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 a temperature 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 while retaining the alloy in a sufficient ductility for fabrication processes.

Other alloying ingredients are known to stabilize β-phase at progressively lower temperatures, until a stable β-phase is obtained at normal temperatures. It is known in the art that the addition of β-stabilisers to a titanium base alloy results in good response to heat treament 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 (..-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, (-•-stabilizers known in the art are Mo, V, Nb, Ta, Mn, Cr, Fe, W, 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 W, 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 created difficulties with regard to ensuring 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 No. 782564 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-l.5Fe-*Al. More recently developed alloys include the Transage series (e.g. T12 ^Q ) 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 the US Patent 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 clearly 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 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 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 base alloy comprising the following constituents in proportions by weight:

vanadium 11.5-12.5 % iron 1.2-1.7 % aluminium 3.5"*+-2 % titanium balance, save for incidental impurities.

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

A preferred composition for the improved titanium alloy comprises:

vanadium 12.0% iron 1.5% aluminium 4.0% titanium balance, save for incidental impurities

(hereafter denoted Ti-12V-1.5Fe-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, of which:

Figure 1 shows the age hardening curves of the Ti-12V-l.5Fe-4Al alloy aged in salt baths at 350-600°C.

Figure 2 shows yield stress against fracture toughness the alloy Ti-12V-1.5Fe-*4Al compared to that of a prior art alloy Ti-10V-2Fe~3Al.

Figure 3 shows yield stress against tensile elongation for the alloy Ti-12V-1.5Fe-*.Al compared to that of a 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 the age hardened (α + β) microstructure are similar to thise reported for similar microstructures in the prior art alloys.

Results have shown that the β-quenched Ti-12V-l.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 and aged in salt baths at 350-600 ^c C. Such alloys attain the maximum hardening and hardness of about 500Hv which is unchanged over a temperature range of 350-*00°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 - 600°C results in a lower level of hardness compared to that attained in the β quenched and aged condition. Age hardening of the alloy annealed at a temperature 50°C below its β-transus (at 750°C) is almost the same for salt bath and furnace ageing.

Mechanical properties

Table 1 shows the tensile properties and fracture toughness values of the age-hardened Ti-12V-1.5Fe-*Al alloy. The table shows the annealing treatment followed by any subsequent age hardening. The duration of each treatment is given in hours(h) or minutes(m). For the high (α + β) annealed and aged condition the Ti-12V-1.5Fe-4Al alloy shows excellent tensile ductility. For example at yield stress levels 1250 -1300 MPa, the alloy shows a tensile elongation of 5-8/.. However for the β-quenched and aged condition, the tensile ductility is reduced and is also affected by heating rate used during ageing (1321 MPa / 0% combination of a furnace aged condition compared to 1298 MPa / % combination of a salt bath aged condition) .

Fracture toughness of the age hardened alloy is affected by the prior annealing temperature. When the fracture toughness of the β-quenched and aged alloy is compared to that of the high (a + β) annealed and aged condition, the latter condition shows better fracture toughness (80 MPam ^1/2 for yield stress of 1248 MPa against 60.3 MPam ^1/2 for 1233 MPa of the β-quenched and aged condition; Table 1).

Salt bath and furnace ageing has little effect on the fracture toughness of the alloy since a precipitate morphology is little affected by the salt bath and furnace ageing. The alloy annealed at 750 °C (also 50°C below the β-transus) and aged at 500°C shows similar fracture toughness values, 63 and 59 MPam ^{1 2} for the salt bath and furnace aged conditions that show the yield stress values respectively 131 and 1321 MPa. Even in the β quenched and aged condition in which furnace ageing deteriorates tensile elongation (1564 MPa/0.8j. combination for the salt bath aged compared with 1321 MPa / 0% combination for the furnace aged Table 1) fracture toughness does not deteriorate (1233 MPa / % / 60.3 MPam ^1/2 combination for the salt bath aged compared with 1350 MPa / 0% / 48.3 MPam ^1/2 combination for the furnace aged).

At a plate thickness up to 29 mm of the fracture toughness specimens age hardening after annealing (in the β or high a + β field) treatment of the alloy was insensitive to the different cooling rates used in the study e.g. the longitudinal section of the alloy shows similar hardness/yield stress values 396 Hv/1269 MPa and 3 6 Hv/1252 MPa respectively for the water quenched and air cooled materials aged at 25°C after annealing at 780°C.

Comparison of mechanical properties of the invention to those of the prior art alloy Ti-10-2--.

The mechanical properties of the invention are compared with those reported for the Ti-10-2-3 alloy in figures 2 and 3- These figures show respectively yield stress versus tensile elongation and yield stress versus fracture toughness values for the (α + β) microstructure and demonstate that the invention has superior yield stress/ ductility/fracture toughness combinations compared to the Ti-10-2-3 alloy. Figure 2 shows that at yield stress of 1200-1500 MPa, the fracture toughness values of the invention developed here are higher than those of Ti-10-2-3 alloy. For example at a yield stress of about 1250 MPa the alloy according to the invention shows fracture toughness values of 80 Mpam ^{1 2} compared to 4l MPam ^{1 2} observed for the Ti-10-2-3 alloy. At a high yield stress value of 1350 Mpa the fracture toughness of the invention is superior at 65 MPam ^{1 2} compared to a value of 30 MPam ^1/2 observed for the Ti-10-2-3 alloy. At this high yield stress level the bimodal structure of the alloys show tensile elongation of 3~1% compared with the tensile elongation of 2-4 % for the Ti-10-2-3 alloy (Fig 3).

Alloys according to the invention, show improved mechanical properties compared to the Ti-10-2-3 alloy because they use the optimum V, Fe and Al combinations to achieve after processing the most attractive (α + β) microstructure at high strength levels. They do not depend on the transitional β > ω + β reaction as in Ti-10-2-3 alloy to achieve the high strength in the (α + β) structure. The uniform a type structure of the Ti-10-2-3 alloy shows poor fracture toughness. In alloys according to the invention a plates of high aspect ratio and of a zig zag morphology are observed at temperatures in which the β > __> + β reactions occur in the Ti-10-2-3 alloy. This enhances fracture toughness in the new alloys.

Table 1 : Mechanical Properties of Ti-12V-1 .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 240 860 28.1 36

850°C/lh WQ + 308 846 30 37.8

850°C/lh IT

750°C/lh WQ +

S 525°C/8h (L) 1190 1235 6.3 7-8

850°C/lh IT

500°C/8h (L) 1225 1281 5-8 7

850°C/lh WQ +

F 590°C/5h (L) 1283 1318 3.4 2.2

S 570°C/4h (L) 1140 1183 4.5 6.5

550°C/2h (L) 1181 1222 6.4 10.4

/4h (T) 1233 1300 60.3 (T-S) 3.0 4

/6h (L) 1161 1207 6.8 18.3

F 550°C/7h (L) 1321 1321 0.0 -

525°C/8h (T) 1350 1350 48.3 (T-L) 0.0 -

S 500°C/8h (L) 1298 1343 3-2 4

430°C/I6h (L) 1564 1587 0.8 1.3

Legend

Salt bath ageing, T-S: Transverse-short, furnace ageing T-L: Transbverse long, Longitudinal direction, WQ: Water quenching Transverse direction AC: Air-cooling

FGC: Furnace gas cooling

IT: Isothermally transformed from the β field (805°C and above)

Table 1 continued : Mechanical Properties of Ti-12V-l .5Fe-4Al Alloy

Heat 0.2% Y. Tensile Fracture Total Reduct.

Treatment Stress Stress Toughness Elong. in Area

(MPa) (MPa) (MPa) ( % ) ( % )

805°C/lh WQ +

S 500°C/8h (L) 1297 1339 3-2 4.4

/I6h (L) 1262 1300 4.1 10.8

780°C/2h WQ +

S 550°C/2h (T) 1248 1293 80 (T-S) 6.9 10.4

525°C/8h (L) 1268 1334 7-6 15.9

F 525°C/8h (L) 1226 1264 9-6 21.2

46θ°C/l6h (L) 1535 1556 2.8 1.8

780°C/2h AC +

F 525°C/8h (L) 1252 1297 7.1 12

780°C/2h FGC +

S 500°C/8h (T) 1361 1420 63.6 (T-S) 6.8 9-9

450°C/l6h (T) 1449 1504 51.7 (T-L) 4.3 6.7

760°C/2h WQ +

S 500°C/8h (L) 1220 1247 7-8 19.6

750°C/4h WQ +

S 550°C/2h (L) 1110 1144 14.1 47.2

F 550°C/4h (L) 1125 1186 11.3 33-2

525°C/8h (L) 1162 1192 11.3 29.3

S 500°C/lh (L) 1218 1322 7-2 11.3

/8h (L) 1190 1251 5-5 17.8

/24h (L) 1154 1230 10.4 31.9

F 500°C/8h (L) 1259 1290 6.4 15.4

460°C/l6h (L) 1347 1405 6 11

S 450°C/l6h (L) 1370 1435 4 7.7

425°C/l6h (L) 1446 1500 5-5 7.5

750°C/6h WQ +

S 500°C/8h (L) 1269 1300 7.1 20.4

750°C/4h AC +

F 500°C/8h (L) 1205 1263 8.6 17.6

S 450°C/l6h (T) 1384 1427 3-6 4.3

750°C/2h FGC +

S 500°C/8h (T) 1315 1390 64.4 (T-S) 5 10

F 500°C/8h (T) 1321 1384 59-3 (T-L) 5-3 6.8

730°C/4h WQ +

S 480°C/8h (L) 1227 1304 7-2 15.2

S 425°C/8h (L) 1232 1322 8.1 25.8