Journal Archive

Platinum Metals Rev., 1989, 33, (3), 106

The Alloying of Titanium Aluminides with Ruthenium

  • By A. Khataee
  • H. H. Flower
  • D. R. F. West
  • Department of Materials, Imperial College of Science, Technology, and Medicine,, London

Titanium is ductile in pure form but hardens considerably on alloying. Titanium alloys have excellent physical and chemical properties, being resistant to oxidation and corrosion in most dilute corrosive environments, including sea water, resulting in their widespread application in chemical plants, oil platforms, underwater engineering and medicine. Furthermore, with a low density of 4.5 × 103 kg/m3 and strengths well in excess of 1000 MPa, coupled with good creep resistance, certain titanium alloys have found extensive use in the aerospace industry.

The physical metallurgy and alloying behaviour of titanium is, however, rather complex due to its allotropy, involving a h.c.p. crystal structure (α ) below 882.5°C and a b.c.c. crystal structure (β ) above this temperature. Alloying elements may dissolve preferentially in α -titanium (α -stabilisers), in β -titanium (β -stabilisers) or in both (neutral), thereby increasing the stability range of these phases. They may also form compounds with titanium. Accordingly, titanium alloys are classed as α -alloys (predominantly α phase), β -alloys and so on, each group having its own characteristic properties. Alpha and near alpha titanium alloys (containing additional small amounts of β -stabilisers) have very good creep resistance, high strength and good weldability. Beta titanium alloys have poor creep resistance, but good workability and formability. It is evident that a combination of useful properties can be achieved in an α + β alloy. Indeed one of the first aerospace alloys developed in the 1950s was an α + β alloy, Ti-6A1-4V (weight per cent), also known in the United Kingdom as IMI318. Capable of operating up to about 350°C, IMI318 is the most widely used of the titanium alloys in aerospace applications.

Over the past 40 years advances in aircraft jet-engine technology have resulted in more efficient and powerful engines operating at increasingly higher temperatures, in order to achieve more thrust. The most advanced titanium alloys today are near-α alloys, such as IMI829 (Ti-5.5Al-3.5Sn-3Zr-1Nb-0.25M0-0.25Si), capable of operating at 550 to 580°C. Above 600°C, however, their oxidation resistance becomes inadequate and they suffer from creep and general metallurgical instability. Future generations of jet-engines will require alloys capable of operating at even higher temperatures, beyond the limit of the conventional alloys currently available. The titanium aluminides, Ti3Al (α 2) and TiAl (γ ), are ordered intermetallic compounds with excellent oxidation resistance and mechanical properties at temperatures up to about 900°C, offering considerable potential in this context. They are, however, brittle at room temperature and most of the research during the past 10 to 15 years has concentrated on overcoming this brittleness. An obvious way to improve ductility is by alloying, with the aim of refining the microstructure and reducing the slip length by producing a second phase dispersion. In this context, β -stabilising elements such as molybdenum, niobium and vanadium have been used to produce aluminide + β structures, with some improvements in ductility, but most workers have ignored the effects of non-conventional β -stabilising elements.

The platinum group metals show extensive solid solubilities in β -titanium. In the past, elements such as palladium have been added to titanium-aluminium, titanium-chromium, titanium-molybdenum and various ternary titanium-aluminium based alloys to enhance their low temperature corrosion properties in aqueous environments (1, 2). IMI260 and IMI262, containing 0.2 weight per cent palladium, are the recommended materials for use below 200°C in chemically reducing environments (3). Studies have shown that it is economically feasible to platinum ion plate some aerospace alloys such as AMS4975, IMI685 and IMI829 in order to increase their operating temperature (4, 5).

Previous work on the titanium-aluminium-ruthenium system (6) showed the presence of a ternary compound in the system, but lacked detailed experimental measurements. The present research was initiated in order to examine the effect of ruthenium on the structure and properties of titanium 17 to 37 atomic per cent aluminium alloys, mainly by establishing the details of constitution and phase equilibria in this section of the titanium-aluminium-ruthenium system.

Experimental Procedure

Alloy preparation and heat treatment entailed precautionary measures to minimise atmospheric contamination, in view of the known effect of impurities on titanium alloys (7, 8). The alloys listed in Table I were prepared by argon arc melting the constituents by a non-consumable tungsten electrode. Prior to melting, the furnace was evacuated and flushed several times with pure argon to reduce contamination. Weight losses after melting were around 0.2 per cent. Alloys, in the form of finger shaped ingots, were wrapped in molybdenum foil, sealed in quartz capsules containing argon and homogenised at 1250°C for 24 hours. Subsequent heat treatments were carried out at 1100 and 1050°C for 24 hours, 950°C for 96 hours and 870 and 770°C for 186 hours. All heat treatments were terminated by breaking the quartz capsules under water, in order to achieve rapid quenching while avoiding atmospheric contamination.


Table I

Nominal Composition of Alloys

Alloy number Composition, atomic per cent
Ti Al Ru
1 Balance 18
2 17 1
3 18 2
4 18 5
5 25
6 24 1
7 23 5
8 37
9 37 1
10 37 5
11 33 29
12 60 18

All the alloys were examined by light microscopy, scanning (SEM) and transmission (TEM) electron microscopy and X-ray diffraction. Elemental analysis was carried out using X-ray energy dispersive spectroscopy (EDS) of both thin foils and bulk samples. The former were prepared in a twin jet electropolishing apparatus employing a 10 to 12 per cent sulphuric acid in methanol electrolyte at −25 to −35°C. X-ray diffraction of bulk samples, powders and extracted particles was also carried out.

Mechanical testing, consisting of compression and tensile tests, was carried out on all the alloys. Selected alloys were isothermally forged at 1100°C to approximately 50 per cent reduction in thickness, before ageing at 770°C. Tensile testing of the forged samples was carried out at 20, 650 and 800°C.

Results and Discussion

Alloys Based on Ti3Al

Composition 17 to 25 atomic per cent Aluminium

Quenching alloys containing 1 to 2 per cent ruthenium from the high temperature β -titanium field results in transformation of b.c.c. β -titanium to martensite, see Figure 1. The martensite is hexagonal in alloys with 1 per cent ruthenium but orthorhombic in alloys with 2 per cent ruthenium. In alloys based on the stoichiometric Ti3Al composition (25 atomic per cent aluminium), the hexagonal martensite contains small ordered domains, shown in Figure 2, with DO19type crystal structure. The finger shaped domains show a positive concentration gradient of ruthenium away from the central rib of the martensite plates, suggesting the involvement of a diffusional growth element coupled with initial martensitic transformations in these alloys.

Fig. 1

Quenching a Ti-24Al-1Ru alloy from 1250°C results in the transformation of b.c.c. β-titanium to martensite, light micrograph

Fig. 2

The ordered finger shaped domains that occur in Ti-24Al-1Ru martensite have a positive concentration gradient of ruthenium away from the central rib of the martensite plates, dark field TEM micrograph

The tempering behaviour of the martensitic structures in these titanium-aluminium-ruthenium alloys is similar to that of other titanium alloys containing β -stabilising elements (9). In short, orthorhombic martensite reverts to β before precipitating low temperature equilibrium α or α 2 phases, whereas such transformation occurs without reversion in hexagonal martensite, depending on the tempering temperature. Details of solid state transformations in titanium-aluminium-ruthenium alloys have been published (10).

The beta phase can be retained at room temperture in alloys containing 5 per cent ruthenium by rapid quenching. Ageing of these alloys at 770 to 950°C results in the precipitation from the beta phase of α, α 2 and/or [Ti,Al]Ru. The latter is a ternary phase based on a solid solution of TiRu and AlRu, in which aluminium and titanium atoms can substitute for each other (6). Ageing in most alloys, however, results in fine scale α 2 + β or α + α 2 + β structures, shown in Figure 3, as can be seen from the partial isothermal sections determined in this work for the titanium-aluminium-ruthenium system at 1250°C and 770°C, and illustrated here in Figures 4(a) and 4(b), respectively. These diagrams serve to show the major constitutional features of this system, namely the extent of the α 2 + β field and the presence of two ternary phases, [Ti,Al]Ru and G-phase (see below). The ternary sections were plotted from phase composition data obtained by EDS analysis of thin foils. For the binary sides, the titanium-aluminium and titanium-ruthenium assessed diagrams by Murray (11, 12) and the aluminium-ruthenium diagram according to Shunk (13) were used.

Fig. 3

Fine scale α 2 + β structures occur in Ti-18Al-5Ru quenched from a temperature of 950°C; an α 2 region is indicated on this SEM micrograph

Fig. 4

Partial isothermal sections of the titanium-aluminium-ruthenium system at a temperature of: (a) 1250°C, (b) 770°C. The numbers refer to the alloys listed in Table 1

An important feature of the titanium-aluminium-ruthenium system is the low solubility of ruthenium (< ∼ 1 per cent) in α -titanium and the titanium aluminides, as evident in the phase diagram of Figure 4(b). The rejection of ruthenium atoms from α, or the aluminides, gives rise to an enriched β phase formation. Thus alloys with as little as 1 per cent ruthenium, contain equilibrium β phase at 770°C, Figure 5. Higher levels of ruthenium and aluminium result in the formation of [Ti,Al]Ru, as in the Ti-23Al-5Ru alloy, shown in Figure 6. Generally, higher levels of aluminium promote compound formation in these alloys, as in the case of the 37 atomic per cent aluminium alloy series, considered below.

Fig. 5

Equilibrium β phase particles are indicated in the α 2 matrix, in Ti-24Al-1Ru, quenched from a temperature of 770°C, TEM micrograph (tensile tested alloy)

Fig. 6

[Ti,Al]Ru particles occur in an α 2 matrix, in Ti-23Al-5Ru, quenched from a temperature of 870°C, TEM micrograph

Alloys Based on TiAl

Composition 37 atomic per cent Aluminium

Investigations of the martensitic transformations in these alloys revealed similar behaviour to that of lower aluminium content alloys, after quenching from the high temperature β field. Quenching of the Ti-37Al-5Ru alloy from 1250°C, however, resulted in only a partial retention of the β phase, see Figure 7. In comparison with lower aluminium alloys (17 to 25 per cent), where 5 per cent ruthenium is sufficient to cause complete retention of β, the Ti-37Al-5Ru alloy with partially retained β demonstrates the α -stabilising effect of aluminium. This is also consistent with the fact that the Ms increases with the aluminium content in titanium-aluminium alloys (14).

Fig. 7

The light micrograph of partially retained β phase in a Ti-37Al-5Ru alloy quenched from a temperature of 1250°C is supplemented by an electron diffraction pattern and key, showing the Burger’s type transformation

The microstructure and equilibria in these alloys are significantly affected by the presence of a ternary compound, the G-phase, in this region of the system. The G-phase is one of a group of compounds with the Th6Mn23 type structure, based on b.c.c., and isotropic with Mg6Cu16Si7, as described by Dwight (15). It is involved in equilibria with α 2, γ, and [Ti,Al]Ru, as indicated in Figure 4(b). In the Ti-37Al-]Ru alloy, the G-phase is present as very small oval particles in equiaxed TiAl (γ ) grains, as shown in Figure 8. The G-phase is, however, very brittle and its presence as large particles in the Ti-37Al-5Ru alloy caused severe embrittlement. Despite this, the compressive plasticity of the Ti-37Al-1Ru alloy showed marked improvement in comparison with the binary Ti-37Al alloy.

Fig. 8

G-phase particles are indicated in a TiAl matrix, in Ti-37Al-lRu quenched from a temperature of 770°C, TEM micrograph (tensile tested alloy)

The data collected in Table II were obtained from a limited amount of mechanical testing of these alloys, and show considerable improvement in ductility and strength for two alloys, namely Ti-24Al-1Ru and Ti-37Al-1Ru. These alloys consist of fine scale microstructures containing small β or G-phase particles in α 2 or α 2 + γ matrix (Figures 5 and 8), after ageing at 770°C. Although room temperature ductility is absent in tension, the results of high temperature tensile testing and sustained high temperature strengths of these alloys has clearly emphasised the potential of ruthenium as an important alloying addition to the Ti3Al and TiAl based alloys. The mechanical testing results given in Table II represent average values of strength and strain obtained by testing two or more alloys in similar conditions.


Table II

Mechanical Properties of Selected Titanium-Aluminium-Ruthenium Alloys

Alloy composition, atomic per cent Test temperature, °C Compressive strength, MPa Compressive strain, per cent Tensile strength, MPa Tensile strain, per cent
Ti-17Al-1Ru 20 1782 4.6
Ti-18Al-2Ru 20 1716 1.7
Ti-18Al-5Ru 20 1825 0.6
Ti-24Al-1Ru* 20 2037 11.5 772 0
  650** 595 0
  800* 651 <1
Ti-23Al-5Ru 20 1625 0
Ti-37Al-1Ru* 20 2249 14.2 807 0
  650* 748 5.7
  800** 491 1.5
Ti-37Al-5Ru 20 0

Alloys tested in tension at Royal Aerospace Establishment; ** only tested in tension


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This research was supported by a Science and Engineering Research Council Case Award, in collaboration with the Johnson Matthey Technology Centre. Isothermal forging and tensile testing of the alloys were carried out by Dr S. Kerry of the Royal Aerospace Establishment.

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