinc layer improves the spreading and wetting of aluminium on steels’ surface. For instance, the same experimental conditions (V = 1 m/min, P = 3 kW and 2 mm spot) generate a 45◦ wetting angle on bare steel versus a 25◦ angle on galvanized surface. This is mostly due to the fact that, on galvanized specimens aluminium spreading occurs on a liquid zinc layer (Tmelt = 419 ◦ C), and not on a solid steel. Under that condition, we cannot talk any more about wetting of a solid surface. On the other hand, if wetting experiments are carried out on a fluxed galvanized surface, wetting angles tend to be strongly increased, but still remain lower than those measured on bare steels (Fig. 7).
Fig. 5. Macroscopic view of the aluminium–steel joints (P = 3 kW and spot diameter = 2.3 mm) for two different scanning speeds: (a) 1.2 m/min and (b) 0.8 m/min. Some geometrical defects (process instabilities) are randomly present but the global aspects are acceptable.
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Fig. 6. Cross-sections of aluminium–steel joints for two different welding speeds (0.6 and 1.2 m/min), on non-galvanized (a and b), and galvanized (c and d) conditions.
5.2. Microstructures of interfaces obtained between non-galvanized steel and aluminium On all the interfaces, the aluminium appears to be bonded to the non-galvanized steel substrate by a thin layer of a very narrow, uniform and defect-free Fe–Al phase.
Fig. 7. Influence of laser scanning speed on the wetting angle α for galvanized and non-galvanized steels with or without fluxing.
The interface thicknesses, estimated with optical analysis (Fig. 8a and b), are nearly constant versus scanning speed (Fig. 11), and remain in the 2–10 m width range, except on specific locations (Fig. 9) where thicker – and most of times cracked – interfaces could be evidenced. These heterogeneities may be due either to some variations of gap between Al and steel sheet, or to local variations of flux thickness prior to wetting. According to previous data on hot-dip coating [2], the thickness of the intermetallic layer depends on the viscosity of the aluminium bath, and decreases with increasing amounts of additive elements in the melt pool. For instance, Si, Cu and Be are known to have the greatest effect in preventing the build-up of the intermetallic layer, and minimizing its thickness. The inter-metallic morphology consists of an inner compact layer, and an outside “composite” layer composed of broken and/or floating particles in the Al matrix (Fig. 8b and Fig. 9). If we have a closer look at these interfaces, it appears that intermetallic layers exhibit serrated interfaces on the aluminium side. This phenomenon is not that obvious between aluminium and steel, and previous investigations using much longer liquid–solid contact times, have found such an aspect in the opposite interface, i.e between steel and intermetallic phase. Under our experimental conditions, only local zones with thicker interfaces (see Fig. 8) exhibit serrated interfaces on both sides. Micro-porosities (a few m), found at various locations along the interfaces, and frequently containing crystallized fluxing
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Fig. 8. Interfaces obtained after laser-induced wetting of non-galvanized (a) and (b), or galvanized (c) and (d) steels at 0.6 and 1.2 m/min. Interfaces obtained on galvanized steels are thicker (37 m vs. 5 m at 0.6 m/min and 10 m vs. 3 m at 1.2 m/min) because of the influence of flux that limits inter-metallic formation.
compounds, were clearly expected to originate from the volatile flux thickness. Their influence on fracture behaviour will be discussed later on. Semi-quantitative EDX investigations carried out on a 0.8 m/min interface indicated a small compositional gradient along the intermetallic layer: nearly 24% Fe 75% Al and 1% Si (wt.%) close to the Al-intermetallic interface, and 26% Fe 72% Al and 2% Si near the steel interface (Fig. 10). Both compositions are closer to Fe2 Al5 (Si) than FeAl3 (Si) sometimes evidenced [3] in Al–steel interfaces. The Si-enrichment near the steel side is not explained (a Si-enrichment would be more likely to occur in the upper part of the melt pool, where solidification ends), but convection movements in the melt pool could modify compositional zones. It also has to be noticed that Si is known to have an inhibiting effect on the formation and growth of Fe2 Al5 intermetallics. Also, FeAl3 needle-like crystals, in-between the intermetallic layer and the aluminium melt-pool (Figs. 9 and 10) are only evidenced at the centre part of reaction layers width, where the aluminium–steel interface experienced the higher temperatures. These crystals, attached to or separated from reaction layers, could be due to iron re-dissolution and diffusion in the aluminium melt pool, possibly favoured by convection movements of the liquid aluminium close to the intermetallic interface.
Fig. 9. Cross-sections of aluminium–steel joints: detail of a local 14 m-thick intermetallic interfacial zone (upper part, solidified aluminium): presence of intermetallic laths and solidification cracks (0.8 m/min, 3 kW).
5.3. Microstructures of assemblies obtained on galvanized steels As indicated in Section 5.1, the zinc layer improves the wetting of aluminium on steel, but this phenomenon is inhibited by a fluxing of aluminium sheets.
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Fig. 10. SEM-EDX analysis of a 5 m intermetallic layer: a compositional gradient is found between the aluminium and the steel side, but the global stoechiometry is close to Fe2 Al5 .
Interfaces obtained on galvanized steels exhibited a Fe2 Al5 Znx stoechiometry [10] with numerous solidification cracks when the interface thickness is above 10 m. Their global aspect is held unchanged compared with those obtained on nongalvanized steels, but their thickness is much higher (up to 30–40 m) than on non-galvanized steels (Fig. 8c and d, Fig. 11) when experiments are performed without flux. This increase seems to be more due to the influence of the flux itself, that limits the intermetallic growth, than to the higher reactivity of the galvanized surface. Indeed, non-galvanized and non-fluxed steels in contact with liquid aluminium exhibit very thick intermetallic even if wetting is close to zero, due to the non-melted oxide layer on aluminium melt pool.
Fig. 12. Microstructure of the molten aluminium generated during reactive wetting on DC04 galvanized steel: formation of pores by zinc vaporisation in the fusion zone and presence of cracks due to Zn local segregation.
The main difference with non-galvanized steel assemblies comes mainly from the ability of Zn to vaporize under laser heating, when the interface temperature reaches 900 ◦ C. This induces large porosities (50–300 m), mostly located in the inner part of the fusion zone, close to the heat-affected zone of aluminium. Combined with the softening effect evidenced in both FZ and HAZ of Al alloy, these porosities are potential locations for cracks initiation. It was also found that Zn dissolution in molten aluminium tends to increase cracking susceptibility inside aluminium melt pool, maybe due to the formation of low melting point Al–Zn compounds at interdendritic spaces (Fig. 12). 6. Microhardness and nano-identation results 6.1. Microhardness results
Fig. 11. Influence of laser scanning speed on the intermetallic thickness, during laser-induced reactive wetting of galvanized or non-galvanized XES steels. Fluxing reduces by a factor 7–8 the intermetallic thickness.
The degree of softening of age-hardened 6016-T4 (100 HV0.025 average hardness) was evaluated by Vickers hardness measurements, with a 25 g loading force, at a 500 m distance from the steel–aluminium interface. Some −35% maximum hardness reductions were obtained, independent on the welding speeds and on the presence of a galvanized layer on steels. The highest degree of softening was obtained in the fusion zone (FZ), and to a lesser extent in the adjacent heat affected zone (HAZ) where precipitates were dissolved or put into solid-solution. Rather wide heat-affected softened zones (more than 5–6 mm) were obtained, with a hardness gradient between the material immediately adjacent to the fusion zone (65 HV = −35%), and a second heat-affected zone (85 HV) where precipitates are not expected to be dissolved, but instead, to coarsen and lose their coherency with the Al matrix. The formation of wide softened regions was associated to a decrease of local strength properties. On the other hand, the steel hardness and microstructure were not affected by the wetting process.
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6.2. Nano-indentation characterization of interfaces
7. Tensile tests
Due to the very thin interface layers, nano-indentation tests, carried out with a Berkovich indenter and 20 mN applied force, were carried-out on different interface thickness, starting with thicker layers (Fig. 13). On all the interfaces investigated so far, results exhibit some rather constant 1100 ± 100 HV20 mN equivalent values in the intermetallic layer, compared with an average 250 HV20 mN value in the steel and 200 HV20 mN in the molten aluminium alloy. These hardness values did not depend on process parameters, and seemed to be in correct agreement with previous data from [6]. They also confirmed a very sharp mechanical gradient close to the interfacial zone. Last, no hardness difference was evidenced between interfaces formed on non-galvanized or galvanized steels, where the chemical composition is close to Fe2 Al5 Znx ( phase) [10]. Also, elastic modulus values in Fe2 Al5 , estimated from the slopes of elastic unloadings, were found to exhibit an average level of 220 GPa.
For ease of processing, 20 mm width lap joint specimens have been used to test aluminium–steel joints, despite tilting of tensile specimens under load, due to bending moments of the lap joints during mechanical testing [18]. The joint strength was evaluated by tensile testing at room temperature, with an initial gauge length of 15 mm (specimen length = 80 mm). Mechanical properties were shown to be rather different, depending on the presence of a galvanized layer or not. On nongalvanized assemblies, maximum resistance of up to 190 N/mm were obtained, with two different failures modes: (1) a brittle interfacial shear fracture or (2) a tensile failure in the fusion zone (Fig. 14). On galvanized steel assemblies, most of the failures occurred in the highly softened fusion zone. Corresponding resistances were evaluated to 130–140 N/mm (Fig. 15). These failures were
Fig. 13. Nano-indentation tests (0.02 N applied force, Berkovich indenter) of a 0.8 m/min specimen (intermetallic thickness = 12 m): (a) localisation of indents and (b) hardness values.
Fig. 14. Tensile curves obtained (a) on non-galvanized joints with 2 failure modes (interfacial brittle failure and failure in the FZ) and (b) on galvanized joints with 2 failure modes (in the HAZ with striction and in the fusion zone, promoted by Zn porosities).
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global properties of dissimilar assemblies. However, our results indicate that this assumption is not valid on age-strengthened materials such as 6016-T4 because most of the failures occur in the solidified aluminium and/or in the softened heat affected zones in the case of galvanized steels. This competition between interfacial failures and fusion zones fractures has been reported already in Ref. [13]. 8. Discussion on aluminium–steel interfaces
Fig. 15. Relationship between mechanical resistance (N/mm) and welding speeds: (a) on non-galvanized steels with two failure modes (1) interfacial failure of Fe2 Al5 and (2) tensile failure in the HAZ and (b) on galvanized steels with 2 failure modes in the HAZ or FZ.
promoted by a high density of pores in the melt-pool (Fig. 12) due to Zinc vaporization above 900 ◦ C, which reduced the section area S (mm2 ) and in turn increased the applied stress F(N)/S (mm2 ) = σ (MPa). In that case, Al–Fe–Zn interfaces were shown to be more resistant than the solidified and annealed 6016 containing pores and cracks (Fig. 12). Interesting results were obtained on galvanized samples with fluxing, even if flux was not shown to be crucial to ensure wettability between aluminium and a zinc-coated steel. In that specific case, maximum resistances of 230 N/mm (resulting in 230 MPa tensile strengths) were shown, with failures located in the first heat affected zone adjacent to the FZ. This result, very close from the fracture strength of 6016-T4 (Table 2), is rather surprising because failure occurred in partially softened zones (rather far away from the fusion zone) where local mechanical properties are higher than those of the FZ. It is tacitly assumed that the mechanical properties of materials adjacent to the joint interface are not important regarding the
Sound Al–steel interfaces could be obtained by the laserinduced reactive wetting assembly mode, with average interface thickness in the 2–15 m range. Rather good mechanical resistance were shown on all the assemblies, whatever the scanning speed and the wetting angles. However, questions remain unanswered about the mechanisms involved in the interface formation, for such short liquid–solid interaction times (0.1–0.5 s). If we first consider the aluminium–steel reactions without zinc layers, we can refer to previous works [4,8,9,19] in isothermal conditions. According to them, the initial growth steps of the interface layer occurs from steel towards aluminium, due to the ease of diffusion of iron atoms into molten aluminium (DFe→liq Al = 8 × 10−8 m2 /s at 1000 ◦ C). Then, a further increase of the Fe2 Al5 thickness modifies diffusion conditions and it becomes necessary to consider the diffusion of Al and Fe atoms through Fe2 Al5 to predict further intermetallic growth. Results presented in our work revealed very thin reaction layers, where the main factors contributing to the growth were expected to be controlled by a reaction instead of diffusion mechanism between steel and aluminium. The mechanisms involved could be as follows: (1) first, iron dissolution takes place towards a thermodynamic metastable equilibrium (with JFe→Al = kdiss (CFe meta − CFe int )), (2) iron diffuses in the aluminium melt pool but remains rather close to the Fe–Al interface (during a 1000 ◦ C–0.15 s thermal cycle, the diffusion path of Fe in Al is less than 200 m), (3) due to the high concentration of dissolved Fe in Al, supersaturation occurs and Fe2 Al5 nucleation occurs, followed by growth. So it can be considered that iron atoms present in Fe2 Al5 come from the dissolved iron occurring in the initial steps [10]. If we consider diffusion coefficients of Fe through Fe2 Al5 from Bouayad et al. [19] at 800 ◦ C (k = 9 × 10−11 m2 /s) or 700 ◦ C (k = 3.9 × 10−11 m2 /s) for low carbon steels, and calculate the diffusion paths, using equivalent times for T above 700 ◦ or 800 √ C (see Fig. 7b), we can compare these calculated values (ecalc 2ktT =700◦ C ) with experimental intermetallic thicknesses (Fig. 16). The agreement is rather correct, especially for the thickness amplitude, even if the evolution with time seems to be underestimated by a simple parabolic approach. Consequently, even if a simple parabolic dependence of the interface thickness against time (d(m) = Kt1/2 ), seems to be rather controversial, due to the complexity of laser-induced interfacial conditions, it seems to reproduce rather well the experimental intermetallic thicknesses. Considering the very specific local conditions induced by laser wetting, we can remind that, due to the small volume of molten aluminium (typically 5 mm3 ), tiny amounts of iron dis-
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Fig. 16. Maximum intermetallic thicknesses on non-galvanized steels vs. the inverse value of the welding speed square root (V−0.5 ). Experimental thickness is estimated with optical microscopy carried out in the middle of each crosssection. Calculated values use the simulated time maintains above 700 or 800 ◦ C with corresponding diffusion coefficients from [19] (9 × 10−11 m2 /s at 800 ◦ C, 3.9 × 10−11 m2 /s at 700 ◦ C).
solution may contribute to the saturation of molten aluminium with iron, and in turn modify (increase) the Fe2 Al5 growth rate. We also could consider that laser-induced reactive wetting involves hydrodynamic flow (see convection paths in Fig. 12) in the liquid aluminium (fluid velocity < 1 m/s), which may affect strongly the reaction mechanism, and possibly promote the dissolution or diffusion rates of the steel as already shown in Ref. [9] on rotating steel samples. Oncoming thermo–diffusional modelling should allow to make more precise predictions of intermetallic thicknesses. As indicated earlier in the paper, high temperature gradients, and high fluid velocities at the center part of interfacial width also promoted the formation of needle-like FeAl3 crystals in aluminium melt pool. However, these crystals are not expected to affect the mechanical resistance of assemblies, as failure usually initiates at the edges of reaction layers, and not on the centre part of layers where those needle-like crystals are located. With respect to other techniques for assembling steel and aluminium, the reactive wetting mode seems to offer a promising potential, because of an interesting balance between metallurgical and mechanical properties on one side, and rather attractive process conditions (high scanning speed . . .) on the other side. If we now consider the presence of the 20 m zinc layer between liquid aluminium and steel, a few comments can be make to explain our results. First, the reactivity between aluminium and zinc (%solubility) is much better than between Al and steel, as confirmed by the Al–Zn diagram. This allows good wetting and generation of intermetallic compounds without fluxing steels. Second, the presence of a 20 m zinc layer which vaporizes under laserirradiation strongly modifies the microstructure of assemblies
337
near the interfacial zones, and in turn the failure modes, mainly located in the fusion zone. No more interfacial failure occurred on galvanized steel assemblies with or without fluxing. The reason why intermetallic layers with Zn are more shear-resistant (or more adherent) than pure Fe2 Al5 is still under investigation, but identical local hardness values seem to indicate that it may not be due to a change in ductility. The influence of fluxing on galvanized steel parts may be explained as follows: during the laser heating, the liquid flux (Tmelt = 560 ◦ C) below the aluminium melt (maximum temperature below 1000 ◦ C: see Fig. 6) acts as a thermal barreer upon the melted zinc (Tboil = 910 ◦ C), and limits the thermal rise and vaporization of the galvanized layer. Consequently, zinc bubbles are avoided, zinc occlusion is limited in the melt pool and mechanical properties are enhanced due to the change of failure mode (from the fusion zone to the adjacent heat affected zone). Last, as mechanical properties of the galvanized joints are more limited by softening effects than by the intermetallic resistance itself, non heat-hardenable aluminium alloys (Al–Mg alloys) could be appropriate candidates for providing higher mechanical resistances with galvanized steels. 9. Conclusions 6016 aluminium alloy and DC04 low carbon steel were joined by laser-induced reactive wetting mode. The wetting assemblies were investigated by process analysis (fast camera, thermocouples), FEM simulations, microstructural observations and tensile testing. The major conclusions of this study can be summarized as follows: 1. Sound beads could be obtained in fillet welding position, by melting and spreading aluminium on solid steel, and creating a 2–10 m thick reaction layer formed by Fe2 Al5 intermetallic phase along the entire interface. 2. The interface layer thickness was estimated between 2 and 25 m on the scanning speed range (0.6–1.4 m/min). 3. The presence of a zinc layer on steel favoured liquid aluminium spreading and reaction kinetics during intermetallic formation, with thicker interfaces for identical (P,V) laser conditions. 4. FEM simulations allowed us to estimate thermal history at the Al–steel interfaces: 750–1000 ◦ C maximum values were obtained for 0.1–1 s time maintains of aluminium in liquid state. 5. Joint strengths obtained on 1 mm thick assembled sheets under tensile loading were found to be close to 110 MPa for non-galvanized steel, and 230 MPa for galvanized steels with flux. The dominant failure mode was interfacial shear failure on non-galvanized assemblies, tensile failure in the fusion zone on galvanized samples without flux, and tensile failure in the heat affected zone for galvanized samples with flux. These mechanical resistances seem to be rather competitive compared with other techniques investigated for assembling steel to aluminium.
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6. The prediction of inter-metallic growth with simple parabolic kinetics provides a rather good fitting with experimental values but cannot be considered as a predictive method for calculating intermetallic thicknesses.
[8] [9] [10]
References
[11]
[1] M. Kreimeyer, F. Wagner, G. Sepold, Proceedings of 23rd International Congress on Applications of Lasers and Electro-Optics (ICALEO’2004), San-Francisco, USA, 2004. [2] Aluminum and Aluminum Alloys Handbook, ASM International, 1993 (Chapter “Aluminum coatings”). [3] A. Bahadur, O.N. Mohanty, Mater. Trans. JIM 36 (9) (1995) 1170– 1175. [4] H.R. Shahverdi, M.R. Ghomashchi, S. Shabestari, J. Hejazi, J. Mater. Proc. Technol. 124 (2002) 344–352. [5] Z. Sun, J.C. Ion, J. Mater. Sci. 30 (1995) 4205–4214. [6] M.J. Rathod, M. Kutsuna, Weld. J. (2004) 16–26. [7] A. Mathieu, S. Matte¨ı, A. Deschamps, M. Su´ery, L. Rodriguez, D. Grevey, “Assemblage Acier/Aluminium par faisceau laser, Proceedings of
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LASERAP’5 Congress, Sarlat, France, Politehnica School of Timisoara, 2003 Rom. K. Bouch´e, F. Barbier, A. Coulet, Mater. Sci. Eng. A 249 (1998) 167–175. V.I. Dybkov, J. Mater. Sci. 25 (1990) 3615–3633. M.L. Giorgi, J.B. Guillot, R. Nicolle, J. Mater. Sci. 40 (9–10) (2005) 2263–2268. S. Fukumoto, H. Tsubakino, K. Okita, M. Aritoshi, T. Tomita, Mater. Sci. Technol. 15 (1999) 1080–1086. K. Kimapong, T. Watanabe, Weld. J. (October) (2004) 277–282. C. Maldonado, T.H. North, J. Mater. Sci. 37 (2002) 2087–2095. S. Katayama, M. Mizutani, Proceedings of 22nd International Congress on Applications of Lasers and Electro-Optics (ICALEO’2004), Scottsdale, USA, 2003. E.A. Brandes, G.B. Brook, Smithells Metals Reference Book, seventh ed., Butterworth & Heineman, 1992. W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (6) (1992) 1564–1583. ABAQUS software, User’s Manual, Version 6.1, Hibbitt, Karlsson and Sorensen Inc., 2001. M. Merklein, A. Giera, M. Geiger, Steel Res. Int. 76 (2–3) (2005) 250–256. A. Bouayad, Ch. Gerometta, A. Belkebir, A. Ambari, Mater. Sci. Eng. 363 (2003) 53–61.
Generation of aluminium–steel joints with laser-induced reactive wetting P. Peyre a,∗ , G. Sierra b,c , F. Deschaux-Beaume c , D. Stuart a , G. Fras c a
GERAILP-LALP, UPR 1578 CNRS, 94114 Arcueil, France b CEA/DRT/DTEN/GERAILP, 94114 Arcueil, France c LMGC, UMR 5508 CNRS, Universit´ e de Montpellier II, 30907 Nˆımes, France Received 1 December 2005; received in revised form 1 July 2006; accepted 1 September 2006
Abstract A new mean of assembling steel to aluminium was developed, following previous work by German workers [1]. In this new method, a laser-induced aluminium melt pool spreads and wets a solid steel, to generate, after solidification a sound and resistant interface layer. Joint properties were investigated, in terms of surface aspects, interface microstructures and mechanical resistances under tensile testing, for non-galvanized and galvanized DC04 steels. Thermal and diffusional finite element (FE) simulations were also carried out to calculate temperature history at interfaces, and reaction layer thickness. The 2–20 m thick reaction layers formed all along the interface were found to be mostly composed of Fe2 Al5 intermetallic compound with a high hardness (1200 HV) and rather low ductility (presence of solidification cracks). The presence of a 10 m thick Zn layer on the steel was shown to have a beneficial influence on the wetting characteristics of the joint, despite the formation of occluded pores in the melt pool due to Zn vaporisation. FE thermal modelling evidenced 760–1020 ◦ C wetting temperatures at the interface between DC04 low carbon steel and 6016 aluminium sheets, with time maintains of the melt pool in the 0.2–0.5 s range, resulting in high-speed reaction kinetics. Using these temperature data, diffusion calculations were shown to provide a rather good prediction of intermetallic thicknesses. Tensile tests were considered on aluminium–steel lap joints and evidenced higher mechanical resistances (220 N/mm linear tensile strength) on galvanized steels, provided that fluxing of the steel surface was carried out prior to welding to avoid zinc vaporisation. Comparatively, non-galvanized assemblies exhibited much lower mechanical resistances (170 N/mm resulting in a 90 MPa interfacial shear strength). It was concluded that the laser-induced wetting technique is a rather effective way for generating Al-steel joints without filler material, and that it should be considered as a competitive technique versus solid assembly modes (friction stir welding . . .). © 2006 Elsevier B.V. All rights reserved. Keywords: Aluminium–steel; Laser; Wetting; Intermetallic compounds
1. Introduction Many processes like hot dip coating against corrosion [2,3], generation of metal matrix composites, or brazing and welding of dissimilar materials [1], are based on the interaction of a solid metal or alloy with a liquid metal [4]. During the interaction time, the solid surface dissolves in the melt, and intermetallic layers are subsequently formed at the interface. Usually, the formation of brittle intermetallic layers limits the mechanical resistance of the transition zone between dissimilar metals [5], when their thickness exceeds a permissible value. This is particularly true for aluminium–steel
∗
Corresponding author. E-mail address: [email protected] (P. Peyre).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.09.111
interfaces which have been investigated already [5–7] because of a huge industrial interest, and which are known to be unfavourable with respect to the mechanical resistance of the joints. The wide diversity of interfacial processes involved in materials joining (diffusion, chemical reaction and fluxing) suggests complicated phenomena during wetting, and wetting kinetics which can be either diffusion controlled or reaction controlled, depending on the systems considered. Several authors have focused on the interface reactions between aluminium and steel in isothermal conditions, to estimate the reaction kinetics and diffusion constants at various temperature, by dipping iron or steel samples in molten aluminium baths at temperatures ranging between 700 and 900 ◦ C [3,4,8]. Microstructure analysis have been carried out on the interface layers, and most of the reported investigations exclu-
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sively identified two distinct intermetallic phases: Fe2 Al5 as the major constituent close to the steel substrate, growing according to a parabolic law, and FeAl3 as the minor phase close to the solidified aluminium [8,9]. After solidification, these intermetallic phases provide bonding of aluminium with steel. Recently, authors [10] have also considered theoretically the interfacial reactions between solid iron and liquid zinc–aluminium alloys. Since brittle intermetallic phases are considered as a critical issue, other techniques, like friction welding [11–13] have been applied successfully to produce butt joints between aluminium and steel without generation of a melt pool, but through the friction of a rotating pin which puts the aluminium in a fluid-like plastic state adhering to the steel surface. In spite of intermetallic issues, several attempts have been made recently (2001–2004) to realize aluminium–steel assemblies with Nd:YAG lasers. Different approaches have been used which can be classified into: (1) a laser brazing technique with filler wire [7], (2) a classical key-hole (vapour capillary) welding mode [14], and (3) a laser roll welding mode where steel remains solid and laser heating provides aluminium melting and subsequent dissolution–diffusion–reaction of steel [6]. All these joining techniques have succeeded in providing acceptable mechanical strengths (in tensile or shear modes), when limiting the interface layer thickness (<10 m) for brazing or reactive wetting, or limiting the aluminium–steel dilution, for the key-hole technique where both metals are melted. Analytical work about diffusion aspects in laser-induced reactive wetting of aluminium with steels have never been published, and seem rather difficult to perform because of: (1) the non-homogeneous spatial profile of lasers, (2) the out-ofequilibrium phenomena involved at high heating and cooling rates, that limit the use of phase diagrams and (3) the strongly non-isothermal nature of laser-heating, and the shorter reaction times than those previously used during isothermal solid steel-liquid aluminium investigations (usually superior to 100 s [3,4,8]). More, the complex geometries of laser assemblies require thermal simulations (with finite element codes), to evaluate the time–temperature histories T = f(t) of interfaces before investigating growth kinetics aspects. In this paper, the generation of aluminium–steel assemblies with a highly non-isothermal laser-induced reactive wetting was addressed. The different aims of the present work were: (1) First, to optimise experimental conditions (welding speeds, overlaps) for the realisation of low carbon steel—6000 aluminium assemblies, using a dedicated experimental set-up, with a specific attention paid to the fast camera analysis of melt pools and to temperature measurements. (2) To make metallurgical and mechanical analysis of assemblies, focusing on interfacial layers. (3) To assess temperature evolution at the aluminium–steel interface with 3D FEM simulations, and to compare them with interfacial thicknesses, in order to envisage further diffusion calculations.
Table 1 chemical composition of DC04 low carbon steel and 6016-T4 aluminium alloy Element (wt.%)
Al
Fe
Mg
Si
Mn
P+S
C
DC04 6016
– Bal.
Bal. 0.5
– 0.3–0.5
– 1–1.3
0.4 0.2
0.06 –
0.08 –
(4) To compare the reactive wetting of galvanized and nongalvanized steels and the mechanical resistance of joints. 2. Materials The materials used for these investigations were a low carbon steel DC04 mostly used for automotive applications, defined by the european normalisation NF EN 10130 and a 6016-T4 aluminium alloy for both aeronautical and automotive issues. Their chemical composition, temperature-dependent physical and mechanical properties are reported in Tables 1 and 2, respectively. The 1.2 mm thick DC04 sheets were provided with a non-galvanized state or a 20 m zinc-coated layer containing nearly 0.2% Al to avoid formation of brittle Fe(Zn)x compounds. On the other hand, 1 mm thick 6016 aluminium sheets were used. Mechanical properties reported in Table 2 were evaluated by quasi-static tensile tests. 3. Experimental procedure 3.1. Reactive wetting laser experiments Respectively, 1.2 and 1 mm thick sheet specimens of DC04 steel and 6016 aluminium alloy were used. Samples were prepared by grade 800 SiC paper polishing, degreasing and (for the non-galvanized steel surface) fluxing with a fluor-base anticorrosive brazing flux. The flux suspension (flux powder dissolved in ethanol), was sprayed rather homogeneously on the surface (10 m < estimated thickness < 50 m), and dried before welding. On galvanized steel sheets, as already shown in Ref. [7], no fluxing was shown to be required to ensure wetting of the zinccoated steel surface, due to the good metallurgical compatibility between the Fe, Al and Zn. However, a few comparative tests have also been carried out with fluxing. Lap wetting tests (Fig. 1) were carried out with a Nd:YAG continuous wave (cw) laser operating in the 2.5–3 kW power range, with scanning speeds between 0.6 and 1.4 m/min. Operating with a 200 mm focal lens, laser was used in defocused condition, with a 30◦ irradiation angle, resulting in a elliptic spot of 2.3 mm × 1.8 mm, and laser intensities in the 70–90 kW/cm2 . All wetting tests were carried out in flat position (using the gravity contribution to the aluminium spreading) on a 100 mm testing length, with the laser beam located on the upper aluminium sheet (Fig. 1a and b), and with an argon shielding gas nozzle (at nearly 10 l/min) to protect heated surfaces from oxidation. For each assembly, a fast camera (Cmos 2000 Hz) recording of the aluminium melt pool (Fig. 2) was done, and temperature
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Table 2 Thermo-physical and mechanical data of DC04 steel and 6016-T4 aluminium alloy [15] Physical properties
ρ (kg/m3 )
λ (W/m/K)
Cp (J/Kg/K)
(20 ◦ C)
(20 ◦ C)
Latent heat (J/kg)
σ Y (MPa)
σ UTS (MPa)
160
280
140
230
DC04 steel
7810
65 45 (400 ◦ C) 28 (800 ◦ C)
470 595 (400 ◦ C) 875 (800 ◦ C)
3.105
6016-T4 alloy
2700
172 (20 ◦ C) 190 (560 ◦ C)
887 (20 ◦ C)
3.8.105 (560 ◦ C–640 ◦ C)
(1450–1510 ◦ C)
Fig. 1. Schematic description of the experimental set-up: (a) global overview and (b) detailed view of the reactive aluminium–steel wetting experiments (α = wetting angle).
measurements near the interaction zone were carried out with 0.2 mm diameter chromel–alumel thermocouples. This allowed us to collect input data for FEM simulations of the process. 3.2. Metallurgical characterizations From the 100 mm-long wetting joints, specimens were cut across the fillet joint seam for macrostructure and microstructure observations. Cross sections of the joints were polished by 220, 500 and 1000 SiC grades, followed by a 3 and 1 m diamond suspension, to a mirror-like surface aspect. Etching with Keller’s reagent (HNO3 + HCL + HF + H2 O) was used to reveal the aluminium microstructure and Nital 4% was used to attack the steel structure. Then, microstructures of the interfaces were observed using: (1) optical microscopy, at magnifications ranging between ×50 and ×200, and (2) scanning electron microscopy (SEM) with an energy dispersive X-ray (EDX) analysis facility to determine the chemical composition of intermetallic phases formed during steel wetting by aluminium melt pool.
Microhardness (10–25 g loading force) and nanoindentation1 (0.1–2 g) tests were also carried out to estimate local mechanical properties of affected zones and interfaces. A more detailed presentation of the nano-indentation technique is available in Ref. [16]. 4. Thermal simulations with the Finite Element Method 4.1. FEM conditions A three dimensional (3D) finite element model was implemented in the finite element software ABAQUSTM Standard [17] with a *Dflux Fortran subroutine to simulate a heat moving source, and a *Heat Transfer card to solve the heat flow equation. The heat was considered to be deposited onto the metal surface with a elliptic intensity distribution, corresponding to a defocused laser irradiation, and a 32% beam absorption
1 The authors wish to thank MC.Ste Catherine and L. Quiniou from CEPArcueil (DGA) for nano-indentation tests.
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used: a uniform 10 W/m2 K value on all the external sides of the body, and h = λ/e = 1500 and 700 W/m2 K values on the lateral vertical edges of, respectively, aluminium and steel sheets. The influence of a 20 m Zn layer was also considered in our FEM description of the thermal history of interfaces. The only effect evidenced was a small temperature decrease (−40 ◦ C maximum value) at the steel–Al interface. 4.2. Process modelling and experimental validation
Fig. 2. Fast camera (1000 Hz) coaxial analysis of the melt pool dimensions during a reactive wetting aluminium–steel assembly (V = 1.2 m/min and P = 3 kW): melt pool length is between 5.3 mm (a) and 6 mm (b).
on aluminium measured by integrated sphere in solid state. Up to 12,000, eight nodes diffusive-conductive hexaedral elements (DC3D8) were used to make a 3D meshing of the overlap aluminium–steel joint. The mesh sizes were chosen in such a way that smaller elements corresponded to the irradiated zone, where flux gradients are maximal. The computation domain was made out of two sheets (20 mm long × 11 mm width × 1 mm thick) with a 5 mm overlap on X-axis (Fig. 3a). Temperature dependent (non-linear) physical properties, reported in Table 2, were taken from Ref. [15]. In the numerical simulation, the latent heat of fusion of each material was taken into consideration, but not the hydrodynamic flow contribution to heating (due to fluid flow in the molten pool, which velocity was estimated to 0.2–0.5 m/s by fast camera analysis). It should be mentioned that in previous work, the thermal conductivity of the liquid state was artificially increased [7] to account for the enhanced heat transport due to surface-tensions-driven convections. The aim of these investigations was to evaluate the thermal history T = f(t) at the aluminium–steel interface, fitting our simulations (and especially the absorption coefficient of the Nd:YAG beam by aluminium or steel) with experimental data from thermocouple measurements near the interaction zone. Data from real-time fast camera analysis of melts (fusion zone dimensions) were also used to optimise our simulations. Concerning the boundary conditions, heat loss coefficients h were used, which traduced the cooling flux due to a laminar flow of protective gas (Qgas = −hgas (T − T0 )), but also the heat dissipation through the clamping device and the lateral sides of the sheets (Qsolid = −hsolid (T − T0 ) with hsolid = λ/e, with λ = conductivity (W/m/K) and e (m) = thickness/length of the part). For our simulations, different h values were
All calculations were carried out with a standard resolution algorithm. The energy supplied by the moving laser head was modelled by a Fortran routine in form of a distributed heat flux (W/m2 ) applied on the external edge of surface elements. It enabled us to define the position of the elliptic laser deposit all along the time range (less than 1 s total time), and for different scanning speeds (between 0.7 and 1.2 m/min). Results of the simulations were obtained in the form of timedependent temperature distributions (local thermal history), and 3D thermal fields. This enabled us to compare T = f(t) simulations with thermocouple data recorded on the steel surface at a “d” distance (between 3 and 6 mm) from the interface. Also, the dimensions of the fusion zone could be compared with fast camera recordings (Fig. 2). Results of calculations, carried out with a 32% absorption of the Nd:YAG light, are shown in Fig. 3 for a 1.2 m/min scanning speed. A 1600 ◦ C peak temperature was obtained in the energy deposit zone of the melt pool, that seemed to be rather high. On the other hand, dimensions of the melt pool (nearly 6.9 mm length and 3.2 mm width) and most of all, thermal history taken
Fig. 3. FEM simulation of a 1.2 m/min–3 kW aluminium–steel laser-induced reactive wetting: (a) thermal fields in the fusion zone in quasi-steady condition (aluminium up, steel down) and (b) upper view of the simulated fusion zone: the bead is 6.9 mm length can be compared with the 5.6 mm experimental value from Fig. 2.
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tion could be mostly controlled by reaction phenomena instead of diffusion mechanisms. 5. Macrostructure and microstructure of joints The results presented hereafter correspond to morphological and metallurgical analysis carried out on Al–steel interfaces produced by the reactive wetting assembly mode. 5.1. Bead aspects
Fig. 4. (a) Comparison between experimental thermocouple values and simulated T = f(t) profiles at 4 mm distance from the interface; (b) thermal histories T = f(t) at the aluminium–steel interface during a reactive wetting test (3 kW, 0.7–1.2 m/min).
at 4 mm distance from the interface, were in correct agreement with experimental data. Following the same procedure, simulations of 0.7–1.2 m/min processes could be carried out and successfully compared with experience (Fig. 4a). The most attractive information from these simulations, is the interfacial temperature history T = f(t), corresponding to the wetting temperature at the molten aluminium–solid steel interface. For the scanning speed range considered, 750–1000 ◦ C peak temperature were obtained (Fig. 4b), for 0.2–1 s time maintain of the wetting. This corresponds to very sharp and short thermal loadings compared to classical steel immersion into aluminium baths (usually 700–900 ◦ C constant temperature during several minutes immersion times) [4,8,9]. Following previous work by Refs. [4,8], it should be expected that, for such short interaction times, intermetallic layer forma-
The effect of travel speed on the appearance of laser-wetted joints is shown in Figs. 5 and 6 on a non-galvanized steel surface (Ra = 0.1 m). In spite of different thermal cycles, the global morphology of beads is kept rather constant, even though the bead widths are reduced with increasing speeds (nearly 4 mm at 0.8 m/min and 3 mm at 1.2 m/min). Also, different wetting (contact) angles are obtained (between 30◦ and 80◦ ), revealing more spreading of aluminium on steel for lower velocities (Fig. 6 and Fig. 7). However, for the same travel speed, contact angle may vary drastically (±20◦ ), due to local changes in heat dissipation (local gaps between upper and lower sheets). When comparing wetting angles obtained on galvanized and non-galvanized steels, one thing is obvious: the zinc layer improves the spreading and wetting of aluminium on steels’ surface. For instance, the same experimental conditions (V = 1 m/min, P = 3 kW and 2 mm spot) generate a 45◦ wetting angle on bare steel versus a 25◦ angle on galvanized surface. This is mostly due to the fact that, on galvanized specimens aluminium spreading occurs on a liquid zinc layer (Tmelt = 419 ◦ C), and not on a solid steel. Under that condition, we cannot talk any more about wetting of a solid surface. On the other hand, if wetting experiments are carried out on a fluxed galvanized surface, wetting angles tend to be strongly increased, but still remain lower than those measured on bare steels (Fig. 7).
Fig. 5. Macroscopic view of the aluminium–steel joints (P = 3 kW and spot diameter = 2.3 mm) for two different scanning speeds: (a) 1.2 m/min and (b) 0.8 m/min. Some geometrical defects (process instabilities) are randomly present but the global aspects are acceptable.
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Fig. 6. Cross-sections of aluminium–steel joints for two different welding speeds (0.6 and 1.2 m/min), on non-galvanized (a and b), and galvanized (c and d) conditions.
5.2. Microstructures of interfaces obtained between non-galvanized steel and aluminium On all the interfaces, the aluminium appears to be bonded to the non-galvanized steel substrate by a thin layer of a very narrow, uniform and defect-free Fe–Al phase.
Fig. 7. Influence of laser scanning speed on the wetting angle α for galvanized and non-galvanized steels with or without fluxing.
The interface thicknesses, estimated with optical analysis (Fig. 8a and b), are nearly constant versus scanning speed (Fig. 11), and remain in the 2–10 m width range, except on specific locations (Fig. 9) where thicker – and most of times cracked – interfaces could be evidenced. These heterogeneities may be due either to some variations of gap between Al and steel sheet, or to local variations of flux thickness prior to wetting. According to previous data on hot-dip coating [2], the thickness of the intermetallic layer depends on the viscosity of the aluminium bath, and decreases with increasing amounts of additive elements in the melt pool. For instance, Si, Cu and Be are known to have the greatest effect in preventing the build-up of the intermetallic layer, and minimizing its thickness. The inter-metallic morphology consists of an inner compact layer, and an outside “composite” layer composed of broken and/or floating particles in the Al matrix (Fig. 8b and Fig. 9). If we have a closer look at these interfaces, it appears that intermetallic layers exhibit serrated interfaces on the aluminium side. This phenomenon is not that obvious between aluminium and steel, and previous investigations using much longer liquid–solid contact times, have found such an aspect in the opposite interface, i.e between steel and intermetallic phase. Under our experimental conditions, only local zones with thicker interfaces (see Fig. 8) exhibit serrated interfaces on both sides. Micro-porosities (a few m), found at various locations along the interfaces, and frequently containing crystallized fluxing
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Fig. 8. Interfaces obtained after laser-induced wetting of non-galvanized (a) and (b), or galvanized (c) and (d) steels at 0.6 and 1.2 m/min. Interfaces obtained on galvanized steels are thicker (37 m vs. 5 m at 0.6 m/min and 10 m vs. 3 m at 1.2 m/min) because of the influence of flux that limits inter-metallic formation.
compounds, were clearly expected to originate from the volatile flux thickness. Their influence on fracture behaviour will be discussed later on. Semi-quantitative EDX investigations carried out on a 0.8 m/min interface indicated a small compositional gradient along the intermetallic layer: nearly 24% Fe 75% Al and 1% Si (wt.%) close to the Al-intermetallic interface, and 26% Fe 72% Al and 2% Si near the steel interface (Fig. 10). Both compositions are closer to Fe2 Al5 (Si) than FeAl3 (Si) sometimes evidenced [3] in Al–steel interfaces. The Si-enrichment near the steel side is not explained (a Si-enrichment would be more likely to occur in the upper part of the melt pool, where solidification ends), but convection movements in the melt pool could modify compositional zones. It also has to be noticed that Si is known to have an inhibiting effect on the formation and growth of Fe2 Al5 intermetallics. Also, FeAl3 needle-like crystals, in-between the intermetallic layer and the aluminium melt-pool (Figs. 9 and 10) are only evidenced at the centre part of reaction layers width, where the aluminium–steel interface experienced the higher temperatures. These crystals, attached to or separated from reaction layers, could be due to iron re-dissolution and diffusion in the aluminium melt pool, possibly favoured by convection movements of the liquid aluminium close to the intermetallic interface.
Fig. 9. Cross-sections of aluminium–steel joints: detail of a local 14 m-thick intermetallic interfacial zone (upper part, solidified aluminium): presence of intermetallic laths and solidification cracks (0.8 m/min, 3 kW).
5.3. Microstructures of assemblies obtained on galvanized steels As indicated in Section 5.1, the zinc layer improves the wetting of aluminium on steel, but this phenomenon is inhibited by a fluxing of aluminium sheets.
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Fig. 10. SEM-EDX analysis of a 5 m intermetallic layer: a compositional gradient is found between the aluminium and the steel side, but the global stoechiometry is close to Fe2 Al5 .
Interfaces obtained on galvanized steels exhibited a Fe2 Al5 Znx stoechiometry [10] with numerous solidification cracks when the interface thickness is above 10 m. Their global aspect is held unchanged compared with those obtained on nongalvanized steels, but their thickness is much higher (up to 30–40 m) than on non-galvanized steels (Fig. 8c and d, Fig. 11) when experiments are performed without flux. This increase seems to be more due to the influence of the flux itself, that limits the intermetallic growth, than to the higher reactivity of the galvanized surface. Indeed, non-galvanized and non-fluxed steels in contact with liquid aluminium exhibit very thick intermetallic even if wetting is close to zero, due to the non-melted oxide layer on aluminium melt pool.
Fig. 12. Microstructure of the molten aluminium generated during reactive wetting on DC04 galvanized steel: formation of pores by zinc vaporisation in the fusion zone and presence of cracks due to Zn local segregation.
The main difference with non-galvanized steel assemblies comes mainly from the ability of Zn to vaporize under laser heating, when the interface temperature reaches 900 ◦ C. This induces large porosities (50–300 m), mostly located in the inner part of the fusion zone, close to the heat-affected zone of aluminium. Combined with the softening effect evidenced in both FZ and HAZ of Al alloy, these porosities are potential locations for cracks initiation. It was also found that Zn dissolution in molten aluminium tends to increase cracking susceptibility inside aluminium melt pool, maybe due to the formation of low melting point Al–Zn compounds at interdendritic spaces (Fig. 12). 6. Microhardness and nano-identation results 6.1. Microhardness results
Fig. 11. Influence of laser scanning speed on the intermetallic thickness, during laser-induced reactive wetting of galvanized or non-galvanized XES steels. Fluxing reduces by a factor 7–8 the intermetallic thickness.
The degree of softening of age-hardened 6016-T4 (100 HV0.025 average hardness) was evaluated by Vickers hardness measurements, with a 25 g loading force, at a 500 m distance from the steel–aluminium interface. Some −35% maximum hardness reductions were obtained, independent on the welding speeds and on the presence of a galvanized layer on steels. The highest degree of softening was obtained in the fusion zone (FZ), and to a lesser extent in the adjacent heat affected zone (HAZ) where precipitates were dissolved or put into solid-solution. Rather wide heat-affected softened zones (more than 5–6 mm) were obtained, with a hardness gradient between the material immediately adjacent to the fusion zone (65 HV = −35%), and a second heat-affected zone (85 HV) where precipitates are not expected to be dissolved, but instead, to coarsen and lose their coherency with the Al matrix. The formation of wide softened regions was associated to a decrease of local strength properties. On the other hand, the steel hardness and microstructure were not affected by the wetting process.
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6.2. Nano-indentation characterization of interfaces
7. Tensile tests
Due to the very thin interface layers, nano-indentation tests, carried out with a Berkovich indenter and 20 mN applied force, were carried-out on different interface thickness, starting with thicker layers (Fig. 13). On all the interfaces investigated so far, results exhibit some rather constant 1100 ± 100 HV20 mN equivalent values in the intermetallic layer, compared with an average 250 HV20 mN value in the steel and 200 HV20 mN in the molten aluminium alloy. These hardness values did not depend on process parameters, and seemed to be in correct agreement with previous data from [6]. They also confirmed a very sharp mechanical gradient close to the interfacial zone. Last, no hardness difference was evidenced between interfaces formed on non-galvanized or galvanized steels, where the chemical composition is close to Fe2 Al5 Znx ( phase) [10]. Also, elastic modulus values in Fe2 Al5 , estimated from the slopes of elastic unloadings, were found to exhibit an average level of 220 GPa.
For ease of processing, 20 mm width lap joint specimens have been used to test aluminium–steel joints, despite tilting of tensile specimens under load, due to bending moments of the lap joints during mechanical testing [18]. The joint strength was evaluated by tensile testing at room temperature, with an initial gauge length of 15 mm (specimen length = 80 mm). Mechanical properties were shown to be rather different, depending on the presence of a galvanized layer or not. On nongalvanized assemblies, maximum resistance of up to 190 N/mm were obtained, with two different failures modes: (1) a brittle interfacial shear fracture or (2) a tensile failure in the fusion zone (Fig. 14). On galvanized steel assemblies, most of the failures occurred in the highly softened fusion zone. Corresponding resistances were evaluated to 130–140 N/mm (Fig. 15). These failures were
Fig. 13. Nano-indentation tests (0.02 N applied force, Berkovich indenter) of a 0.8 m/min specimen (intermetallic thickness = 12 m): (a) localisation of indents and (b) hardness values.
Fig. 14. Tensile curves obtained (a) on non-galvanized joints with 2 failure modes (interfacial brittle failure and failure in the FZ) and (b) on galvanized joints with 2 failure modes (in the HAZ with striction and in the fusion zone, promoted by Zn porosities).
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global properties of dissimilar assemblies. However, our results indicate that this assumption is not valid on age-strengthened materials such as 6016-T4 because most of the failures occur in the solidified aluminium and/or in the softened heat affected zones in the case of galvanized steels. This competition between interfacial failures and fusion zones fractures has been reported already in Ref. [13]. 8. Discussion on aluminium–steel interfaces
Fig. 15. Relationship between mechanical resistance (N/mm) and welding speeds: (a) on non-galvanized steels with two failure modes (1) interfacial failure of Fe2 Al5 and (2) tensile failure in the HAZ and (b) on galvanized steels with 2 failure modes in the HAZ or FZ.
promoted by a high density of pores in the melt-pool (Fig. 12) due to Zinc vaporization above 900 ◦ C, which reduced the section area S (mm2 ) and in turn increased the applied stress F(N)/S (mm2 ) = σ (MPa). In that case, Al–Fe–Zn interfaces were shown to be more resistant than the solidified and annealed 6016 containing pores and cracks (Fig. 12). Interesting results were obtained on galvanized samples with fluxing, even if flux was not shown to be crucial to ensure wettability between aluminium and a zinc-coated steel. In that specific case, maximum resistances of 230 N/mm (resulting in 230 MPa tensile strengths) were shown, with failures located in the first heat affected zone adjacent to the FZ. This result, very close from the fracture strength of 6016-T4 (Table 2), is rather surprising because failure occurred in partially softened zones (rather far away from the fusion zone) where local mechanical properties are higher than those of the FZ. It is tacitly assumed that the mechanical properties of materials adjacent to the joint interface are not important regarding the
Sound Al–steel interfaces could be obtained by the laserinduced reactive wetting assembly mode, with average interface thickness in the 2–15 m range. Rather good mechanical resistance were shown on all the assemblies, whatever the scanning speed and the wetting angles. However, questions remain unanswered about the mechanisms involved in the interface formation, for such short liquid–solid interaction times (0.1–0.5 s). If we first consider the aluminium–steel reactions without zinc layers, we can refer to previous works [4,8,9,19] in isothermal conditions. According to them, the initial growth steps of the interface layer occurs from steel towards aluminium, due to the ease of diffusion of iron atoms into molten aluminium (DFe→liq Al = 8 × 10−8 m2 /s at 1000 ◦ C). Then, a further increase of the Fe2 Al5 thickness modifies diffusion conditions and it becomes necessary to consider the diffusion of Al and Fe atoms through Fe2 Al5 to predict further intermetallic growth. Results presented in our work revealed very thin reaction layers, where the main factors contributing to the growth were expected to be controlled by a reaction instead of diffusion mechanism between steel and aluminium. The mechanisms involved could be as follows: (1) first, iron dissolution takes place towards a thermodynamic metastable equilibrium (with JFe→Al = kdiss (CFe meta − CFe int )), (2) iron diffuses in the aluminium melt pool but remains rather close to the Fe–Al interface (during a 1000 ◦ C–0.15 s thermal cycle, the diffusion path of Fe in Al is less than 200 m), (3) due to the high concentration of dissolved Fe in Al, supersaturation occurs and Fe2 Al5 nucleation occurs, followed by growth. So it can be considered that iron atoms present in Fe2 Al5 come from the dissolved iron occurring in the initial steps [10]. If we consider diffusion coefficients of Fe through Fe2 Al5 from Bouayad et al. [19] at 800 ◦ C (k = 9 × 10−11 m2 /s) or 700 ◦ C (k = 3.9 × 10−11 m2 /s) for low carbon steels, and calculate the diffusion paths, using equivalent times for T above 700 ◦ or 800 √ C (see Fig. 7b), we can compare these calculated values (ecalc 2ktT =700◦ C ) with experimental intermetallic thicknesses (Fig. 16). The agreement is rather correct, especially for the thickness amplitude, even if the evolution with time seems to be underestimated by a simple parabolic approach. Consequently, even if a simple parabolic dependence of the interface thickness against time (d(m) = Kt1/2 ), seems to be rather controversial, due to the complexity of laser-induced interfacial conditions, it seems to reproduce rather well the experimental intermetallic thicknesses. Considering the very specific local conditions induced by laser wetting, we can remind that, due to the small volume of molten aluminium (typically 5 mm3 ), tiny amounts of iron dis-
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Fig. 16. Maximum intermetallic thicknesses on non-galvanized steels vs. the inverse value of the welding speed square root (V−0.5 ). Experimental thickness is estimated with optical microscopy carried out in the middle of each crosssection. Calculated values use the simulated time maintains above 700 or 800 ◦ C with corresponding diffusion coefficients from [19] (9 × 10−11 m2 /s at 800 ◦ C, 3.9 × 10−11 m2 /s at 700 ◦ C).
solution may contribute to the saturation of molten aluminium with iron, and in turn modify (increase) the Fe2 Al5 growth rate. We also could consider that laser-induced reactive wetting involves hydrodynamic flow (see convection paths in Fig. 12) in the liquid aluminium (fluid velocity < 1 m/s), which may affect strongly the reaction mechanism, and possibly promote the dissolution or diffusion rates of the steel as already shown in Ref. [9] on rotating steel samples. Oncoming thermo–diffusional modelling should allow to make more precise predictions of intermetallic thicknesses. As indicated earlier in the paper, high temperature gradients, and high fluid velocities at the center part of interfacial width also promoted the formation of needle-like FeAl3 crystals in aluminium melt pool. However, these crystals are not expected to affect the mechanical resistance of assemblies, as failure usually initiates at the edges of reaction layers, and not on the centre part of layers where those needle-like crystals are located. With respect to other techniques for assembling steel and aluminium, the reactive wetting mode seems to offer a promising potential, because of an interesting balance between metallurgical and mechanical properties on one side, and rather attractive process conditions (high scanning speed . . .) on the other side. If we now consider the presence of the 20 m zinc layer between liquid aluminium and steel, a few comments can be make to explain our results. First, the reactivity between aluminium and zinc (%solubility) is much better than between Al and steel, as confirmed by the Al–Zn diagram. This allows good wetting and generation of intermetallic compounds without fluxing steels. Second, the presence of a 20 m zinc layer which vaporizes under laserirradiation strongly modifies the microstructure of assemblies
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near the interfacial zones, and in turn the failure modes, mainly located in the fusion zone. No more interfacial failure occurred on galvanized steel assemblies with or without fluxing. The reason why intermetallic layers with Zn are more shear-resistant (or more adherent) than pure Fe2 Al5 is still under investigation, but identical local hardness values seem to indicate that it may not be due to a change in ductility. The influence of fluxing on galvanized steel parts may be explained as follows: during the laser heating, the liquid flux (Tmelt = 560 ◦ C) below the aluminium melt (maximum temperature below 1000 ◦ C: see Fig. 6) acts as a thermal barreer upon the melted zinc (Tboil = 910 ◦ C), and limits the thermal rise and vaporization of the galvanized layer. Consequently, zinc bubbles are avoided, zinc occlusion is limited in the melt pool and mechanical properties are enhanced due to the change of failure mode (from the fusion zone to the adjacent heat affected zone). Last, as mechanical properties of the galvanized joints are more limited by softening effects than by the intermetallic resistance itself, non heat-hardenable aluminium alloys (Al–Mg alloys) could be appropriate candidates for providing higher mechanical resistances with galvanized steels. 9. Conclusions 6016 aluminium alloy and DC04 low carbon steel were joined by laser-induced reactive wetting mode. The wetting assemblies were investigated by process analysis (fast camera, thermocouples), FEM simulations, microstructural observations and tensile testing. The major conclusions of this study can be summarized as follows: 1. Sound beads could be obtained in fillet welding position, by melting and spreading aluminium on solid steel, and creating a 2–10 m thick reaction layer formed by Fe2 Al5 intermetallic phase along the entire interface. 2. The interface layer thickness was estimated between 2 and 25 m on the scanning speed range (0.6–1.4 m/min). 3. The presence of a zinc layer on steel favoured liquid aluminium spreading and reaction kinetics during intermetallic formation, with thicker interfaces for identical (P,V) laser conditions. 4. FEM simulations allowed us to estimate thermal history at the Al–steel interfaces: 750–1000 ◦ C maximum values were obtained for 0.1–1 s time maintains of aluminium in liquid state. 5. Joint strengths obtained on 1 mm thick assembled sheets under tensile loading were found to be close to 110 MPa for non-galvanized steel, and 230 MPa for galvanized steels with flux. The dominant failure mode was interfacial shear failure on non-galvanized assemblies, tensile failure in the fusion zone on galvanized samples without flux, and tensile failure in the heat affected zone for galvanized samples with flux. These mechanical resistances seem to be rather competitive compared with other techniques investigated for assembling steel to aluminium.
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6. The prediction of inter-metallic growth with simple parabolic kinetics provides a rather good fitting with experimental values but cannot be considered as a predictive method for calculating intermetallic thicknesses.
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