Aluminium – Molybdenum – Titanium

Introduction

Structural materials based on Ti are well-established in industry. The stability and properties of the high-temperature bcc (βTi) and low-temperature hcp (αTi) solid solutions can be controlled by the addition of alloying elements such as Al and Mo (see also ‘Notes on Materials Properties and Applications’). Therefore, the phase relations in the Ti-rich corner of the ternary Al-Mo-Ti system were already studied in the 1950s to 1980s [1954Mar, 1962Ere, 1962Min, 1963Min1, 1963Min2, 1963Min3, 1965Wil, 1969Cro, 1969Fed, 1969Nar, 1971Ham, 1972Ham, 1973Ham, 1974Zwi, 1975Ham, 1975Zan, 1977Zan, 1981Ere, 1981Tre, 1986Gro, 1988Gro].

During the 1980s, studies on TiAl-based alloys revealed their potential as light-weight structural materials for applications in the aerospace and automotive industries. As Mo was found to be one of the most promising alloying elements for improving the mechanical properties of the brittle binary Ti aluminides (see, e.g. [2001Dja]), many investigations focused on the complex phase relations in this Mo-poor part of the ternary Al-Mo-Ti system [1978Ban, 1980Ban1, 1980Ban2, 1992Kim, 1993Das1, 1993Das2, 1994Mor, 1997Che, 1997Sin1, 1997Sin2, 1998Has, 1998Joh, 1999Kim, 2000Kai, 2002Jun, 2003Sin, 2004Kim, 2006Aza, 2010Sch, 2011Kab, 2011May1, 2011May2, 2011Sch1].

Phase relations on the Al-Mo side of the ternary system have rarely been investigated. Experimental results on high-temperature phase equilibria in this composition region were reported by [2003Nin, 2017Kri1, 2019Kri].

The phase relations in the ternary system are very complex and difficult to study experimentally. The main reasons for this are:

- the occurrence of various metastable phases (see sections ‘Solid Phases’ and ‘Miscellaneous’)

- extended solid solubility ranges as, e.g., in (βTi,Mo), (αTi), and TiAl₃, which, moreover, strongly depend on temperature (see sections ‘Solid Phases’)

- phases originating from the binary boundary systems and existing as isolated single-phase fields in the ternary system in certain temperature ranges (such as (αTi), TiAl, TiAl₃, and Al₆₃Mo₃₇; see section ‘Solid Phases’)

- crystallographic similarities of some equilibrium phases ((αTi)/Ti₃Al, (βTi,Mo)/(βTi,Mo)_o, TiAl/TiAl₂/TiAl₃; see section ‘Solid Phases’ and Table 2).

Since the previous MSIT evaluation of the Al-Mo-Ti system by [2004Tre], a large number of new experimental investigations and thermodynamic modelling has been performed on the ternary system [2009Cup, 2010Cup, 2017Hua, 2017Kri1, 2018Wit, 2019Kri]. On the basis of these data, the system description of [2004Tre] including reaction scheme, liquidus surface, and isothermal sections can be significantly improved. Therefore, the present report represents a complete re-assessment of the system.

Investigations regarding the phase relations, structures, and thermodynamics of the Al-Mo-Ti system are summarized in Table 1.

Binary Systems

The Al-Mo and Al-Ti phase diagrams are accepted from [2020Wal] and [2020Pal], respectively.

The Mo-Ti phase diagram is accepted as described by [Mas2]. This version of the phase diagram contains a stable miscibility gap in the bcc (βTi,Mo) solid solution following clear experimental evidence from [1977Ter, 1978Ter, 1982Ter]. It should be noted that the existence of a miscibility gap was also confirmed in later experimental work [1998Fur], but is not taken into account in some recent thermodynamic descriptions [2018Hu, 2018Mar, 2020Jia].

Solid Phases

Crystallographic data and stability ranges of all solid phases are summarized in Table 2. The system contains two ternary and 14 binary intermetallic equilibrium phases as well as the solid solutions (Al), (βTi,Mo), and (αTi). Some information about the individual phases and the solubility of the ternary elements is given in the following.

(Al): The solubility of Ti and Mo in the fcc (Al) solid solution is very low, maximum values given in the literature for the binary subsystems are 0.8 at.% Ti [2006Sch] and 0.07 at.% Mo [2016Sch].

(βTi,Mo): The bcc (βTi,Mo) solid solution extends far into the ternary system and its phase field connects to the bcc MoAl phase at temperatures above 1470°C [2003Nin]. The (βTi,Mo) phase dominates the system at high temperatures covering about 80% of the liquidus surface (see, e.g., [2018Wit]). At low temperatures, its phase field shrinks and finally contracts in the Mo-rich corner.

(αTi): According to the binary Al-Ti phase diagram, the hcp (αTi) solid solution exists in two separated phase fields and the same holds true for the ternary system. At temperatures up to (1110±10)°C, (αTi) only exists along the Al-Ti boundary with Al contents up to >20 at.% and low Mo contents of a maximum of ~1 at.% [2017Hua]. At higher temperatures, another phase field of (αTi) appears at high Al contents. At 1500°C, the phase field of (αTi) has detached from the Al-Ti boundary and exists in the ternary system with compositions ranging from about 3 to 8 at.% Mo and 49 to 53 at.% Al [1998Joh]. The maximum temperature of (αTi) is estimated as (1515±10)°C; see section ‘Liquidus Surface’. The extension of the phase field into the ternary system is largest at 1350 to 1400°C reaching Mo contents of ~13 at.% [1997Che, 2018Wit].

Ti₃Al: The crystal structure of Ti₃Al can be regarded as an ordered version of hcp (αTi). Therefore, Ti₃Al is frequently designated as α₂ phase. It does not form from the melt but in a peritectoid reaction from the (αTi) and (βTi,Mo) solid solutions. The solubility for Mo reaches a maximum of about 2 at.% [2017Hua].

TiAl (also designated as γ phase) reaches a solubility of 12.6 at.% Mo [2017Hua]. Thermodynamic calculations indicate that, similar to the (αTi) phase, the maximum of the melting temperatures is located in the ternary system (1463°C) [2018Wit].

TiAl₂ forms congruently in the solid state from TiAl. A Mo content of 2.7 at.% was measured at 1200°C decreasing to a solubility of 1.2 at.% Mo at 800°C [2017Hua].

TiAl₃ (sometimes designated as ε phase) has a very peculiar phase field. Its homogeneity range is < 1 at.% in the binary Al-Ti system, but strongly widens in the ternary system. Mo can replace about 80% of Ti along (Ti_1–xMo_x)Al₃ (corresponding to about 20 at.% Mo) and about 20% of Al along Ti(Al_1–xMo_x)₃ (corresponding to about 15 at.% Mo) [1970Han, 1987Ere, 1990Ere]. The liquidus temperature increases with increasing Mo content, and thermodynamic calculations give a maximum of 1492°C for about 20 at.% Mo [2018Wit]. In the binary system, a low-temperature variant TiAl₃(l) exists, which is a superstructure of TiAl₃ (same tetragonal space group, but c axis four times as long). However, the temperature of the sluggish transformation is not known (735 to about 950°C according to [2001Bra], <600°C according to [1973Loo]) and the low-temperature polymorph is not taken into account in the accepted phase diagram of the binary system from [2020Pal]. Therefore, it is not considered in the present assessment.

MoAl₁₂, MoAl₅, Mo₅Al₂₂, Mo₄Al₁₇, MoAl₄, and MoAl₃: There is only scarce information on the solubility of Ti in the Al-rich Al-Mo intermetallic phases MoAl₁₂, MoAl₅, Mo₅Al₂₂, Mo₄Al₁₇, MoAl₄, and MoAl₃. [1994Sok] report solubilities of 2, 4, and 2 at.% Ti in MoAl₁₂, MoAl₅, and MoAl₃, respectively, at 500°C. However, these values are estimates resulting from metallographic observations and no composition analyses were performed. In a more recent experimental study by [2017Hua], much lower values of 0.6, 0, 0, and 0.4 at.% Ti were measured for the phases MoAl₅, Mo₅Al₂₂, MoAl₄, and MoAl₃, respectively. These low values are accepted for the present assessment. [1991Sch] found three different polymorphs of the binary phase MoAl₅. Besides the before-known variant, which has a hexagonal WAl₅-type crystal structure, they found another hexagonal variant MoAl₅(h’), which should exist in an intermediate temperature range, and a low-temperature variant MoAl₅(r) with a rhombohedral structure. Because the transformation temperatures are not known and since it was also suggested that the intermediate phase is only metastable [1991Sch, 2006Eum], MoAl₅ is treated as one single phase in the present assessment.

Mo₃Al₈ dissolves about 3 to 4 at.% Ti in the temperature range 800 to 1200°C [2017Hua]. The shape of the phase field indicates that Ti preferentially replaces Al.

Mo₃₇Al₆₃ is a high-temperature phase existing between 1490 and 1570°C in the binary system. Addition of Ti stabilizes the phase down to lower temperatures. At 1400°C, the solubility for Ti exceeds 16 at.% [2017Kri1]. The phase decomposes at 1317°C in a ternary eutectoid reaction [2019Kri].

MoAl: As mentioned above, the congruently melting bcc MoAl phase, which is another high-temperature phase in the Al-Mo system (1470 to 1707°C), forms a continuous solid solution with the (βTi,Mo) phase.

Mo₃Al: The only Mo-rich Al-Mo intermetallic phase Mo₃Al has a very high and nearly temperature-independent solubility for Ti in the range of 20 to 25 at.% (see, e.g., [2017Hua]).

(βTi,Mo)_o, (Ti₂AlMo): The occurrence of B2 (CsCl-type) ordering in the bcc (βTi,Mo) phase in a wide range of compositions centered around Ti₂AlMo was discovered already by [1958Boe]. Ordering occurs below a critical temperature, which depends on composition. This temperature is highest for the composition Ti₂AlMo (1418°C according to [2003Das]). Assuming a first-order transformation, [1993Bud] suggested the existence of a two-phase A2 + B2 field in their assessment. Ab initio calculations instead indicated the possible existence of an A2 + L2₁ two-phase field and predict L2₁ to be the lowest energy state of Ti₂AlMo [2000Alo, 2004Alo]. However, experimental investigations neither could prove such a two-phase field nor found any L2₁ ordering. Instead, the observations indicate the transformation from A2 (βTi,Mo) to B2-ordered (βTi,Mo)_o to be a second order reaction, which was also accepted for the assessments of [2004Tre] and [2018Wit] as well as for the present assessment.

σ, Ti₃Mo₃Al₄: A second ternary phase with the approximate composition Ti₃Mo₃Al₄ was found by [1970Han] in alloys containing 25 to 36 at.% Mo, 25 to 33 at.% Ti, and 42 to 48 at.% Al after annealing at 900°C for one week. Its existence was later on confirmed in several studies [1988Ere2, 1990Ere, 2017Hua, 2018Wit]. Since the crystal structure of this phase is of the σCrFe polytype, it is usually designated as σ. The phase forms in a peritectoid reaction at 1166°C from Mo₃Al and the ordered (βTi,Mo)_o solid solution.

Metastable phases: In Ti-based alloys, quenching from the (βTi,Mo) phase often results in martensitic transformations to metastable αʹ (hexagonal), αʺ (orthorhombic), or athermal ω (hexagonal) phases, which significantly affect the mechanical properties, see ‘Notes on Materials Properties and Applications’ and ‘Miscellaneous’.

Metastable phases were also found in a ternary alloy with 50 at.% Al and 15 at.% Mo, which was annealed at various temperatures between 800 and 1400°C and water-quenched [1997Che]. A sample quenched from 1400°C contained a phase with a tetragonal L6₀-type structure, which is closely related to the L1₀ structure of TiAl. It was interpreted as an intermediate state in the transformation from TiAl to another new phase, which possesses a D0₂₂-type structure (similar to TiAl₃) and was termed γʹ phase. [1997Che] describes the transformation from the L1₀ γ phase to the D0₂₂ γʹ phase as a diffusionless, massive transformation and concludes that this phase should be an equilibrium phase at 800°C. In addition, prolonged annealing at 800°C was found to result in the formation of a γʺ phase, which is another tetragonal structure resulting from further ordering of the γʹ phase but having a different composition. [1997Che] suggests that both γʹ and γʺ are equilibrium phases at 800°C. However, as their occurrence was not confirmed in any other experimental study, they were not taken into account in the present assessment.

A metastable phase with an orthorhombic D0₂₂-like structure (designated as D0₂₂ʹ) was found in a series of as-cast Al-rich Al-Mo-Ti alloys containing approximately 72-74 at.% Al, 21-24 at.% Mo, and 5-6 at.% Ti. It forms a checkerboard-like microstructure with the tetragonal D0₂₂ TiAl₃ equilibrium phase, which has a slightly lower Mo content. Due to both its structural similarity to the tetragonal TiAl₃ structure and the common structural features with the binary Al₈Mo₃ phase, the metastable D0₂₂ʹ phase is regarded as a precursor to the formation of Al₈Mo₃, which occurs as equilibrium phase during annealing at elevated temperatures [2020Lei].

Invariant Equilibria

The invariant equilibria are tabulated in Table 3 and the reaction scheme including the binary systems is shown in Figs. 1a to 1d. The schematic is mainly based on the modeling work of [2018Wit] which took into account the experimental results of [1958Boe, 1963Min3, 1970Han, 1972Ham, 1980Ban2, 1988Ere1, 1990Ere, 1992Kim, 1997Che, 1997Sin2, 1998Joh, 2000Kai, 2003Nin, 2010Cup, 2017Hua, 2017Kri1, 2018Wit]. In addition, the present assessment considers the experimental results from [1973Ham, 1975Ham, 1982Ere, 1987Ere, 1993Has, 1994Mor, 1997Sin1, 1998Has, 1999Kim, 2010Sch, 2019Kri], and the binary boundary systems have been adjusted according to the accepted binaries [Mas2, 2020Pal, 2020Wal]. The differences between the current assessment and the work of [2018Wit] regarding invariant equilibria involving the liquid phase are discussed in the next section as they are directly related to the liquidus surface. Therefore, the following text focuses on changes pertaining to the solid-state invariant equilibria.

The transition reaction (αTi) + TiAl₃ ↔ (βTi,Mo) + TiAl (U₄) must take place in a temperature range between 1400°C and 1452°C, even though the value resulting from modeling is much lower. [2018Wit] experimentally found tie-lines (βTi,Mo) + TiAl and TiAl + TiAl₃ at 1400°C indicating the existence of the tie-triangle (βTi,Mo) + TiAl + TiAl₃ at this temperature. Since 1452°C is the upper limit given by the eutectic reaction E₁, the reaction temperature is estimated as (1425±25)°C.

At (1340±20)°C the critical eutectoid reaction (βTi,Mo) ↔ (βTi,Mo)_o + TiAl (ec₁) takes place. This is the point at which the (βTi,Mo)_o phase-field, that formed at 1418°C at a composition of Ti-25Al-25Nb [2003Das], connects with the phase boundary (βTi,Mo)/(βTi,Mo)+TiAl. The “shift” of the (βTi,Mo)/(βTi,Mo)_o+TiAl line to lower Mo-contents is determined by the experimental results of [2010Sch], who obtained disordering temperatures above 1200°C for two alloys (Ti-45Al-3/7Mo) via in situ XRD measurements. Moreover, the critical transition reaction (βTi,Mo) + TiAl ↔ (αTi) + (βTi,Mo)_o (Uc₂) has to take place above 1200°C, since the alloy Ti-45Al-3Mo lies within the reported tie-triangle (αTi)+TiAl+(βTi,Mo) [2000Kai] and is ordered at this temperature [2010Sch]. As a result of this reaction and the fact that ordering of (βTi,Mo) is not accepted in the binary Al-Ti system [2020Pal], the critical transition reaction (αTi) + (βTi,Mo) ↔ (βTi,Mo)_o + Ti₃Al (Uc₃) has to take place at (1185±15)°C due to the formation of the tie-triangle (αTi) + (βTi,Mo) + Ti₃Al in the binary Al-Ti system at 1200°C. The resulting tie-triangle (αTi) + (βTi,Mo)_o + Ti₃Al reacts in a ternary eutectoid reaction (E₃) at (1110±10)°C. The modelling work of [2018Wit] had resulted in a value for this reaction temperature of 1098°C, but the experimental results of [2017Hua] clearly show that the tie-triangle (βTi,Mo)_o + TiAl + Ti₃Al is present at 1100°C. Therefore, the reaction must take place at a temperature in the interval 1100 to 1120°C.

Another difference from the modelling work of [2018Wit] results from the fact that ordering of (βTi,Mo) does not occur in the binary Al-Ti system according to the accepted phase diagram of [2020Pal]. Therefore, the tie-triangle (αTi) + (βTi,Mo) + Ti₃Al resulting from the binary reaction p₁₁ at 1170°C has to undergo ordering of (βTi,Mo) in the ternary system. The corresponding ordering reaction Ti₃Al + (βTi) ↔ (βTi)_o + (αTi) (Uc₄) must take place below 1170°C (p₁₁) and above 1107°C, which is the ordering temperature from the modelling of [2018Wit] in the binary subsystem. Therefore, the reaction temperature is estimated to be (1140±30)°C.

At 1166°C the ternary phase σ forms in a peritectoid reaction (βTi,Mo)_o + Mo₃Al ↔ σ (p₁₂) [2018Wit]. The involvement of (βTi,Mo)_o in this reaction requires the critical eutectoid reaction Ec₁ to finish above the formation temperature of the σ-phase and is estimated to take place at (1180±10)°C. The resulting tie-triangle Mo₃Al + (βTi,Mo)_o + Ti₃Al reacts in a transition-type reaction U₈ with the tie-triangle (βTi,Mo)_o + Mo₃Al + σ (which results from the formation of the ternary σ-phase) at an estimated temperature of (1080±10)°C. The calculated reaction temperature of 1139°C [2018Wit] is too high, since [2017Hua] found a tie-line σ + Mo₃Al at 1100°C which is clear experimental evidence that the reaction must take place below 1100°C. For the transition-type reaction U₉ there is no clear experimental evidence whether this reaction takes place above or below 1000°C. However, following the 1000°C isothermal section proposed by [2017Hua], the reaction temperature must lie below 1000°C and is estimated as (980±10)°C.

The proposed ternary eutectoid decomposition of Mo₃₇Al₆₃ [2018Wit] has been determined by [2019Kri] with DSC measurements to take place at 1317°C, and this value is accepted here.

As a result of the accepted congruent formation of TiAl₂ at 1215°C in the binary Al-Ti system [2020Pal], there is a TiAl₂ phase field growing from the binary Al-Ti system towards the phase boundary TiAl/TiAl+TiAl₃ with decreasing temperature. At 1205°C, the reaction e₆ (TiAl ↔ TiAl₂ + TiAl₃) occurs, splitting the TiAl single-phase region into two parts. The Ti-rich of the two resulting TiAl + TiAl₂ + TiAl₃ tie-triangles is stable down to room temperature, while the Al-rich vanishes at 977°C in the binary Al-Ti system (e₈).

Liquidus Surface

The liquidus surface has been newly established based on the modelling of [2018Wit] and experimental results from [1971Rex, 1998Joh, 2017Kri2, 2019Kri, 2020Lei] and is displayed in Fig. 2. The main differences compared to the version proposed by [2018Wit] are discussed in the following.

The invariant line separating the primary solidification fields of (βTi,Mo) and Mo₃Al has to change its character from peritectic (p₁) to eutectic (e₁) due to the accepted congruent formation of MoAl [1971Rex, 2017Kri2, 2020Wal]. This change in character is estimated to take place between 1700°C and 1800°C.

Directional solidification experiments [1998Joh] showed the existence of a single-phase (αTi) island at 1500°C which forms in a peritectic reaction L + (βTi,Mo) ↔ (αTi) (p₃). The modelling work of [2018Wit] had indicated a lower maximum reaction temperature of only 1490°C which is even lower than the binary reaction temperature (1491°C). Therefore, he had to introduce an additional peritectic minimum. However, since there is no experimental evidence for the suggested peritectic minimum and since the experimentally found maximum temperature clearly is above 1500°C [1998Joh], it is concluded that there is only a peritectic maximum (p₃) on the invariant line between (βTi,Mo) and (αTi). Since a reaction temperature of 1490°C was calculated by [2018Wit] and the experimental results of [1998Joh] prove this temperature to be >1500°C, a reaction temperature of (1515±10)°C is tentatively estimated in the present evaluation based on the data for the monovariant line p₅-E₁ and isotherm at 1500°C in the liquidus surface projection.

Furthermore, the character of the reaction L + TiAl ↔ (αTi) + TiAl₃ (U₃) is changed from eutectic [2018Wit] to transition type. There is no experimental evidence for a ternary eutectic, and, moreover, by assuming this invariant reaction to be U-type, there is no longer a need for adding a eutectic maximum on the monovariant line, which had been suggested merely by modelling [2018Wit].

The proposed peritectic reaction (p₄) [2018Wit] on the monovariant line Mo₈Al₃/TiAl₃ is supported by the extension of the TiAl₃ phase field across this monovariant line at 1400°C [2020Lei].

The reactions in the Al-rich corner of the liquidus surface are accepted from [2018Wit] and are displayed in Fig. 3.

Isothermal Sections

The isothermal sections at 1700°C and 1600°C are shown in Fig. 4 and Fig. 5, respectively. The phase boundary L/L+(βTi,Mo) is taken from the liquidus surface (Fig. 2). [2003Nin] showed that MoAl is connected with (βTi,Mo) at temperatures above 1470°C. Therefore, the proposed tie-triangle L + MoAl + (βTi,Mo) [1988Ere1, 1988Ere2, 1990Ere] is not necessary. The congruent formation of MoAl along with the resulting eutectic reaction L ↔ Mo₃Al + MoAl at 1597°C [2020Wal] leads to the formation of a second liquid field in these isothermal sections and the tie-triangle L + Mo₃Al + (βTi,Mo) is required.

At 1500°C (Fig. 6), there is an (αTi)-island [1998Joh] (cf. section ‘Liquidus Surface’) resulting in two tie-triangles L + (βTi,Mo) + (αTi). In addition, the binary Mo₃₇Al₆₃ phase occurs and its homogeneity range significantly extends into the ternary system as was shown experimentally [2017Kri1, 2019Kri] and by thermodynamic calculations [2018Wit]. At 1400°C, the Mo₃₇Al₆₃ phase is isolated from the binary system (Fig. 7) and vanishes at 1317°C in a ternary eutectoid reaction [2019Kri]. The resulting tie-triangle Mo₃Al₈ + (βTi,Mo) + TiAl₃ is observed at 1300°C [1990Ere] (Fig. 8).

At 1418°C, (βTi,Mo) begins to order in an alloy with composition Ti-25Mo-25Al [2003Das] and forms an isolated (βTi,Mo)_o-field that grows with decreasing temperature as shown in Fig. 7 and the figures which follow. In addition, at 1400°C the (αTi) phase field is connected to the binary boundary system as was shown by [1997Sin1, 1997Sin2] and extends far into the ternary system to approximately 13 at.% Mo [2018Wit]. With these results the position of the tie-triangle (αTi) + (βTi,Mo) + TiAl is estimated at 1400°C. The same tie-triangle has been experimentally determined at 1300°C (Fig. 8) and 1200°C (Fig. 9) [2000Kai] showing a significant decrease of the solubility of Mo in (αTi).

At 1300°C, the (βTi,Mo)_o phase field is connected with the two-phase field (βTi,Mo)_o + TiAl as a result of the eutectoid reaction (βTi,Mo) ↔ (βTi,Mo)_o + TiAl (ec₁). This two-phase field broadens with decreasing temperature as can be seen at 1200°C (Fig. 9). At 1253°C, the (βTi,Mo)_o phase field also reaches the boundary to the (βTi,Mo) + Mo₃Al two-phase field resulting in the eutectoid reaction (βTi,Mo) ↔ (βTi,Mo)_o + Mo₃Al [2018Wit]. Therefore, the (βTi,Mo) field is divided into two separate parts at 1200°C by the (βTi,Mo)_o field [2018Wit]. This separation ends at 1180°C with the reaction Ec₁ ((βTi,Mo) ↔ (βTi,Mo)_o + TiAl₃ + Mo₃Al) by which the central (βTi,Mo) field vanishes as is also visible in the 1100°C section (Fig. 10).

The formation of a TiAl₂ phase field is visible at 1200°C (Fig. 9) as a result of its congruent formation at 1215°C in the Al-Ti binary system [2020Pal] (as described in section ‘Invariant Equilibria’). The resulting tie-triangle TiAl + TiAl₂ + TiAl₃ was experimentally observed at 1100°C and 1000°C [2017Hua].

At 1166°C, the ternary σ-phase forms [2018Wit] and is visible in the isothermal sections at 1100°C and lower (Figs. 10 to 13). The isothermal sections from 1100°C to lower temperatures (Figs. 10 to 13) mostly rely on the experimental results of [2017Hua]. These results and the respective discussion in the above section ‘Invariant Equilibria’ indicate that a transition-type reaction U₈ ((βTi,Mo)_o + Mo₃Al ↔ TiAl₃ + σ) takes place below 1100°C. As a consequence, a tie-triangle σ + Mo₃Al + TiAl₃ must exist at lower temperatures, which was indeed experimentally observed at 1000°C [2017Hua]. Down to lower temperatures in the Al-rich corner, phase equilibria involving the liquid phase, TiAl₃, and the series of Al-rich, binary Al-Mo phases can be observed.

Between 900 and 800°C, the miscibility gap in the binary (βTi,Mo) solid solution ((βTi,Mo) ↔ (βTi,Mo)_Ti + (βTi,Mo)_Mo) opens at 850°C and extends into the ternary system with decreasing temperature. At about the same temperature (842°C), the second-order transformation (βTi,Mo)_o ↔ (βTi,Mo) changes its character to a first-order type transformation at the boundary to the two-phase field with Mo₃Al. This causes the formation of the tie-triangle (βTi,Mo)_o + (βTi,Mo)_Mo + Mo₃Al [2018Wit]. Both effects are visible in the 800°C isothermal section in Fig. 13.

Temperature – Composition Sections

There is a variety of different temperature-composition sections in the literature [1958Boe, 1963Min1, 1963Min3, 1969Kor, 1975Ham, 1980Ban2, 2003Nin, 2010Sch, 2011Kab, 2019Kri]. Starting from the binary Mo-Ti systems, [1963Min1, 1963Min3] show a series of temperature-composition sections for various fixed Al-contents as well as two sections parallel to Al-Ti with 0.5 and 5 at.% Mo. However, in these sections the miscibility gap in (βTi,Mo) [Mas2] as well as the existence of the phase TiAl₂ were not taken into account. Therefore, these temperature-composition sections are not considered here.

[1958Boe] present a section that connects the binary compositions Al₅₀Ti₅₀ and Mo₅₀Ti₅₀. It is solely based on their own data (light-optical microscopy and XRD) as [1958Boe] were the very first to study phase relations in this region of the ternary system. The existence of the Ti₃Al phase was not yet known at that time, and in addition the authors have assumed that the (βTi,Mo) + TiAl two-phase field is only stable above 1200°C, which is not correct as can be seen in the isothermal sections at 1100°C and below (Figs. 9 to 13). Therefore, their section is not accepted here.

The partial temperature-composition section of [1969Kor], which shows the effect of Mo additions to Ti₃Al, contains incorrect phase equilibria above 1100°C and, therefore, is not used in this assessment.

The partial temperature-composition sections presented by [1980Ban2] (at Ti-31 mass% Al and at Ti-9 mass% Mo) do not consider (βTi,Mo)_o and show incorrect phase equilibria for example at 1400°C. These sections are not accepted for the present assessment.

The temperature-composition sections shown in [1975Ham] and [2003Nin] disagree with the phase equilibria accepted for the present isothermal sections at 1200 and 1300°C (Figs. 8 and 9).

On the basis of experimental data in the literature, [2010Sch] calculated temperature-composition sections with constant Al-contents of 43 and 45 at.% in the range 0 to 10 at.% Mo. From these sections and the information coming from the isothermal sections of the current assessment, a revised temperature-composition section for 45 at.% Al and 0 to 10 at.% Mo has been constructed and is shown in Fig. 14.

[2019Kri] calculated a temperature-composition section at 65 at.% Al with Ti-contents in the range 0 to 10 at.%. This section qualitatively is accepted with, however, adding the missing high-temperature L + (βTi,Mo) and L + (βTi,Mo) + Mo₃₇Al₆₃ phase fields and adjusting the positions of the boundary lines to the respective isothermal sections in Figs. 5 to 8 (Fig. 15).

Thermodynamics

Total energy calculations for the disordered A2 and ordered B2 structure types of alloys with Ti₅₀ Mo₂₅Al₂₅ composition revealed slightly lower formation energy for the ordered variant [2000Alo]. Thermodynamic calculations to model phase relations in the Ti-rich corner of the system were reported by [1977Zan, 1986Gro, 1988Gro] and respective calculations for the TiAl + Mo region of the ternary system were performed by [1998Has, 1998Joh, 1999Kim, 2010Sch]. A rough estimation of the liquidus surface was obtained by thermodynamic calculations based on Gibb’s energy data of the binary subsystems available at that time [1982Dan]. CALPHAD modelling of high-temperature phase equilibria mainly focusing on the Al-Mo side of the ternary system was performed by [2009Cup, 2010Cup]. Finally, a comprehensive thermodynamic study including CALPHAD modelling of the entire ternary system was reported by [2018Wit]. The resulting isothermal sections, liquidus surface projection, and reaction scheme were used as basis for the present assessment; see also the respective sections above.

Notes on Materials Properties and Applications

Mechanical properties: Ti-based alloys are attractive for many industries due to their high strength, good corrosion resistance and low density. Mo stabilizes the high-temperature bcc (βTi) phase, while Al stabilizes the low-temperature hcp (αTi) phase. The potential of a combined addition of Al and Mo to improve the mechanical behavior of Ti has been studied since the 1950s [1954Cro, 1954Kes, 1954Mar, 1956Cro, 1957Cro, 1959Kno1, 1959Kno2, 1962Ere, 1962Min, 1963Min2, 1972Luz, 1973Ham, 1975Moi] and today is still a topic of scientific work [2019Jej] (Table 4). The hardness of the α phase increases with both Al and Mo content with Mo having a much stronger solid-solution-hardening effect than Al [2018Che]. Quenching of Ti-Mo alloys from the (βTi,Mo) phase can result in martensitic transformations to metastable αʹ (hexagonal), αʺ (orthorhombic), or athermal ω (hexagonal) phases, which can be suppressed by the addition of Al [1971Wil, 1975Moi, 1980Sas, 1993Cui, 2000Ike, 2013Dev, 2016Zha]. [1971Kho] studied a series of Ti-xMo-3Al (x = 0.5 to 30, in mass%) alloys. After quenching from the (βTi,Mo) region, alloys with 0.5-10 mass% Mo consisted of martensitic phases, while alloys with 15-30 mass% Mo retained the cubic (βTi,Mo) structure. Cold working after quenching significantly increased the strength, and the best mechanical properties were achieved for a Mo content of 15 mass% [1971Kho].

Extending beyond just the (αTi) and (βTi,Mo) phases, Ti aluminides based on TiAl (γ) and Ti₃Al (α₂) have attracted much interest. Very few studies have focused on the effect of Mo additions to single-phase Ti₃Al [1969Kor, 1976Zel] or TiAl, e.g., [1991Mae, 1993Has, 2013Zho]. Most work is related to the effect of Mo additions on lamellar or duplex two-phase TiAl + Ti₃Al alloys, which are the basis for alloys applied today as light-weight, high-temperature resistant blade materials in jet engines or car turbo chargers. Mo is the most potent β stabilizing alloying element (see, e.g. [2001Dja], which is a very beneficial effect at high temperatures as the existence of the cubic (βTi,Mo) phase allows hot working of the material, i.e., forming by hot rolling or forging. However, at low temperatures ordering of the β phase or transformation to ω-type non-equilibrium phases can be detrimental as these phases have a strongly embrittling effect, see, e.g., [2011May1, 2011May2, 2014Gup, 2019Erd].

Other two- or three-phase combinations investigated with respect to their mechanical behavior are (βTi,Mo) + Ti₃Al alloys containing 7 at.% Mo [1973Ham, 1975Ham, 1975Hid], three-phase (βTi,Mo)_o + Ti₃Al + metastable ω phase alloys with 3-4 at.% Mo [1991Dja, 1992Dja1, 1992Dja2], and Ti₃Al + TiAl + (βTi,Mo)_o alloys containing 2 at.% Mo [1994Li, 1994Mor].

Some research has focused on single-phase, B2-ordered (βTi,Mo)_o alloys with compositions around Ti-25Mo-25Al (at.%) [1992Nak, 1993Tho, 1994Tho, 2017Lu, 2019Pat, 2020Lu]. The yield stress of stoichiometric Ti-25Mo-25Al increases from 330 MPa to 570 MPa in the temperature range from room temperature to 900°C [1992Nak, 1993Tho, 1994Tho]. A recent investigation into the compressive stress–strain response of stoichiometric, Mo-enriched and Ti-enriched single-phase B2-ordered Al-Mo-Ti alloys in the temperature range 700 to 1100°C revealed a strong compositional influence. While the off-stoichiometric alloys showed high yield strength at low temperatures, which continuously decreases with increasing temperature, the stoichiometric alloy showed an apparent positive temperature dependence of the yield strength thus exceeding the values of the off-stoichiometric alloys at 1000°C. TEM analysis revealed that the alloys deform primarily via <111> slip at elevated temperatures [2020Lu]. First-principles calculations for B2-ordered Ti-25Mo-25Al give values for the bulk modulus B, shear modulus G, and Young's modulus E of 118, 45, and 120 GPa, respectively [2019Pat].

Two-phase TiAl₃ + Mo₃Al alloys with either equiaxed or lamellar microstructure were produced by heat-treatments of a Ti-36Mo-47Al [2003Nin, 2004Miu, 2005Miu]. Heat treatments and high-temperature compression tests in the temperature range 1000 to 1066°C proved the equiaxed microstructure to be very stable [2004Miu, 2005Miu].

Al-based ternary alloys containing small additions of Ti and Mo for strengthening were studied by [2010Kaz, 2014Zai]. Rapid solidification of an Al alloy with 0.1 at.% and 0.3 at.% Mo resulted in a single-phase supersaturated solid solution, whereas higher alloying contents of 0.30-0.75 at.% Mo and 0.10-0.25 at.% Ti caused the crystallization of a metastable MoAl₃ phase [2010Kaz]. Addition of 0.08 at.% Ti and 0.03 at.% Mo to Al metal resulted in significant grain refinement, and deforming such an alloy by ECAP (equal channel angular pressing) enhanced the mechanical properties [2014Zai].

Density: A density contour map covering the entire Al-Mo-Ti system was calculated by [2015Zhu] from the molar volumes and fractions of the constituent phases using a simple thermodynamic model for the molar volumes.

Corrosion behavior: Regarding the high-temperature structural application of Ti-Al-based alloys, oxidation resistance plays a very important role. Binary Ti-Al alloys with up to 50 at.% Al are known to have a very poor oxidation resistance due to the formation of non-protecting titanium oxide scales on the alloy surfaces. However, Mo additions to Ti₃Al [2010Ram] and TiAl [1994Shi, 1995Ana, 2002Fer] were found to have a very beneficial effect on the oxidation resistance leading to the formation of dense, protective aluminium oxide layers. It was also shown that the so-called ‘fluorine treatment’, which is another chemical concept to improve the oxidation behavior of Ti-Al alloys, is applicable to Mo-containing alloys [2014Pfl].

The corrosion behavior of Ti-based Al-Mo-Ti alloys in nitric acid at 25-75°C was studied for two-phase (αTi) + (βTi,Mo) alloys with compositions Ti-10Mo-10Al and Ti15Mo-5Al (in mass%, corresponding to Ti-5Mo-17Al and Ti-8Mo-9Al in at.%). The experiments revealed good corrosion resistance with a larger passive potential range than pure Ti in nitric acid [2005Pop]. The alloys were also proven to be resistant to concentrated chloride solutions [1996Rad].

The sulfidation properties of TiAl containing 2 at.% Mo was studied at 900°C and 1.3 Pa sulfur pressure in an H₂S-H₂ gas mixture [2000Izu]. The resulting surface scale was found to be a multi-layer structure composed of alloy substrate, layers of TiAl₂, TiAl₃, Mo-Al alloy, and a sulfide scale, which consisted of various mixtures of Ti-sulfides and Al₂S₃.

In biomedical applications, the corrosion behavior of implants plays a crucial role. Al-based Al-Mo-Ti alloys with up to 10 at.% Mo and 1 at.% Ti were electrodeposited on a Cu substrate resulting in dense, compact, and well-adhering layers. The corrosion resistance in simulated body fluid, so-called Ringer’s solution, was investigated by potentiodynamic pitting corrosion measurements. It was found to be better than that of pure Ni but somewhat inferior to 316 L stainless steel [2010Tsu].

Functional properties: Application-related research on Al-Mo-Ti alloys focuses nearly exclusively on structural applications, while only very few studies deal with functional properties. One example refers to the effect of Al additions on the superconductivity of Ti_0.75Mo_0.25. The superconduction transition temperature of (Ti_0.75Mo_0.25)_1–xAl_x alloys with x = 0 to 0.06 was found to decrease approximately linearly from 3.9±0.1 K at x = 0 to 1.9±0.1 K at x = 0.06 [1985Ho]. Another example is related to shape memory behavior. Ti-based Al-Mo-Ti alloys containing 6-6.5 at.% Mo and 3 at.% Al can reversibly change their structure on heating and cooling from the high-temperature cubic (βTi,Mo) phase to the low-temperature orthorhombic αʺ martensite. Therefore, their potential for shape memory applications was studied. Since these Al-Mo-Ti alloys have a higher cytocompatibility than Ni-Ti alloys, they were highlighted as potential Ni-free replacements in biomedical applications [2007Yam].

Miscellaneous

Non-equilibrium phases and microstructures: Several publications discuss metastable phases in the Al-Mo-Ti system, see also above in ‘Notes on Materials Properties and Applications’ and Table 5. [1993Cui] showed that valence electron concentration, electronegativity, and atomic radius of the elements are factors controlling the occurrence of the metastable ω phase. First-principles calculations of the lattice stabilities of (βTi,Mo), (βTi,Mo)_o, and the hexagonal variants ω₀ and ωʺ for a composition Ti₄Al₃Mo revealed that the most stable structure should actually be ωʺ followed by the ordered structure (βTi,Mo)_o [2020Ye]. In contrast to that, first-principles calculations of [2014Hu] for the same composition predicted (βTi,Mo)_o to be more stable than the ω-related phases. Comparing the stability of the ordered (βTi,Mo)_o versus the disordered (βTi,Mo) phase in TiMo_xAl_1–x alloys by DFT (density functional theory) calculations, [2015Hol] found that the B2-ordered variant is stable for Mo contents above 4 at.% (x = 0.04). Moreover, they proved that there is a barrier-less transformation path from (βTi,Mo)_o to TiAl allowing for a spontaneous transformation from the symmetry-stabilized (βTi,Mo)_o to the ground state TiAl phase. TEM experiments of [1993Li] had already indicated the existence of this diffusionless transformation. A high number of antiphase domains (APDs) were found in the TiAl phase of a Ti-2Mo-48Al alloy after quenching from 1350°C. These APDs originate from the transformation of (βTi,Mo)_o to TiAl proving that at high temperatures the (βTi,Mo)_o phase is the equilibrium phase [1993Li].

Non-equilibrium microstructures in Ti-(1-2)Mo-(45-47)Al (at.%) alloys were obtained by rapid solidification of liquid droplets revealing a very strong dependence on both Al and Mo content [2016Ken]. The evolution of the lamellar Ti₃Al + TiAl microstructure in a Ti-2Mo-48Al alloy was studied as a function of aging temperature (1000 to 1100°C) and time (1 to 16 h) and compared to binary Ti-48Al and ternary Ti-4Nb-48Al alloys. Aging resulted in the transformation of the lamellae into an equiaxed microstructure, as well as precipitation of a small amount of (βTi,Mo)_o phase. Compared to the binary and the Nb-containing ternary alloy, the microstructural morphology change proceeded much faster in the Mo-containing alloy, and the volume fraction of equiaxed microstructure was 88% after 16 h aging at 1100°C [2005Sre].

The parameters for the technological manufacture of Al-Mo-rich Al-Mo-Ti alloys (composition range 48-52 mass% Mo and 6-9 mass% Ti, corresponding to typical compositions in at.% near Ti-23Mo-70Al) by an aluminothermic smelting process were estimated by thermodynamic simulations [2013Udo]. Such non-equilibrium foundry alloys are used as alloying material for the production of Al- and Mo-containing titanium alloys. Their crystallization process and phase compositions were investigated by [2014Sel]. The alloy contained a certain amount of Si. Therefore, a Si-containing phase with an average composition given as Ti_2.3Mo_2.6Si_0.6Al_4.2 was found in addition to the phases Mo₃Al, Mo₃Al₈, and Ti₃Al.

Orientation relationships: The interfacial structures in a Ti-5Mo-50Al alloy, which had been annealed for 150 h at 1240°C and subsequently furnace-cooled, were studied by conventional and high-resolution TEM. The alloy consisted of the three phases Ti₃Al + TiAl + (βTi,Mo)_o. All orientation relationships were such that the close-packed planes and directions in all of the phases were parallel [1996Das].

Diffusion: From the analysis of diffusion couples in the (βTi,Mo) solid solution region, interdiffusion coefficients and the diffusion mechanism were derived at 1100°C [2021Cha] and 1250°C [2014Che]. The interdiffusivity was found to depend strongly on the Mo content, while this was not the case with respect to the Ti and Al contents [2021Cha]. Diffusion was found to occur primarily via a vacancy mechanism [2014Che].

Site occupation: The site occupation of Mo has been analyzed for the Al-Ti intermetallic phases Ti₃Al [1988Nan, 1999Hao, 2000Yan], TiAl [1990Nan, 1998Woo, 2000Son, 2000Yan, 2001Kan, 2009Ayk], and the B2-ordered (βTi,Mo)_o-phase [1988Nan, 1995Che, 1996Sik, 2007Sin, 2008Sin, 2019Pat]. For Ti₃Al alloys, measurements by the ‘atom location by channeling-enhanced microanalysis’ (ALCHEMI) method show that Mo atoms occupy Ti sites [1999Hao]. In the case of the L1₀ lattice of the TiAl phase, very contradicting results were reported. According to first principles calculations by [2000Son] and [2009Ayk], Mo preferentially occupies the Ti sublattice sites, whereas according to [2001Kan] Mo prefers the Al sites for both Ti-rich and Al-rich TiAl. In contrast to these results, [1990Nan, 1998Woo] find that Mo can fill either sublattice, i.e., for Al-lean compositions Mo can occupy the left-over sites on the Al sublattice and, similarly, for Ti-lean compositions Mo can occupy the Ti sublattice. For the B2-ordered (βTi,Mo)_o phase, it was shown that the Ti and Al atoms tend to never share a common sublattice site, while Mo can fill either sublattice depending on composition [1988Nan, 1995Che, 2019Pat]. For a composition Ti₂AlMo, one of the B2 sublattice sites is entirely filled up with Ti atoms, while the Al and Mo atoms randomly share the other site [1996Sik, 2007Sin].