Alloying Elements in Intermetallic γ‐TiAl Based Alloys – A Review on Their Influence on Phase Equilibria and Phase Transformations

Intermetallic γ‐TiAl based alloys are innovative structural materials dedicated to applications in the automotive and aeronautic industries. Especially their low density, high specific yield strength, and excellent resistance against creep and oxidation make them a suitable choice for structural high‐temperature components in combustion engines. However, further improvement of their properties and processability is required to conquer new application areas and facilitate cost‐effective production. While the incorporation of additional alloying elements is a promising possibility, their effects on the phase transformations and phase equilibria need to be considered rigorously. In this context, this review provides a detailed survey of the research work on the influence of technically important alloying elements on the Ti–Al phase diagram. First, an introduction to the fundamentals of the phase transformations in γ‐TiAl based alloys and the consequences of changed cooling conditions, relevant for example to the latest developments within the field of processing, is given. Afterward, the alloying elements, categorized with respect to their stabilization effect, are discussed and their particularities are highlighted. Topics covered include established ternary phase diagrams and isoplethal sections, the stabilization of additional phases as well as the influence of alloying elements on the microstructure of modern engineering intermetallic γ‐TiAl based alloys.


Introduction
Recent regulations on the allowed emissions of harmful greenhouse gases combined with the ever-growing demand of personal and economic mobility and surge in fuel prices sparked a search for methods to increase the efficiency of combustion engines in the aeronautic and automotive sector.Since improved engine concepts only present one part of the solution and often necessitate an enhancement of the material properties, novel design and processing strategies for the utilized material systems are required from a material scientist's point of view.[3] Among the many different alloying concepts for γ-TiAl based alloys, especially the GE-4822 alloy (Ti-48Al-2Nb-2Cr, in at%) and the TNM alloy (Ti-43.5Al-4Nb-1Mo-0.1B, in at%) have been successfully implemented by globally leading engine manufacturers on the industrial scale.The former has been included in the GEnX engine by General Electrics since 2011, [2,3] while the latter has been introduced by Pratt and Whitney in their Geared Turbofan engine in 2014. [4,5]he main benefit of intermetallic γ-TiAl based alloys, when compared to their main competitors, namely, the formerly used heavy Ni-base superalloys, can be found in their low density of roughly 4 g cm À3 (depending on the alloy composition).In combination with their high yield strength and stiffness, this gives rise to excellent specific mechanical properties perfectly suited for the application of high-temperature materials and especially the improvement of the efficiency of combustion engines.[8][9] The weight-saving potential due to the utilization of γ-TiAl based alloys has also been recognized by the automotive sector, in particular in the form of valves and turbocharger turbine wheels of high-performance automotive diesel engines, where this class of materials has the potential to replace heavier Ni and Ti-base alloys. [6,10,11]14] A crucial aspect contributing to the success of intermetallic γ-TiAl based alloys in the industrial sector is the exploitation of the beneficial effects of additional alloying elements.][17] In particular, alloy design concepts belonging to the latest generation of γ-TiAl based alloys are often developed to satisfy the requirements of a certain envisaged manufacturing route, e.g., by improving the hot workability by stabilization of ductile phases. [17,18]However, the resulting alterations of the phase transformation behavior and associated changes in the microstructural characteristics of the material have to be equally considered when adding additional alloying elements to γ-TiAl based alloys to harness their full potential.Especially the knowledge of the phase stability regions and associated transition temperatures as well as of the volume fractions of the individual phases as a function of temperature and alloy composition are of crucial importance for the development of suitable heat treatment strategies and sustainable processing routes.An excellent example, for which this procedure has been applied successfully, is the TNM alloying concept. [7,17,19]Due to changed phase equilibria at high temperatures associated with the addition of Nb and Mo, the hot workability of this alloy family could be improved significantly with respect to other γ-TiAl based alloys and, thus, production cost could be lowered immensely. [17]To minimize the adverse effects of said alloying elements on the application-relevant properties, special heat treatments were derived utilizing the characteristics of the phase transformations in the alloy system on the basis of a profound experimentally generated knowledge of the influences of Nb and Mo. [19,20] tool, which has been proven valuable for the development of the TNM alloy as well as for alloy development programs in general, is the prediction of binary and multicomponent phase diagrams via thermodynamic calculations. [7,17,21,22]In particular, while the choice of type and amount of additional alloying elements has been guided mainly by the expertise of the alloy designer and, essentially, involved an exhaustive trial and error process in the past, such an application of thermodynamic calculations and the availability of ever-growing databases have enabled researchers to adopt a more knowledge-based approach. [17]Fast predictions of the phase transformation behavior for a certain alloy composition, e.g., with regard to the solidification behavior and phase fraction evolution, can save time and lower the costs of alloy development programs.However, for this procedure to function properly, the alloying elements of interest need to be covered in the thermodynamic databases in the relevant composition ranges.To further optimize the quality of the results of such thermodynamic databases, which to date still lack some accuracy, especially for high amounts of additional alloying elements and the associated cointeractions, reliable experimental data are a necessity. [23,24]These involve, for example, phase transformation temperatures as a function of alloy compositions and phase composition as a function of temperature.
This review is dedicated to the question as to how alloying elements influence the phase transformations and phase equilibria of intermetallic γ-TiAl based alloys.It shall provide the reader with a fundamental overview on both the thermodynamics of γ-TiAl based alloys and the associated effects of technically important alloying elements.The alloying elements in the focus of this work, categorized according to the corresponding stabilized phases, are: 1) the β-stabilizing elements Nb, Mo, W, Ta, Mn, V, Cr, Fe, Zr, and Hf; 2) the α-stabilizing elements Al, Si, C, O, and N; and 3) further elements such as B and rare-earth metals.The first part of this review establishes a profound background on the Ti-Al system and the relevant phase transformations occurring in γ-TiAl based alloys upon cooling, i.e., β !α, α !γ, and α !α 2 þ γ.Furthermore, changes to these transformations associated with different cooling rates are addressed.The second part presents a detailed compilation of the literature on the influence of the aforementioned alloying elements on these phase transformations, the stabilization of additional phases and the occurring microstructure.
For previous reviews dealing with alloying elements in γ-TiAl based alloys, the reader is referred to the previous studies. [25,26]owever, while Duan et al. [25] focused on the history and development of γ-TiAl based alloys and Raji et al. [26] mainly investigated the effects of additional elements in β-solidifying alloys, this review is dedicated specifically to phase transformations and the associated influence of alloying elements in γ-TiAl based alloys.This work rigorously elucidates the fundamentals of the phase transformations in the Ti-Al system, e.g., with regard to the influence of the cooling rate, and describes in great detail the thermodynamic impact of the investigated alloying elements, both quantitatively and qualitatively, addressing all γ-TiAl based alloys regardless of their solidification type.

Fundamental Phase Transformations in Intermetallic γ-TiAl Based Alloys
The fundamental basis for all titanium aluminide alloys is the binary Ti-Al system, which is presented in Figure 1a based on the assessment of Schuster and Palm. [27]While some research has been dedicated to single-phase α 2 -Ti 3 Al and single-phase γ-TiAl alloys in the past, their inherent low ductility at room temperature has led to the development of two-phase alloys based mainly on the γ phase with minor volume fractions of α 2 phase. [28,29]30] This culminates in the presence of several different phases depending on temperature.The crystal structures and the corresponding crystallographic data of the occurring phases are given in Figure 1b and Table 1, respectively.At room temperature, the main constituents are the ordered γ-TiAl phase and the ordered α 2 -Ti 3 Al phase.The crystal structure of the former can be described as a stacking of individual Ti and Al atom layers along the [001] direction and is, from this point of view, closely related to the face-centered cubic crystal structure.However, the γ phase is slightly distorted along the c direction resulting in a tetragonal structure with a c/a ratio of 1.02 at the stoichiometric 50Ti:50Al composition.This ratio, which is of course affected by the amounts of incorporated Ti and Al, plays an important role, since an increased tetragonality negatively impacts the dislocation motion and, thus, the ductility of the material. [28,31,32]enerally, in the Al range of γ-TiAl based alloys, neither the γ phase nor the α 2 phase is present in their ideal stoichiometry or, synonymously, their perfectly ordered state even in thermodynamic equilibrium, as shown in Figure 1a.Due to the applied technical cooling rates as well as the incorporation of additional elements, their composition may change even further. [33,34]The ratio between the γ and α 2 phase, which is in favor of the γ phase in advanced alloying concepts, is mainly governed by the respective Al content, but also influenced by additional alloying elements and the material's thermal history. [3,6,7,16]For example, a cast binary Ti-45Al (at%) alloy contains 72 m% γ and 28 m% α 2 , while a binary Ti-48Al (at%) alloy consists of 92.3 m% γ and only 7.7 m% α 2 . [35]In the case of binary γ-TiAl based alloys, an Al content of 48 at% has been found to yield the highest strain to rupture at room temperature due to an optimal volume ratio between γ and α 2 phase. [32]dditionally, the size of microstructural features has been found to depend on the Al content. [36,37]Further essential effects of Al include a significant influence on the solidification behavior, as discussed in detail in Section 2.1, as well as an improved oxidation resistance. [6]e phase transformation pathway typically observed during the heating of a binary γ-TiAl alloy starts with the disordering of the hexagonal α 2 phase into the high-temperature hexagonal α-Ti(Al) phase, once the temperature exceeds the eutectoid temperature.This temperature as well as other important phase transition temperatures are marked by colored lines in Figure 1a.Upon further heating, the γ phase continuously transforms into the α phase.The temperature at which the complete dissolution of the γ phase is achieved is called the γ-solvus temperature.While this temperature coincides in the binary Ti-Al phase diagram with the α-transus temperature, i.e., the temperature at which the material enters an α single-phase field region, this is not guaranteed in the case of ternary or more complex alloying systems, e.g., see Section 3.2.At the highest temperatures, the disordered body-centered cubic β phase is present.Similar to the α-transus temperature described above, the β-transus temperature marks the transition into the β single-phase field region.While both the α and β phase are also prominent in conventional Ti-base alloys, their individual phase compositions are characterized by significantly higher amounts of dissolved Al in the case of γ-TiAl based alloys.A more detailed description of the phase transformations associated with these transition temperatures and the influence of the cooling rate and additional alloying elements is given in Section 2.2 and 3, respectively.
Generally, relevant phase transformations in thermodynamic systems fall into one of two categories depending on the aggregation states of the involved phases: liquid-solid and solid-solid.In the case of γ-TiAl based alloys, the multitude of phases present yields the occurrence of several distinct phase transformations.A look at the binary Ti-Al phase diagram in Figure 1a reveals that, in addition to three transformations involving the liquid phase (L), i.e., precipitation of β from L, precipitation of α from L and L þ β !α, three solid-solid phase transformations may arise for the Al contents of intermetallic γ-TiAl based alloys during cooling, i.e., β !α, α !γ and α !α 2 þ γ.The respective manifestation of both the liquid-solid and solid-solid transformations substantially depends on the cooling rate and alloy composition.Consequently, a variety of microstructures with  different microstructural features, e.g., in terms of phase morphology, grain size and texture, and, thus, different mechanical properties are achievable. [6,7]In the following, the phase transformations occurring during the cooling of γ-TiAl based alloys are discussed in detail.Pathways of microstructural evolution associated with the different phase transformations during solidification and subsequent cooling are presented in Figure 2. The schematic provides an overview of the potential liquid-solid as well as solid-solid phase transformations in binary γ-TiAl based alloys depending on their Al content.

Liquid-Solid Phase Transformations
][40] Even powder metallurgical processing techniques utilizing only solid-solid phase transformations, such as spark plasma sintering, rely on liquid-solid phase transformations during the atomization of molten feedstock material into powder. [38,41]For the relevant Al range from 42 to 48 at%, the section of the binary Ti-Al phase diagram in Figure 1a reveals the occurrence of three different transformations involving the liquid phase: precipitation of β from the liquid, precipitation of α from the liquid and L þ β !α.Once the material enters the β þ L phase field region, the first solidifying phase for Al contents below 49 at% is the β phase. [29]The β dendrites grow with the 〈100〉 directions parallel to the heat flow direction, which is generally true for cubic phases. [6,42,43]While the primary solidification via the β þ L phase field region is typically observed in connection with conventional casting methods, thermodynamic calculations suggest that this may change in favor of the α phase for Al contents above 46 at% under very high cooling rates. [44]Such a behavior was indeed observed in a binary Ti-48Al (at%) alloy for the high cooling rates of additive manufacturing, i.e., in the magnitude of %10 3 °C s À1 . [45]owever, Graf et al. [46] reported for a Ti-48Al-2Cr-2Nb (at%) alloy that the β phase is the first solid phase during an adapted selective laser melting process for a similar cooling rate, suggesting a substantial influence of additional alloying elements as well as the interplay of temperature gradient and solidification rate.
Figure 2 shows that, depending on the exact Al content of the alloy, three different solidification pathways are possible after the precipitation of β phase from the liquid.For Al contents below 44 to 45 at%, the solidification occurs entirely via the β phase, e.g., see route I in Figure 2. The other two pathways, corresponding to routes II and III in Figure 2, are characterized by the peritectic formation of α phase in accordance to the L þ β !α transformation, which occurs for higher Al contents up to 49 at%. [27,29]oute II and III differ in the subsequent phase transformation, i.e., β !α on the Al-lean side of the peritectic point and the precipitation of α from the remaining liquid phase on the Al-rich side.Accordingly, γ-TiAl based alloys are divided into β-solidifying and peritectically solidifying alloys in literature, because the solidification pathway has a significant influence on the appearance and homogeneity of the resulting microstructure.For example, a comparison between the solidification structures of a binary Ti-45Al and Ti-48Al (at%) alloy is given in Figure 3.While the β-solidification results in an equiaxed microstructure with only small microsegregations (Figure 3a), the peritectic solidification yields extensive macrosegregations as well as a coarse columnar grain structure (Figure 3b). [6,47]Although the orientation relationship between the β and the α phase (see Section 2.2.1) gives rise to 12 possible α variants, strong temperature gradients allow only one α orientation per β dendrite during the peritectic transformation, thus, causing these coarse columnar grains and sharp solidification textures. [16,43]Contrarily, β-solidification results in an isotropic and texture-free material due to the subsequent solid-solid transformation between the β and α phase, allowing all 12 orientation variants to occur. [47]he third liquid-solid phase transformation observable during the solidification of γ-TiAl based alloys is the precipitation of the α phase from the liquid for Al contents above the peritectic point marked with "P" in Figure 2.For sufficient amounts of Figure 2. Schematic representation of the microstructural evolution during solidification and subsequent cooling to room temperature for three compositions with differing Al contents: I) β-solidification, II) peritectic solidification followed by cooling through the α þ β phase field region, III) peritectic solidification followed by cooling through the α þ L phase field region.The orange point labeled P marks the peritectic point.Note that the microstructure formation during processing may differ from the idealized one, for example due to chemical segregations associated with solidification or the applied cooling rate (see text).
remaining liquid phase after the peritectic transformation, α dendrites associated with this transformation are detectable in the solidification microstructure. [29]These dendrites grow depending on the cooling rate either with their [0001] or their 1120 direction parallel to the heat flow. [6,42,48]Ultimately, it is worth mentioning that the primary solidifying phase is the α phase for Al contents above 49 at%, which can participate in a second peritectic reaction with the liquid into γ phase.This transformation typically occurs for Al contents above the composition range of γ-TiAl based alloys, i.e., above 50 at%, but may arise at lower Al contents due to pronounced segregations. [29]2.Solid-Solid Phase Transformations

Transformation β ! α
As shown in Figure 1, the α þ β phase field region, in which the β !α transformation occurs upon cooling, is situated at the highest temperatures when compared to the other solid-solid phase transformations in intermetallic γ-TiAl based alloys.However, it is only present for Al contents below the peritectic composition of 47 at%.Because the β !α transformation determines the grain size of the α phase, it is of utmost importance for the mechanical properties of the material. [6]Generally, this transformation bears many similarities with the β !α transformation observed in Ti alloys.Especially, the same orientation relationship, the so-called Burgers orientation relationship as presented in Equation (1), holds true for both transformations from the body-centered cubic to the hexagonal crystal structure [49] ð0001Þ α k f110g β and h1120i α kh 111i β (1) A graphical representation of this orientation relationship is given in Figure 4a.Essentially, the parallel arrangement of the most closely packed lattice planes of the β and the α phase, the {110} β and the (0001) α planes, as well as of the closed-packed directions, the 111 h i β and 1120 α direction, gives rise to 12 independent orientation variants.
As shown in Figure 4b, the β !α transformation can manifest itself in three different ways depending on the actual cooling rate and exact alloy composition: 1) diffusion-controlled nucleation and growth of α grains, 2) composition-invariant massive transformation, or 3) diffusion-less martensitic transformation.Slow cooling from the β single-phase field region results in the precipitation and diffusion-controlled growth of α grains obeying the Burgers orientation relationship (Equation (1)), which yields a refinement of the final microstructure when compared to peritectically solidifying alloys (Figure 3). [6,7,50]While an increase of the cooling rate to a moderate level does not affect the underlying mechanism of nucleation and growth, the microstructural appearance of the formed α grains is shifted toward Widmannstätten plates, which are expected to exhibit higher transformation kinetics when compared to the growth of equiaxed α grains. [51]A corresponding microstructure of a Ti-44Al-3Mo-0.1B (at%) alloy water-quenched from the β single-phase field region, see Figure 4c, consists of colonies of several such Widmannstätten plates.Similar microstructures are often observed in αand (α þ β)-Ti alloys under the name basket-weave microstructure. [52]These Widmannstätten colonies are found to nucleate at boundaries between β grains, at grain boundary allotriomorphs (essentially heterogeneously formed α phase), or intragranularily.They grow via a diffusion ledge mechanism in accordance to the Burgers orientation relationship with their long axis roughly in the 335 h i β direction. [6,52,53]Each α plate is often separated by retained β phase, which enriches in Ti and β-stabilizing elements in the course of this transformation.
Increasing the cooling rate to levels at which the massive β !α phase transformation occurs results in a significantly different microstructure. [6]This becomes apparent when comparing the Widmannstätten microstructure in Figure 4c with the massive one in Figure 4d, belonging to an alloy whose composition promotes the massive transformation under water quenching.[56] A complete massive transformation was observed in a Ti-42Al-5 V (at%) alloy after water quenching, [54] while an incomplete massive transformation was reported in gas-atomized powder of a Ti-43.5Al-4Nb-1Mo-0.1B(at%) alloy due to elemental segregation during the atomization process. [57][61] The nuclei of these massive α grains form with the same Adapted with permission. [47]2004, Wiley.composition as the parent β phase and subsequently grow with rates several orders of magnitude higher than the diffusioncontrolled transformation described above, e.g., 10 4 μm s À1 . [6,56]hese fast-growth kinetics are achieved by the migration of grain boundaries by short-range diffusion and jumping of atoms across the interface. [6,58]While the massive α grains have been found to disobey the Burgers orientation relationship with their related parent β grains, [57,58] Hu and Jiang [58] found a geometric correlation between neighboring massive α grains.In particular, some of the massive grains originating from the same parent β grain shared a 60°rotation relationship around a common a axis, which is a configuration that allows a reduction of elastic energy concentration. [58]t the highest cooling rates, the β !α transformation manifests itself in the form of a martensitic transformation. [6]The formation of hexagonal martensitic laths or plates, characterized by a high dislocation density, is achieved by a shear deformation of the body-centered cubic crystal via the β-½111ð112Þ and β-½111ð101Þ shear systems.Contrarily to the massive transformation, the martensitic grains obey the Burgers orientation relationship with respect to their parent β grain. [52]An example for a fully martensitic microstructure is presented in Figure 4e.][64][65] The α 00 martensite, which is mainly observed in β-Ti alloys, [65][66][67][68] is not within the scope of this review and the interested reader is referred to the cited references.In the case of binary γ-TiAl based alloys, the martensitic transformation from β to α has only been observed under the extremely high cooling rates achievable during the solidification of fine powder particles via gas atomization. [69,70]For the martensitic transformation to occur in bulk material of binary γ-TiAl based alloys, even ice-brine quenching was found to be insufficient and a massive transformation was observed instead. [50]In literature, this behavior is attributed to the high T 0 temperature of the β !α transformation. [50]owever, similar to steels, the critical cooling rate to obtain martensite can be lowered by additional, often slow-diffusing, alloying elements, e.g., Nb, Mo, V, Mn. [54,57,58,63,64] The formation of disordered α 0 martensite occurred in a gas-atomized TNM powder as well as in an ice-brine quenched Ti-44Al-4Nb-4Hf-0.1Si (at%) alloy in the previous studies, [57,58] respectively.The formation of ordered α 2 0 martensite via the combination of a local diffusion-controlled and a purely displacive transformation was evidenced by Mayer et al. [62] in a Mo-bearing γ-TiAl based alloy.Furthermore, this type of martensite has also been observed in V-and Mn-containing alloys. [63,64]Takeyama and Kobayashi [54] showed that the martensitic transformation is favored over the massive one by either increasing the amount of additional alloying elements, e.g., compare the microstructures for two different V contents in Figure 4, or by reducing the Al content.

Transformation α ! γ
The second solid-solid phase transformation of particular importance for intermetallic γ-TiAl based alloys during cooling is the α !γ transition.It occurs once the temperature falls below and (e) are adapted with permission. [54]005, Elsevier.
the γ-solvus temperature, e.g., see green line in Figure 1b, which is generally situated at lower temperatures than the β !α transformation.The crystallographic orientation relationship between the hexagonal α/α 2 phase and the tetragonal γ phase according to Blackburn is presented in Equation ( 2). [71]Essentially, this orientation relationship describes a parallel arrangement of the most closely packed lattice planes and directions of both these structures, as illustrated in Figure 5a, in a similar manner to the Burgers orientation relationship.
The α !γ phase transformation yields different microstructural features depending on the applied cooling rate and, thus, can substantially influences the mechanical properties of the material. [6,7]A schematic continuous time-temperaturetransformation diagram for the cooling from the α single-phase field region as well as micrographs of the respective microstructures are presented in Figure 5b.Going from low to high cooling rates, a lamellar γ formation (Lam.), a massive γ formation (γ m ) or even a suppression of the γ precipitation (only α 2 ) may occur.These different formation types of γ phase are described in the following.It is worth mentioning that the precipitation of globular γ grains during very slow cooling has been observed in literature, e.g., see the work by Appel et al. [6] and the references therein.However, globular γ grains are more frequently observed as a product of isothermal heat treatments within the α þ γ phase field region rather than as a result of cooling. [7]Furthermore, the α !γ transformation is characterized by a high undercooling of 50-100 °C even for relatively low cooling rates, which further depends on the alloy composition and the α grain size, [72] suggesting a high nucleation barrier for this transformation.
For low cooling rates, such as furnace cooling, the α !γ transformation results in the formation of γ lamellae inside the α grains in accordance to the Blackburn orientation relationship, as shown in Figure 5b. [6,7,73,74]In the microstructure, one such grain with lamellae is known as a lamellar colony.A microstructure consisting solely of such lamellar colonies is known as a fully lamellar microstructure in literature. [6,7]While such a microstructure can be adjusted by cooling from the α single-phase field region, also other microstructural types are commonly observed.In particular, heat treatments within the α þ γ phase field region result in microstructures containing lamellar colonies, which form during cooling, and globular single-phase γ grains, which form during dwelling.These microstructures, organized in accordance with their amount of globular γ grains (γ globular ) with respect to the fraction of lamellar colonies, are called nearly lamellar (γ globular < α 2 /γ colonies), duplex (γ globular % α 2 /γ colonies), and globular (γ globular > α 2 /γ colonies). [6,7]he nucleation of the individual γ lamellae within the α 2 /γ colonies happens heterogeneously at α grain boundaries or grain boundary allotriomorphs.[77] The particular arrangement of these stacking faults, consisting of two Shockley partial dislocations, changes the ABABAB stacking sequence of the hexagonal structure into the ABCABC sequence of the face-centered cubic lattice. [75]The following change of the composition via diffusion and ordering of this face-centered cubic structure into the L1 0 structure finally results in the nucleation of the γ lamella. [71,77,78]Depending on the cooling rate, subsequent growth of the γ phase occurs either by the motion of partial dislocations, by twinning or by continuous nucleation of new lamellae. [73,79]In most cases, all lamellae within one individual colony are arranged in parallel, which can be explained by the  [7] 2013, Wiley; c) Gibbs free energy considerations for the massive α !γ transformation upon quenching from a temperature T 1 in the α single-phase field region to a temperature T 2 in the α þ γ phase field region for an alloy with an Al content x Al 0 .
Blackburn orientation relationship and the fact that the hexagonal crystal structure of the α phase only offers one set of basal planes, i.e., (0001) α planes. [6,71]Another consequence of the Blackburn orientation relationship with respect to the interface between adjacent γ lamellae is the occurrence of different interface types.In particular, the three possible arrangements of the (111) γ plane on the (0001) α plane, i.e., rotations of multiples of 60°around the [0001] α direction, yield a pseudo-twin boundary at 60°, an order domain related interface at 120°and a true twin boundary at 180°misorientation between neighboring γ lamellae. [80]][87][88] Widmannstätten colonies, which occur at lower cooling rates than feathery structures, consist of packets of α 2 and γ lamellae misoriented approximately 64°around the 1100 α 2 direction with respect to the α 2 phase of the surrounding α 2 /γ colony. [73,85]However, the lamellae inside the Widmannstätten colony are finer than their counterparts in the normal α 2 /γ colony, and the α 2 phase is present in lower amounts with a sparse, discontinuous and lens-shaped morphology. [86]Two different explanations exist on how these Widmannstätten colonies are formed based on the observed misorientation: 1) twinning of the α grain during cooling and subsequent lamellae formation; 2) direct nucleation at an already existing γ lamella orientated in accordance with the Blackburn orientation relationship (Equation ( 2)) with regard to the parent α grain. [73]Feathery structures appear more complex when compared to Widmannstätten colonies and consist of multiple small lamellar packets, which are in contact with each other but slightly misaligned.The resulting characteristic fanning-out morphology, similar to feathers, gives this microstructural feature its name.The feathery structures are misoriented with respect to the surrounding colony by about 15°, but the contained α 2 and γ lamellae are aligned in accordance with the Blackburn orientation relationship. [72]Furthermore, a continuous change of the orientation of the individual lamellar packets is found, which is attributed to the elastic stresses arising during their formation. [73]More precisely, these structures are thought to form via nucleation of γ lamellae during the growth of an already existing lamellar packet. [89]The typical cooling rates, at which Widmannstätten colonies and feathery structures can be obtained, are in the range of air cooling to water quenching, but also depend on the chemical composition of the material. [73,90]or example, Hu et al. [91] observed that alloying with Ru promotes the precipitation of a feathery microstructure due to a reduction of the interfacial energies of γ/γ boundaries.
Increasing the cooling rate when cooling down from the α single-phase field region, decreases the average width of the formed γ lamellae. [7,83]However, once a critical cooling rate, which depends on the alloy composition and α grain size, is exceeded, the γ phase does not precipitate in a lamellar manner, but rather via a massive transformation, as shown in Figure 5b. [6,7,73]This manifestation of the α !γ phase transformation is compositioninvariant and results in γ grains with an irregular morphology and a high density of internal defects and boundaries. [92]It is promoted by the use of weakly β-stabilizing and slow-diffusing elements that show an unequal partition behavior between the α and γ phase.For example, additions of the elements Nb, Ta, and Ru have been found to favor the formation of massive γ phase under suitable cooling rates. [91,93,94]By contrast, a high O content is able to suppress the massive α !γ transformation due to the α-stabilizing effect of this element. [95]Furthermore, additions of B also affect the occurrence of the massive transformation as a result of changes of the initial microstructure. [96,97]he composition invariance causes an Al depletion in the formed γ phase when compared to its ideal Ti:Al stoichiometry and, consequently, yields a decreased tetragonality as well as a c/a ratio close to unity. [31,98,99][102] This predominant nucleation at α grain boundaries results in a dependence of the amount of massive γ phase on the α grain size and, thus, also on the heat treatment time in the α single-phase field region.Especially, longer heat treatments and the associated α grain growth yield lower amounts of massive γ grains in the final microstructure.Interestingly, during the nucleation at an α grain boundary, the Blackburn orientation relationship is obeyed with one of the parent α grains, but the growth of the massive γ grain, which happens via short-range diffusion across the migrating incoherent interface, then continues into the other adjacent α grains. [103,104]Consequently, no simple orientation relationship between the massive γ grains and the surrounding α grain has been observed in literature. [92,105]owever, some research has shown that the massive γ grain may also grow into the parent α grain, but the Blackburn orientation relationship is destroyed due to successive twinning during growth. [106] representation of the massive formation of γ phase with respect to thermodynamics is given in Figure 5c, which presents the Gibbs free energy curves G α and G γ in dependence on the chemical composition x, i.e., the Al content, after quenching from the α single-phase field region (T 1 ) to a temperature inside the α þ γ phase field region (T 2 ).The massive transformation can only occur when the lamellar or Widmannstätten α !γ transformations are suppressed.Furthermore, a temperature must be reached upon cooling from the α single-phase field region at which G γ is lower than G α for the material's composition x Al 0 .Consequently, the driving force for the massive transformation ΔG massive is essentially proportional to the difference of the two Gibbs free energies G α and G γ , and increases even further for higher cooling rates and, thus, higher undercoolings ΔT. [93,107] The Gibbs free energy curves in Figure 5c also show that the massive transformation results in a thermodynamically unstable material condition as its Gibbs free energy is higher than the one of the mixture of α and γ phase with their respective equilibrium compositions x Al α;eq: and x Al γ;eq: .Technologically, this metastable material condition, which is only retainable at low temperatures due to a lack of thermal activation and diffusion, can be used to produce microstructures with interesting mechanical properties.[98,108] During annealing heat treatments subsequent to the massive γ formation, the material approaches thermodynamic equilibrium, i.e., a mixture of α 2 and γ phase.Consequently, the α 2 phase precipitates during such heat treatments in the former massive γ grains, which change their chemical composition due to diffusional processes.The precipitated α 2 phase obeys the Blackburn orientation relationship with respect to the γ phase and forms lamellar structures.[6] However, contrarily to the strictly parallel arrangement of γ lamellae forming during the α !γ transformation due to only one set of (0001) α planes, structures of crossed α 2 lamellae can be formed in the massive γ phase, since it possesses four different {111} γ planes.Such crossed lamellar microstructures have been found to exhibit high strength and creep resistance as well as good thermal stability.[98,105,108] Finally, even the massive formation of γ phase can be suppressed during cooling from the α single-phase field region once a critical cooling rate is exceeded.As shown in Figure 5b, this complete suppression of the γ phase formation results in a microstructure consisting only of supersaturated α 2 grains.It is worth mentioning that the ordering of α into α 2 even occurs for the high cooling rate of water quenching.[109] 2 The eutectoid point in the binary Ti-Al phase diagram, at which the α phase transforms into α 2 and γ phase, is situated at an Al content of 39 at% and a temperature of 1120 °C. [27]In contrast to eutectoid transformations in other material systems, e.g., steels, the α !α 2 þ γ transformation has never been observed as a discontinuous transformation, i.e., the growth of the two product phases in the parent phase. [6,110]This is attributed to the fact that the parent α phase is structurally closely related to the α 2 phase via an ordering transformation only requiring short-range diffusion and that the composition of the eutectoid point is very close to the composition of the resulting α 2 phase. [111,112]Especially, even in case of equilibrium conditions, only a slight excess of Al atoms has to diffuse to form the small amount of γ phase required.The formation of γ phase during this phase transformation has been found to rely on the same mechanisms as described above in Section 2.2.2.However, the sudden increase of the amount of γ phase, which should occur in the course of the eutectoid transformation in thermodynamic equilibrium, is suppressed to large extent for technical cooling rates. [6,110]With respect to the influence on the microstructure, this phase transformation only plays a minor role when compared to the β !α and α !γ phase transformations as highlighted in Figure 2.

Alloy Generations
Early investigations on binary TiAl alloys revealed that these alloys do not meet the required properties for applications at high temperatures, such as sufficient creep strength and oxidation resistance. [7]Consequently, a vast amount of research effort was put into the investigation of the impact of different alloying elements on the properties of γ-TiAl based alloys. [7,28,30]his ultimately led to the definition of particular TiAl alloy generations.Therein, alloys are categorized regarding their development time and their specific composition. [8]For information about the 1 st alloy generation, the reader is referred to the work by Kim. [30]In general, alloys of the 2 nd generation, such as the GE-4822 alloy, have a chemical composition (in at%) that can be described as follows Ti-ð45-48ÞAl-ð1-3ÞX-ð2-5ÞY-ð< 1ÞZ (3) where X = Cr, Mn or V, Y = Nb, Ta, W or Mo, and Z = Si, B or C. Broadening the application area of γ-TiAl based alloys even further required an increase in service temperature and, therefore, in creep resistance.The conducted research work led to the definition of the 3 rd alloy generation, which includes, for example, the TNB alloy family.The chemical constitution of 3 rd generation alloys can be summarized as follows [7] Ti-ð42-48ÞAl-ð0-10ÞX-ð0-3ÞY-ð0-1ÞZ-ð0-0:5ÞRE where X = Cr, Mn, Nb or Ta, Y = Mo, W, Hf or Zr, and Z = C, B or Si; RE = rare-earth metals.
The so-called process-adapted γ-TiAl alloys are a recent development and correspond to the 4 th alloy generation. [8]While their alloying concept is similar to 3 rd generation alloys, the chemical compositions of these latest generation alloys are adapted to optimize the material for a specific manufacturing route, e.g., nearconventional forging in the case of the TNM alloy [7] or additive manufacturing. [113]For more details on the history and development of γ-TiAl based alloys, the reader is referred to the work by Duan et al. [25]

β-Stabilizing Elements
The success of 3 rd generation intermetallic γ-TiAl based alloys can be attributed to a large extent to the incorporation of β-stabilizing elements.Among the multitude of different β-stabilizing elements, Nb, Mo, and W have experienced extensive use in several different alloying concepts.However, also other β-stabilizing elements, such as Ta, Mn, V, Cr, Fe, Zr, and Hf, have attracted some research interest in the past.
The effect of β-stabilizing elements on the Ti-Al phase diagram is exemplified in Figure 6a for the TNM alloy system, i.e., Ti-xAl-4Nb-1Mo-0.1B(at%).Compared to the binary Ti-Al phase diagram, an expansion of the β single-phase field region to higher Al contents, accompanied by a reduction of the α single-phase field region, occurs in this isoplethal section upon alloying with the given amounts of β-stabilizing elements Nb and Mo.Especially, while the α phase field region is present over the whole Al range of binary γ-TiAl based alloys, it only occurs above 44 at% in the TNM-based system.Simultaneously, β-stabilizing elements can heavily impact the solidification behavior of the material by shifting the region of the peritectic transformation toward higher Al contents, thus, promoting a β-solidification, e.g., see Section 2.1.However, certain β-stabilizing elements, e.g., W, are prone to pronounced segregations upon solidification, because of their low diffusivity and distinct partitioning behavior. [6,114]he most prominent feature of the system shown in Figure 6a and generally most systems with sufficient amounts of β-stabilizing elements, though, is the presence of an ordered B2 phase at low temperatures when compared to the binary Ti-Al system.This so-called β o phase, whose crystal structure is depicted in Figure 6b, corresponds to the ordered counterpart of the high-temperature β phase. [6][117] One example of such a complex microstructure is shown in Figure 6c, which presents the microstructure of a heat-treated TNM alloy consisting of lamellar α 2 /γ colonies, β o phase, and globular γ grains.
A crucial aspect that needs to be considered with respect to the incorporation of β-stabilizing elements is that the stability of the β/β o phase over a large temperature range bears both advantages and disadvantages.The presence of small amounts of the soft and ductile β phase at forging temperatures in the α þ β phase field region allows the forging of certain γ-TiAl based alloys using conventional forging equipment. [17,18]This stands in contrast to most other γ-TiAl based alloys, which can only be forged isothermally by encapsulation of the material, culminating in a more complex and expensive processing. [6,7]However, the presence of β o phase at service temperature has been proven detrimental for the creep resistance, especially if present in large amounts. [19,116]This can be attributed to its less densely packed B2 crystal structure resulting in a higher diffusivity compared to the hexagonal and tetragonal lattices of the α 2 and the γ phase, respectively.Furthermore, literature suggests that the β o phase exhibits a negative impact on high-temperature strength and room temperature ductility. [6,34,118]In the case of the TNM alloy, a special two-step heat treatment has been designed to reduce the amount of β o phase in the final microstructure to a level innocuous for the creep performance. [20,119]Especially, the first heat treatment step exploits certain characteristics of the phase fraction evolution of the TNM alloy, i.e., a temperature range above 1200 °C in which the amount of β phase exhibits a local minimum and thermal activation is still high.For more details see the study by Schwaighofer et al. [19] Additions of β-stabilizing elements bring several beneficial consequences with respect to the mechanical properties.These elements generally increase the material's high-temperature strength due to solid solution hardening of the individual phases.[122] However, both of these mechanisms require a certain amount of the respective alloying element to be uniformly distributed within the phases, which significantly depends on the element's partitioning behavior. [41]The partitioning behavior as well as the preferred sublattice in the α 2 , β o and γ phase of the alloying elements in the focus of this review are presented in Table 3.It shows that almost all of the discussed β-stabilizing elements tend to occupy the Ti sublattice in the α 2 and γ phase, while the Al sublattice is preferred in the β o phase.The other elements either occupy solely the Al sublattice, as in the case of Si, or are dissolved interstitially, as in the case of C, O, N, and B.
Regarding the important β-stabilizing elements Nb, Mo and W, the latter exhibits the most pronounced partitioning behavior.In particular, W favors the β phase, followed by the α phase and at last by the γ phase at high temperatures.The partitioning of Mo behaves essentially the same, but is less pronounced when compared to W. While Nb also prefers the β phase over the other two phases, albeit to a lesser extent than Mo and W, it is almost equally distributed between the α and γ phase. [33]In general, the solubility of Nb in γ-TiAl based alloys is significantly higher than that of Mo and W. For example, while the α þ γ phase field region extends up to approximately 9 at% of Nb at 1200 °C, it only reaches 1 to 1.5 at% in the case of Mo and W. [33] Concerning the other β-stabilizing elements shown in Table 3, their partitioning behavior between the major TiAl-related phases is, in many cases, qualitatively similar to the one of Mo and W, i.e., Based on the data reported in the study by Schwaighofer et al. [298] ; b) B2 crystal structure of the ordered β o -TiAl phase.Red atoms correspond to Ti, blue ones to Al; c) Microstructure of a heat-treated TNM alloy showing a complex multi-phase microstructure consisting of α 2 , β o and γ phase.Adapted with permission. [255]Copyright 2022, Elsevier.
β > α > γ or β o > α 2 > γ in the case of V and Hf.One exception is the element Zr, which favors the γ phase compared to the other two phases.Similar to W, the elements Cr and Fe show a reversed partitioning behavior between the α 2 and γ phase at low temperatures and the α and the γ phase at higher temperatures.Especially, while the α phase is preferred over the γ phase, the α 2 phase is less favored compared to the γ phase.
[125][126] In particular, the addition of W in the IRIS alloy has been found to yield the highest improvement of the oxidation resistance when compared to the Nb-containing GE-4822 alloy and the Mo-and Nb-containing TNM alloy during static as well as cyclic oxidation conditions at 900 °C.This was attributed to the formation of a compact continuous alumina layer, which is enhanced by W in the IRIS alloy and which is not so dominant in the other two alloys. [126][130] In the following Section 3.2.1-3.2.10, the available research work on the β-stabilizing elements Nb, Mo, W, Ta, Mn, V, Cr, Fe, Zr, and Hf is reviewed individually and discussed with respect to the influence on the Ti-Al phase diagram, associated microstructural changes and the possible formation of additional intermetallic phases.
of the Ti-Al-Nb ternary system in the composition and temperature range relevant to γ-TiAl based alloys.[142][143] These phases are described in more detail in the following two paragraphs, but, as can be seen in Figure 7, they are only thermodynamically stable at lower temperatures and dissolve upon heating of the material.Concerning the solid-solid phase transformations inherent to the Ti-Al system, Chladil et al. [133] found that Nb significantly increases the eutectoid temperature, while the γ-solvus temperature remains mostly unaffected in TNB-based alloys in the investigated composition range of Ti-45Al-(5-7.5)Nb(at%).The same behavior was also observed for a higher Al content of 46 at%. [144]Furthermore, Nb moves the α-transus line to higher Al contents.With respect to the liquidus and the solidus temperature, this element has been found to shift the phase transitions involving the liquid phase to higher temperatures. [145,146]he ω o phase is generally observed in Nb-containing γ-TiAl alloys after cooling from any β/β o -containing phase field region with moderate or lower cooling rates or after isothermal aging below its dissolution temperature. [115,139,147,148]The crystallographic data of this phase are given in Table 2.In accordance with the original article by Bendersky et al., [148] a recent phase diagram study by Distl et al. [137] suggested that a (Ti,Nb) 2 Al stoichiometry may be more applicable than the Ti 4 Al 3 Nb designation often found in literature.Generally, this phase is found inside the β o phase due to the underlying β o !ω o transformation in γ-TiAl based alloys, but its formation within lamellar colonies via a direct α 2 !ω o transformation has also been reported in literature. [20,115,119,149]The β o !ω o transformation involves two individual diffusion-controlled steps. [139]It is possible to suppress the second transformation step by applying a sufficiently high cooling rate, which results in the occurrence of the ω 00 phase, i.e., the precursor variant of the ω o phase. [115,139]As mentioned above, the ω o phase is thermodynamically unstable at higher temperatures.In particular, dissolution temperatures of 790 and 870 °C for the ω 00 and the ω o phase, respectively, were reported for a quenched Ti-45Al-10Nb (at%) alloy during heating with 20 °C min À1 [139] as well as of 830-880 °C for the ω o phase in Ti-(34-40)Al-10Nb (at%) alloys. [150]In the case of the TNM alloy, the ω 00 and the ω o phase have also been detected in the microstructure after suitable heat treatments, but they dissolved at 775 and 825 °C upon heating. [115,151]Finally, although the ω o phase is able to form nanometer-sized precipitates within the β o phase via aforementioned transformation and, thus, can significantly increase the hardness of this phase, [115] it has been found to be detrimental to the overall mechanical properties of γ-TiAl based alloys. [146,152]nother ternary intermetallic phase relevant to γ-TiAl based alloys is the orthorhombic O phase with an ideal stoichiometry of Ti 2 AlNb. [6]While research work in the past sometimes distinguished between two different types of O phase, i.e., O1 and O2, which differ in their sublattice occupations, a recent work by Xu et al. [153] has shown that only one of these variants actually  Based on the data presented in the studies by Distl et al. [137,138] corresponds to the O phase, namely O2, from a crystallographic point of view.[156][157] Sadi et al. [157] observed three different phase transformations involving the O phase upon cooling from 1200 °C in alloys close to its ideal stoichiometry: 1) Additionally, the β o !O transformation changed to a massive-type transformation for high cooling rates. [157]The peritectoid transformation α 2 þ β o !O was also observed by Muraleedharan et al. [155] Furthermore, the occurrence of an Al-rich and Nb-poor orthorhombic phase at low temperatures has also been observed in γ-TiAl based alloys with compositions of engineering relevance, e.g., TNB alloys. [117,143,158]In particular, the Nb content threshold for the occurrence of this phase is situated around 5-7.5 at%, while the Al content must not exceed 46-47 at% at the same time. [143]][160] Finally, it is notable that the presence of this orthorhombic phase is not tied to the Ti-Al-Nb system, as demonstrated in a recent study on several different γ-TiAl based alloys by Rackel et al. [143]

Molybdenum
An alloying concept that relies on additions of Mo is the TNM alloying concept with a nominal chemical composition of Ti-43.5Al-4Nb-1Mo-0.1B(at%). [17,21]Generally, the β-stabilizing effect of Mo is considered four times stronger with respect to the amount of stabilized β phase when compared to alloying with Nb, [7,15,161] which gives a Mo-equivalent of approximately 2 at% for the TNM alloy.As shown in Figure 8a, the strong β-stabilizing effect of Mo manifests itself in a decreasing β-transus temperature as well as a pronounced reduction of the α single-phase field region, which disappears completely for Mo contents above 2 at%.Furthermore, the increase of the eutectoid temperature for low Mo contents indicates a stabilization of the ordered α 2 phase against its disordered counterpart, the α phase, upon alloying with Mo.Concerning the γ-solvus temperature, only a slight variation as a function of the Mo content is observable in Figure 8a, although the involved phase field regions change.Consequently, the described influence of Mo on the eutectoid and γ-solvus temperature is similar to the one of Nb mentioned in Section 3.2.1. [23,162]The disordering of the β o phase into the β phase upon heating is shifted to higher temperatures by Mo additions, hence, indicating a more pronounced stabilization of the former against the latter phase. [112]This can also be seen from the blue triangles in Figure 8a corresponding to the disordering temperature of the β o phase in alloys containing 3 and 7 at% Mo, respectively, as determined by in situ neutron diffraction.Contrarily, the green triangle in Figure 8a, i.e., the disordering temperature for 5 at% Mo, was determined by in situ high-energy X-ray diffraction, which is less sensitive with respect to ordering phenomena in Ti-Al-related phases.Consequently, the superstructure peaks associated with the ordered crystal structure of the β o phase become undetectable at lower temperatures during heating experiments and the disordering temperature appears to be lower compared to the ones measured by neutron diffraction. [112,163]][166][167][168][169] For instance, Singh and Banerjee [165,166] investigated the Ti-(44-50)Al-(2, 4, 6)Mo (at%) alloy system with respect to its microstructural evolution during solidification and subsequent heat treatments.In addition to the phase transformations inherent to the Ti-Al system, they observed several other transitions depending on the Mo content, especially involving the β phase, such as a eutectoid α !β þ γ, a peritectoid α þ β !γ and a β !β þ γ transformation, resulting in particular microstructural features. [165,166]An α !β þ γ reaction was also observed by Musi et al. [163] in an as-cast Ti-44Al-5Mo (at%) alloy.This reaction resulted in coral-like structures, which occurred in the former α grains next to lamellar structures arising from the α !γ transformation, as demonstrated in Figure 8b.Additionally, indications of a direct precipitation of the γ phase from the β phase during cooling were also found in the microstructure. [163]In the case of a β-homogenized and waterquenched Ti-44Al-7Mo (at%) alloy, Erdely et al. [169] traced such a β o !γ transformation from the metastable to the stable state heating as determined by neutron diffraction (blue, 3 and 7 at% Mo) and high-energy X-ray diffraction (green, 5 at% Mo), respectively.Adapted with permission. [163]Copyright 2022, Elsevier; b) Microstructure of an as-cast Ti-44Al-5Mo (at%) alloy, investigated in ref. [163], showing features of different transformations: The contrasts of the individual phases are dark gray (γ), gray (α 2 ), and light gray (β o ).
upon heating from room temperature to 900 °C without signs of the occurrence of any intermediate phase via in situ diffraction and scattering techniques.Small-angle X-ray scattering combined with atom probe tomography allowed to study the early stages of the formation of disc-shaped, partially coherent precipitates in the β o parent phase, while the results based on high-energy X-ray diffraction confirmed the Kurdjumov-Sachs orientation relationship between β o and γ. [169] With respect to ω o -related intermetallic phases, which are frequently observed in Nb-containing γ-TiAl based alloys such as TNB and TNM under the correct circumstances, e.g., see Section 3.2.1,Mo is characterized by a destabilizing effect and, consequently, is able to completely suppress their formation upon sufficient alloying. [151]Furthermore, as seen in Figure 8a, no other phases than the Ti-Al-related phases and their derivatives occur in thermodynamic equilibrium in the Mo range feasible for γ-TiAl based alloys.However, starting from a nonequilibrium material condition, i.e., after water-quenching from the β single-phase field region, the occurrence of an orthorhombic phase upon reheating was observed in a Ti-44Al-3Mo-0.1B (at%) alloy. [141,170]Schmoelzer et al. [141] described this orthorhombic phase as the B19 phase, whose crystallographic data are given in Table 2.It has to be noted that the differentiation between the B19 phase and the O phase is generally not a trivial task, due to their very similar crystal structures.In particular, a representation of both phases in terms of their common space group Cmcm, e.g., see ref. [117], shows that they actually only differ in the occupation of certain sublattices. [117,170]The orthorhombic phase observed in the Ti-44Al-3Mo-0.1B(at%) alloy was suggested to act as a metastable transition phase between the α 2 and the γ phase upon the transition of the nonequilibrium material condition toward thermodynamic equilibrium.In the temperature range from 600 to 720 °C it formed modulated structures within the α 2 phase and the α 2 0 martensite. [170]It is worth mentioning that other authors have also observed this B19 phase in other intermetallic γ-TiAl based alloys, e.g., see refs.[171-174].

Tungsten
Tungsten is utilized in the ABB-2 alloy (Ti-46Al-2 W-0.5Si, in at%) and in the IRIS alloy (Ti-48Al-2 W-0.1B, in at%), both containing 2 at% of this element in their nominal chemical composition. [41,175]Compared to Mo, W acts as an even more potent β-stabilizing element with respect to the amount of β phase stabilized in γ-TiAl based alloys. [15,33]However, the Ti-Al-W system is only scarcely treated in literature when compared to the information available on the influence of Nb and Mo on the Ti-Al phase diagram.Consequently, open questions still remain regarding the influence of the alloying element W on the transition temperatures in γ-TiAl based alloys.A study by Kainuma et al. [33] clearly predicted that for γ-TiAl based alloys containing 2 at% W only the β o phase and the γ phase should be thermodynamically stable for Al contents above 44 at% and temperatures below 1200 °C. [176]However, the α 2 phase has been observed in both the ABB-2 alloy and IRIS alloy with 46 and 48 at% Al, respectively, after long-term creep exposure and heat treatments. [175,177]Especially, α 2 phase was still present in the microstructure of an IRIS alloy heat-treated at 750 °C for 8200 h. [178]Similarly, the microstructure of a Ti-46Al-2 W-0.2B (at%) alloy was found to consist of α 2 , β o, and γ phase at room temperature after an in situ high-energy X-ray diffraction experiment adapted to reproduce a spark plasma sintering (SPS) densification. [179,180]Furthermore, the conducted in situ experiment granted insights into the equilibrium phase transformation pathway during cooling from 1325 °C, i.e., [179] Especially, no α single-phase field region was present in this Ti-46Al-2 W-0.2B (at%) alloy, despite the relatively high content of α-stabilizing Al.Interestingly, the isoplethal section presented in Figure 6 predicts an extensive α single-phase field region for the same amount of Al and a similar amount of β-stabilizing elements Mo and Nb, i.e., a Mo-equivalent of 2 at% compared to the 2 at% W. Thus, further evidence for the strong β-stabilizing effect of W in comparison to Mo is given. [19,179]

Tantalum
A 4 th generation γ-TiAl based alloy, which relies on the additions of only Ta and which has generated some interest in the past, is the Ti-46Al-8Ta (at%) alloy. [181,182]This alloy has been in the focus of research due to the possibility of microstructural refinement via the so-called "massive transformation technique". [183]s mentioned in Section 2.2.2, Ta is able to lower the cooling rate necessary for the formation of massive γ phase during the α !γ transformation.In particular, the incorporated 8 at% of Ta are sufficient to yield an almost completely massively transformed microstructure after air cooling, while for the same amount of Nb only 5% of the microstructure had transformed into massive γ phase under the same cooling conditions. [181]uring a subsequent annealing treatment, the α 2 phase precipitated from the massive γ phase, resulting in a nearly or fully convoluted microstructure depending on the fraction of initial massive γ phase. [181,183,184]Such a convoluted microstructure, which is presented in Figure 9a, is characterized by the presence of α 2 lamellae in the former massive γ grains.However, due to this γ massive !α 2 transformation, which occurs above 980 °C in the Ti-46Al-8Ta (at%) alloy upon heating at 5 °C min À1 , [184] the Blackburn orientation relationship (see Equation ( 2)) allows four different lamella orientations instead of only one orientation as in the case of the α !γ transformation.
An isopleth of the Ti-Al-Ta ternary system containing aforementioned Ti-46Al-8Ta (at%) alloy is shown in Figure 9b.Notable is the presence of the ternary Ta-rich τ phase (space group P6 3 /mmc) at temperatures below 900 °C, which is structurally identical to the hexagonal ω o -Ti 4 Al 3 Nb phase reported by Witusiewicz et al. [145] in the Ti-Al-Nb system, as well as the predicted absence of α 2 above 42 at% Al at low temperatures.This τ phase was indeed observed by Lapin et al. [182,185,186] at grain and lamellae boundaries in the Ti-46Al-8Ta (at%) alloy after longterm aging treatments and creep experiments in the temperature range from 700 to 800 °C.More specifically, the τ phase formed via the decomposition of α 2 laths into α 2 /γ colonies. [186]oncerning the predicted γ þ τ phase field region at low temperatures, these long-term annealing treatments showed that α 2 phase is still present in the microstructure after 10 000 h at 750 °C, thus contradicting the presence of this phase field region. [182]Furthermore, for even higher contents of Ta, the presence of an orthorhombic O phase similar to the one observed in Nb-containing alloys has been reported in literature. [187]ith respect to the phase transformation pathway at high temperatures, i.e., melting and solidification, Shuleshova et al. [188] conducted in situ high-energy X-ray diffraction experiments on a Ti-45Al-7Ta (at%) alloy.During solidification and cooling to approximately 1050 °C, they observed the phase transformation pathway which reversed during the following reheating and melting of the alloy.While this is generally in good agreement with the isoplethal section shown in Figure 9b, slight discrepancies observable at the lowest temperatures within the investigated range, e.g., the absence of the α 2 þ β o þ γ phase field region, may be attributed to an insufficient undercooling to provoke the corresponding phase transformations.

Manganese
Manganese, besides other alloying elements such as Mo and Nb, is one of the β-stabilizing elements which have already been demonstrated to positively influence the hot workability of γ-TiAl based alloys, e.g., the Ti-42Al-5Mn (at%) alloy. [18,33] illustrated in Figure 10a as well as in the isothermal sections shown in the study by Huang et al., [189] Mn extends the β singlephase field region to higher Al contents.The addition of this element also stabilizes the β phase at lower temperatures and shifts the α single-phase field to higher Al contents compared to the binary Ti-Al phase diagram. [18,27,33][191][192] In the case of the Ti-42Al-5Mn (at%) alloy, the β/β o phase is stabilized from about 900 °C up to the liquidus temperature. [189,193]This circumstance permits conventional hot forging, since the presence of the cubic β phase provides enough independent slip systems for dislocations in the same manner as in the TNM alloy system. [18,194]Consequently, the use of expensive equipment for isothermal forging is not required.As Mn is also a less costly alloying element than, for example, Mo or Nb, the Ti-42Al-5Mn (at%) alloy has been in the focus of recent research as an inexpensive alternative to the established β-stabilized γ-TiAl based alloys such as TNM and TNB.However, in this Ti-42Al-5Mn (at%) alloy, the brittle Ti(Al,Mn) 2 .Adapted with permission. [184]2010, Wiley; b) Calculated isopleth of the Ti-Al-Ta system for 8 at% Ta.The red line marks the composition of the Ti-46Al-8Ta (at%) alloy.The black dashed line indicates the disordering temperature of the β o phase.The letter τ marks the existence range of a ternary Ta-rich phase (see text).Based on the data reported in ref. [182].
(a) (b) Figure 10.a) Ti-rich part of the isothermal section of the Ti-Al-Mn ternary system at 800 °C.The red point indicates the composition of the Ti-42Al-5Mn (at%) alloy.Based on the data reported in the work by Huang et al. [189] ; b) Isopleth of the Ti-Al-V system for a constant Al content of 45 at%.Based on the data given in the study by Takeyama and Kobayashi. [54]hase has been reported to occur below 900 °C. [189,195,196]This phase is a hexagonal C14 Laves phase, which causes a significant reduction of the room temperature ductility and strength. [197]uring annealing at 800 °C, it was found to form already after 15 min inside the β o phase as well as within the lamellar colonies. [195,196]Since the stability range of the Laves phase is especially important for applications at elevated temperatures, efforts have been made to prevent its formation. [198]Tang et al. [198] reported that alloying with only 0.5 at% Mo or W can suppress the formation of the brittle Laves phase, with the effect of Mo being stronger than the one of W. [198] Furthermore, Mn has been found to act as a ω o -destabilizing element in Nb-containing γ-TiAl based alloys. [199,200]In particular, the effect of Mn on the ω o phase is stronger than that of Mo, e.g., by substituting Mo with Mn, the formation of ω o phase in a TNM alloy could be inhibited. [200]2.6.Vanadium Vanadium is a very prominent alloying element in conventional Ti-base alloys and, due to its incorporation in the Ti64 alloy (Ti-6Al-4 V, in wt%), most likely the quantitatively most utilized β-stabilizing element.The isopleth of the Ti-45Al-V (at%) system depicted in Figure 10b as assessed by Takeyama et al. [54] shows that the effect of this element is qualitatively quite similar to the one of Mo.However, with respect to the occurring phase transformations, V-containing γ-TiAl based alloys often exhibit a sequential α !α þ γ !β þ γ transformation sequence in contrast to the eutectoid α !β þ γ transformation in Mo-containing alloys. [54]Quantitatively, V is a weaker β-stabilizing element than Mo.For example, the α single-phase field region extends up to 15 at% V in the Ti-Al-V system compared to only 2 at% Mo in the Ti-Al-Mo system for similar Al contents.However, since its β-stabilizing effect is stronger compared to Nb, V can be considered as an intermediate β-stabilizing element. [34]Assessments of the Ti-Al-V system by different authors have shown that no ternary intermetallic phases exist regardless of the temperature, [201,202] which is different from other ternary systems containing β-stabilizing elements, such as Nb, Ta and Mn.Concerning solidification, Zhai et al. [203] reported that the primary solidifying phase changes from β to α for a Ti-45Al-4 V (at%) alloy upon rapid solidification during free fall in drop tube experiments under an estimated cooling rate of 10 3 to 10 5 K s À1 .However, by increasing the V content of the material, the β phase could be retained even for such high cooling rates. [203]Another consequence of an increasing V content with respect to the solid-solid phase transformations was the suppression of the massive α !γ transformation in favor of the α !α 2 transformation, which is preferred for Al-leaner alloys. [204]Zhai et al. [204] argued that this can be attributed to the strong similarities between Ti and V, which are horizontal neighbors in the periodic table, thus making the alloy effectively Al-lean upon alloying with V. Furthermore, as mentioned in Section 2.2.1, this element is able to promote the martensitic β !α transformation over the massive one upon quenching from high temperatures, if a sufficient amount of this element is present. [54]With respect to the phase distribution, Wang et al. [205] reported that the alloying of 2 at% V to Nb-and Cr-containing γ-TiAl based alloys increased the amount of γ phase present in the microstructure.Additionally, V has been found to negatively impact the stability of the ω o phase in Nb-containing alloys. [206]2.7.Chromium Chromium is incorporated in the GE-4822 alloy and the so-called K5 alloy family (Ti-46Al-3Nb-2Cr-0.2W, in at%) with 2 at% each, which has been found to be an optimal amount to enhance the room temperature ductility. [6,207]Its effect concerning the stabilization of the β phase is considered stronger than the one of Nb and V. [34] With respect to the GE-4822 alloy, the contained 2 at% of Cr may be a reason from a compositional point of view, why small amounts of β o phase can be found in the microstructure of this alloy, [44,208] since 2 at% of Nb are not sufficient to stabilize said phase in a ternary Ti-Al-Nb alloy at room temperature, e.g., see Figure 7.A prediction of the stabilization of the β o phase at low temperatures by additions of Cr can be seen in Figure 11.
Figure 11.Calculated isopleths of the Ti-Al-Cr system after the assessment by Witusiewicz et al. [211] for a) 2 at% Cr and b) 45 at% Al.Panels (a) and (b) are based on the data reported in ref. [211]; Isothermal sections of the Ti-Al-Fe system at c) 800 °C and d) 1200 °C.Single-and triple-phase field regions are colored in gray and light gray, respectively.Note that the two single-phase field regions at high Al and low Fe contents correspond to the TiAl 2 and the so-called 1d-APS phase (not labeled).Panels (c) and (d) are based on the data given in ref. [213].
The α-transus temperature has been found to decrease by about 60 °C due to alloying of 4 at% Cr to a Ti-48Al based (at%) material.Simultaneously, the amount of α 2 phase was reduced in favor of the γ phase. [209]Interestingly, phase field regions containing the disordered α phase are shifted to lower temperatures at the expense of α 2 -containing regions for increasing Cr contents, i.e., the eutectoid temperature is lowered as exemplified by Figure 11b.Consequently, Cr tends to stabilize the α phase over the α 2 phase in contrast to, for example, Mo and V. [210,211] Several authors have reported that Cr inhibits the formation of the ω o phase (Table 2) to some extent. [199,200]However, as demonstrated by Xiong et al., [200] this element is not able to completely prevent the ω o phase formation in a derivative of the TNM alloy, in which the 1 at% Mo was replaced by 1 at% Cr.Since these authors observed a complete suppression of the ω o phase for the substitution of Mo with Mn, it can be concluded that Cr is not as strong as Mn with respect to this inhibition of the ω o phase. [200]Furthermore, at the highest Cr contents, as shown in Figure 11b, the intermetallic C15 Laves phase can occur, i.e., it is present above approximately 5 at% of Cr at low temperatures.

Iron
The alloying element Fe has only seen limited use in intermetallic γ-TiAl based alloys when compared to the β-stabilizing elements discussed so far.On the one hand, this may be attributed to its tendency to strongly segregate during solidification. [212]n the other hand, while Fe is a β-stabilizing element considered stronger than V with respect to the amount of β o phase stabilized at room temperature, [34] it also yields the precipitation of several Fe-based intermetallic phases even at relatively low alloying amounts.These phases may be detrimental to the ductility of the material. [212,213]Especially, as shown in Figure 11c,d, the phases τ 2 and τ 2 * as well as the Fe 2 Ti phase occur for different temperatures in the composition range relevant to γ-TiAl based alloys. [213]For example, Tokar et al. [214] reported on the presence of the τ 2 phase in a Ti-50Al-2Fe (at%) alloy, which formed via the decomposition sequence of a Fe-and Al-rich β o phase upon cooling from 1300 °C, i.e.,

Zirconium
][217] For instance, Kawabata et al., [215] who studied the influence of ternary additions in alloys based on the Ti-50Al (at%) composition, showed that Zr additions increase the yield and fracture stress at room temperature and 600 °C more effectively than V, Nb, Cu, and Mn.][217][218] In contrast to many other alloying elements, Zr prefers to enrich in the γ phase when compared to the α and the β phase (see Table 3).At temperatures at which the γ phase is not present, a partitioning tendency toward the β phase has been found.The solid solution strengthening by Zr additions is caused by a distortion of the crystal structure of the γ phase. [6]Especially, an asymmetric influence of Zr on the a and c axis of the L1 0 structure has been observed. [31,215,217,219]However, individual pieces of research have reported contradicting results with respect to the exact influence of Zr.Kawabata et al. [215] mentioned a decreasing length of both axes as well as an increased tetragonality, while several other authors found a decreasing c/a ratio for increased additions of Zr. [31,217,219] In the literature on conventional Ti-base alloys, Zr is generally considered as a neutral alloying element with the main purpose of strengthening the α phase. [52,220]Examination of the binary Ti-Zr system, e.g., see ref. [221], reveals that Zr additions reduce the β-transus temperature on the Ti-rich side, thus, indicating a β-stabilizing effect in the composition range of Ti alloys. [221]owever, the effect of Zr on the β/α equilibrium can be considered lower than one of the β-stabilizing elements discussed above, especially since a full miscibility of Ti and Zr in the α and β phase exists in this system due to the chemical similarities between these two elements. [221]Contrarily, the binary Zr-Al phase diagram is characterized by the presence of several intermetallic phases. [222]In the part of the Ti-Al-Zr system relevant to γ-TiAl based alloys, Zr has been found to drastically reduce the melting temperature of the material, i.e., additions of 5 at% can lower it by about 100 °C, as demonstrated in Figure 12. [223,224] Zirconium tends to stabilize the liquid phase over the β phase and, especially, the α phase.In particular, the α single-phase field region, which spans over 150 °C in a binary Ti-46Al (at%) alloy, vanishes for Zr contents above approximately 4.5 at%.][227] The contraction of the α single-phase field region is further promoted by the increase of the γ-solvus temperature.In particular, Musi et al. [228] reported that Zr stabilizes the γ phase at temperatures below this transition temperature, i.e., the amount of γ increases at the expense of α/α 2 .Additionally, the eutectoid temperature increases upon alloying with Zr, as demonstrated in Figure 12, indicating the stabilization of the α 2 compared to the α phase similar to the cases with Mo and V. [228,229] At intermediate temperatures, i.e., the temperature range of the α 2 þ γ phase field region, sufficient alloying with Zr may cause the formation of the (Zr,Ti) 5 Al 3 phase. [225]ith respect to ternary intermetallic compounds, contradicting results can be found for this system.Some authors claim that no ternary phases are thermodynamically stable in the Ti-Al-Zr system. [218,226]Other research work reported on the presence of the phases Ti 2 ZrAl and Zr 2 TiAl, which, however, could also be ordered substructures of the β phase as argued by some authors. [225,230,231]

Hafnium
Hafnium can be considered a β-stabilizing element, since it has been found to lower the β-transus temperature of Ti for additions up to about 20 at%. [232]Similarly, Hf acts as a weak β-stabilizing element in γ-TiAl based alloys. [233,234]For example, Chen et al. [233] reported on the presence of β o phase in a Ti-48Al-0.5W (at%)/Ti 2 AlN composite with 1 at% Hf compared to a Hf-free alloy variant.Additionally, Hf seems to increase the amount of γ phase at the expense of the α 2 phase in the microstructure. [233]Siahbuomi et al. [234] observed no change in the primary solidifying phase in alloys based on the GE-4822 alloy with up to 6 at% Hf.However, for Hf contents above 4 at%, an additional eutectic phase transformation occurred, resulting in the formation of eutectic structures in the interdendritic parts of the solidification microstructure containing Al 3 Hf 2 and TiAl 2 phase. [234,235]In accordance to this, Imayev et al. [216] only observed α 2 and γ phase in the microstructure of a Ti-44Al-2.5Zr-2.5Hf-0.2B(at%) alloy and defined the phase transforma- γ for this alloy.Furthermore, these authors suggested that the combined amount of 2.5 at% Zr and 2.5 at% Hf has a similar β-stabilizing effect as 5 at% Nb on the resulting microstructure. [216]However, Huang [236] observed a higher amount of β o phase in a Ti-44Al-8Nb-0.1B (at%) alloy when compared to a Ti-44Al-4Nb-4Hf-0.2Si(at%) after long-term heat treatments at 700 °C and concluded that Hf may be a weaker β-stabilizing element than Nb.

α-Stabilizing Elements
The following Section 3.3.1 and 3.3.2deal with the influence of α-stabilizing elements on the thermodynamics of γ-TiAl based alloys.Compared to the large amount of β-stabilizing elements, the group of technically relevant α-stabilizing elements only comprises Si, C, O, and N. Their preferred sublattices and partitioning behavior in the Ti-Al-related phases are given in Table 3.It is noteworthy that also Al acts as an α-stabilizing element.For example, it shifts the β-transus temperature to higher temperatures in the binary Ti-Al phase diagram in Figure 1a.However, since this element represents one of two fundamental constituents of γ-TiAl based alloys, it has already been discussed in Section 2.

Silicon
][239][240][241][242][243][244][245] The amount of Si in γ-TiAl based alloys is typically well below 1 at%, e.g., TNMþ with 0.3 at%, [116] K5S with 0.2 at%, [207] and the ABB-2 alloy with 0.5 at%, [175] to prevent a detrimental effect on the tensile properties and, especially, on the ductility. [237]This can be attributed to the fact that the solubility of the Ti-Al-related phases for the substitutionally incorporated Si is rather low, thus, promoting the precipitation of silicides. [6,246]Klein et al. [246] observed that the α 2 phase can dissolve the highest amount of approximately 0.5 at% Si, followed by the γ phase with 0.22 at% in the lamellar colonies of a heat-treated TNMþ alloy.The β o phase was found to exhibit the lowest solubility in this study.However, at high temperatures, this trend reverses, i.e., the β phase is able to incorporate more Si than the α phase. [246]Furthermore, it is suggested in literature that this element's overall solubility within an alloy, i.e., the amount of an alloying element that can be added without stabilizing new phases, may also depend on additions of quaternary alloying elements, e.g., strong silicide-forming elements such as Hf and Zr reduce the solubility of Si. [247] When the Si amount exceeds the solubility limit of the alloy, ζ-Ti 5 Si 3 silicides are formed, whose crystallographic data and orientation relationships with the Ti-Al-related phases are presented in Table 2. [175,237,240] The beneficial influence of Si on strength can be attributed to precipitation hardening by these fine ζ-Ti 5 Si 3 silicides as well as to solid solution hardening by dissolved Si. [240,242] The enhanced creep resistance can be explained by interfacial drag mechanisms as well as the hindering of dislocation movement by the ζ-Ti 5 Si 3 silicides. [243,248,249]Noda et al. [237] proposed two possible phase transformations causing the formation of silicides: a eutectic L ! β þ ζ-Ti 5 Si 3 transformation during solidification and a eutectoid α 2 !γ þ ζ-Ti 5 Si 3 transformation at lower temperatures.Especially, the eutectoid transformation results in the formation of silicides at the α 2 /γ and γ/γ interfaces in the lamellar α 2 /γ colonies and has been Based on the data reported in the work by Musi et al. [228] observed in many different alloys during/after creep exposure. [116,237,241,250]After the nucleation at the interface, the growth of the silicides proceeds into the α 2 lamellae. [237]owever, the formation of silicides inside the colonies significantly depends on the existing ratio of α 2 and γ phase due to their inherently different solubilities for this element. [246]Especially, if the amount of γ phase is too low, for example, owing to an Al-lean alloying concept, all Si can remain in a solid solution and no ζ-Ti 5 Si 3 silicides are able to precipitate.Another transformation involving silicides, i.e., supersaturated β o !β o þ ζ-Ti 5 Si 3 , was proposed by Klein et al. [246] to occur in the β o phase of β-stabilized γ-TiAl alloys due to its very low solubility for Si. [175][253][254] A portion of the liquidus projection is shown in Figure 13a.While the primary solidification still occurs via the β phase in the relevant composition range, some other phase transformations are detectable in this part of the ternary system.Especially, the two following liquid-solid phase transformations, labeled "U" and "E" in Figure 13a, can be extracted from the liquidus projection of the Ti-Al-Si system: L þ β !α þ ζ-Ti 5 Si 3 ("U") and L ! α þ γ þ ζ-Ti 5 Si 3 ("E"). [251]In the work by Noda et al., [237] who proposed a L ! β þ ζ-Ti 5 Si 3 transformation, coral-shaped precipitate skeletons were found in the interdendritic regions of the solidified microstructure. [237]Taking into account the high Al content of the studied alloy, which was 47.5 at% and, thus, promoted a peritectic solidification, the phase transformation labeled as "U" may also seem plausible to have occurred in the material of the respective work instead. [237]sothermal sections of the Ti-Al-Si system at elevated temperatures predict that the phase equilibria in the Al range of engineering γ-TiAl based alloys, i.e., 42 to 48 at%, remain mostly unchanged when compared to the binary Ti-Al system, with the only exception being the presence of ζ-Ti 5 Si 3 silicides. [252,253]urthermore, these isothermal sections support the observation that additions of Si stabilize the γ phase over the α 2 phase, i.e., [ 237,250] These findings were confirmed recently also for alloys based on the TNM alloying concept.Not only was Si found to increase the amount of γ phase at the expense of the α 2 phase at low temperatures while leaving the amount β o phase constant, it also increased the temperatures of solid-solid phase transformations, e.g., see Figure 13b. [255]

Interstitial Elements C, O, and N
The alloying element C is often intentionally alloyed to engineering γ-TiAl based alloys in quantities below 1 at% to improve the high-temperature capabilities, tensile and creep strength as well as the microstructural stability. [6,7,116]In contrast to C, the elements O and N are considered as impurities picked up during processing or high-temperature application in most alloying concepts.260] The alloying elements C, O and N, which are dissolved interstitially in the Ti-Al-related phases (Table 3), are characterized by a strong α-stabilizing effect in γ-TiAl based alloys.This results in a shift of the peritectic L þ β !α phase transformation to lower Al contents, thus, promoting the peritectic solidification over the β-solidification.This can be seen in Figure 14, which depicts the Ti-43.5Al-4Nb-1Mo-0.1B-xC(at%) system according to Schwaighofer et al. [261] In particular, these authors found that additions of 0.45 at% of C already change the solidification pathway of the TNM alloy from β-solidification to a peritectic one, e.g., see circle labeled "1" in Figure 14.Furthermore, Zhang et al. [262] have reported that sufficient additions of N can also change the primary solidification phase from β to α in high Nb-containing alloys.Simultaneously, the α-stabilizing effect of these elements impacts the α/β phase equilibrium below the solidus temperature and, consequently, the associated β-transus temperature.In particular, C, O, and N shift this particular transition temperature to higher values due to an enlargement of the α single-phase field region.Additionally, in the case of the TNM alloy, which does not exhibit such a phase field region at any Si 3 phase transformation, respectively.Based on the data presented in the work by Bulanova et al. [251] ; b) Experimentally determined isopleth of the TNM-Si system.Important phase transition temperatures are labeled.Red data points were determined by a combination of in situ high-energy X-ray diffraction (HEXRD), differential scanning calorimetry (DSC) and heat treatments (HT), while green ones were obtained by DSC and heat treatments.Based on the data given in the work by Musi et al. [255] temperature, alloying with more than 0.1 at% C stabilizes the α phase sufficiently to reintroduce this phase field region, e.g., see circle labeled "2" in Figure 14.][263] Additionally, N was able to prevent the massive formation of γ phase during air cooling of a GE-4822 alloy. [263]n the case of O, the α-transus temperature decreases with increasing amounts of this particular element. [257,264,265]dditionally, the α/α þ γ phase field region boundary in the ternary Ti-Al-O system is shifted to the Ti-rich side at 1200 °C for higher O contents. [266]A recent study on the Ti-Al-O ternary alloying system showed that O, while generally being described as an α-stabilizer, also promotes the formation of the ordered α 2 phase at the expense of the α phase.In particular, adding 1 at% O to a Ti-43Al (at%) alloy caused the ordering of the α phase into the α 2 phase at 1200 °C. [267]At higher temperatures, i.e., 1300 °C, the increasing amount of O absorbed during heat treatments in air was found to significantly affect the α/β phase equilibrium in a TNM alloy.As demonstrated in the study by Musi et al., [255] the β phase dissolved during heat treatments at 1300 °C in air, while it was stable at the same temperature in an inert Ar atmosphere.Such a β-destabilizing effect of O has also been observed at lower temperatures in the case of the β o phase.270] Similarly, the β o phase transformed into α 2 and a Cr 2 Ti-type Laves phase in the GE-4822 alloy via a eutectoid decomposition in a temperature interval comparable to the one the TNM alloy was subjected to. [269]Ultimately, the increase of the material's overall O solubility due to the high solubility within the formed α 2 phase was identified as a contributing factor to the driving force of these transformations. [271]he influence of O on the phase equilibria of γ-TiAl based alloys is also a crucial aspect concerning the production of these materials.Generally, individual processing routes tend to result in different amounts of dissolved O.For example, Gerling et al. [272] obtained an O content of around 500 m.-ppm after hot-isostatic pressing and sintering compared to 1500 m.-ppm after metal injection molding of the same alloy.The associated changes of the phase distribution for a given nominal alloy composition have to be considered when designing heat treatment strategies and dealing with different impurity contents.For example, Nakashima and Takeyama [273] showed that adding O to ternary Ti-Al-V and Ti-Al-Cr alloys shifts the α 2 þ β þ γ phase field region toward higher contents of V/Cr, which affects the volume fractions of the individual phases and, ultimately, the mechanical properties.Furthermore, work by Distl et al. [273] indicated that the solubility of Nb in the α phase at 1200 °C increases from about 12 to 16 at% when the amount of O is increased from 200 to 400 m.-ppm, which also means that the α þ β þ γ phase field region shifts in a similar manner.
The solubilities for C, O, and N in the Ti-Al-related phases show similar tendencies, i.e., the solubility in the α 2 phase is significantly higher than the one in the γ phase.Consequently, the material's overall solubility for these elements substantially depends on the Al content governing the phase distribution between α 2 and γ.In the case of binary TiAl alloys, the γ phase can only dissolve 200-300 at.-ppm C. [274] Interestingly, the C solubility of the γ phase can be increased by Nb additions causing the formation of Ti octahedrons, which are energetically advantageous for C within the L1 0 structure. [274,275]The same effect can be achieved if the Ti/Al ratio in the γ phase is increased. [275]he higher C solubility of the α 2 phase compared to the γ phase is also attributed to the presence of such Ti octahedral sites, occurring naturally in the D0 19 crystal structure.In the case of the TNM alloy, the C solubility of the γ phase and the α 2 phase are 2500 and 15 000 at.-ppm, respectively. [261]or comparison, the solubility for O in binary two-phase α 2 þ γ TiAl alloys is 300 at.-ppm in the γ phase and 8000-22 000 at.-ppm in the α 2 phase. [274,276]The same order of magnitude was observed by Perdrix et al. [256] for the N solubility of the α 2 phase, which was found to be higher than 23 200 at.-ppm.
Exceeding the alloy's overall solubility limit for C, O, or N yields the precipitation of carbides, oxides or nitrides, respectively, in the microstructure.In the case of C, the cubic Ti 3 AlC perovskite carbides (P-type) and hexagonal Ti 2 AlC carbides (H-type) can occur depending on the composition of the material, the thermal history and the applied processing route. [6,7,277,278]The H-type carbides are stable at higher temperatures, but have less of an impact on the dislocation mobility and, thus, creep resistance. [6,277]In particular, this type of carbides is often observed in the as-cast microstructure of C-containing γ-TiAl based alloys. [261,279]The formation of P-type carbides from the γ phase within lamellar α 2 /γ colonies during aging at 750 °C was reported by Schwaighofer et al. [280] in a C-containing TNM alloy.The dissolution temperature of these carbides during subsequent heating was found to be 1185 °C. [280]In the case of O and N, the formation of oxides and nitrides is commonly observed as a consequence of oxidation under normal atmosphere. [126,245]owever, such phases have also been reported in literature to occur as a part of the microstructure of bulk material.In the case of N, the Ti 3 AlN and Ti 2 AlN nitrides with the same crystal  [261] Phase field boundaries associated with important transition temperatures are labeled.The open red circles show: 1) the vanishing of the β single-phase field region and, thus, the transition from β-solidification to peritectic solidification; 2) the emergence of an α single-phase field region.Adapted with permission. [261]Copyright 2014, Elsevier.
structures as their respective carbide counterparts (see Table 2) can precipitate under suitable circumstances. [259,260,263]With respect to O, Menand et al. [274] reported that excess O causes the formation of fine-scaled oxides inside single-phase γ alloys.

Further Elements
This section focuses on elements for which a classification into αor β-stabilizing may not be applicable, either due to an extremely low solubility or insufficient available data in literature.The elements in question are B as well as selected rare-earth metals such as Y, Sc, La, Ce, Nd, Gd, Ho, and Er.In the case of B, information on the partitioning behavior and occupied sublattices in the Ti-Al-related phases is given in Table 3.

Boron
While the typical amounts of B added to intermetallic γ-TiAl based alloys are only in the range of 0.1-1 at%, a significant impact on the mechanical properties has been found by extensive research work. [281]For example, Hu [282] observed an increase in the room temperature ductility of a GE-4822 alloy from 0.2% to 1% when alloying with 1 at% of B. The beneficial effect of B is attributed to its role as grain refinement agent, regardless of the type of solidification, i.e., β-solidification or peritectic solidification.Depending on the exact alloyed amount, B yields a fine as-cast microstructure either via increased heterogeneous nucleation of grains at borides in the melt or via constitutional undercooling promoted by its high partition coefficient favoring the liquid phase. [282]However, the efficiency of the B effect seems to depend on the alloy's Al content, as Al-lean alloys require less B to achieve effective grain refinement. [283]dditionally, during the solid-solid phase transformations subsequent to solidification, i.e., the β !α transformation, the TiB 2 borides (see Table 2) reduce the α grain size by heterogeneous nucleation. [16,281]introducing strong borideforming elements like W, Ta and Nb to γ-TiAl based alloys may change the prevalent boride from TiB 2 to TiB (see also Table 2) and reduce the efficiency of the added B as grain refinement agent due to its removal from the melt and the related decreased constitutional undercooling. [36]But also other processing methods, such as gas atomization and powder densification by spark plasma sintering, benefit from the positive influence of B on the final size of the microstructural constituents. [284,285]enerally, it is worth mentioning that the B solubility in the solid Ti-Al-related phases is well below 0.1 at% at room temperature. [286,287]Furthermore, an effect of these small B additions on the equilibrium phase transformation temperatures has never been observed.

Rare-Earth Metals
[293][294][295][296] In particular, such precipitates have been found to refine the size of microstructural features.However, compared to the alloying elements discussed previously, only a few studies are currently available in the literature regarding the influence of rare-earth metals on the phase transformations in γ-TiAl based alloys.In the case of Y, which may be the most-studied rare-earth metals-based alloying element in γ-TiAl based alloys up to this day, additions of 2 at% have been found to increase the amount α 2 phase and yield the precipitation of Al 2 Y in a cast Ti-46Al-2Cr-2Nb (at%) alloy. [290]ther authors have reported on the formation of Y 2 O 3 precipitates through internal oxidation in alloys based on the composition Ti-47Al-1Mn-2Mo-0.3C-(0-0.6)Y(at%). [296]hese precipitates were able to promote the dislocation activity in the γ phase by reducing the O content and, thus, improved the ductility. [296]A study by Yang et al. [293] investigating the influence of small additions of Sc and La-rich misch metal to a Ti-40Al-16Nb (at%) alloy showed that these elements decrease the α-transus temperature and stabilize the α 2 phase.Contrarily, Choi et al. [297] observed an increased amount of γ phase upon alloying Ce and La to a Ti-45Al-3Mo-2Nb (at%) alloy.Summarized, more dedicated studies are required to fully understand the effect of the individual rare-earth metals on the phase transformations, since comparisons are often difficult due to low overall amounts of rare-earth metals and significantly different base alloy compositions.

Summary and Outlook
A significant amount of research work has been dedicated to the investigation of the influence of alloying elements on phase transformations and properties of intermetallic γ-TiAl based alloys.Often, such additional elements increase the complexity of the material by introducing new stable and metastable phases in addition to changing the occurring phase equilibria.The purpose of this work was to summarize the results of the many different alloying elements that have been studied in past.First, an overview on the important phase transformations associated with the Ti-Al system, i.e., β !α, α !γ and α !α 2 þ γ, and on how these are affected by different cooling rates were given.Afterward, the different alloying elements were categorized with respect to their influence on the β/α phase equilibrium into β-stabilizing (Nb, Mo, W, Ta, Mn, V, Cr, Fe, Zr, Hf ), α-stabilizing (Si, C, O, N) and other elements (B, rare-earth metals), and the current state of knowledge on how these elements affect the Ti-Al system was discussed.Ultimately, future research dedicated to phase transformations in γ-TiAl based alloys may be directed toward the cointeraction of several alloying elements, which can be assumed to alter the thermodynamic properties with respect to the individual ternary systems.Such experimental data represent an important aspect for future alloy development programs as well as the optimization and validation of thermodynamic databases.

Figure 1
Figure 1.a) Section of the binary Ti-Al phase diagram according to the assessment in the study by Schuster and Palm.[27]Phase field region boundaries associated with important transition temperatures, i.e., the peritectic (purple), the β-transus (orange), the α-transus (=γ-solvus in the binary system) (green), and the eutectoid (blue) temperature, are highlighted in different colors.The red region marks the Al content of engineering γ-TiAl based alloys.Based on the data reported in the study by Schuster and Palm;[27] b) Crystal structures of the ordered and disordered phases in binary γ-TiAl based alloys.Red atoms represent Ti, blue atoms Al.Purple atoms indicate that these lattice points are randomly occupied by either Ti or Al.

Figure 4 .
Figure 4. a) Graphical representation of the Burgers orientation relationship connecting the body-centered cubic lattice of the β phase (green) and the hexagonal lattice of the α phase (purple).The green dashed lines illustrate an extension of the (110) β plane and highlight the fact that the atomic positions of the β and α phase overlap almost completely in a configuration according to the Burgers orientation relationship; b) schematic time-temperaturetransformation diagram for cooling from the β phase field region (MS: microstructure).Note that only β !α solid-solid phase transformations discussed in the text are shown and further phase transformations are possible; Panels c) through e)show examples of a Widmannstätten, a massive and a martensitic microstructure in different γ-TiAl based alloys, i.e., (c) Ti-44Al-3Mo-0.1B (at%), d) Ti-42Al-5 V (at%) and e) Ti-42Al-10 V (at%).Note that ordering of the phases α and β may occur subsequent to the β !α transformation for the cooling rate of water quenching depending on the material's composition, thus, resulting in the presence of α 2 and β o in the final microstructure.Panels (d) and (e) are adapted with permission.[54]2005, Elsevier.

Figure 5 .
Figure 5. a) Graphical representation of the Blackburn orientation relationship connecting the α/α 2 and γ phase; b) time-temperature-transformation diagram for the cooling from the α single-phase field region showing the manifestation of the α !γ transformation for different cooling rates.Adapted with permission.[7]2013, Wiley; c) Gibbs free energy considerations for the massive α !γ transformation upon quenching from a temperature T 1 in the α single-phase field region to a temperature T 2 in the α þ γ phase field region for an alloy with an Al content x Al 0 .

Figure 6 .
Figure 6.a) Isoplethal section of the Ti-xAl-4Nb-1Mo-0.1B(at%) system.The red line corresponds to the nominal chemical composition of the TNM alloy.Based on the data reported in the study by Schwaighofer et al.[298] ; b) B2 crystal structure of the ordered β o -TiAl phase.Red atoms correspond to Ti, blue ones to Al; c) Microstructure of a heat-treated TNM alloy showing a complex multi-phase microstructure consisting of α 2 , β o and γ phase.Adapted with permission.[255]Copyright 2022, Elsevier.

Figure 7 .
Figure 7. Isothermal sections of the Ti-Al-Nb system relevant to γ-TiAl based alloys at a) 700 °C and b) 1200 °C.Single-and triple-phase field regions are colored in gray and light gray, respectively.Based on the data presented in the studies by Distl et al.[137,138]

Figure 9 .
Figure 9. a) Convoluted microstructure of an aged Ti-46Al-8Ta (at%) alloy showing α 2 lamellae with different orientations inside individual γ grains (see text).Adapted with permission.[184]2010, Wiley; b) Calculated isopleth of the Ti-Al-Ta system for 8 at% Ta.The red line marks the composition of the Ti-46Al-8Ta (at%) alloy.The black dashed line indicates the disordering temperature of the β o phase.The letter τ marks the existence range of a ternary Ta-rich phase (see text).Based on the data reported in ref.[182].

Figure 12 .
Figure 12.Several isopleths of the ternary Ti-Al-Zr system for a) a fixed Zr content of 2 at%, b) for a fixed Zr content of 4 at%, and c) for a fixed Al content of 46 at%.Phase field region boundaries associated with important transition temperatures are labeled.Based on the data reported in the work by Musi et al.[228]

Figure 13 .
Figure 13.a) Section of the liquidus projection of the ternary Ti-Al-Si system in the composition range of γ-TiAl based alloys.C 1 and C 2 mark the maximum temperatures on the corresponding blue liquidus lines.P, U and E denote the points of the binary peritecticL þ β !α, the ternary quasi-peritectic L þ β !α þ ζ-Ti 5 Si 3 and the eutectic L ! α þ γ þ ζ-Ti 5Si 3 phase transformation, respectively.Based on the data presented in the work by Bulanova et al.[251] ; b) Experimentally determined isopleth of the TNM-Si system.Important phase transition temperatures are labeled.Red data points were determined by a combination of in situ high-energy X-ray diffraction (HEXRD), differential scanning calorimetry (DSC) and heat treatments (HT), while green ones were obtained by DSC and heat treatments.Based on the data given in the work by Musi et al.[255]

Figure 14 .
Figure 14.TNM-C phase diagram after Schwaighofer et al.[261] Phase field boundaries associated with important transition temperatures are labeled.The open red circles show: 1) the vanishing of the β single-phase field region and, thus, the transition from β-solidification to peritectic solidification; 2) the emergence of an α single-phase field region.Adapted with permission.[261]Copyright 2014, Elsevier.

Table 1 .
Crystallographic data of important solid phases occurring in binary γ-TiAl based alloys.

Table 3 .
Preferred sublattices in the α 2 , β o and γ phase as well as the partitioning behavior of the alloying elements in the focus of this work.Note that the presented partitioning behavior distinguishes between high temperatures (α, β, γ) and low temperatures (α 2 , β o , γ).
a)n/a = not available (to the authors' best knowledge); int.= interstitial; -= no preference.