Titanium aluminides

8.6.3 Titanium aluminides

Because of the limited scope for improvements in the properties of conventional titanium alloys above 650 ◦

C, either by alloy development or by TMP, increased attention is being given to the titanium intermetallics, Ti 3 Al (α 2 -phase) and TiAl (γ-phase). With low density, high modulus and good creep and oxidation resistance up to 900 ◦

C, they have considerable potential if the poor ductility at ambient temperatures could be improved. A comparison of Ti 3 Al- and TiAl-based materials with conventional Ti alloys is given in Table 8.7.

Advanced alloys 471 Table 8.7 Comparison of α 2 and γ superalloys with conventional titanium alloys.

Property

(α 2 + β) (γ +α 2 ) Density (g cm −3 )

Titanium alloys

145 176 RT tensile strength (MN m −2 )

E, stiffness (GN m −2 )

10 620 HT (760 ◦

C) tensile strength (MN m −2 ) 620 550 Max. creep temp. ( ◦ C)

730 900 RT ductility (%)

3 Service temp. ductility (%)

Electron microscopy studies of Ti 3 Al or α 2 have shown that deformation by slip occurs at room temperature by coupled pairs of dislocations with b planes and by very limited glide on

The ductility increases at higher temperatures due to climb of the increased glide mobility of 1/6 Only limited activity of the {0 0 0 1}

The most successful improvements in the ductility of Ti 3 Al have been produced by the addition of β -stabilizing elements, particularly niobium, to produce α 2 alloys. An addition of 4 at.% Nb produces

significant slip on on

for Ti with a consequent reduction in the Peierls–Nabarro friction stress. Alloys based on α 2 are Ti– (23–25)Al–(8–18)Nb, of which Ti–24Al–11Nb has excellent spalling resistance. Most Ti 3 Al + Nb alloys, such as super α 2 , also contain other β-stabilizers including Mo and V, i.e. Ti–25Al–10Nb– 3V–1Mo, which exhibits about 7% room temperature elongation. Alloying Ti 3 Al with β-stabilizing elements to produce two-phase alloys significantly increases the fracture strength. These Ti 3 Al-based alloys can be plasma-melted and cast followed by TMP in the (α 2 + β) or β-range. The improved ductility of Ti 3 Al alloys has led to aerospace applications in after-burners in jet engines, where it compares favorably in performance with superalloys and gives a 40% weight saving. Developments are taking place in rapid solidification processing to include a second phase (e.g. rare-earth precipitates) and to provide powders, which may be consolidated by HIPing, to produce fully dense components with properties comparable to wrought products. There are also developments

in intermetallic matrix composites by the addition of SiC or Al 2 O 3 fibers ( ∼10 µm). These have some attractive properties, but the fiber/intermetallic interface is still a problem. The γ-phase Ti–(50–56)Al has an ordered fc tetragonal (L1 0 ) structure up to the m.p. 1460 ◦

C, with c/a = 1.02 (Figure 8.16). Deformation by slip occurs on {1 1 1} planes and, because of the tetragonal- ity, there are two types of dislocations, namely ordinary dislocations 1/2

At room temperature, deformation occurs by both ordinary and superdislocations. However, [0 1 1] and [1 0 1] superdislocations are largely immobile because segments of the trailing superpartials 1/6 stacking faults are largely sessile because of the Peierls–Nabarro stress. Some limited twinning also occurs. The flow stress increases with increasing temperature up to 600 ◦

C as the superpartials become mobile and cross-slip from {1 1 1} to {1 0 0} to form K–W-type locks, the 1/2 increases and twinning is promoted.

The two-phase (γ +α 2 ) Ti–Al alloys have better ductility than single-phase γ with a maximum at 48 at.% Al. This improvement has been attributed to the reduced c/a with decreased Al, further promotion of twinning and the scavenging of O 2 and N 2 interstitials by α 2 . The combination of high

472 Physical Metallurgy and Advanced Materials [1 1 2]

(b) Figure 8.16 Structure of TiAl (Ll 0 ) (a) and (1 1 1) plane (b) showing slip vectors for possible

dissociation reactions, e.g. ordinary dislocations 1/2[1 1 0], superdislocations [0 1 1] and 1/2[1 1 2], and twin dislocations 1/6[1 1 2] (after Kim and Froes, 1990).

stiffness (E = 175 GPa at 20 ◦

C to 150 GPa at 700 ◦ C), density-normalized strength similar to cast

Ni-based alloys, high temperature strength and reasonable oxidation resistance to 750 ◦

C, low thermal expansion coefficient and high thermal conductivity have led to a high level of interest in TiAl-based alloys. The major limitations to their application are the intrinsic low room temperature ductility (no better than 2–3%), the low fracture toughness (between 10 and 20 MPa m 1/2 at 20 ◦

C) and the high growth rate of fatigue cracks. Alloy and process development have resulted in some successful applications of these alloys over the last 10 or so years. Cast turbochargers are now manufactured in Japan for cars and wrought exhaust valves were used in Formula 1 cars for some years. Major applications for these alloys are still awaited despite the success of these two applications; the high cost of processing is holding commercial developments back. In the case of thermomechanical processing the costs are high because the alloys are strong at normal hot-working temperatures and because some sort of protection (such as canning with steel) from oxidation must be used during working. In the case of castings the efficiency of material usage is very low, both because casting technology is not efficient and because melting and casting are difficult because of the reactivity of molten TiAl-based alloys.

The compositions of TiAl-based alloys which are of commercial interest lie within the range of about Ti/45–48 Al (at.%) but all alloys contain other elements in attempts to improve the properties of the binary alloy. Additions of Nb between about 5 and 8 at.% are important in improving oxidation resistance and also imparting some solid solution strengthening.

An understanding of the microstructures which can be obtained in cast or in wrought products of TiAl-based alloys requires knowledge of the phase changes that occur over the temperature range from the melting point to room temperature. The relevant part of the binary phase diagram between

Ti and Al is shown in Figure 8.17, which also indicates schematically the influence of some alloying additions on phase boundaries.

The various phase transformations in the Ti–Al system offer the possibility of microstructural control both for the wrought route and for the casting route. Thus, cooling samples containing less than about 44 at.% Al the solidification will take place through the formation of β, which may or may not be removed via the peritectic reaction. Subsequent cooling of the α-phase results in precipitation of

Advanced alloys 473

Atomic percent aluminum

Figure 8.17 Partial Ti–Al phase diagram showing the influence of ternary additions on the position of the various boundaries (courtesy M. H. Loretto).

γ - which, under typical cooling rates encountered with castings, results in the formation of a lamellar structure. This ‘fully lamellar’ structure consists of parallel lamellae of γ and α and of twinned γ.

These lamellae are formed on the (0 0 0 1) plane of the α-phase and thus their length is defined by the pre-existing α grain size. Somewhat slower cooling results in some of the γ lamellae coarsening at colony boundaries to form γ grains through local growth of the lamellae, to form a ‘near fully lamellar’ structure. Hot working in the two-phase region results in the formation of equiaxed γ and α grains; the ratio of the amounts being defined by the average alloy composition and the hot-working temperature. Subsequent cooling results in the equiaxed α grains forming lamellae to yield a duplex microstructure, or if extensive hot working is used, a structure consisting of γ and α grains is formed, termed ‘near γ’.

If the cooling rate is increased, as in oil or water quenching, the α-phase transforms massively to γ if the Al content is above about 44 at.% (below this Al-content, α is retained) and this transformation offers a further opportunity for microstructural control in cast samples by heat treating in the two- phase field so that α can precipitate on all four {1 1 1} planes, throughout the γ grains. The tendency to transform massively is strongly dependent upon the composition, which has important consequences upon the choice of alloy composition in cast samples.