Precipitation Behavior of ω o Phase in Ti-37 . 5 Al-12 . 5 Nb Alloy

Mutual transformation between α2 andωo phases has been an interesting topic in recent years. In this study, martensitic α2 was obtained by air-cooling from 1250 ◦C in Ti-37.5Al-12.5Nb (at%) alloy while fourωo variants formed in the βo phase matrix during the cooling process. Nonetheless, only oneωo variant was observed at the periphery of the α2 plates in the βo phase and the orientation relationship between these two phases was [0001] α2//[1210] ωo; (1120) α2//(0002) ωo. Thin γ plates precipitated within the α2 phase and were thought to be related to the appearance ofωo phase. The redistribution of the compositions during the phase transformations was studied by energy dispersive X-ray spectroscopy analysis. The corresponding mechanisms of the phase transformations mentioned above are discussed.


Introduction
High Nb-containing TiAl (Nb-TiAl) alloys have been considered as potential materials for high-temperature applications due to their low density, high strength, good oxidation resistance, and creep properties [1][2][3].Recently, Stark et al. showed that the amount of ω o phase increased with the content of Nb in high Nb-TiAl alloy [4,5].Meanwhile, Nb is a β phase stabilizer that extends the β phase field and facilitates the ordered ω (ω o ) phase transitions in the β o phase in high Nb-TiAl alloys [6][7][8][9].Numerous studies have reported the ω o phase transformations in high Nb-TiAl alloys, indicating that the ω o phase is stable at 700-900 • C [10,11].However, these studies have mainly focused on the transition process between the β o and ω o phases [6][7][8][9][10][11][12][13] or the α 2 to β o phase [14][15][16].Recently, some reports concentrated on the precipitation of the ω o phase in α 2 laths during aging and studied the relationship between the α 2 and ω o phases [10,[17][18][19][20].To summarize, there are two different thoughts regarding the precipitation of ω o from α 2 phase.First, Huang et al. reported the perpendicular decomposition of coarse α 2 laths in Ti-44Al-8Nb-B alloy and suggested that the α 2 to β o (ω) transformation occurred after exposing at 700 • C in air for up to 10,000 h, indicating the occurrence of α 2 →β o →ω o transformation [17].Similar cases of α→β→ω transformation in titanium alloys had been reported by Vohra et al. [18] and Gupta et al. [19].Second, Bystrzanowski et al. observed that the applied stress could enhance the ω o precipitation and suggested that the ω o precipitation was directly transformed from the α 2 phase [20].Furthermore, Song et al. observed the direct α 2 to ω o phase transformation in Ti-45Al-9Nb alloy after aging at 900 • C [10].Although these studies discussed the transformation process and orientation relationships (ORs) between the ω o and α 2 phases, few reports focused on the nucleation sites of the ω o phase and the preferential ORs between the ω o and α 2 phases.Moreover, the nucleation behavior of ω o particles associated with α 2 phase in β o phase has scarcely been reported.In this work, the precipitation of ω o phase in Ti-37.5-12.5Nballoy was examined.The preferential OR between the ω o and α 2 phases was evaluated.The corresponding mechanisms were also discussed.

Materials and Methods
An ingot of the Ti-37.5Al-12.5Nb(at%) used in this study was prepared using induction levitation melting.The ingot was flipped and remelted three times to ensure compositional homogeneity.Table 1 lists the chemical compositions measured via wet chemical analysis.Specimens with sizes of 10 × 10 × 10 mm were cut from the center of the ingot by electric-discharge machining.The specimens were heat treated at 1250 • C for 2 h followed by air cooling with a cooling rate of approximately 20 K/s.The microstructures after heat treatments were examined using a Zeiss Supra 55 scanning electron microscope (SEM) in back-scattered electron (BSE) mode.Thin foils used for transmission electron microscopy (TEM) observation were prepared by twin-jet electro-polishing in a solution of 65 vol% methanol, 30 vol% butanol, and 5 vol% perchloric acid at 30 V and −30 • C. TEM analysis was conducted on a Tecnai G 2 F30 field emission transmission electron microscope operating at 300 kV.The compositions were obtained by energy dispersive X-ray spectroscopy (EDS) on TEM.Each parameter was an average value of more than five results measured at different locations.

Results
The actual composition of the alloy is Ti-37.0Al-13.0Nb, as obtained by chemical analyses, which is close to the nominal composition.Figure 1a shows the BSE image of the Ti-37.5Al-12.5Nballoy after air-cooling.The microstructure is composed of α 2 plates and β o matrix which can be identified by TEM as in Figures 1b and 2b.The well-defined dark lines (arrowed in Figure 1a) observed in α 2 plates are believed to be the midribs of the martensite, which is similar to the result for the α 2 phase form from the β phase by iced-brine quenching in Ti-44Al-4Nb-4Hf-0.1Si [21].It is difficult to distinguish whether the ω o phase exists in the β o region or not from the SEM image.Thus, the precipitation behavior of the ω o phase can be studied by using TEM. Figure 1b shows the bright-field TEM image of the air-cooled sample.Some particles with sizes of tens of nanometers distribute uniformly in the β o region.The corresponding selected area diffraction (SAD) pattern of the β o region is shown in Figure 1c, indicating that β o phase can readily transform to ω o phase during air cooling in this alloy.
Commonly, the observed ω o phase in high Nb-TiAl alloys can form from the "ω-collapse" in β o phase [6], i.e., the {111} β o layers "collapse" and the "-A-B-A-B-A-B-A-" stacking sequence in β o phase transforms into "-A-B/A-B-A/B-A-".There are four equivalent {111} β o layers so that four possible ω o variants can form in one β o grain.Considering the ORs between these four ω o variants and β o phase, the {0110} ω o diffraction spots of two ω o variants (denoted as "ω o1 " and "ω o2 " in Figure 1c) can be observed at 1/3{112} β o under the zone axes: <110> β o //<2110> ω o1 , ω o2 .However, the <0112> zone axes of "ω o3 " and "ω o4 " are parallel with <110> β o thus the diffraction patterns of these zone axes are overlapped completely.As a result, the intensities of the superposition spots of the ω o and β o phases are significantly increased in Figure 1c. Figure 2a shows the bright-field image of the α2 plates.The circled area in Figure 2a demonstrates an almost precipitate-free region except for some particles that precipitate at the boundary of the α2 phase.The corresponding SAD pattern of this area is shown in Figure 2b.Only one {0110} ωo spot exists at 1/3{112} βo under the same beam direction as Figure 1c, which indicates that only one ωo variant exists at the α2/βo boundary.The OR between these phases is obtained as: [110] βo//[0001] α2//[1210] ωo; (111) βo//(1120) α2//(0002) ωo Figure 2c is the corresponding dark field image taken by using the diffraction spot of the ωo variant, as circled in the SAD pattern in Figure 2b.It is indicated that the precipitates at the α2/βo boundary are of one kind of ωo variants.The compositions of the different phases (denoted in Figure 2a) were obtained by EDS equipped on TEM in Table 2.The results show that the ωo precipitates at the periphery of the α2 phase are more concentrated in Nb than βo-martix and α2 phase.This case can be interpreted in that Nb is a stabilizing element of the ωo phase rather than of the βo and α2 phases [22][23][24].Thus, ωo precipitations at the α2/βo boundary are more concentrated in Nb than βo and α2 phases because of the expulsion of Nb in the βo and α2 phases during the cooling process.Figure 2a shows the bright-field image of the α 2 plates.The circled area in Figure 2a demonstrates an almost precipitate-free region except for some particles that precipitate at the boundary of the α 2 phase.The corresponding SAD pattern of this area is shown in Figure 2b.Only one {0110} ω o spot exists at 1/3{112} β o under the same beam direction as Figure 1c, which indicates that only one ω o variant exists at the α 2 /β o boundary.The OR between these phases is obtained as: Figure 2c is the corresponding dark field image taken by using the diffraction spot of the ω o variant, as circled in the SAD pattern in Figure 2b.It is indicated that the precipitates at the α 2 /β o boundary are of one kind of ω o variants.The compositions of the different phases (denoted in Figure 2a) were obtained by EDS equipped on TEM in Table 2.The results show that the ω o precipitates at the periphery of the α 2 phase are more concentrated in Nb than β o -martix and α 2 phase.This case can be interpreted in that Nb is a stabilizing element of the ω o phase rather than of the β o and α 2 phases [22][23][24].Thus, ω o precipitations at the α 2 /β o boundary are more concentrated in Nb than β o and α 2 phases because of the expulsion of Nb in the β o and α 2 phases during the cooling process.Further studies on the interior of the α 2 phase reveal that there are a few stacking faults in it.Figure 4a shows some fine-scale planar defects within the α 2 phase, indicating certain phase transformations occur, as also suggested by the distortions at the boundary of the α 2 plate.Figure 4b is the HRTEM image of the interface between the α 2 and β o phase.The β o region and an ω o precipitate nucleated at the α 2 boundary are observed (the corresponding FFT images are shown in Figure 4c,d).It is worth pointing out that although the beam direction is [1120] α 2 and the ω o precipitate is under [0001] ω o direction, the OR between these two phases is as same as those obtained in Figures 2 and 3.The magnified image of Figure 4b is shown in Figure 4e, the atomic stacking sequence of the α 2 phase changes from "-ABAB-" to "-ABCABC-" (FCC-stacking) due to Shockley partial dislocations moving on alternate basal plane (0001) α 2 planes.Repeating this mechanism every two basal planes of the hexagonal matrix leads to the crystal structure change, thus, the stacking faults can act as the nucleus of the γ phase [25,26].Moreover, it has also been reported that γ phase precipitated in this alloy after annealing at 700 • C for 26 days [6].However, the γ phase observed in [6] consists of γ grains precipitated directly from the matrix and not thin γ laths transformed from the α 2 phase.Despite the different morphologies, these facts suggest that the γ phase is an equilibrium phase.Further studies on the interior of the α2 phase reveal that there are a few stacking faults in it.Figure 4a shows some fine-scale planar defects within the α2 phase, indicating certain phase transformations occur, as also suggested by the distortions at the boundary of the α2 plate.Figure 4b is the HRTEM image of the interface between the α2 and βo phase.The βo region and an ωo precipitate nucleated at the α2 boundary are observed (the corresponding FFT images are shown in Figure 4c,d).It is worth pointing out that although the beam direction is [1120] α2 and the ωo precipitate is under [0001] ωo direction, the OR between these two phases is as same as those obtained in Figures 2 and 3.The magnified image of Figure 4b is shown in Figure 4e, the atomic stacking sequence of the α2 phase changes from "-ABAB-" to "-ABCABC-" (FCC-stacking) due to Shockley partial dislocations moving on alternate basal plane (0001) α2 planes.Repeating this mechanism every two basal planes of the hexagonal matrix leads to the crystal structure change, thus, the stacking faults can act as the nucleus of the γ phase [25,26].Moreover, it has also been reported that γ phase precipitated in this alloy after annealing at 700 °C for 26 days [6].However, the γ phase observed in [6] consists of γ grains precipitated directly from the matrix and not thin γ laths transformed from the α2 phase.Despite the different morphologies, these facts suggest that the γ phase is an equilibrium phase.

Single ωo Variant Nucleated at the α2 Boundary
Transformation matrices are useful for calculating the ORs between the precipitation and matrix phase, which has been widely used in calculating the habit-planes and misorientations [27][28][29][30].As described above, four possible ωo variants exist in the βo phase.The ORs between the ωo variants and βo phase and between the ωo and α2 phases are calculated by using the transformation matrices.The transformation matrices B for βo to ωo phases are shown in Table 3 (see Appendix A.1 for details).According to the Burgers OR between the α2 and βo phases: {110} βo//( 0001 ) α2; <111> βo//<1120> α2, six βo variants can form from the α2 phase and the transformation matrices C for these variants are shown in Table 4 (see Appendix A.2 for details).The hypothesis is that the L is the transformation matrix from the crystallographic coordinate system to the orthogonal coordinate system (see Appendix A).By using Equation (1), the matrices T for α 2 to ω o phases are obtained, deriving the arbitrary parallel crystallographic directions between the α 2 and ω o phases.According to these results, only four ω o variants have good lattice matching and two ORs between the α 2 and ω o phases are obtained (see Appendix A.2 for details): ORII can be also expressed as <1010>α 2 //<2423>ω o ; {0002} α 2 //{0112} ω o by using a superimposed stereographic projection.Thus, both ORs calculated from the transformation matrices are the same as the results obtained by edge-to-edge matching calculation in Ti-45Al-9Nb alloy [10].According to the results, the selection of OR between the α 2 and ω o phases is essentially based on which ω o variant has good lattice matching with α 2 phase.It was reported that the misfit between the α 2 and ω o phases had a minimum value if ORI was formed [10].Thus, it is believed that only one ω o variant nucleated at the α 2 boundary.
Moreover, according to the EDS results in Table 2, the composition of the ω o particle nucleated at the boundary of the α 2 phase is more concentrated in Nb than the β o matrix.It is suggested that the ω o phase primarily nucleates at the boundary of the α 2 phase and enriches in Nb during growth.Thus, the untransformed area of the periphery of the α 2 phase is depleted in Nb so that the precipitate-free regions are observed.

Thin γ Plates Precipitated within the α 2 Phase
Because of the misfit matching between the α 2 and ω o phases, distortions at the interface are expected.The interplanar spacings of the (0002) α 2 and (1120) ω o planes measured by HRTEM and SAD software are 0.233 nm and 0.229 nm, respectively.That is a 1.7% mismatch in the interplanar spacing when ORI is formed between the α 2 and ω o phases.It means that α 2 phase may have an extra (0002) α 2 plane after successive stacking of 59 pairs of (1120) ω o and (0002) α 2 planes.The interplanar spacing of (111) γ is 0.232 nm is smaller than that of the (0002) α 2 but larger than (1120) ω o .Thus, the fine γ plates may relieve the distortion at the interface of the ω o and α 2 phases (Figure 4e).It has been reported that the precipitation of γ in α 2 is simply a HCP→FCC structure change which can be brought about if a/6 <1010> type Shockley partials move on alternate basal plane (0001) α 2 planes [31,32].As mentioned above, an extra (0002) α 2 plane exists after successive stacking of 59 pairs of (1120) ω o and (0002) α 2 planes.This case may cause the distorted-region separated along the α 2 and ω o interface.As a consequence, the sliding of the partial dislocations in the distorted-region can produce separated fine γ plates at certain intervals.Moreover, the interplanar spacing of 59 (0002) α 2 planes is approximate 13.7 nm, which is consistent with the average spacing of the γ laths, which is approximate 12 nm as obtained from Figure 4a.

Conclusions
In this work, the precipitation of ω o phase in Ti-37.5Al-12.5Nballoy was examined mainly by TEM.The ORs between different phases were calculated.The main results are summarized as follows: 1.
Only one ω o variant preferentially nucleates at the α 2 boundaries.This is because the minimum misfit exists at the α 2 /ω o interface if the OR between these two phases is: <2110> Precipitate-free regions are observed at the α 2 boundaries.EDS results indicate that the ω o precipitates are more concentrated in Nb than β o -matrix.The preferred nucleation of the ω o variant causes solute depletion surrounding the α 2 plates, which inhibits the nucleation and growth of new ω o precipitates in the un-precipitated regions.

3.
Thin γ plates precipitate within the α 2 phase.These fine γ plates can relieve the distortion caused by the mismatch at the α 2 /ω o interface.

Figure 1 .
Figure 1.Microstructure of the Ti-37.5Al-12.5Nballoy after annealing at 1250 °C for 2 h followed by air cooling: (a) back-scattered electron (BSE) image; (b) bright-field transmission electron microscopy (TEM) image; (c) the corresponding selected area diffraction (SAD) pattern of the βo region in (b).

Figure 1 .
Figure 1.Microstructure of the Ti-37.5Al-12.5Nballoy after annealing at 1250 • C for 2 h followed by air cooling: (a) back-scattered electron (BSE) image; (b) bright-field transmission electron microscopy (TEM) image; (c) the corresponding selected area diffraction (SAD) pattern of the β o region in (b).

Figure 2 .
Figure 2. TEM images of (a) bright-field image of the α2/βo boundary; (b) the corresponding SAD pattern at the α2/βo boundary; (c) dark-field image of the same area obtained by taking the spot of the ωo phase circled in the SAD pattern in (b).

Table 2 . 8 Figure 3
Figure 3 shows the High-resolution TEM (HRTEM) image of the α2/βo interface obtained under [0001] α2 direction.The fast Fourier transformation (FFT) images of the ωo and βo areas are shown in Figure b,c respectively.It is demonstrated that ωo phases sized about a few tens of nanometers nucleate at the α2 boundary in Figure 3a.This indicates that the preferential ωo phase can nucleate at the boundary of α2 phase and keep a certain OR with α2 phase.

Figure 2 .
Figure 2. TEM images of (a) bright-field image of the α 2 /β o boundary; (b) the corresponding SAD pattern at the α 2 /β o boundary; (c) dark-field image of the same area obtained by taking the spot of the ω o phase circled in the SAD pattern in (b).

Figure 3
Figure 3 shows the High-resolution TEM (HRTEM) image of the α 2 /β o interface obtained under [0001] α 2 direction.The fast Fourier transformation (FFT) images of the ω o and β o areas are shown in Figure 3b,c respectively.It is demonstrated that ω o phases sized about a few tens of nanometers nucleate at the α 2 boundary in Figure 3a.This indicates that the preferential ω o phase can nucleate at the boundary of α 2 phase and keep a certain OR with α 2 phase.Further studies on the interior of the α 2 phase reveal that there are a few stacking faults in it.Figure4ashows some fine-scale planar defects within the α 2 phase, indicating certain phase transformations occur, as also suggested by the distortions at the boundary of the α 2 plate.Figure4bis the HRTEM image of the interface between the α 2 and β o phase.The β o region and an ω o precipitate nucleated at the α 2 boundary are observed (the corresponding FFT images are shown in Figure4c,d).It is worth pointing out that although the beam direction is [1120] α 2 and the ω o precipitate is under [0001] ω o direction, the OR between these two phases is as same as those obtained in Figures2 and 3.The magnified image of Figure4bis shown in Figure4e, the atomic stacking sequence of the α 2 phase changes from "-ABAB-" to "-ABCABC-" (FCC-stacking) due to Shockley partial dislocations moving on alternate basal plane (0001) α 2 planes.Repeating this mechanism every two basal planes of the hexagonal matrix leads to the crystal structure change, thus, the stacking faults can act as the nucleus of the γ phase[25,26].Moreover, it has also been reported that γ phase precipitated in this alloy

Figure 3 .
Figure 3. (a) High resolution TEM (HRTEM) image of the α2/βo interface, the corresponding fast Fourier transformation (FFT) images of ωo and βo areas denoted in (a) are shown in (b,c) respectively.

Figure 4 .
Figure 4. (a) TEM image of the α2 plate; (b) HRTEM image of the α2 boundary; (c,d) the corresponding FFT images transformed from the ωo and βo regions in (b); (e) the magnified image of (b) at the α2/ωo interface.

Figure 4 .
Figure 4. (a) TEM image of the α 2 plate; (b) HRTEM image of the α 2 boundary; (c,d) the corresponding FFT images transformed from the ω o and β o regions in (b); (e) the magnified image of (b) at the α 2 /ω o interface.

Table 1 .
Chemical composition of the as-cast material.

Table 2 .
Energy dispersive X-ray spectroscopy (EDS) results of different phases.