Phase Transformation Behavior of a β -Solidifying γ -TiAl-Based Alloy from Different Phase Regions with Various Cooling Methods

: The phase transformation behavior of Ti-42Al-5Mn (at.%) alloy from different phase regions with various cooling rates was investigated based on electron probe micro analyzer-backscattered electrons (EPMA-BSE). It is shown that β → α 2 (cid:48) takes place when this alloy is cooled at a high rate, such as water quenching (WQ), oil cooling (OC), from β single phase. With the decreasing cooling rate to air cooling (AC), β → α 2 (cid:48) is restrained and β → γ is promoted by forming γ platelets. The room-temperature microstructure is β o + α 2 when alloy cooled (WQ and OC) from ( β + α ) dual-phase. However, under AC, β → γ occurs and γ platelets form. It should be noted that α 2 → γ happens when this alloy cooled from 1180 ◦ C (> T eut ) by OC and AC, forming an incomplete lamellae ( α 2 / γ ) structure in the α 2 phase. However, when the alloy cooled from 1100 ◦ C ( <T eut ), α 2 / γ → β o,sec occurs and complete lamellae generates in α 2 phase.


Introduction
In general, Ti and Ti alloys are quite attractive for aerospace, medical, and industrial applications [1][2][3][4], of which gamma-titanium aluminide (γ-TiAl) has been considered as one of the topmost candidates for high-temperature structural material due to its low density and high specific modulus and strength [5][6][7][8][9]. However, the intricate processing behavior, particularly the poor ductility and the low hot workability appears to be the main obstacle for a wide industrial application of γ-TiAl alloys [9]. In 1997, β-solidifying γ-TiAl alloys (via β-phase solidification) were proposed by Naka [10], also called beta solidified gammalloys (BSG) by Kim in 2018 [11], arousing increasing attention due to their excellent workability and the appropriate solidification process without any peritectic segregation [12]. These BSG alloys usually have one or more β stabilizing elements, such as Mo, Cr, Mn, V, Nb and a lower Al concentration (42-44%) when compared to conventional gammalloys (CG) [11], which finally results in their solidification pathway going through the single β-phase region.
As for the CG alloys, typically four types of microstructures can be obtained by heat treatment around the gamma dissolution temperature (T γsolv ), namely fully lamellar (FL), nearly lamellar (NL), duplex (DP) and near gamma (NG) [13,14]. The solidification procedure from the liquid to the stable solid state for the CG alloys can be given as L→L + β→(β + α) or (L + α)→α→α + γ→α 2 + γ. In the case of BSG alloys, since they contain a sufficient amount of β-stabilizing elements, the solidification pathway appears to be more complicated, and the typical structures of CG alloys are also no longer applicable. Clemens et al. [15,16] defined three types of structures for Ti-43.5Al-4Nb-1Mo-0.1B (TNM) alloy (BSG), including NL + β 1 , NL + β 2 , and NL + γ. The three microstructure consist of α 2 /γ + β o + γ, but there are significant differences in their morphology. From the aspect of the γ phase morphology, the γ phase in the first two kinds of structure is platelet, and the γ phase in the latter structure is block. Compared with NL + β 1 and NL + β 2 , the difference between the two microstructures is that there are very few β o,sec and γ phases around the colony boundary for NL + β 1 , while there are plenty of β o,sec and γ phases for NL + β 2 .
Ti-42Al-5Mn, developed by Tetsui et al. in 2002 [17], is a typical BSG alloy, and it has the advantage of low cost and great workability, and seems to be a promising material for automotive industrial applications. Since then, Tetsui et al. reported some studies on this alloy from 2002 to 2013 [18][19][20], but their investigations were still limited to smelting and forging. Our team has also conducted more in-depth research on Ti-42Al-5Mn in recent years. Our previous research confirmed its solidification pathway from the liquid to room temperature as the following: liquid→liquid + β→β→β + α→β + α + γ→β o + α 2 + γ→β o + γ + α 2 /γ→β o + γ + α 2 /γ + β o,sec , with the T β = 1311 • C, T γsolv = 1231 • C, T β→βo/α→α2 = 1168 • C, T eut = 1132 • C and T α2/γ→βo,sec ≈ 1120 • C [21]. On this basis, the continuous cooling transformation curve (CCT) of the forged alloy was examined and a CCT diagram of the alloy was established [22]. It has been found that five kinds of phase transformation exist during continuous cooling from 1300 • C (forging temperature) with the cooling rates of 0.1 • C/s, 0.5 • C/s, 2 • C/s, 10 • C/s, 50 • C/s, and 200 • C/s, including β→α, β→γ, β→β o /α→α 2 , α 2 →α 2 /γ and α 2 /γ→β o,sec . Despite this, the study of the alloy is still at an initial stage; the corresponding microstructure evolution and phase transformation are not well understood. In order to exert the advantages of the alloy, it is overwhelmingly indispensable to grasp their phase transformation behavior.
In the present study, the phase transformation behavior of as-casted Ti-42Al-5Mn alloy from different phase regions (β, α + β, α + β + γ) with various cooling rate was further investigated. The work aimed to clarify the phase transition characteristics of Ti-42Al-5Mn in order to provide fundamental data, knowledge and information (DKI) for BSG alloy development.
The samples with size of Φ8 mm × 8 mm were cut along the edge of the foot section of the as-cast ingot. Subsequently, each sample was put into a separate sealing tube filled with argon atmosphere to avoid oxidation, and the sealing tube was then placed in the muffle furnace. In order to ensure the consistency of each experiment, the samples were first heated to 1340 • C to ensure the α 2 /γ lamellae and γ grain could be completely dissolved into the β phase. After holding for 30 min, samples were cooled by water quenching (WQ), oil cooling (OC), and air cooling (AC), and the decreasing ordering in cooling speed was WQ, OC, AC [23]. The other specimens were then furnace cooled (FC, valued about 5 • C·min −1 [24]) to the target temperatures of 1260 • C (α + β), 1180 • C (α + β + γ, close to T eut ), and 1100 • C (α + β + γ). After holding for 30 min, these samples were also cooled with the WQ, OC, and AC methods, as illustrated in Figure 1. The microstructures of the specimens were characterized by electron probe micro analyzer (EPMA, JXA-8230, Northeastern University, Shenyang, China) with back scattered electron (BSE) and wave dispersive spectroscopy (WDS) mode.  Figure 2 shows the microstructure characteristics of the samples subjected to 1340 • C with different cooling methods. In this paper, β o (ordered body-centered cubic (bcc) B2 structure) and α 2 (ordered hexagonal D0 19 structure) are used to describe the room-temperature phase composition. The reason is that the disorder-order phase transformation shows the least sensitivity to the cooling rate in γ-TiAl alloy, namely a high cooling rate could not suppress the nucleation for the transformation [25]. Based on our previous research [21], there should be only pure β phase at 1340 • C since the Ti-42Al-5Mn is in the β single-phase region. However, from Figure 2a, some needlelike martensitic structures, namely α 2 [26] was observed in the above WQ microstructures. As described in reference. [22], the martensitic phase is considered to be the decomposition of β phase by β→α 2 . Moreover, upon quenching from the single β-phase field region a martensitic transformation was reported in "high-β" bearing γ-TiAl based alloys, such as TiAl-Mo [26], TiAl-V [27] or TiAl-Nb-Hf [28].

Microstructure Cooled from 1340 • C
When cooled with the OC method (Figure 2b), the martensitic phase also existed. This means that the β→α 2 cannot be restrained even though the cooling rate decreases to OC. It can also be seen that the size of α 2 is obviously coarsened due to the relatively lower cooling rate than WQ. Besides, in this situation, some fine γ (ordered face-centered tetragonal L1 0 structure) phase named γ-platelets (γ p ) [29] are formed within the β o phase, which is the decomposition of the β phase by β→γ.
With the cooling rate decreasing to AC (Figure 2c), it can be found that the needlelike martensitic structure disappeared, which suggests that the β→α 2 is completely restrained at this cooling rate. The β o phase is mainly distributed in parallel strips which are quite different from the above samples.
The α 2 phase, defined as the supersaturated α 2 -grain [30], was found between the parallel lines. It should be mentioned that the nucleated γ phase from β o was also very fine, which is similar to the OC sample.
3.2. Microstructure Cooled from 1260 • C Figure 3 presents the microstructure characteristics of the samples subjected to 1260 • C with different cooling methods. As can be seen in Figure 3a, in comparison with Figure 2a, plenty of supersaturated α 2 -grains are precipitated from the single β phase when the sample was cooled from 1340 • C to 1260 • C by furnace cooling. Both the WQ and OC samples consisted of β o and α 2 , and no martensitic α 2 phase was detected. This indicates that the β→α 2 transformation was not only affected by the cooling rate but also influenced by the temperature before cooling.
When it cooled with the AC method (Figure 3c), the microstructure consisted of β o , α 2 and γ phases. Lots of fine γ p were nucleated from the β phase, and the γ phase grew in a staggered manner in the β phase, not parallel, as in Figure 2c.  Figure 4 shows the microstructure characteristics of the samples subjected to 1180 • C with different cooling methods. As can be seen in Figure 4a, in comparison with Figure 3a, plenty of globular γ grains named γ g [29] with a median size of 15 µm were nucleated from the β phase. In the α 2 region, a few dispersive γ laths named γ l also formed. With decreasing cooling rates to OC and AC, the sizes of the γ g and β o phases were all intensively increased to a median size of 40 µm. Moreover, the amount of γ l phase in the supersaturated α 2 -grains increased as well. The incomplete α 2 /γ lamellae called (α + γ) still formed for these three samples, which might be due to the 1180 • C being relatively higher than the T eut (~1140 • C) [21]. Figure 5 shows the microstructure characteristics of the samples subjected to 1100 • C with different cooling methods. As can be seen in Figure 5a, in comparison with Figure 4a, the complete α 2 /γ lamellae formed. It should be mentioned that some fine β o,sec phases were also nucleated in the lamellae due to the 1100 • C in the temperature range of the α 2 /γ→β o,sec transformation. When the cooling rates decreased to OC and AC (Figure 5b,c), the amount of β o,sec phases increased. The lamellar spacing and γ grain size also significantly increased with the decreasing cooling rate.

Discussion
As stated above, the phase transformation behavior of the Ti-42Al-5Mn alloy from different phase regions with various cooling rate was obtained, as shown in Table 1. It can be seen that the influence of the temperature and cooling rate on the phase transformation and microstructure of this alloy at room temperature was very significant. For the same phase region, with the decrease of the cooling rate, the microstructure of the sample coarsened, and the types of phase transitions increased. This was mainly due to the greater time for the lower cooling rate, which favors more phase transitions. It should be mentioned that the γ morphology at room temperature was also closely related to the temperature and cooling rates. Table 2 shows the corresponding relationship between temperature, cooling rate and γ' morphology. As can be seen from this table, three types of γ' morphology can be identified as stated in Section 3, including γ grain or globular (γ g ), γ platelet (γ p ), and γ lath (γ l ), which are illustrated in Figure 6. The γ g with block morphology, was directly nucleated from the β phase. It was found that when the holding temperature was below T γsolv , the γ g could be obtained under the present cooling methods, such as WQ, OC and AC. In the case of γ p , it was also directly formed in the β phase. However, for this structure, the sample only cooled by AC could be obtained when the holding temperature was above T γsolv . These fine γ p phases are usually dispersed in the β region. The third one is γ l , which was precipitated from the α 2 /γ lamellae, and its formation condition was similar to γ g , i.e., the γ l phase was almost accompanied by the formation of the γ g phase. It was ascertained that the treating temperature below T γsolv can promote the formation of γ g and γ l , and inhibit the appearance of γ p . Whereas, the temperature above T γsolv can restrain the nucleation of γ g and γ l , and when combined with a relative low cooling rate, for example, the AC method, it will promote the appearance of γ p .
Although the γ g and γ l are all nucleated from the β phase, the size was different between these two phases. It can be seen from Figures 2-5, that the size of γ g was significantly larger than that of γ p , which might be the result of their different precipitation conditions. In this study, the fine γ p is often formed when cooled from a temperature that does not contain the γ phase region, such as the β or (β + α) phase region. In this situation, the γ p is precipitated from the β phase during the continuous cooling with a relatively slow cooling rate. Such a phase would have a smaller size compared to the γ g . However, once the temperature is lowered to the phase region including the γ phase, such as the (β + α + γ) phase region, the pre-precipitated γ phase furnace cooling from 1340 • C to the (β + α + γ) phase region would be further coarsened during the following cooling to form the γ g . Moreover, it can be also noted that the size of γ g would be obviously coarsened with a further decrease of the cooling rate. Hence, it is suggested that γ g and γ l can be clarified into the same phase, but the grain size between them is different due to their different precipitation behavior.
Note: "×" is not found; " √ " is found. Figure 6. The typical morphology of γ phase in this study.
In fact, the different sizes of γ at the α 2 /γ colony boundary have a significant effect on the properties of the γ-TiAl alloy. For the globular γ phase (γ g ), Mayer et al. [15] suggested that the γ g at the α 2 /γ colony boundary is a ductile phase itself due to its high slip lines, and it will help improve the ductility of the γ-TiAl alloy. In addition, Schwaighofer et al. [30] proposed that the presence of the γ phases at the colony boundary can intensively retard grain growth, but a high volume fraction of globular γ-grains would lead to a significant decrease of the yield strength at room temperature (RT) and creep strength at elevated temperature. Recently, in our research [31], we concluded that the γ phase at the α 2 /γ colony boundary benefit in terms of improving the ductility of Ti-42Al-5Mn alloy depends on its grain size. This means that the γ phase can only significantly improve the plasticity of the alloy if it has a suitable size, i.e., too large or too small will have more or less adverse effects on the ductility of γ-TiAl alloy plasticity. Hence, it is necessary to effectively control the grain size of the γ phase to ensure that the alloy has good comprehensive performance in the regulation of the alloy microstructure and mechanical properties.
In particular, the fully lamellar microstructure usually exhibits a superior combination of strength, creep resistance, fracture toughness and ductility [32], i.e., the γ phase at α 2 /γ colony boundary might be detrimental to the creep strength, then various heat treatments were applied to reduce the volume fraction of γ phase at the α 2 /γ colony boundary. For instance, for Ti-44.5Al-6.25Nb-0.8Mo-0.1B (BSG), Bolz et al. [33] found that when treating at 1270 • C, which is just slight below the T γsolv (1280 • C), and finally in combination with a same annealing system (800 • C/6 h/FC), the volume fraction of the globular γ-grain sharply reduced from 61.4% to 12.9% with the increased cooling rate from FC to OC, while that of the lamellar α 2 /γ colonies increased intensively from 23.5% to 84.8%. Moreover, they also clarified that the volume fraction of the globular γ-grain can decrease to nearly zero when treating at 1300 • C holding for 1 h and cooling with the AC method, in combination with the 800 • C/6 h/FC annealing system. This means that treating at a higher temperature (normally above T γsolv ) and combining with a higher cooling rate can help to suppress the globular γ-grain, which inversely promotes γ l formation under 800 • C/6 h/FC annealing system. Based on the present research, for Ti-42Al-5Mn, it was also seen that the γ' morphology can gradually transition from globular to platelet, and both γ g and γ p can disappear completely with an increased treating temperature above T γsolv combined with a high cooling rate (WQ, OC). Due to the absence of annealing treatment, for the sample accompanied with no γ g and γ p phases, large amounts of supersaturated α 2 phase would in turn form.
For the binary TiAl alloys (CG), it should be clarified that the globular γ-grain can emerge when the treating temperature is in the (α + γ) region. With temperature increases to above T γsolv , the full lamellar microstructure without any globular γ-grains can be formed. This suggests that the γ phase at the α 2 /γ colony boundary can be completely restrained if treating with a relatively high temperature above T γsolv both for the CG and BSG alloy.
In addition, it should be mentioned that the parent phase of the γ phase α 2 /γ colony boundary is not consistent in different γ-TiAl alloys. For the CG alloys, it has been claimed that this γ phase would form in the α-phase when the cooling rate is sufficiently high to suppress the formation of lamellar (α 2 /γ) microstructure [34,35]. For BSG alloys, however, this γ phase would only nucleate directly from the β phase [36]. In general, the transformation paths of γ either from β or from α are closely related to the solidification pathway of γ-TiAl alloy. For the CG alloys, the solidification procedure is: L→L + β→β + α or L + α→α→α + γ→α 2 + γ. At low temperatures, the β phase will completely transform into the α phase. In this situation, the γ phase can only emerge from α phase. In contrast, for the BSG alloys the solidification pathway is: L→L + β→β→β + α→β + α + γ→β o + α 2 + γ. The (β + α) phase will turn to (β + α + γ) with decreasing temperature. Based on the present results, the γ phase nucleates directly from the β phase while the α phase transforms into α 2 /γ lamella with a γ lath inside. The phase transformation from α 2 -Ti 3 Al(D0 19 ) of γ-TiAl(L1 0 ) not only includes the obvious change of the atom stacking order but also the re-adjustment of the localized chemical composition by long-distance diffusion [37]. This phase transformation is controlled by thermal activation and the reaction rate is relatively slow. In fact, Bolz et al. [33] have concluded that a higher driving force is apparently necessary for α→α 2 + γ (γ l ) than that of β→γ (γ g and γ p ). On the contrary, both β o -TiAl(B 2 ) [38] and γ-TiAl (L1 0 ) belong to the cubic structure and have the structure correlation shown in Figure 7. Moreover, the local chemical atom composition of β o is much closer to the γ phase. Hence, it can be seen that the γ phase is more likely to be precipitated from the β phase with the bcc structure than the α phase with the hcp structure.

Conclusions
The phase transformation behavior of Ti-42Al-5Mn alloy from different phase regions with various cooling rates was studied. The basic conclusions are the following: I.
The β→α 2 takes place when this alloy is cooled at a high rate (WQ and OC) from β single phase. With the decreasing cooling rate to AC, β→α 2 is restrained and the β→γ is promoted by the formation of γ platelets. II.
The room-temperature microstructure is β o +α 2 when this alloy is cooled by WQ and OC from (β + α) the dual-phase. However, under AC, the β→γ takes place and γ platelets form. III.
The α 2 →γ occurs when this alloy is cooled from the (β + α + γ) temperature, slightly higher than T eut (1132 • C), by WQ, OC and AC, forming incomplete lamellae (α 2 /γ) structures in the α 2 phase. In this situation, plenty of globular γ grains rather than γ platelets are nucleated from the β phase, and the size of γ grains increased intensively with the decreasing cooling rate. IV.
When the alloy cooled from 1100 • C (<T eut ), the α 2 /γ→β o,sec occurs and complete lamellae are generated in the α 2 phase.
Funding: This research was funded by National Natural Science Foundation of China for providing financial support during this study (Project No. 51801210).