Experimental Phase Equilibria and Isopleth Section of 8Nb-TiAl Alloys

Abstract

The 8Nb isopleth section of a Ti-Al-Nb system is experimentally determined based on thermal analysis and thermodynamic calculation methods to obtain the phase transformation and equilibrium relations required for material design and fabrication. The phase transus and relations for the 8Nb-TiAl system show some deviations from the calculated thermodynamic results. The ordered β_o phase transforms from the disordered β/α phases at 1200–1400 °C over a large Al concentration range, and this transformation is considered to be an intermediate type between the first- and second-order phase transitions. Moreover, the β_o phases are retained at the ambient temperature in the 8Nb-TiAl microstructures. The ω_o phase transforms from the highly ordered β_o phase, rather than from α₂ or β_o with a low degree of atom ordering B2 (LOB2) structure, with Al concentration of 32–43 at% at approximately 850 °C. From the experimental detection, the transition of the ω_o phase from the β_o phase is considered to be a further ordering process.

1. Introduction

In the past decade, TiAl-based alloys have been considered promising candidates for high-temperature materials in aerospace and automotive applications because of their excellent properties of low density, high specific yield strength and stiffness, and favorable oxidation resistance and creep properties up to high temperatures [1,2,3]. TiAl alloys containing high amounts of Nb, based on the γ-TiAl and α₂-Ti₃Al intermetallic, exhibit excellent high-temperature strength and oxidation resistance [4] and have attracted significant attention [3,5]. Advanced materials based on the Ti-Al-Nb alloys can be used at temperatures above 800 °C [6]. As reported by Appel et al. [3] and Erdely et al. [7], both the hot-workability and the ductility can be effectively promoted by the β phase [5,8], which provides a sufficient number of independent slip systems to act as a ductile constituent in the final microstructure [9,10]. Moreover, the solid-state transformation pathway and the microstructure of TiAl-based alloys can be manipulated through β/β_o phases [3,7,11]. According to reports by Cheng et al. [12], Kobayashi et al. [13], and Takeyama et al. [14], a multitude of solid-state transformations and resulting microstructural morphologies can be achieved by stabilizing the β phase [3].

The phase diagram and phase equilibria of the Ti-Al-Nb system play important roles in material design and fabrication, especially in the manipulation of the solid-state transformation pathway and microstructure [5,15,16]. Phase transition behaviors can be precisely detected by some thermal detection methods, such as differential scanning calorimetry (DSC), differential thermal analysis (DTA), and in-situ X-ray/neutron diffraction technologies. These determination methods generally show good agreement with each other [17,18,19,20,21,22,23]. Phase diagrams of the three constituent binary systems in the Ti-Al-Nb system—Ti-Al [24], Al-Nb [25], and Ti-Nb [26]—have been sufficiently delineated. The isothermal sections at 1000 °C [27,28,29,30], 1100 °C [30,31], 1150 °C [28,29,30], 1200 °C [27,32], 1300 °C [33], and 1400 °C [28,29,34,35] from experimental data are available in the literature. In addition, isopleth sections show the development of phase equilibria, such as the Ti_72.5Al_27.5-Nb [36], Ti₇₈Al₂₂-Nb [37], TiAl-TiNb [37,38], 45Al [25], 47Al [25], 8Nb [25,39], and 10Nb sections [39]. Thermodynamic methods have also been adopted to assess the diagrams of Ti-Al-Nb systems [25,40,41]. Recent reviews of literature on the Ti-Al-Nb system up to 2011 have been provided by Witusiewicz et al. [25], Cupid et al. [41], and Raghvan et al. [42]. However, insufficient experimental data contributed to some differences between these assessments, including the equilibria of the α₂, β_o, O, and ω_o phases.

The 8Nb-TiAl alloys, as a new family of TiAl-based alloys, have been widely investigated owing to their higher strength and good oxidation resistance [43,44]. For addition of 8 at% Nb in the TiAl alloys, the transus temperatures of α and β, as well as the α phase field, are generally decreased, inducing the significant refinement of the structure and the improvement of the yield strength [45].

Table 1. Crystallographic data of phase designations occurring in this work.

Designation	Pearson	S. G. (#)	Strukt.	Prototype	Lattice Parameters (nm)
β-Ti(Al,Nb)	cI2	$Im\overline{3}m$ (229)	A2	W	a = 0.331	-	-
α-Ti(Al,Nb)	hP2	$P{6}_3/mmc$ (194)	A3	Mg	a = 0.295	-	c = 0.468
γ-TiAl	tP4	$P4/mmm$ (123)	L10	AuCu	a = 0.283	-	c = 0.408
α2-Ti3Al	hP8	$P{6}_3/mmc$ (194)	D019	Ni3Sn	a = 0.577	-	c = 0.463
βo-TiAl(Nb)	cP2	$Pm\overline{3}m$ (221)	B2	CsCl	a = 0.322	-	-
ωo-Ti4NbAl3	hP6	$P{6}_3/mmc$ (194)	B82	InNi2	a = 0.458	-	c = 0.552
O-Ti2NbAl	oC16	$Cmcm$ (63)	---	NaHg	a = 0.616	b = 0.973	c = 0.470

presents the phases most frequently used for the Ti-Al-Nb system in this work, as well as their crystallographic data [20,41,46,47]. However, a lack of knowledge regarding the constitution of such multicomponent systems is one of the major obstacles to alloy development and significant work has to be expended to obtain reliable information on phase relationships [46].

In this work, the 8Nb isopleth section, as well as the phase equilibria, of the Ti-Al-Nb system is experimentally determined by DSC, scanning electron microscopy (SEM), and X-ray diffraction (XRD). The samples were annealed sufficiently in high vacuum and then slowly cooled to achieve the equilibrium state as far as possible. However, considering the difficulty of achieving thermodynamic equilibrium in practice, the experimental phase equilibria may have some deviations from thermodynamic equilibria. This work aims to establish the experimental isopleth and phase equilibria of the 8Nb-Ti-Al system, which have the treating conditions closer to the actual production, to obtain the phase transformation and equilibrium relationships required for material design and fabrication.

2. Materials and Methods

Eighteen ternary 8Nb-TiAl alloy buttons were prepared by non-consumable arc melting (with a tungsten electrode) from high-purity initial components (Ti, Al, Nb: 99.99 wt%) using a water-cooled copper hearth in an ultra-high-purity argon atmosphere. Only those samples with less than 1.0 wt% weight losses were adopted for further analysis. Each button (with a mass of approximately 30 g) was melted, turned over, and re-melted five times to ensure homogeneity. The diffusion annealing was performed in high vacuum at 1400 °C for 20 h to equilibrate and coarsen the phase. The samples were then slowly cooled to 50 °C in the furnace (cooling rate: ~1.0 °C/min). Thermodynamic equilibrium calculations were performed by the Calculation Phase Diagrams (CALPHAD) method using the commercial software Pandat^® (v2017.1, Compu Them, Middleton, WI USA) [48], following the thermodynamic description from Witusiewicz’s investigation [25]. The phase structures were detected by XRD (D8 Advance, Bruker, Bremen, Germany) using powder samples, prepared in a metallic miller, and operated at 40 kV and 40 mA with Cu Kα. For the morphology analysis, sample sizes of 8 mm × 8 mm × 8 mm were cut from the centers of the buttons with an electron discharge cutting machine (New Sparkie, Shaoxing, China), and the samples were then polished following standard mechanical polishing procedures. The microstructures of the annealed samples were obtained by SEM in back-scattered electron (BSE) mode (JSM-6380, JEOL, Musashino, Japan), operated at 20 kV and at a working distance of 10 mm. The compositions in the samples were analyzed by the optical emission spectrometer (OES, Q4 TASMAN 130, Bruker, Bremen, Germany).

The solid-state transformation temperatures were determined by DSC (404C, NETZSCH, Selb, Germany) with an Al₂O₃ crucible. Discs of 4 mm diameter and 3 mm thickness (150–180 mg), which produced sufficient signals in the argon atmosphere, were machined from the centers of the buttons. Temperature calibrations for the equipment were conducted by melting high-purity Cu, Al, Ag, and Ni at scan rates of 5 and 10 °C/min to determine the correction equation. In this work, the heating rate for all test samples (excluding the calibration samples) was 10 °C/min, using a covered Al₂O₃ crucible with a maximum temperature of 1450 °C, and within a dynamic Ar atmosphere with a flow rate of 30 mL/min. To account for the influence of the containers, the DSC signals of the empty crucibles were subtracted from the detected traces. As a standard comparison sample, thermal data were acquired from sapphire, using the same parameters as those applied to the 8Nb-TiAl samples, to determine the specific heat capacity (C_p) by Equation (1).

$$C_p^{Sample}=\frac{DS{C}_{sample}}{DS{C}_{sapphire}}·C_p^{Sapphire}$$

(1)

In this work, to improve the accuracy and uniformity of the data acquisition process, a new method was developed for determining and calibrating the onset transformation temperatures for the different cooling/heating rates in DSC detection. The first derivatives of the DSC traces (DDSC) were employed to determine the transformation temperatures. The temperature difference, ΔT = 38 °C, was adopted to calibrate the dynamic transformation temperatures. Full details are given in Supplementary Materials.

3. Results

3.1. Phase Evolution

The chemical compositions of the 8Nb-TiAl alloys (S01–S18) determined by OES, with the nominal compositions, are presented in

Table 2. Nominal and determined compositions (at. %) of the 8Nb-TiAl alloys.

Sample	Nominal Composition	Chemical Composition (OES)
		Al	Ti	Nb
S01	Ti73.6Al18.4Nb8	18.34	73.69	7.97
S02	Ti70.8Al21.2Nb8	21.13	70.89	7.98
S03	Ti68.1Al23.9Nb8	23.80	68.23	7.97
S04	Ti65.3Al26.7Nb8	26.65	65.37	7.98
S05	Ti62.6Al29.4Nb8	29.37	62.66	7.97
S06	Ti59.8Al32.2Nb8	32.16	59.87	7.97
S07	Ti58.4Al33.6Nb8	33.52	58.51	7.97
S08	Ti57.0Al35.0Nb8	34.89	57.11	8.00
S09	Ti55.7Al36.3Nb8	36.24	55.75	8.01
S10	Ti54.3Al37.7Nb8	37.63	54.35	8.02
S11	Ti52.9Al39.1Nb8	39.01	52.96	8.03
S12	Ti51.5Al40.5Nb8	40.42	51.56	8.02
S13	Ti50.1Al41.9Nb8	41.83	50.14	8.03
S14	Ti48.8Al43.2Nb8	43.13	48.84	8.03
S15	Ti47.4Al44.6Nb8	44.53	47.44	8.03
S16	Ti46.0Al46.0Nb8	45.93	46.04	8.03
S17	Ti44.6Al47.4Nb8	47.35	44.63	8.02
S18	Ti43.2Al48.8Nb8	48.72	43.27	8.01

. The specific reflections of certain phases—α₂(202), γ(002)/(200), O(310), and β_o(200)/ω_o(202)—were selected to describe the phase evolution, as shown in Figure 1. The reflection in S01–S12 at approximately 38.5° was much stronger than the reflection at 40.8°, showing the overlap of the diffraction peaks with the β_o phase. For this β_o phase, the (110)/(220) reflections of the disordered β structure were preserved, whereas the (200)/(211) reflections were suppressed, indicating a LOB2 structure, as shown in Figure 1b,c. However, as shown by the XRD pattern of Sample S06–S12, the reflection at approximately 39.3° was considered to indicate overlapping of the β_o and ω_o phases, as shown in Figure 1c. The β_o phase was fully depressed in S14–S18, as shown in Figure 1d. Check out the Supplementary Materials for more details.

The evolution of the fraction of each phase in the samples can be roughly evaluated by the reflection peak height, where unstacked reflections should be selected first. Figure 2 shows the peaks of the α₂(202), γ(002)/(200), O(310), and β_o(200)/ω_o(202) reflections in the 8Nb-TiAl samples. The α₂ phase was detected as the main phase in all the samples, as shown in Figure 1. The fraction of the α₂ phase increased with increased Al content in Samples S01–S10, then decreased with continuing Al increases in Samples S11–S18, as shown in Figure 2a. The fraction of γ phase was small in Samples S01–S10, and then the γ phase fraction increased continuously as the Al content increased in Samples S11–S18, as shown in Figure 2c. The presence of the O phase could be identified by the (310) reflection in the Al-lean S01–S05 samples, as shown in Figure 2b, and the ω_o and β_o phases were identified in Samples S06–S12, as shown in Figure 1b and Figure 2d. The fraction of β_o+ω_o phases increased dramatically in Samples S06–S09, then decreased in S10–S13. In the Al-rich S13–S18 samples, only two phases formed: α₂ and γ.

3.2. Morphology

As shown in Figure 3a,b, the morphologies of Samples S01–S05 (18.4–29.4 at% Al) comprised coarse α₂ laths and O phase (dark regions); flaky α₂ particles in the low order β_o matrix; interwoven β_o structures (bright regions). The β_o phase mainly appeared along the α₂ grain boundaries and only a small volume fraction of the β_o phase was present within the α₂ laths. As Al increased in Samples S01–S05, the fraction of α₂ increased while the O and β_o phases decreased, as observed by the morphology analysis. The presence of ω_o in Samples S06–S12 was identified by XRD, as shown in Figure 1, whereas this phase was not observed in the microstructure, as shown in Figure 3 and Figure 4. For Samples S06–S09, the morphology mainly consisted of the coarse α₂ laths and interwoven β_o+ω_o structures, as shown in Figure 3c,d. The β_o phase was mainly located along colony boundaries and only a small volume fraction was present within the colonies [5]. The morphologies of Samples S10–S12 consisted of lamellar colonies, (α₂+γ)_L, and interwoven β_o+ω_o structures, as shown in Figure 3e,f. Isolated γ equiaxed grains were found on the boundaries of α₂ and β_o, as shown in Figure 3 and Figure 4. The morphologies of Samples S13–S17 were fully lamellar, as shown in Figure 4a–d. The size of the lamellar colonies and the interlamellar spacing in these samples increased greatly. Figure 4c,d show that both the proeutectoid α₂ and the γ were largely situated along boundaries, at the triple lines of (α₂+γ)_L, and around the β_o grains. However, the presence of β_o and γ was extremely restrained as the Al content increased to 44.6 at.%, as shown in Figure 4c,d. The morphology of Sample S18 consisted of coarse γ laths and interwoven eutectoid α₂+γ microstructures, as shown in Figure 4e,f.

3.3. Thermal Analysis

The phase transformation temperatures of 8Nb-TiAl alloys during experimental continuous heating were determined by thermal analysis. The phase transitions were simultaneously determined by DSC and C_p data, with the calculated C_p for comparison. Figure 5 shows the thermodynamically calculated C_p data and experimentally determined DSC and C_p data for the selected samples. The abrupt changes in the calculated C_p, shown by the narrow peaks in Figure 5a, correspond to the formation of an ordered β_o phase from α/β; the order-disorder transitions of β ↔ β_o and the eutectoid reaction of α ↔ β_o + γ can also be determined from the experimental C_p data, as shown in Figure 5b,c. However, as shown in Figure 5c, there is a particularly broad peak at approximately 900 °C in the experimental C_p curve, indicating an additional ordering transition in Samples S06–S12, which was considered to be the β_o ↔ ω_o transition. For Samples S15–S18, no β ↔ β_o transition was found at elevated temperatures, as shown in Figure 5d. Based on the DSC, SEM, and XRD analyses, the phase transitions in 8Nb-TiAl alloys were determined, as shown in Figure 5b–d, where the phase regions and isopleth sections are schematically identified in the DSC/C_p curves.

3.4. Isopleth

Based on this work’s thermal analyses, the experimental 8Nb-TiAl isopleth section was deduced and compared with a thermodynamic isopleth, as shown in Figure 6. More details about the transformation temperatures determined by DSC detection in this work and by the literatures are provided in Supplementary Materials. The experimental transus between phase transitions was approximated by a three-order polynomial with a least-squares solution, as shown in Figure 6. The experimental isopleth showed some differences from the thermodynamic calculations.

The phase transition behaviors found in this work were in good agreement with those previously reported in the literature, as shown in Figure 6. Schmoelzer et al. [11] and Erdely et al. [7] determined the phase transition temperatures in Ti-43.5Al-(4–5)Nb-(Mo, B) alloys using DSC and in-situ high-energy X-ray diffraction (HEXRD) methods. The eutectoid transition (α ↔ β_o + γ and β_o ↔ α₂ + γ) temperature has been considered to be T_eut = 1160–1180 °C, whereas the α ↔ γ temperature has been T_α = 1246-1255 °C; these results agreed well with the transition temperatures taken from the experimental transus, T_eut = 1151–1178 °C and T_α = 1230 °C, respectively. The order-disorder β ↔ β_o transition temperatures have been detected at 1225 °C for Ti-43.5Al-4Nb-(Mo, B) [7,11], and 1200 °C for Ti-(42–43.5)Al-Nb [7,17], in agreement with the transus temperature, 1250–1270 °C in the experimental isopleth, as shown in Figure 6. The α ↔ α₂ + γ transformation temperatures in Ti-(42–43.5)Al-Nb have been detected at 1169 °C by in-situ HEXRD and DSC [7,17], located exactly in the phase regions of the experimental isopleth. Meanwhile, the α ↔ β transition temperature has been considered to be approximately 1300 °C [17], corresponding to the experimental value, approximately 1310 °C, in the isopleth. For the Ti-45Al-8.5Nb-(W,B,Y) alloy [49], the α ↔ γ and eutectoid transition temperatures have been considered to be in the range of 1150–1310 °C and 1100–1220 °C, and these were determined to be 1280 °C and 1150–1176 °C in the experimental isopleth, respectively.

The β_o ↔ ω_o transition temperatures in Ti-45Al-10Nb have been determined to be approximately 780 °C with a slow cooling process (10 °C/min), and approximately 850–870 °C with a heating process (10 °C/min) [20,50]. In addition, during a slow furnace cooling process, the β_o ↔ ω_o transition temperature in Ti-42Al-8.5Nb has been detected at 928 °C by in-situ HEXRD and DSC with a heating rate of 20 °C/min [17]. These experimental results are in good agreement with the transus determined in this work. In contrast, at the temperatures near 800 °C, the α₂ ↔ ω_o transition has been seen to occur in only a few minutes with applied stress, although the transition took several hundred hours without applied stress [51]. Moreover, the β_o ↔ ω_o transition is promoted by the addition of a β/β_o-stabilizing element, such as Nb or Mo [20,25,52,53]. The α₂ ↔ ω_o transition has been observed to be restrained by the addition of an α₂-stabilizing element, such as C [51].

4. Discussion

4.1. Phase Equilibria

As shown in Figure 6, the experimental 8Nb-TiAl alloys exhibited the transformation pathway β → β + β_o → β_o + α₂ → β_o + α₂ + O for Al concentrations lower than 30 at%, and β → β + β_o → β_o + α₂ → β_o + α₂ + γ → β_o + ω_o + α₂ + γ for Al concentrations of 30–39 at%. The eutectoid transitions, α ↔ β_o + γ and β_o ↔ α₂ + γ, as well as the β_o ↔ γ transition, had to be considered a single region in this work because of the lack of sufficient experimental data. For alloys with 39–42 at% Al, a pathway with the α phase emerged, β → β + β_o/α → β + β_o + α → β + β_o + α + γ → β_o + ω_o + α₂ + γ. For alloys with 42–47 at% Al, the pathway was β → β + α → α → α + γ→ α + β_o + α₂ + γ → α₂ + γ. For alloys with Al concentrations greater than 47 at%, the pathway was β + α → α → α + γ → α₂ + γ.

In the 8Nb-TiAl system, the β ↔ β_o transition occurred over a large range of Al concentrations, as shown in Figure 6. The transitions of samples with Al ranges of 18.4–43.2 at% at 1200–1400 °C, indicated by the #1 transus line in Figure 6, exhibited higher transformation temperatures than those calculated, which increased as the Al concentration became lower than 33.5 at%, with a maximum transformation temperature of 1380 °C at 32 at% Al, while the transition temperature decreased over the range of 33.5–42.5 at% Al and disappeared entirely as the Al concentration increased above 44.0 at% Al.

For alloys with 18–36 at% Al, following the β ↔ β_o transus, there was a high-temperature phase transition, determined as the ordered α₂ phase transformation from β_o phases (the #2 and #3 transus lines in the isopleth), as shown in Figure 6. The preliminary α₂ phase formed the coarsening laths among the β_o phase, and the subsequent α₂ phase forms the flaky particles in the β_o matrix, as shown in Figure 1.

For alloys with 18–30 at% Al, the formation of the O phase was determined by XRD, although not observed in the microstructures because of the small amount of the phase fraction in the alloy matrix, as shown in Figure 1. The fraction of O phase in the samples decreased and the α₂ phase increased as the Al concentration increased, as shown in Figure 2c,d. The O phase can be formed in a large temperature range between 800–1000 °C in Ti-Al-Nb ternary alloys [54,55], with two different pathways in the Ti-Al-Nb ternary system—α₂ ↔ O and β_o ↔ O—depending on the composition and treatment process [17,56]. One pathway occurs in the α₂ phase in low-Nb-containing TiAl alloys (≤12.5 at% Nb), and the other pathway occurs in the β_o phase in high-Nb-containing alloys (~25 at% Nb) [17,56,57]. The transformation of the O phase from the β_o phase is considered to occur through a martensitic transition [57]. Therefore, in this work, the transus line in the range of 18–30 at% Al at approximately 850 °C in the 8Nb-TiAl system was confirmed to be the α₂ ↔ O transition, in agreement with the thermodynamic calculation.

For alloys with 32–36 at% Al, there was an extra γ precipitation from α₂, forming the ultrafine lamellar structure. As the Al concentration increased above 36 at%, the lamellar structure mainly formed through eutectoid reactions at 1150–1200 °C, with α ↔ β_o + γ and β_o ↔ α₂+γ, forming the lamellar colonies as shown in Figure 3, Figure 4 and Figure 6. Therefore, as mentioned in the thermal analysis, the experimental phase transformation behaviors within α-containing regions were in good agreement with the thermodynamic calculations, as shown in Figure 5a,b,d, indicating that the influence of experimental conditions on these regions is small. The dominant activity under these conditions is suggested to be the main attribution.

For alloys with Al concentrations lower than 30 at% (Samples S01–S05) or greater than 43 at% (Samples S14–S18), there were no β_o (LOB2 structure) and ω_o phases found at the ambient temperature, as indicated by XRD and SEM detections, while for Al concentrations in the range of 32–42 at% (Samples S06–S12), the ω_o phase and highly ordered β_o phase were detected simultaneously. As a result, the ω_o phase was preferentially formed in the highly ordered β_o phase.

The experimental phase equilibrium relations and transformation pathways of the 8Nb-TiAl alloys deviated significantly from the thermodynamic calculations, as shown in Figure 6. These deviations may have been associated with the treatment conditions, such as the heating/cooling rates, and/or insufficiencies in the thermodynamic database for high-Nb-containing TiAl systems [5,11,58]. Because of the actual preparation process, the experimental isopleth showed metastable phase behavior, such as the formation of ω_o from the retained β_o, rather than from α₂ as expected by the thermodynamic equilibrium behavior.

4.2. β_o Transformation

The β ↔ β_o transition temperatures and ordering procedure were in agreement with the studies of Clemens et al. [5] and Schmoelzer et al. [11], and were strongly affected by the compositions, as in [11,25]. The transition temperature differences between the experimental and calculation results may also have been related to the ordering of the β phase [59]. Crystallographic analysis has shown that the structure of the ordered β_o phase contains two sublattices: one occupied by Al (B site) and the other occupied by Ti atoms (A site). The two sublattices are randomly occupied with all species in the disordered β phase. As previously reported [60], Nb atoms should replace the Ti atoms in the structure. However, Leonard et al. [61] suggested that the site occupancy of Nb depends on the composition. The atomic occupancy and lattice parameter are slightly changed during the β ↔ β_o order-disorder transition.

The phase transition type is evidently reliant on the order of the isobaric-isothermal potential derivative (free Gibbs energy) [62], which exhibits an abrupt change at the transition position. In the first-order phase transitions (first derivatives), determined by the heat release or endothermic peaks in the DSC curve, there are abrupt changes in the specific volume and entropy. In the second-order phase transitions, such as order-disorder transitions, there are continuous changes in the specific volume and entropy, with no transition heat shown in the DSC curve, although there are abrupt changes in the second derivatives of the isobaric-isothermal potential, such as in the specific heat capacity, thermal expansion coefficient, and isothermal compressibility [62].

In this work, the ordered β_o phase formed from disordered β/α phases as the temperature decreased from the elevated temperatures. As mentioned in reference to the calculated C_p patterns, shown in Figure 5a, there was one narrow peak in Sample S03, indicating the order-disorder β ↔ β_o transition, and there were two narrow peaks in Samples S06–S15, corresponding to two types of order-disorder transitions, β ↔ β_o and α ↔ β_o + γ. The thermal analysis unambiguously detected the transformations of β_o from β/α, as shown in Figure 5b,d; these transformations are commonly considered second-order phase transitions. However, in addition to the peaks found in the C_p patterns, endothermic peaks were also detected in the DSC patterns for β ↔ β_o and α ↔ β_o + γ transitions, as shown in Figure 5b,d, indicating that the order-disorder transitions, β ↔ β_o and α ↔ β_o + γ, in the 8Nb-TiAl system should be defined as an intermediate transition type, between the first- and second-order phase transitions.

4.3. ω_o Transformation

The transformation of β_o phase to ω_o phase in Nb-containing TiAl alloys at 700–900 °C has been frequently observed [20,63,64,65]. The ω_o phase has been found to be transformed within the β_o matrix by homogeneous nucleation, or on the boundaries of β_o/γ phases by heterogeneous nucleation, of particles with nano- or micro-scale sizes during continuous heating/cooling processes [17,20,51,65]. The formation of ω_o phase is associated closely with the treatment process [51,66] and the chemical composition [20], and Nb has been suggested to act as the ω_o stabilization element [67].

As shown in Figure 5c, there was an endothermic peak and an abrupt change in the DSC and C_p patterns, respectively, indicating that, to some extent, the β_o ↔ ω_o transition was an ordering process. As mentioned by Stark et al. [20], the β_o phase exhibits a crystallographic orientation with the ω_o phase: $\left\{111\right\}{\mathsf{\beta }}_{\mathrm{o}}\parallel \left\{0001\right\}{\mathsf{\omega }}_{\mathrm{o}}$ or $\langle 1\overline{1}0\rangle {\mathsf{\beta }}_{\mathrm{o}}\parallel \langle 11\overline{2}0\rangle {\mathsf{\omega }}_{\mathrm{o}}$. According to this relationship, Figure 7 shows the atomic configuration of the ω_o phase with the Wyckoff positions and site occupancies, as well as the atom distribution in β_o phase as seen along its three-fold axis. During the β_o→ω_o transformation, the four stacking atomic layers in the β_o phase, Layers 1–4 in Figure 7a, merge into double layers in the ω_o phase, Layers 1–2 in Figure 7b. Furthermore, the site occupancy preference also changed, such that the Nb atoms moved into the 2a site, whereas the Ti and Al atoms move almost into the 2c and 2d sites, as shown in Figure 7. The lattice of the ω_o phase is slightly compressed along the $\langle 111\rangle {\mathsf{\beta }}_{\mathrm{o}}/\langle 0001\rangle {\mathsf{\omega }}_{\mathrm{o}}$ direction, whereas it stretches significantly along the $\langle 1120\rangle {\mathsf{\omega }}_{\mathrm{o}}$ direction [20].

5. Conclusions

(1) The isopleth section of the 8Nb-TiAl system (with 18–48 at% Al) was determined experimentally by DSC, SEM, and XRD methods, and compared with thermodynamic calculations (from CALPHAD). The experimental isopleth showed some deviations from the calculations, especially for β_o containing regions.

(2) The ordered β_o phase transformed from the disordered β phase in the range of 18.4–43.2 at% Al at 1200–1400 °C. The β_o phase also transformed through a eutectoid transition, α ↔ β_o + γ, from the α phase. The transition type of β ↔ β_o was considered to be an intermediate type between the first- and second-order phase transitions. In particular, the β_o phases were retained at the ambient temperature in the 8Nb-TiAl alloys under experimental conditions.

(3) The ordered ω_o phase transformed experimentally from the highly ordered β_o phase rather than from the α₂ phase or the β_o phase with a LOB2 structure. The experimentally detected ω_o formation occurred within 32–43 at% Al concentrations at approximately 850 °C, having a smaller composition range and higher transforming temperature than the calculated thermodynamic transition. The formation of the ω_o phase from the β_o phase was considered to be a further ordering process.