New alloy group under the microscope - tomorrow's lightweight joining solutions

TWI Bulletin, May/June 1999

Lee Smith joined TWI at the beginning of 1996 and is a Senior Project Leader in the Materials Department. He received his first degree in Metallurgy and his PhD in the development of intermetallic alloys from the University of Birmingham.

Philip Threadgill gained his BSc and PhD in Physical Metallurgy from University College, Swansea. He has worked at TWI since 1976, and is currently R&D Manager for the Friction and Forge Processes group.

A new class of material, the intermetallic _γ -TiAl, has emerged in the past decade for use in lightweight elevated temperature applications. Lee Smith and Philip Threadgill describe the known welding metallurgy of this group of alloys.

Intermetallic alloys have been the subject of intense research over the past ten or so years, principally to increase operating temperatures and/or decrease the weight of future aeroengines. The _γ -TiAl alloys in particular are assured a future for high temperature applications where weight is a critical concern (see the Table for a summary of the key properties). Engine trials have been performed with _γ -TiAl compressor and low pressure turbine blades and it seems likely that increasing competition between the aeroengine manufacturers will eventually provide the impetus for their introduction into service.

Table: Selected properties of _γ -TiAl alloys compared with those of conventional titanium alloys

Increasingly, the automotive, especially motor sports, sector is proving to be an early adopter of cutting edge technology. This has proven so with the _γ -TiAl alloys; valve heads and turbo rotors for automotive engines have been developed and have already entered service. Many of the automotive and aeroengine applications will require joining processes, from repair welding of castings to joining _γ -TiAl to dissimilar materials. Progress in alloy development and in the understanding of welding metallurgy has been considerable, to the point where a review would be beneficial to those working in applications requiring high temperature lightweight materials. This work describes in detail the physical and welding metallurgy of this important group of alloys.

Base materials

Most _γ -TiAl alloys contain approximately 48%Al and are actually two phase ( α ₂-Ti ₃Al and _γ -TiAl) under equilibrium conditions below 1125°C ^[1] ( Fig.1). Both phases have ordered lattices; the _γ phase exhibits the L1 ₀ crystallographic structure, while the α ₂ phase has the D0 ₁₉ structure ( Fig.2). As with most multiphase materials, mechanical properties are strongly dependent on microstructure ( Fig.3) and, by careful processing and heat treatment, microstructures can be tailored to suit particular property requirements. ^[2]

The fully lamellar microstructure ( Fig.4b) consists of colonies of α 2 and γ plates. It is typically exhibited by as-cast material, but can be obtained by heat-treating for short periods above the alpha transus (1320°C for Ti-48Al-2Cr-2Nb). In both instances the grain size tends to be quite large and such alloys can have extremely limited ductility and poor strength. Ductility and strength can be improved by minimising the grain size through selective heat treatments or thermomechanical processing. Even so, the lamellar microstructure typically gives the best fracture toughness.

Finally, 'duplex' microstructures ( Fig.4c) exhibit a mixture of fine equiaxed γ grains and lamellar colonies and are typically formed after heat treatment within the α+ γ phase field, where the equilibrium composition consists of approximately equal proportions. Such alloys exhibit superior ductility and strength to a lamellar microstructure, but at the expense of fracture toughness.

Effect of alloying additions and impurities

Contaminants absorbed during welding and base metal impurities can also present welding problems. Oxygen and nitrogen embrittle γ -TiAl; alloys with impurity levels greater than approximately 200ppm typically exhibit brittle behaviour at room temperature. The enhanced ductility of two phase alloys compared to single phase gamma alloys may be related, in part, to interstitial segregation. Interstitial elements tend to segregate strongly to the α 2 phase, [4] increasing the purity, and consequently ductility, of the γ phase. In most cases this may prove beneficial, but if excessive oxygen contamination occurs, large embrittling α 2 grains may form. Furthermore, nitrogen and oxygen may segregate to the solidification front in a manner analogous to that of P and S in austenitic stainless steels, although in this case promoting predominantly solid state cracking, rather than solidification cracking ( Fig.5). Quite small levels of hydrogen can lead to hydride formation and embrittlement in conventional titanium and α 2-Ti 3Al alloys, but pose less of a problem for γ -TiAl alloys. Although hydrides can form, the hydrogen solubility in a typical lamellar γ -TiAl alloy is much greater (>360ppm) than either of the other alloy types. [5]

Overall, the susceptibility of the γ -TiAl alloys to embrittlement by O, C, N and H is no worse than that of conventional titanium alloys, for which the precautionary measures necessary to avoid detrimental contamination during fabrication are well documented. Thus, no difficulties should be encountered if good practice is followed.

Fusion zone and HAZ microstructures

For most alloys, equilibrium solidification will follow the sequence L → β +L → α. The presence of retained β can degrade mechanical properties, and is more likely in Cr-containing alloys. However, high undercoolings, such as typically experienced during welding, may favour the sequence L → α +L → α, particularly for Cr-free alloys. Transformations which occur during cooling from the α phase can be categorised into five types, producing, in order of increasing cooling rate; lamellar gamma ( γ l), Widmanstätten gamma ( γ w), 'feathery' gamma ( γ f), massively transformed gamma ( γ m) and retained alpha ( α 2r).

For cooling rates typical of, say, TIG welding, lamellar α 2+ γ microstructures are commonly observed. [6,7] The formation of the lamellar morphology in the Al-lean ( ie <50 atomic % Al) alloys follows the sequence α → α2 → α 2+ γ or α → α+ γ → α 2+ γ , the latter occurring at greater cooling rates. The γ laths grow on the basal plane of the α (or α 2) phase, nucleating predominantly at grain boundaries with the lath interface parallel to {111} γ . More than one lath orientation or colony can exist in each prior- α grain, although each colony tends to be quite large. There are six γ / α orientation variants so adjacent laths can be either γ / α 2 or γ / γ . The mechanical performance of weldment γ l morphologies can compare favourably with as-cast lamellar structures. For example, the lath spacing tends to be finer, aiding ductility. It is likely, however, that postweld heat treatment (PWHT) is necessary to relieve residual stresses and to promote the interlocking of neighbouring colonies, improving fatigue performance.

The Widmanstätten morphology ( Fig.6a) consists of acicular colonies of γ and α 2 lamellae. Although the γ /prior- α orientation relationship is identical to that of γ l ( ie four possible lath orientations), the colonies have a basket-weave appearance. Several mechanisms have been suggested, including intragranular nucleation on lattice defects in the α phase and growth via a shear transformation, [8] similar to that observed in ferrous materials. There is, however, compelling evidence against such a mechanism: the γ w morphology is typically encountered only in re-heated material, rather than on cooling from the liquid state. A plausible mechanism that might account for these observations is the formation of acicular α subgrains on {111} γ within lamellar regions of the original prior- α grains. [9] On cooling, these transform to a lamellar structure, following the transformation sequence for γ l. It is speculated that the toughness of γ w structures might compare quite favourably with that of as-cast lamellar alloys, due to grain refinement through recrystallisation and a more frequent change in lath orientation.

The morphology of γ f consists of fine colonies of lamellar gamma with no consistent rational orientation relationship with the prior α, often appearing to fan out from an original colony ( Fig.6b). It should be noted that some observers make no distinction between γ f and γ w, but these microstructures are indicative of quite different thermal histories. The γ f morphology has been observed in material cooled at a wide variety of rates, but is only observed upon cooling from above the α transus ( eg fusion zone and high temperature HAZ). Whilst the γ f colonies do contain α 2 lamellae, the volume fraction of α 2 depends on cooling rate, with very little being observed in rapidly cooled material. Transformation is by a displacive reaction, with γ f colonies initiating from γ l and possibly γ w primary colonies. [9] The growth rate of the γ l morphology is quite sluggish, whereas it is postulated that the growth of γ f laths on non-basal α planes might be more favourable at greater undercoolings, due to a higher energy interface. The toughness of the γ f morphology may benefit from the non-aligned laths, but as yet no quantitative studies have been published.

At very high cooling rates, such as is possible for laser and electron beam welding, massively transformed gamma ( γ m) may form ( Fig.6c). The evolution of this morphology has been explained by incoherent growth. [10] Initially, during cooling from above the α-transus, the α phase decomposes to form γ l laths, nucleating preferentially at favourably orientated grain boundaries. The transformation front proceeds into the primary grain, but also impinges into and along the grain boundary close to the nucleation zone. Growth continues incoherently into the neighbouring grain via a massive transformation, proceeding at an accelerated rate due to the high energy interface. At even greater cooling rates the α → γ transformation may be suppressed entirely, leading to mostly α 2 microstructures (the α → α 2 transformation occurs very rapidly since it is a second order transformation, requiring only short-range diffusion). Both γ m and α 2r morphologies are particularly sensitive to solid state cracking. The avoidance of these morphologies represents a significant challenge to the practical application of power beam processes such as laser and electron beam welding.

Solid state weld microstructures

Both diffusion bonding (DB) and friction welding have been applied successfully to γ -TiAl alloys. Diffusion bonding can be performed at a range of temperatures, although those between the eutectoid temperature and the alpha transus are perhaps more commonly used. Under these conditions, a 'double necklace' of γ grains forms typically at the bond line ( Fig.7) and is most probably a consequence of the deformation induced on the joint surfaces during preparation. [11] This microstructure can, however, be avoided by polishing (rather than grinding) the joint surfaces. Even so, the mechanical properties of the joints with the 'double necklace' are comparable to those of typical matrix microstructures. At higher DB temperatures microstructural transformations in the matrix are likely to occur; ie the formation of equiaxed γ grains at grain boundaries and other high interfacial energy sites. Bonds made at temperatures below the alpha transus show no significant change in microstructure or hardness away from the bond line. Indeed, in circumstances where new γ grains do not nucleate at the bond line, grain growth across the boundary occurs and the bond line is almost impossible to detect.

Although no fusion occurs during friction welding, profound microstructural changes occur due to the combination of temperature and severe deformation. [12] At higher forge forces material near the joint undergoes recrystallisation, producing very small equiaxed γ grains.

At lower forge forces, refinement of the microstructure is much less marked, and a lamellar microstructure is formed at the bond line, which is typically finer than the parent metal. Further away from the bond line and the more massively deformed material, but within the zone heated close to, or above, the α transus, secondary lamellae may form within prior lamellar colonies. These secondary γ lamellae are thought to nucleate on basal defects introduced into the α 2 laths by the thermomechanical cycle. The effects of friction welding are particularly evident in lamellar γ -TiAl alloys; the lamellae are extensively twinned and deformed either side of the recrystallised zone ( eg Fig.8). No mechanical property studies of friction welds have been published, but it is likely that the weld properties will be dominated by those of the recrystallised zone. Although fully γ matrix microstructures show typically poor mechanical performance, the very fine grain structure of the recrystallised zone may prove beneficial.

Postweld heat treatment

Postweld heat treatment may be essential to achieve desired weld properties, achieving stress relief, reducing hardness and promoting the formation of beneficial microstructures. Unfortunately, microstructural alteration will require high temperatures (significantly greater than 1000°C) to achieve any marked changes. In most instances it will be impractical to perform heat treatments above the alpha transus and, thus the welding process and parameters should ideally be selected to achieve a desirable microstructure in the as-welded condition. In most instances this will preclude parameters which produce the brittle γ m and α 2r morphologies, so only the γ l and γ f microstructures are considered here for fusion welds.

No studies have been published showing the effect of PWHT on friction welds, but it is likely that the fine equiaxed γ grains of the recrystallised zone will transform to a fine grained duplex or lamellar structure, depending on the heat-treatment temperature chosen. The double necklace structure present at the bond line of diffusion bonds can be eliminated by PWHT above the α transus, but this is at the expense of extensive grain growth in the matrix.

Conclusions

References

Further reading
Smith L S and Threadgill P L: 'The physical and welding metallurgy of titanium aluminides: A review.' TWI Members Report 633/1998.