Welding of rapidly solidified alloy 8009

TWI Bulletin, November/December 1995

Mike Ellis has a BSc(Eng) in Metallurgy from Imperial College and a PhD on the subject of fracture mechanisms in a 2 ¹/ ₄ Cr-1Mo pressure vessel steel, from Cambridge University. He has worked as a Research Associate for the National Research Council of Canada (NRC) and as a Senior Metallurgist for Alcan International Ltd in Banbury. Mike joined TWI's Materials Department in 1992, and has since completed an MBA at Warwick Business School. He is currently employed carrying out Business Development activities and securing funding for European Collaborative projects.

Having been awarded a BA in Metallurgy and Material Science, Martin Strangwood stayed on at Cambridge University following PhD research into the prediction and assessment of ferrous weldmetal microstructures, concentrating on the formation of acicular ferrite. In 1990, he joined the School of Metallurgy and Matertials/IRC in Materials for High Performance Applications at the University of Birmingham as a lecturer in phase transformations and materials joining. His major research effort is directed towards quantitative phase transformation theory and its application in microstructural modelling.

Alloy 8009 is a rapidly solidified Al-8.5%Fe alloy designed for high temperature (up to 400°C) aerospace applications. Both fusion and solid state joining techniques have been shown to produce welds. Variations in the weld microstructures are assessed by Mike Ellis of TWI and Martin Strangwood of the University of Birmingham.

Aluminium alloys are widely used by the aerospace industry, in particular because of their fairly high specific strength to specific modulus ratio. The wrought aerospace alloys are largely age-hardenable and can be used only at relatively low temperatures, ie up to 100°C. Any length of time at temperatures greater than this will lead to overageing of the alloy, and hence loss of strength.

At operating temperatures above those acceptable for Al alloys, titanium alloys may be used, but at some loss in weight savings. The need for a new generation of Al alloys with high temperature properties approaching those of titanium has been identified. Allied Signal Inc have developed a series of rapidly solidified (RS) Al-Fe-V-Si alloys which have properties which remain stable up to approximately 400°C.

Most materials applications require joining technology, and this article presents initial studies in the welding of Alloy 8009. The laser beam, tungsten inert gas (TIG) and resistance (spot) fusion welding processes and solid state friction stir welding were used. The welds were metallurgically assessed using light microscopy, scanning electron microscopy (SEM) and transmission electron microscopy (TEM).

Experimental practice

Material

The Alloy 8009 was supplied in rolled sheet approximately 2mm thick with the composition shown in Table 1.

Table 1 Chemical composition of Alloy 8009

Base metal microstructure

The parent material microstructure is shown in Fig.1a. The grain size was fine, about 0.2-0.3µm, and a high volume fraction of fine, Fe-rich dispersoids was present. Overall, the structure was reasonably homogeneous, but some banding was observed with variations in the size and number density of the Fe-rich dispersoids. Regions with a lower density of larger dispersoids exhibited larger grain sizes, up to 0.5µm. The dispersoids had a typical size range of 50-100nm, but some were as large as ~150nm ( Fig.1b). These larger dispersoids were identified as either Al ₅M ₂ (M = Fe, V) or Al ₃M, with the former pre-dominating. The finer dispersoids were found to be Al ₆M, although some metastable Al ₁₃M ₂ and AlM ₂ was noted. The larger grains exhibited very fine precipitates ( eg the darker grain seen in Fig.1b) which were probably Mg and Si-rich. These were extremely metastable, both growing and dissolving under electron beam heating in the TEM.

Fig. 1. The as-received parent material microstructure of the Alloy 8009:
a) SEM micrograph;

b) TEM micrograph showing low particle density (large grain size) region with Al ₅M ₂ (M = Fe, V); note the fine precipitates in the dark grain

Fusion joining

Tungsten inert gas (TIG) welding of Alloy 8009

The TEM study showed the heat affected zone gave a discernible gradation in microstructure from base to fused material, Fig.2c. Increased diffusion and thermal activation at higher temperatures close to the fusion boundary caused a shift from AlM ₂ to Al ₃Fe, giving a lower density of larger dispersoids. This resulted in a solute-depleted, structure with a coarsened grain size (1.5-3µm). The amount of possible coarsening increased as the fusion boundary was approached.

Table 2 Summary of tungsten inert gas (TIG) bead on plate welding trials on Alloy 8009

Fig. 2. TIG melt runs on Ally 8009:

a) Optical micrograph showing bead profile;

b) SEM micrograph showing intermetallic needles in the fusion zone, nominal;

c) TEM micrograph showing coarsening of particles through the HAZ, A - fine particles, B - coarse particles

As the fusion boundary was crossed, large Fe ₃Al intermetallics were formed. A process of multiple nucleation or sympathetic nucleation was seen, but developed to such an extent that 'star and cross' ( Fig.2b) features were observed by SEM. Near the centre of the fused zone, the size of the intermetallic particles became appreciably coarser. These large features were not observed in the TEM, as they were too coarse to electropolish successfully. Using the TEM, fragments of primary Fe ₃Al were observed, the formation of which (taking V and Si into solution) had decreased the amount of 'interdendritic/grain boundary' FeAl ₃ and Si-Al intermetallic. These latter were present as a discontinuous series of isolated lenticular particles. The solute-depleted Al grains had grown to >5µm in diameter in the weld bead. No fine precipitates were observed.

Laser welding of Alloy 8009

A total of four welds was made using a continuous wave Laser Ecosse 5kW CO ₂ laser ( Table 3). Laser welding produced a narrow weld. This meant that there was no detectable transition zone between base metal and fused material ( Fig.3a) and, hence, there was insufficient material to carry out a thorough TEM examination of the HAZ.

Table 3 Summary of laser beam welding trials on Alloy 8009

However, the fused material showed great changes in microstructure compared with the base material. Typical laser weld microstructures are shown in Fig.3. As with TIG-welded Alloy 8009, the fine dispersoids reacted to form coarse intermetallics. In this case, they were not 'lath-like' but had a 'blocky' geometry ( Fig.3b). An idea of the scales of the parent material and weld metal microstructures is given if Fig.1a is compared with Fig.3b. At high magnification, the microstructure ( Fig.3c) contained large (1.5-2.5µm) particles of Fe ₃Al. These particles are the primary solidification phase as shown by the hexagonal morphology noted in Fig.3d.

Fig. 3. Laser beam welds made in Alloy 8009:

a) Optical micrograph showing profile of weld;

b) SEM micrograph showing the 'blocky' nature of the microstructure;

c) TEM micrograph showing large primary Fe ₃Al with secondary solidification phases on interface (arrowed). The light regions are solute-depleted Al and surrounded by interdendritic FeAl ₃ and Si, V rich intermetallic phase;

d) TEM micrograph showing large, multiple Fe ₃Al particles and grain boundary network of FeAl ₃ and Al-Si-V intermetallics

The material grain size had increased to 1-2µm in diameter. The final phases to form were an 'interdendritic/grain boundary' structure, composed of FeAl ₃ and Si, V-rich metallics.

Resistance welding of Alloy 8009

Resistance welding was carried out using a Schlatter 350kVA DC welding machine. Electrodes were made from a Cu-1%Cd alloy with a dome radius of 75mm. Coupons of dimensions 30 x 30mm were welded together. Surface pre-treatment involved abrading with 600 grit wet and dry paper. A total of thirteen welds was made, some were sectioned for metallography and some joints were peel-tested to reveal the failure mode. A summary of the conditions used is shown in Table 4. The lower peak currents (~25kA) produced spot welds which failed in the fused parent material and at the interface. At higher current (31kA), the bonds were relatively good, and failure occurred around the fusion boundary leaving a proud piece of material equivalent to the size of the fused zone. If currents were increased further (39kA), excessive 'flash' was produced giving rise to a poor bond and it was considered that optimum conditions occur at a welding current of ~30kA. A typical spot weld is shown in Fig.4a. In the centre of the weld, there was slight clumping of the intermetallic particles, and small shrinkage cavities.

Table 4 Summary of resistance welding trials on Alloy 8009

Typical microstructures are shown in Fig.4b and 4c. Welds No 9 and 13 (see Table 4) were examined in the SEM and TEM respectively. As with the laser weld, the HAZ of the resistance weld was too small to examine by TEM. However, a sharp change in the microstructure was noted at the fusion boundary into the weld metal and acicular intermetallics were observed ( Fig.4c).

As with the TIG and laser welds, the weld metal microstructure was characterised by large (2-50µm) Fe ₃Al particles as the primary solidification phase within a solute-depleted Al matrix. This matrix contained 'interdendritic/grain boundary' intermetallics of FeAl ₃ and the Si, V-enriched phase, Fig.4c. In terms of the grain size and continuity of grain boundary intermetallics, the resistance weld fell between the laser and TIG welds, (roughly 2.5µm in size).

Fig. 4. Resistance welds made in Alloy 8009:

a) Optical micrograph showing resistance weld nugget;

b) SEM micrograph of Weld No 9;

c) TEM micrograph of weld No 13 showing faceted Fe ₃Al matrix in the centre of the nugget

Solid state joining

Friction stir welding of Alloy 8009

The friction stir welding technique joins aluminium alloys by using a rotating tool to plasticise and consolidate the material about the joint line. ^[1] The operation is achieved by drilling a hole at the start of the joint, equal to the joint depth required, and inserting a close fitting pin. The pin is rotated and moved forward in the direction of welding. When the pin is rotated it frictionally heats an annular region of aluminium alloy, rapidly producing a plasticised tubular shaft of metal around itself. As it is moved in the direction of welding, the pressure provided by the leading face forces plasticised material to the back of the pin, where it cools and consolidates. The process requires no filler material or bevelled edge preparation.

A friction stir weld was made by producing a lap joint from two pieces of Alloy 8009 and pushing a rotating tool through the top sheet into the bottom sheet at a linear speed of 160mm/min. Bonding parameters were not optimised. To make a second bond, a sheet of Alloy 8009 was placed on top of a piece of 6082 (Al-Mg-Si) extruded alloy and a rotating tool was pushed through the surface of the rapidly solidified alloy. The stirring motion of the tool mixed both the alloys together to produce a solid bond. The microstructures of the bonds for Alloy 8009 to Alloy 8009 and Alloy 8009 to 6082 are shown in Fig.5 and Fig.6 respectively.

Fig. 5. Friction stir weld in Alloy 8009:

a) SEM micrograph showing weld metal microstructure;

b) TEM micrograph showing Al ₅M ₂ and MAl ₃ near bond line

Fig. 6. Friction stir weld between Alloy 8009 and 6082:

a) Optical micrograph showing mechanical mixing;

b) TEM micrograph showing l ₅M ₂ and MAl ₃ near bond line in Alloy 8009

In contrast to the fusion techniques, the microstructure of the Alloy 8009 material in the friction stir welds was very similar to that of the base material (compare Fig.5a with Fig.1a). Some increase in the intensity of particle banding, presumably formed by strain concentration in the pre-existing coarse grain bands in the as-received material, was observed. The TEM micrographs, Fig.5b and 6b, indicate that there was a slight increase in dispersoid size, reduction in number density and increase in Al grain size when compared to the parent material. This was associated with an increase in the proportion of the dispersoids that were of the Al ₃M type rather than Al ₅M ₂. Some alignment of the dispersoids along the flow lines of the material occurred during joining. There was also apparently a loss of solute to the dispersoid matrix interface, inferred by the reduction in ageing precipitates observed forming and dissolving under the electron beam of the TEM.

Discussion

The results of this work have shown that both fusion and solid state techniques can be used to weld Alloy 8009. Obviously, more detailed studies would be necessary to optimise welding conditions, and to ensure freedom from defects such as the longitudinal cracking in some TIG welds, and the shrinkage cavities in TIG and resistance welds. The fracture face of the longitudinal cracks in wide TIG beads showed no evidence of a solidification structure, and the cracking is presumably a reflection of the low inherent material ductility, probably reduced further by intermetallic formation in the molten zone. Nevertheless, it is evident that substantially sound joints can be produced by a range of processes. The success of friction stir welding will be noted, in particular that joining was obtained between Alloy 8009 and 6082 alloys, despite the marked difference in tensile properties and stiffness between the materials.

As would be expected, fusion and solid state welding led to considerable variations in the bond microstructure. There has been a limited amount of work carried out on the fusion welding of rapidly solidified Al-Fe alloys, ^[2-5] and even less work has been reported for Alloy 8009. Due to the high temperatures during fusion welding, microstructural changes will occur. Analogous with the present results, the microstructure of TIG welds in an Al-Fe-V-Si alloy can transform from one containing fine spherical dispersoids in an Al matrix to one having coarse, acicular intermetallics likely to be Fe ₃Al, thus reducing the mechanical properties of this alloy. ^[6] It has also been noted that residual hydrogen, originating from the aluminium powder, ^[6] can be released and form unacceptable levels of porosity in the weld metal.

The microstructure of the laser weld was different to that of the TIG weld. Even though coarse intermetallics were produced during laser welding, these were not lath-like, but blocky ( Fig.3), presumably reflecting the rapid thermal cycle associated with this welding technique. Similar microstructures in the fusion and heat affected zones have been reported in other RS alloys for pulsed Nd:YAG laser welding ^[2] and electron beam welding. ^[3]

Some encouraging results were obtained from the spot welded material ( Table 4). Nugget failure was observed in a number of cases when currents were of the order of 31kA. Too low a current (~25kA) gave rise to failure either in the parent material or at the interface of the weld. Too high a current (~39kA) led to flash and a relatively poor weld. Although there are limited data on the resistance welding of RS Al alloys, work has been carried out using the capacitor discharge welding technique on a RS Al-Fe-Ce alloy. ^[4,5] The microstructure in that case was fine and gave high strength.

A number of the problems associated with fusion welding can be alleviated by the use of solid state joining. For the present family of alloys, this has the beneficial effect of avoiding the formation of coarse intermetallics, eg Fe ₃Al. A considerable amount of work on inertia and linear friction welding has been carried out, ^[6,7] joining a rapidly solidified FVS1212 to itself and to the high strength 2024-T351 monolithic alloy. This study showed that linear and inertia friction welding were effective methods for joining these alloys. The weld region in the similar alloy joints had negligible dispersoid and alpha grain coarsening with minimum hardness degradation across the joint. Joint efficiencies achieved were approximately 85% of parent material strength at room temperature when using a rapid thermal cycle. When the thermal cycle was slower, increased dispersoid coarsening and softening at the weld interface occurred and joint efficiency fell to only about 80%. In the dissimilar alloy welds, nearly all the deformation occurred in the 2024-T351 material.

All of the above findings are confirmed by the results of the friction stir welding trials. The base material dispersoids appeared relatively unaffected, other than a small increase in size and a reduction in number density. The dispersoids seemed to align themselves along the flow lines associated with the joining deformation, while the smaller size distribution at the bond line may have an effect of increasing bond line ductility.

Conclusions

Substantially sound fusion and solid state welds were successfully made in Alloy 8009, without extended optimisation trials. The TIG, laser, resistance and friction stir methods were used. The use of fusion techniques destroyed the base metal structure with primary solidification of Fe ₃Al, loss of solute, formation of larger Al grains and the formation of grain boundary FeAl ₃ and Si, V-enriched intermetallics. In contrast, friction stir welding caused minimal microstructural change with some modification of dispersoid alignment, a small amount of coarsening and loss of solute to dispersoid/matrix interfaces.

Acknowledgements

The authors would like to thank Dr L Kallman of Saab Scania AB (Saab Military Aircraft) for providing the Alloy 8009 sheet.

References