This invention relates to an improved titanium base alloy and more particularly to titanium base alloys including vanadium, aluminium and iron as essential constituents.

In many applications it is important to combine high strength with good ductility and fracture toughness, as well as good hot and cold workability. High strength and toughness as well as good hot workability are important in thick section titanium forgings such as compressor discs and helicopter rotor heads. Applications for titanium alloys where good cold formability is required include springs, fasteners and honeycomb structures. In order to obtain high strength the alloy must have sufficient fracture toughness to prevent failure in the presence of microscopic flaws.

Elemental titanium at ambient temperatures, is of a close packed hexagonal structure commonly called the α-phase, whilst at temperatures of 883°C or higher it transforms to a body-centered cubic structure known as the β-phase. The α-phase is stabilised by aluminium, preferably in the range 2.5 to 8% by weight, which leads to an increase in the mechanical strength and hardness of the resulting alloy whilst retaining a sufficient ductility for fabrication processes.

- Other alloying ingredients are known to stabilize the β-phase at progressively lower temperatures, until a stable β-phase is obtained at normal temperatures. It is known in the art that addition of β-stabilisers to a titanium base alloy results in good response to heat treatment and imparts high strength to the alloy. It is also known that the heat treatability and depth of hardening of an alloy increase in proportion to the amount of retained β-phase after quenching the alloy to room temperature, β-phase titanium alloys exhibit better forgeability

than titanium alloys of other conformations as a consequence of their body centered cubic crystal lattice structure, although as stated above, α-phase titanium alloys still have good formability properties. Shear and diffusionable transformations which take place during cooling can be inhibited by alloying to produce so called deep hardenable materials. In this context hardenability is related to the capacity for retention of β-phase, through thick sections, coupled with high strength, β-stabilizers known in the art are Mo, V, Nb, Ta, Mn, Cr, Fe, , Co, Ni and Cu. Combining both a- and β-microstructures enables optimisation of a number of mechanical properties. However, only the elements Mo, V, Nb, Ta, Mn, Cr, Fe and , from the list above are suitable for producing mixed-phase (α and β) alloys or stable β-alloys. β-titanium alloys are potentially very attractive to the aerospace industry. However early alloys suffered problems arising from the use of high melting point components which compromised alloy homogeneity. In addition, the use of compound forming elements produced undesirable intermetallic phases. The alloying additions also increase density compared with the titanium itself.

UK Patent Specification No. 782 64 describes improvements relating to titanium-aluminium base alloys. This document lists the mechanical properties of titanium alloys including three containing additions of aluminium, vanadium or iron having the following point compositions: Ti-5V-4A1, Ti-10V-4Al and Ti-5V-1.5Fe-4Al. More recently developed alloys include the Transage series (e.g. T12Q) and Ti-10-2-3 (a titanium alloy comprising 10% vanadium, 2% iron, 3% aluminium, proportions being in % weight) . The latter of these two alloys is described in US Patent Specification No. 3802877- Vanadium is employed as the principal β- stabilising element and small additions of aluminium reduce density and increase hardenability. Good combinations of strength and toughness may be obtained in the aged condition. There is a significant market for these alloys in the aerospace field and the alloy Ti-10-2-3 appears to offer the best combination of properties currently available. However this alloy has limited cold workability and the age hardening levels are sensitive to cooling rates of the annealed material.

It is therefore an object of the invention to provide a titanium alloy which can easily be processed to give various combinations of high strength ductility, fracture toughness and fatigue strength superior to those reported for the commercial Ti-10V-2Fe-3Al alloy.

It is a further object of the invention to provide a titanium alloy with good hot processing properties and in which age hardening levels are less sensitive to cooling rates of the annealed material compared to the commercial Ti-10-2-3 alloy.

The invention is a near β-phase titanium alloy comprising the following constituents in proportions by weight:

vanadium 7•6-8.2 % iron 2.1-2.7 % aluminium 3-5-4.2 % titanium balance, save for incidental impurities.

In the following text, all compositions are given in proportions by weight.

A preferred composition for the improved titanium alloy comprises:

vanadium 8.0 % iron 2.5 % aluminium 4.0 % titanium balance, save for incidental impurities.

(hereafter denoted Ti-8V-2.5Fe-4Al)

A near β-phase alloy is one in which a significant proportion of the structure is β-phase.

It is especially advantageous if alloys according to the invention are heat treated to obtain the requisite microstructure for maximum fracture toughness and fatigue strength. Preferably the alloys are annealed at 800°C, water quenched and then aged in a salt bath at 525°C over a period of 8 hours.

The invention will now be described by way of example only with reference to the following drawings in which:

Figure 1 shows the age hardening curves of the Ti-8V-2.5Fe-4Al alloy aged in salt baths at 350-550°C.

Figure 2 shows the age hardening curves of the Ti-8V-2-5Fe-4Al alloy aged in salt baths (a) and furnaces (b) after annealing at 800°C.

Figure 3 shows yield stress against fracture toughness for the alloy Ti-8V-2.5Fe-4Al compared to that of the prior art alloy Ti-10V-2Fe-3Al

Figure 4 shows yield stress against tensile elongation for the alloy Ti-8V-2.5Fe-4Al compared to that of the prior art alloy Ti-10V-2Fe-3Al.

Investigations into the structure/property relationships of the inventive alloys have helped to explain why the strength/ductility/fracture toughness combinations shown by their age hardened (α + β) microstructures are superior to those reported for similar microstructures in prior art alloys.

Results have shown that the β-quenched Ti-8V-2.5Fe-4Al alloy contains nearly 100% retained β and a very small amount of orthorhombic α" martensite.

Age Hardening

Figure 1 shows the age-hardening curves of alloys having compositions falling within the invention aged in salt baths at 350-550°C. Such alloys attain the maximum hardening and hardness of about 500Hv which is unchanged over a temperature range of 350-450°C. Age-hardening at temperatures above these temperature ranges is lower and decreases with the rise in ageing temperature. Annealing at temperatures high in the

(a + β) phase field and ageing at 350-550°C results in a lower level of hardness compared to that attained in the β-quenched and aged material.

In the alloy it is observed that the level of hardening attained after annealing at a temperature high in the α and β range, e.g. at 50°C below the β-transus (at 800 ^C C) , is lower when the ageing is carried out in a salt bath rather than in a furnace (Fig. 2) .

Mechanical Properties

Table 1 shows the tensile properties and fracture toughness values of an age-hardened Ti-8V-2.5 ^_ 4A1 alloy. The table shows the annealing treatment followed by subsequent hardening treatment. The duration of each treatment is given in hours (h) or minutes (m) . For the high (a + β) annealed and aged condition the alloy shows excellent tensile ductility. For example, at yield stress levels 1250 - 1300 MPa it shows tensile elongation of 6-12%. It was also observed that when aged in a furnace, the tensile ductility of the alloy is reduced.

Fracture toughness of the age-hardened alloy is affected by the prior annealing temperature. The alloy annealed at 800°C (50°C below the β transus) shows, after subsequent ageing in a salt bath at 25°C, a fracture toughness value of 7 MPam ^1/2 at a yield stress of 1290 MPa, while the alloy annealed at the higher temperature of 830°C (20°C below the β transus) and aged in a salt bath at 550°C shows a smaller fracture toughness value of 57 MPam ^{1 2} at a lower yield stress of 1253 MPa.

Fracture toughness of the aged alloy is also affected by the heating rate in the ageing process. For example, the alloy annealed at 800°C (50°C below the β transus) and aged in a salt bath at 25°C has a yield stress of 1290 MPa, tensile elongation of 12% and a fracture toughness value of 72 MPam ^1/2 , whereas the furnace aged condition with similar yield stress of 131 MPa and comparable tensile elongation of % , shows a fracture toughness value of only 48 MPam ^1/2 .

Mechanical properties following age hardening after annealing (in the β or high (α + β) field) of the alloy depended on the different cooling rates used in in the study e.g yield stress values of 1290 and 1207 MPa are observed for the water quenched and air cooled alloy respectively aged at 525°C after annealing at 800°C. In addition, age hardening level is affected by specimen size; e.g. the furnace gas cooled specimen annealed at 800°C and aged at 450°C showed yield stress/ductility combination of 1357MPa/5.6% and 1170MPa/12.1 % respectively for the 4.54mm thick section specimen and the larger 29mm thick section specimen. At plate thickness of 29 m isothermal reactions are suppressed during cooling from the annealing temperature only by water-quenching and the water quenched

4.54mm thick section specimen and the larger 29mm thick section specimen shows similar levels of age-hardening (e.g. specimen annealed at 800°C water quenched and aged at 525°C in a furnace shows yield stress levels of about 1320 MPa for both specimen sizes) .

Comparison of the mechanical properties of Ti-8V-2. ^[: iFe-4Al to those of the prior art allov Ti-10-2-3

The mechanical properties of the invention are compared to those reported for the Ti-10-2-3 alloy in figures 3 and 4. These figures show respectively yield stress versus tensile elongation and yield stress versus fracture toughness values for the (a + β) microstructure and demonstrate that the invention has superior yield stress/ductility/fracture toughness combinations compared to the Ti-10-2-3 alloy. Figure 3 shows that at yield stress 1000-1300 MPa, the fracture toughness values of the invention developed here are higher than those of Ti-lO-2-3 alloy. For example at a yield stress of about 1100 MPa the alloy according to the invention shows fracture toughness values of 83 ^_ 97 Mpam ^1/2 compared to 63-73 MPam ^1/2 observed for the Ti-10-2-3 alloy. At the high yield stress level of 1300 MPa the invention shows a tensile elongation of 5 ^_ 12% and a fracture toughness of 48-72 MPam ^{1 2} compared to the tensile elongation of 2 - 10% and fracture toughness of 28 - 32 MPam ^1/2 shown by the Ti-10-2-3 alloy (Figs. 3 and 4).

Alloys according to the invention, show improved mechanical properties compared with the Ti-10-2-3 alloy, because they use the optimum V, Fe and Al combinations to achieve, after processing, the most attractive (a + β)

microstructure at high strength levels. They do not to depend on the transitional β > ω + β reaction as in the Ti-10-2-3 alloy to achieve the high strength in the (α + β) structure. The uniform α type structure of the Ti-lO-2-3 alloy shows poor fracture toughness. In alloys according to the invention α plates of high aspect ratio and having a zig zag morphology are observed at temperatures in which the β > w + β reactions occur in Ti-10-2-3 alloy. This enhances fracture toughness in the new alloys.

Table 1 : Mechanical Properties of Ti-8V-2.5Fe-4Al Alloy

Heat 0.2% Y. Tensile Fracture Total Reduct.

Treatment Stress Stress Toughness Elong. in Area

(MPa) (MPa) (MPam ^{1 2} ) (%) (%)

-890°C/30m WQ 172 957 25.6 25

~890 ^o C/30m WQ +

S 570°C/4h (L) 1158 1208 7 12.7

550°C/4h (L) 1240 1287 4.8 5.3

550°c/4h (T) 1245 1289 3-5 5.3

525°C/8h (T) 1270 1320 3-6 6.5

475°C/l6h (T) 1455 1487 1.6 4

-890°C/30m IT

800°C/lh WQ +

S 525°C/8h (L) 1286 1343 4.1 6.3

-830°C/40m WQ +

S 570°C/4h (L) 1168 1226 12.4 31.3

* 550°C/4h (T) 1253 1293 56.6(T-L) 7.3 .14.6

525°C/8h (L) 1296 1356 6.6 11.4

F 525°C/8h (L) 1296 1337 6.1 8.7

'800°C/100m WQ +

S 570°C/4h (L) 1153 1194 14.6 33.4

550°C/4h (T) 1200 1253 16 51

* 525°C/8h (T) 1290 1341 72.4(T-S) 12.1 38.8

*F 525°C/8h (T) 1315 1373 47-7(T-S) 9.4 19

^* 525°C/8h (L) 1318 1351 9-3 22.8

S 475°C/l6h (L) 1367 1439 7-3 12.6

430"C/l6h (T) 1603 1708 4.4 8.6

^* 800°C/100m AC +

S 525°C/8h (T) 1207 1263 14.4 40.4

F 525°C/8h (L) 1265 1309 6.1 15.6 s 475°C/l6h (L) 1283 1349 10.4 33.7

430°C/l6h (L) 1482 1610 5.6 6.5

"800°C/100m FGC +

S 550°C/4h (T) 1196 1244 10.4 22.4

525°C/8h (T) 1261 1298 8 16.7

450°C/l6h (T) 1357 1417 5.6 7.8

*800°C/100m FGC +

S 550°C/2h AC 1061 1113 106.8(T-S) 17.6 53.8

500°C/8h AC 1105 1152 97 (T-S) 16.0 57-8

F 500°C/8h AC 1102 1156 83.3 (T-S) 13.7 46.5

S 450°C/l6h AC 1170 1253 62.9 (T-L) 12.1 36.0

^* 770°C/100m WQ +

S 525°C/8h (T) 1175 1216 9-9 31.4

Legend

^* small section specimens (4.54mm diameter by l6mm guage length specimens)

* large section specimens (29mm thick fracture toughness specimens) S: Salt bath ageing, F: Furnace ageing L: Longitudinal direction, T: Transverse direction T-S: Transverse-short, T-L: Transverse-long WQ: Water quenching, AC: Air-cooling, FGC: Furnace Gas Cooling

IT: Isothermally Transformed from 890°C, 4θ°above β transus.