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Article

Detrimental Effects of βo-Phase on Practical Properties of TiAl Alloys

by
Toshimitsu Tetsui
1,* and
Kazuhiro Mizuta
2
1
National Institute for Materials Science, Tsukuba 305-0047, Ibaraki, Japan
2
AeroEdge Co., Ltd., Ashikaga 329-4213, Tochigi, Japan
*
Author to whom correspondence should be addressed.
Metals 2024, 14(8), 908; https://doi.org/10.3390/met14080908
Submission received: 15 July 2024 / Revised: 2 August 2024 / Accepted: 7 August 2024 / Published: 9 August 2024
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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Abstract

:
The TNM alloy, a βo-phase-containing TiAl alloy, has been withdrawn from use as a last-stage turbine blade in commercial jet engines as it suffered frequent impact fractures in service, raising doubts regarding the necessity of the βo-phase in practical TiAl alloys. Here, we evaluate the practical properties required for jet engine blades for various TiAl alloys and investigate the effects of the βo-phase thereupon. First, we explore the influence of the βo-phase content on the impact resistance and machinability for forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys; the properties deteriorate significantly at increasing βo-phase contents. Subsequently, two practical TiAl alloys—TNM alloy and TiAl4822—were prepared with and without the βo-phase by varying the heat treatment temperature for the former and the Cr concentration for the latter. In addition to impact resistance and machinability, the creep strength is significantly reduced by the presence of the βo-phase. Overall, these findings suggest that the βo-phase is an undesirable phase in practical TiAl alloys, especially those used for jet engine blades, because, although the disordered β-phase is soft at high temperatures, it changes to significantly more brittle and harder βo-phase after cooling.

1. Introduction

Elements such as Nb, Cr, Mo, Mn, and W are the main additives in TiAl alloys and function as β-stabilizers. These additive elements are present in practical TiAl alloys that were developed early, such as TiAl4822 (Ti–48Al–2Nb–Cr (at. %; this unit is omitted hereafter) [1,2,3], 45, 47XD (Ti–45, 47Al–2Nb–2Mn–0.8vol% TiB2) [4,5,6], and DAT-TA2 (Ti–46.5Al–3.2Nb–0.8Cr–0.7Si–0.1C) [7]. Notably, the β-phase content in these initial TiAl alloys is minimal because of the low added amount.
From a practical standpoint, the β-phase has been actively used for ~20 years to forge TiAl alloys, which was previously considered impossible. At high forging temperatures, the β-phase transforms into a disordered bcc phase that significantly improves deformability and enables forming by forging. By contrast, after cooling, it changes into the ordered βo(B2)-phase of an intermetallic compound. Ti–42Al–5Mn [8,9,10], which is a representative alloy with high β/βo-phase content, can be readily deformed by hot forging during which the heated material is removed from the furnace and pressed at a high speed in the air using a die at a relatively low temperature, enabling the production of large parts [11].
In contrast, the TNM alloy (Ti–43.5Al–4Nb–1Mo–0.1B) [12,13,14], which is a representative forged TiAl alloy with low β/βo-phase content, is not conducive to high-speed deformation while cooling owing to its low β-phase levels. Therefore, hot forging cannot be readily used to manufacture products that require no defects; instead, these are manufactured by isothermal forging, in which a high-temperature die is used in a high-temperature chamber, with deformation conducted at an extremely low speed. The TNM alloy was once used in Pratt & Whitney’s PW1100G commercial aircraft jet engine [15]; however, it was discontinued and replaced with a common heat-resistant alloy owing to frequent impact fractures caused by collisions with flying debris inside the engine [16]. In contrast, other jet engines such as General Electric’s (GE) GEnx [15] and CFM International’s LEAP [17,18] use large amounts of TiAl4822 cast material with virtually no βo-phase and do not experience impact fracture or other problems. This extreme difference is undoubtedly related to the fact that the PW1100G engine, which is of a geared turbofan variety, exhibits greater collision energy with debris owing to its faster rotational speed than that of regular turbofan engines such as LEAP. Nevertheless, this variation was naturally anticipated in the design process; in other words, the differences in the material properties between TNM alloy and TiAl4822 are potentially responsible for their extreme disparity in practical use. The most significant difference between the two is the presence or absence of the βo-phase.
Therefore, the present study was aimed at evaluating the need for the β/βo-phase from a practical viewpoint and clarifying whether the β/βo-phase offers any value for TiAl alloys other than improving forgeability. The assumed product is the last-stage turbine blades of a jet engine. The critical strength properties for this product include creep strength and high-/low-cycle fatigue strength at elevated temperatures. Concerning high-cycle fatigue strength, Sallot et al. [19] have already reported that the βo-phase preferentially oxidized during prior atmospheric heating reduces fatigue strength. In terms of low-cycle fatigue strength, Nakatani et al. [20] reported that the preferential oxidation of the βo-phase during testing reduces fatigue strength.
In addition to these strength properties, two other attributes are essential for fabricating high-performance and reliable jet engine blades. The first characteristic, impact resistance, is vital because of the disastrous experience with the TNM alloy. In essence, high impact resistance is critical to ensure that TiAl blades are reliable during operation and not easily damaged by collisions with debris. The second aspect, machinability [21], is important to reduce manufacturing costs, given that TiAl blades for jet engines are currently manufactured from simply shaped ingots cast in permanent molds [22] or from investment cast blade materials with large amounts of excess wall [23]. Even forged materials (TNM alloy) that have been used in the past have used blade-like materials with large amounts of excess wall [24]. Therefore, they must be processed into thin product blades by high-volume machining, and excellent machinability is required to reduce costs.
Considering these aspects, at first, Ti–43.5Al–xCr ternary forged alloys and Ti–46.0Al–xCr ternary cast alloys with different βo-phase contents were prepared in this study by changing the amount of Cr, and the effects of the βo-phase on impact resistance and machinability were investigated. Subsequently, two practical TiAl alloys—TNM alloy and TiAl4822—with and without the βo-phase were prepared by varying the heat treatment temperature for the former and the Cr concentration for the latter, and the effects of βo-phase on creep strength were evaluated in addition to impact resistance and machinability. Finally, the obtained results were scrutinized to determine whether the β/βo-phase in practical TiAl alloys, especially those used for jet engine blades, offers benefits other than improving forgeability.

2. Materials and Methods

2.1. Materials

Ti–Al–xCr ternary alloys with varying βo-phase contents were prepared by altering the Cr concentration to 2.0, 2.5, 3.0, 3.5, and 4.0 in the 43.5Al forged alloys and 46.0Al cast alloys. Furthermore, modified TNM alloys with and without the βo-phase were prepared by varying the post-forging heat treatment temperature with the nominal composition of the TNM alloy (Ti–43.5Al–4.0NB–1.0Mo–0.1B). For the modified TiAl4822 alloys, the nominal Al concentration is 48; however, the material commonly used today has a slightly reduced Al concentration. Therefore, samples of modified TiAl4822 with and without the βo-phase were prepared by changing the content of Cr slightly (1.79, 1.91, and 2.14) in Ti–47Al–2.0Nb.
The raw materials—Ti sponge, Al pellets, Cr grains, Nb thin plates, and Mo and TiB2 powders—were induction melted in a CaO crucible in an atmosphere that was replaced by Ar after vacuuming. The weight of one batch was ~850 g. After all the raw materials were melted, the resulting molten metal was kept for 3 min with the melting power applied and then poured into a metal mold that was divided into two parts. For the Ti–43.5Al–xCr ternary forged alloys and modified TNM alloys, cylindrical ingots with approximate dimensions of ø55 × 90 mm were produced by casting the molten metal into a metal mold with a cylindrical cavity. For the Ti–46.0Al–xCr ternary cast alloys and modified TiAl4822 alloys, the molten metal was cast in the shape of a material with a flat-plate-type specimen collection section measuring 60 × 90 × 16 mm with a feeder head above it.
Melting in a CaO crucible cannot prevent the increase in oxygen concentration caused by oxygen contamination from the crucible [25]; therefore, the oxygen concentration will be higher than that of the existing TiAl alloy materials used for jet engine blades, which are typically melted in a water-cooled Cu crucible [22]. Therefore, in this study, 0.15 mass% Ca was added in the form of an Al–10mass% Ca alloy for deoxidization. The added Ca combines with oxygen in the molten metal to form CaO, which is released from the molten metal surface in the form of fumes during melting. Furthermore, excess Ca is discharged as Ca fumes. Consequently, the oxygen concentration of the TiAl alloy material after casting is less than 0.1 mass%, and the Ca concentration is less than 0.05 mass%. The details of this process will be provided in a separate report. As an example, the results of the chemical analysis of the modified TiAl4822 alloy produced in this study are shown in Table 1. The impurity levels are not significantly different from those of the TiAl4822 material melted in a water-cooled Cu crucible.
Hot forging was used in this study instead of isothermal forging for cost reasons. To manufacture the Ti–43.5Al–xCr ternary forged alloys and modified TNM alloys, the cylindrical ingots were coated with ATP610 [26] (Advanced Technical Products Supply Co., Inc., West Chester, Ohio, USA) for lubrication and insulation during forging. Each ingot was held in a furnace heated to 1330 °C for ~1 h, removed from the furnace, and then pressed in the direction of the ingot height in a single forging operation using a 300-ton hydraulic press, forming a pancake with an approximate diameter and thickness of 130 mm and 16 mm, respectively. Cracking occurred at the periphery of the forged material in the Ti–43.5Al–xCr ternary forged alloys with low Cr concentrations and in the modified TNM alloys; however, this did not affect the subsequent evaluations as they were conducted using the interior parts.
For the Ti–43.5Al–xCr ternary forged alloys, a heat treatment protocol of 1280 °C/5 h/furnace cooling (FC) was applied. For the TNM alloy, two-stage heat treatment is typically performed after forging [12]: air cooling after high-temperature holding, followed by FC after low-temperature holding. In the present study, samples (TNM alloys of nominal composition) with and without the βo-phase were prepared by changing the heat treatment temperature of the first stage, with values of 1177, 1207, and 1237 °C selected based on preliminary heat treatment tests. After forging, the TNM alloy pancakes were maintained at each of these temperatures for 3 h and then air-cooled. Subsequently, in the second stage, all pancakes were subjected to the standard 850 °C/6 h/FC protocol used for TNM alloy. For the Ti–46.0Al–xCr ternary cast alloys and modified TiAl4822 alloys, casting defects were eliminated by performing hot isostatic pressing (HIP) [27] using the 1200 °C/4 h/186 MPa protocol commonly employed for TiAl4822 cast materials.

2.2. Methods for Evaluating Various Properties

2.2.1. Impact Resistance

The Charpy impact test is the simplest and most realistic method for evaluating the impact resistance of industrial materials and can be easily conducted at high temperatures; therefore, it was used in the present study to determine impact resistance. After heat treatment and HIP, the forged and cast materials were machined on their front and back surfaces to achieve a uniform thickness of 10 mm. Subsequently, prismatic specimens with dimensions of approximately 10 × 10 × 55 mm were produced by further processing.
The impact resistance of TiAl alloys is significantly lower than that of normal metallic materials; therefore, unnotched specimens and a small 30J hammer were used to measure accurately the differences between the alloys. Room temperature (RT) and 700 °C, which is reportedly close to the maximum service temperature of TiAl alloys intended for jet engine blades, were selected as the test temperatures. For the 700 °C tests, an electric furnace was installed next to the testing machine. Each specimen was placed in the furnace at 700 °C for ~1 h, removed from the furnace, and then quickly placed in the Charpy impact testing device. Notably, the sample removed from the electric furnace was tested within 5–10 s. All the alloys were tested approximately 10 times at RT and 700 °C, and the impact resistance of each alloy was compared with the mean value of the absorbed energy. Moreover, the standard deviation was calculated.

2.2.2. Machinability

Materials with less tool damage allow faster machining speeds and therefore reduce costs. Therefore, the machinability of each alloy was evaluated based on the magnitude of the tool damage. The details of the cutting tests are provided elsewhere [28]. Briefly, square K10 carbide inserts were mounted on a seven-blade cutter (ø100 mm diameter between circumferentially mounted insert tips), and dry milling was performed using a face milling machine. Although the tool coating and lubrication conditions (dry, wet, etc.) have a significant effect on the machinability, the tests were aimed at comparing the relative machinability of the different alloys; therefore, uncoated carbide inserts were used under dry conditions.
The cutting tests were performed after the complete removal of the alteration layer on the surfaces of the forged and cast materials, which were formed by heat treatment and HIP. The milling conditions were constant with a cutter rotation speed of 130 revolutions per minute (peripheral speed of the insert tip of 0.680 m/s), feed rate of 4.97 mm/s, and cutting depth of 0.2 mm per cut. This was repeated until approximately 100 g of sample volume was removed.
Wear-induced tool weight loss was evaluated by measuring the change in weight for the sum of seven inserts before and after the cutting tests using a high-resolution analytical balance with a minimum measurement value of 0.00001 g. The obtained tool weight loss was normalized by the weight of the sample removed to compare the relative effect of each alloy on tool damage (that is, the machinability of each alloy).

2.2.3. Other Properties

The microstructure of each alloy was examined using backscattered electron images acquired in compositional mode. In addition, the area ratio of each phase was determined by analyzing these images with image-processing tools. To this end, a section corresponding to the center of the plate thickness of the forged and cast materials was cut and polished. For modified TNM and TiAl4822 alloys, creep tests were conducted at 750 °C/225 MPa using round bar specimens (diameter = 4.0 mm) with a 20.0 mm long gauge section, and the relationship between test duration and creep strain was measured. Additionally, certain alloys were subjected to Vickers hardness tests using a 294 N load. The five-point average was obtained, and brittleness was simply assessed by observing cracks formed near the indentation.

3. Results

3.1. Ti–Al–xCr Ternary Forged and Cast Alloys

In the five forged materials—that is, Ti–43.5Al–xCr, with x = 2.0, 2.5, 3.0, 3.5, and 4.0—the circumferential cracks decreased with increasing Cr content (Figure 1). This was due to the increase in the amount of the disordered β-phase with increasing Cr incorporation at heating temperature (1330 °C) before forging. Therefore, the beneficial effect of the disordered β-phase in improving forgeability was confirmed.
The microstructures of the forged Ti–43.5Al–xCr alloys (x = 2.0, 2.5, 3.0, 3.5, and 4.0) were examined after the 1280 °C/5 h/FC heat-treatment protocol (Figure 2). Three phases were identified: the main phase is the γ-phase, the second phase is the α2-phase, and the third phase is the βo-phase. The alloys with 2.0Cr and 2.5Cr exhibited a fully lamellar structure and lacked the βo-phase, whereas the remaining three alloys contained the βo-phase, whose magnitude escalated with increasing Cr addition.
The microstructures of the cast Ti–46.0Al–xCr alloys (x = 2.0, 2.5, 3.0, 3.5, and 4.0) subjected to the 1200 °C/4 h/186 MPa HIP treatment were subsequently analyzed (Figure 3). The same types of phases were present as in the forged alloys. The 2.0Cr-doped alloy exhibited a fully lamellar structure and contained a small amount of the α2-phase but lacked the βo-phase. Contrarily, the remaining four alloys contained the βo-phase, whose extent increased when the incorporated Cr increased.
The relationship between the area ratio of the α2-phase (the second phase) and that of the βo-phase (the third phase) vs. the Cr concentration—obtained by processing the backscattered electron images of the forged Ti-43.5Al-xCr and cast Ti-46.0Al-xCr alloys—is plotted in Figure 4. For both alloys, the area ratio of the α2-phase decreases and that of the βo-phase increases with increasing Cr concentration. The α2-phase is mostly present in the lamellar structure, which indicates that the ratio of the lamellar structure decreases with increasing Cr concentration. A comparison of the forged and cast alloys reveals that the ratio of the βo-phase in the former is higher due to the lower Al concentration.
The mean absorbed energy and the corresponding standard deviation were determined for the forged and cast Ti–Al–xCr ternary alloys from the Charpy impact tests conducted at RT and 700 °C (Table 2). In addition, the relationship between the area ratio of the α2-phase (the second phase) and that of the βo-phase (the third phase) and the mean absorbed energy is shown in Figure 5. These results suggest that the factors responsible for the decrease in the mean absorbed energy at both RT and 700 °C in the forged and cast Ti–Al–xCr ternary alloys are the increase in the βo-phase and the decrease in the α2-phase (decrease in the lamellar structure ratio). The decrease in the impact resistance due to the decrease in the lamellar structure ratio was already confirmed in our previous study [21,28]. In addition to this, it is clear from this study that as the βo-phase increases, the impact resistance decreases. Regarding the test temperatures used, the mean absorbed energy was higher at 700 °C than at RT for both the forged and cast alloys, possibly owing to the increased ductility while maintaining strength at 700 °C.
The cutting test results of the forged and cast Ti–Al–xCr ternary alloys were analyzed to establish a relationship between tool weight loss and Cr concentration (Figure 6); the tool weight loss was normalized to the weight of the sample removed by cutting. The findings, in combination with the results shown in Figure 2, Figure 3 and Figure 4, suggest that the extent of tool wear is independent of the α2-phase (lamellar structure) ratio and increases with increasing βo-phase content, confirming the negative effect of the βo-phase on the machinability of the TiAl alloys.

3.2. Modified TNM and TiAl4822 Alloys

The microstructures of the modified TNM alloys in nominal composition were examined using different heat treatment temperatures in the first heat treatment stage after hot forging at 1330 °C (Figure 7). In the alloy heat treated at 1177 °C, the amount of the βo-phase was high, whereas, in the alloy heat treated at 1207 °C, the microstructure was a mixture of lamellar structure and the α2-, γ-, and βo-phases, which resembles the previously reported standard microstructure of TNM alloy [12]. Moreover, in the alloy heat treated at 1237 °C, the βo-phase almost disappeared and changed to a mixed microstructure containing a lamellar structure and the α2- and γ-phases.
The microstructures of modified TiAl4822 alloys (Ti–47Al–2.0Nb-xCr) with varying Cr contents (1.79, 1.91 and 2.14) were examined after subjecting them to HIP at 1200 °C/4 h/186 MPa (Figure 8). The 1.79Cr and 1.91Cr alloys exhibited a duplex structure with lamellar structure and γ-grains, in addition to a small amount of the α2-grains. The 2.14Cr alloy exhibited a similar microstructure but with a βo-phase.
Subsequently, the average absorbed energy and standard deviation were determined for the modified TNM and TiAl4822 alloys from the Charpy impact tests conducted at RT and 700 °C (Table 3). Among the modified TNM alloys, the samples heat treated at 1177 and 1237 °C exhibited the lowest and highest absorbed energy, respectively, at both RT and 700 °C. In other words, the decrease in impact resistance with increasing βo-phase content was estimated by a comparison of these findings with those shown in Figure 7. Among the modified TiAl4822 alloys, the absorbed energy of the alloy with 2.14Cr was the lowest both at RT and 700 °C, whereas that of the alloys with 1.79Cr and 1.91Cr did not differ significantly. These results, in combination with those shown in Figure 8, indicate that the impact resistance decreases when the βo-phase is present.
Regarding the test temperatures, the increase in the absorbed energy of the modified TNM alloys at 700 °C was smaller than that at RT. Conversely, this increase was considerably greater for the modified TiAl4822 alloys presumably because the amount of the γ-phase in the TiAl4822 alloys was significantly higher than in the TNM alloys, resulting in improved ductility with maintained strength at 700 °C. However, the relationship of the heat treatment temperature for the modified TNM alloys or the amount of Cr added in the modified TiAl4822 alloys with the superiority in terms of impact resistance was the same as that at RT. Thus, the βo-phase continued to exert a detrimental effect on the impact resistance of practical TiAl alloys at the operating temperature of jet engine blades.
The results of the cutting tests performed on the modified TNM and TiAl4822 alloys were subsequently analyzed (Figure 9). For the modified TNM alloys, the tool weight loss (normalized to the weight of the sample removed) was plotted against the first-stage heat treatment temperature. The results, in combination with those shown in Figure 7, indicate that the tool wear magnitude increases at increasing amounts of the βo-phase, confirming the undesirable influence of the βo-phase on the machinability of the modified TNM alloys.
For the modified TiAl4822 alloys, the relationship between the Cr concentration and the scale of the tool weight loss was investigated. The tool wear of the modified TiAl4822 alloys was significantly lower than that of the modified TNM alloys. This was probably due to the presence of significantly higher and lower proportions of the softer γ-phase and the harder α2-phase, respectively, in the modified TiAl4822 alloys than those in the modified TNM alloys. Among the modified TiAl4822 alloys, the tool wear of the 2.14Cr alloy was slightly higher than that of the 1.79Cr and 1.91Cr alloys. Essentially, these results, in combination with those shown in Figure 8, imply that the machinability decreases in the presence of the βo-phase.
The relationship between test duration and creep strain was examined by subjecting the modified TNM and TiAl4822 alloys to creep tests at 750 °C/225 MPa (Figure 10). Among the modified TNM alloys, the alloys heat treated at 1177 and 1237 °C exhibited the lowest and highest creep strengths, respectively. These results demonstrate that the creep strength decreased significantly at increasing βo-phase content. Among the modified TiAl4822 alloys, the alloy with 2.14Cr exhibited a lower creep strength than those of the alloys with 1.79 and 1.91Cr. Based on this result, we could conclude that a reduction in creep strength was observed owing to the presence of the βo-phase.
The microstructures near the ruptured positions in the creep test specimens were examined for the modified TNM alloy heat treated at 1177 °C (which exhibited the lowest creep strength among the modified TNM alloys), and the modified TiAl4822 alloy integrated with 2.14Cr (which showed the lowest creep strength among the modified TiAl4822 alloys) (Figure 11). The results reveal creep voids concentrated in the βo-phase in both alloys. In other words, the microstructural analysis indicated that the creep strength of the βo-phase was lower than that of the lamellar structure, γ-phase, and α2-phase. Overall, these findings indicate that the βo-phase adversely affected the creep strength of practical TiAl alloys, in addition to impact resistance and machinability. Notably, as mentioned earlier, the preferentially oxidized βo-phase degrades high- and low-cycle fatigue properties at high temperatures [19,20]. In other words, the βo-phase can have a significant detrimental effect on the strength properties that are critical for jet engine blades.

4. Discussion

4.1. Summary of the Influence of Each Phase on the Practical Properties of TiAl Alloys

The basic phase of TiAl alloys is the γ-phase (TiAl phase). In practical TiAl alloys, the second phase, i.e., the α2-phase, is formed by reducing the Al concentration. Because most of the α2-phase exists in the form of a lamellar structure, the ratio of the α2-phase is proportional to that of the lamellar structure. On the other hand, the third phase, i.e., the β/βo-phase, is formed by adding more than a certain amount of β-stabilizing elements such as Cr and Mn, which are commonly added to practical TiAl alloys. Other factors affecting the properties of TiAl alloys include finer grain size and finer lamellar spacing. However, with regard to the former, it is difficult to obtain fine grains, especially in practical TiAl alloys for jet engine blades, because they must be heat treated or HIPed at high temperatures for long periods to ensure microstructural stability in service. On the other hand, the latter requires rapid cooling after high-temperature heat treatment, which requires special processes such as gas-fan cooling, resulting in higher costs. Therefore, the ratio of the α2-phase (lamellar structure) and that of the β/βo-phase is considered to have the greatest influence on the properties of practical TiAl alloys.
Based on the results obtained in this study and related references, we summarized the effects of the α2- and βo-phases ratios on each practical property in Table 4. These results indicate that the α2-phase (lamellar structure) often has a positive effect, whereas the βo-phase has only a negative effect.

4.2. Causes of Negative Effects of βo-Phase on Impact Resistance and Machinability

The reasons underlying the undesirable effects of the βo-phase on impact resistance and machinability were then investigated using a simple method. Essentially, in the Ti–Al–Cr ternary system, single-phase alloys were prepared based on the phase diagram at 1200 °C [31,32] and evaluated. Alloys with compositions of Ti–40.5Al–10.0Cr, Ti–41.0Al–1.0Cr, and Ti–50.0Al–2.0Cr, which were designed to contain only the βo-, α2-, and γ-phases, respectively, were prepared. The melting method was identical to that described above, and microstructural observations and Vickers hardness tests at 294 N were performed on the cast materials at RT after heat treatment; these involved maintaining the samples at 1200 °C for 10 h followed by water quenching.
The microstructures, cracks formed near the indentation caused during the Vickers hardness test, and Vickers hardness values at RT were investigated for these alloys (Figure 12). A small amount of the γ-phase coexisted in the alloy designed to exhibit a single βo-phase; however, the βo-phase was predominant. The Vickers hardness of this alloy (HV 493) indicated that it was extremely hard. Moreover, the observed cracking around the indentation implied that this alloy was remarkably brittle. The alloy designed to exhibit a single α2-phase contained a single α2-phase; its hardness (HV 369) was lower than that of the aforementioned alloy with an almost single βo-phase, and its cracking behavior was minor. The alloy intended to exhibit a single γ-phase also contained a single γ-phase; it exhibited the lowest hardness value (HV 200) among these alloys and showed only microscopic cracks around the indentation. These results indicate that although the disordered β phase is softer at high temperatures (forging temperature), it transforms into the βo-phase, which is significantly more brittle and harder than the α2- and γ- phases after cooling.
In the preceding discussion on impact resistance, the impact resistance of forged and cast Ti–Al–xCr ternary alloys, as well that of the practical TiAl alloys (modified TNM and TiAl4822 alloys), decreased at increasing βo-phase content, presumably owing to the brittleness of the βo-phase, as corroborated above. Furthermore, the machinability deteriorated at increasing βo-phase content for different alloys; this was also related to the fact that the βo-phase was significantly harder than the other phases, as verified above.
The α2-phase is less brittle and less hard than the βo-phase but more brittle and harder than the γ-phase; thus, it may cause negative effects similar to those of the βo-phase. Nonetheless, because the α2-phase is essentially present in TiAl alloys as a phase that forms lamellar structures, and a softer γ-phase exists between the thin α2-phase plates, the combined effect of the microstructure seemingly mitigates the brittleness and hardness of the α2-phase to a certain extent. However, unlike the α2-phase, the βo-phase cannot co-exist as a composite structure with the γ-phase but exists independently, suggesting that its brittleness and hardness cannot be mitigated by the microstructural factors.

4.3. Assessment of Various Existing TiAl Alloys Assuming the Harmfulness of βo-Phase

4.3.1. Cast TiAl Alloys

For cast TiAl alloys, there is no need for the disordered β-phase that is required for the forging process; therefore, the composition should be designed to prevent the formation of the detrimental βo-phase. Among the known TiAl cast alloys, TiAl4822 (Ti–48Al–2Nb–2Cr) [1,2,3]—which is practically used to fabricate jet engine blades—contains 2% Nb and 2% Cr, which exhibit small and large β-stabilizing effects, respectively. As confirmed in the present study, a slight increase in Cr concentration will lead to the formation of the harmful βo-phase. However, this rarely occurs with the current compositional specification ranges; therefore, this is not considered a major problem. Furthermore, it would be perfectly fine to limit the Cr content a little lower than the current specification. The 45, 47XD alloy (Ti–45, 47Al–2Nb–2Mn–0.8vol% TiB2) [4,5,6], which had been considered for fabricating jet engine blades for Rolls-Royce [33], is perfectly safe because it contains 2% Nb and 2% Mn, which has a lesser β-stabilizing effect than that of Cr. Additionally, DAT-TA2 (Ti–46.5Al–3.2Nb–0.8Cr–0.7Si–0.1C) [7]—which is being used to fabricate turbine wheels for passenger car turbochargers—has a slightly increased Nb content than that in the two aforementioned practical alloys; notably, this is perfectly acceptable given the low Cr content. In essence, the TiAl cast alloys developed and used in the distant past are excellent in that the addition of a small amount of β-stabilizing elements improves their properties and suppresses the formation of the harmful βo-phase.
In contrast, modern TiAl alloys such as TNB-V2 (Ti-45Al-8Nb-0.2C) [34] and IRIS (Ti-48Al-2W-0.08B) [35] contain many β-stabilizing elements. Only Nb, which has a small β-stabilizing effect, is added in the former; however, the amount incorporated is sufficiently large for forming the βo-phase. Meanwhile, 2% W—which has an extremely large β-stabilizing effect—is added to the latter, resulting in an abundant βo-phase content. Newer TiAl alloys, including these alloys, contain large amounts of additives mainly to improve strength and oxidation resistance; however, the practical properties required for jet engine blades, such as impact resistance and machinability, are considered to be significantly inferior to those of the practical TiAl alloys developed earlier owing to the presence of the βo-phase. These observations underscore the superiority of the previously developed TiAl alloys.
Regarding the additive elements, Nb was added mainly to improve the oxidation resistance of TiAl alloys; however, no significant differences in oxidation resistance have been observed between TiAl alloys in which 5% and 10% Nb were incorporated [36]. Therefore, large amounts of Nb need not be added; this can prevent βo-phase formation and reduce material costs. Additionally, W and Mo exhibit a large β-stabilizing effect; therefore, even small levels of incorporation will produce the βo-phase. Similar to Nb, W was added to improve oxidation resistance [37,38]; however, because W induces the same improvement in oxidation resistance as that of Nb but at considerably lower levels, it could act as a beneficial additive when incorporated within the range where no βo-phase is produced. Furthermore, Mo does not need to be incorporated into cast alloys as it does not provide any property-enhancing benefits other than disordered β-phase stabilization for improved forgeability. Additionally, Cr and Mn act as beneficial additives, improving properties such as ductility and impact resistance [39] when added in the range in which the βo-phase does not form.

4.3.2. Forged TiAl Alloys

As confirmed for the forged Ti–43.5Al–xCr ternary alloys in this study, the βo-phase could be eliminated by heat treatment when the Cr levels incorporated in them were ≤2.5%; but the forgeability was significantly reduced (Figure 1). At Cr levels of ≥3.0 at. %, which led to moderate forgeability, the βo-phase remained after the heat treatment. From these results, it can be said that the impact resistance, machinability, and creep strength of the forged TiAl alloys with a certain degree of forgeability were inevitably inferior to those of the cast TiAl alloys without the βo-phase. These findings reveal that the application of forged TiAl alloys is limited to large parts that cannot be produced using current TiAl alloy casting technologies. However, in such cases, the deterioration of various properties is inevitable.
Conversely, the TNM alloy is very interesting. In the present study, the βo-phase was almost eliminated by heat treatment at 1237 °C in the modified TNM alloys with nominal composition. The melting and forging methods used in this study are different from those used for the TNM alloy used in jet engine blades, and thus do not faithfully reproduce the material in practical use. Contradictorily, several cases have been reported where the βo-phase of TNM alloy has been eliminated by heat treatments [40,41,42]. Therefore, it is believed that the βo-phase could be almost eliminated in the TNM alloy used in actual jet engine blades. In such cases, the impact resistance, machinability, and creep strength can be significantly improved, as shown in this study. In addition, fatigue strength, which has been reduced by preferential oxidation of the βo-phase [19,20], should also be improved. It is very interesting why the βo-phase was not removed in the TNM alloy commercialized for jet engine blades.

5. Conclusions

In this study, the influence of the βo-phase on the practical properties of TiAl alloys was investigated for the use case of a last-stage turbine blade in a jet engine using 43.5Al forged alloys and 46.0Al cast alloys in a Ti–Al–xCr ternary system and two practical TiAl alloys (TNM alloy and TiAl4822). The salient results are outlined below:
  • For the forged and cast Ti–Al–xCr ternary alloys, the impact resistance at RT and 700 °C, as well as the machinability, decreased significantly as the βo-phase content increased.
  • The results of modified TNM and TiAl4822 alloys show that the βo-phase, present within the material, reduced the creep strength, impact resistance, and machinability.
  • Although the disordered β-phase is soft at high temperatures (forging temperature), it changes to the ordered βo-phase, which is significantly more brittle and harder after cooling. Thus, the detrimental effect of the βo-phase on the impact resistance and machinability can be attributed to this change.
  • An evaluation of various existing TiAl alloys developed to date with respect to the presence or absence of the βo-phase suggests that the practical properties of βo-phase-free TiAl alloys (that were developed in the distant past) are superior to those of more recent TiAl alloys that contain the βo-phase.
  • Finally, it can be concluded that the βo-phase should not be included in practical TiAl alloys, especially those used for jet engine blades.

Author Contributions

Conceptualization, T.T. and K.M.; methodology, T.T.; software, T.T.; validation, T.T.; formal analysis, T.T.; investigation, T.T.; resources, T.T. and K.M.; data curation, T.T.; writing—original draft preparation, T.T.; writing—review and editing, K.M.; visualization, T.T.; supervision, T.T.; project administration, K.M.; funding acquisition, K.M. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Japan Science and Technology Agency, grant number AS0216001.

Data Availability Statement

Data will be made available on request.

Conflicts of Interest

Kazuhiro Mizuta was employed by the AeroEdge Co., Ltd. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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Figure 1. Appearance of Ti–43.5Al–xCr (at. %) alloys hot-forged after heating at 1330 °C.
Figure 1. Appearance of Ti–43.5Al–xCr (at. %) alloys hot-forged after heating at 1330 °C.
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Figure 2. Backscattered electron images showing microstructures of forged Ti–43.5Al–xCr ternary alloys subjected to the 1280 °C/5 h/furnace cooling heat treatment protocol after hot-forged at 1330 °C: x = (a) 2.0Cr, (b) 2.5Cr, (c) 3.0Cr, (d) 3.5Cr, and (e) 4.0Cr.
Figure 2. Backscattered electron images showing microstructures of forged Ti–43.5Al–xCr ternary alloys subjected to the 1280 °C/5 h/furnace cooling heat treatment protocol after hot-forged at 1330 °C: x = (a) 2.0Cr, (b) 2.5Cr, (c) 3.0Cr, (d) 3.5Cr, and (e) 4.0Cr.
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Figure 3. Backscattered electron images showing microstructures of cast Ti–46.0Al–xCr ternary alloys subjected to hot isostatic pressing (HIP) using the 1200 °C/4 h/186 MPa protocol: x = (a) 2.0Cr, (b) 2.5Cr, (c) 3.0Cr, (d) 3.5Cr, and (e) 4.0Cr.
Figure 3. Backscattered electron images showing microstructures of cast Ti–46.0Al–xCr ternary alloys subjected to hot isostatic pressing (HIP) using the 1200 °C/4 h/186 MPa protocol: x = (a) 2.0Cr, (b) 2.5Cr, (c) 3.0Cr, (d) 3.5Cr, and (e) 4.0Cr.
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Figure 4. Relationship between Cr concentration and ratio of α2- and βo-phase in (a) forged Ti–43.5Al–xCr alloys and (b) cast Ti–46.0Al–xCr alloys.
Figure 4. Relationship between Cr concentration and ratio of α2- and βo-phase in (a) forged Ti–43.5Al–xCr alloys and (b) cast Ti–46.0Al–xCr alloys.
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Figure 5. Relationship between the area ratio of α2- and βo-phase and the mean absorbed energy obtained in the Charpy impact test at RT and 700 °C for (a) forged Ti–43.5Al–xCr alloys and (b) cast Ti–46.0Al–xCr alloys.
Figure 5. Relationship between the area ratio of α2- and βo-phase and the mean absorbed energy obtained in the Charpy impact test at RT and 700 °C for (a) forged Ti–43.5Al–xCr alloys and (b) cast Ti–46.0Al–xCr alloys.
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Figure 6. Results of machining tests conducted on forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys, showing the relationship between Cr concentration and tool weight loss at each Cr incorporation stage.
Figure 6. Results of machining tests conducted on forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys, showing the relationship between Cr concentration and tool weight loss at each Cr incorporation stage.
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Figure 7. Backscattered electron images showing microstructures of modified TNM alloys hot-forged at 1330 °C and then subjected to the first heat treatment stage at (a) 1177, (b) 1207, and (c) 1237 °C for 3 h followed by air cooling and subjected to the second heat treatment stage at 850 °C for 6 h followed by furnace cooling.
Figure 7. Backscattered electron images showing microstructures of modified TNM alloys hot-forged at 1330 °C and then subjected to the first heat treatment stage at (a) 1177, (b) 1207, and (c) 1237 °C for 3 h followed by air cooling and subjected to the second heat treatment stage at 850 °C for 6 h followed by furnace cooling.
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Figure 8. Backscattered electron images showing microstructures of modified TiAl4822 alloys after HIP at 1200 °C/4 h/186 MPa: (a) Ti–47.0Al–2.0Nb–1.79Cr, (b) Ti–47.0Al–2.0Nb–1.91Cr, and (c) Ti–47.0Al–2.0Nb–2.14Cr.
Figure 8. Backscattered electron images showing microstructures of modified TiAl4822 alloys after HIP at 1200 °C/4 h/186 MPa: (a) Ti–47.0Al–2.0Nb–1.79Cr, (b) Ti–47.0Al–2.0Nb–1.91Cr, and (c) Ti–47.0Al–2.0Nb–2.14Cr.
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Figure 9. Results of machining tests conducted on the modified TNM and TiAl4822 alloys, showing the relationship between the (a) first-stage heat treatment temperature and tool weight loss for modified TNM alloys, and (b) Cr content and tool weight loss for modified TiAl4822 alloys.
Figure 9. Results of machining tests conducted on the modified TNM and TiAl4822 alloys, showing the relationship between the (a) first-stage heat treatment temperature and tool weight loss for modified TNM alloys, and (b) Cr content and tool weight loss for modified TiAl4822 alloys.
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Figure 10. Creep curves of (a) modified TNM alloys and (b) modified TiAl4822 alloys subjected to a creep test performed at 750 °C/225 MPa.
Figure 10. Creep curves of (a) modified TNM alloys and (b) modified TiAl4822 alloys subjected to a creep test performed at 750 °C/225 MPa.
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Figure 11. Microstructures near ruptured position in creep specimens, showing the locations of creep voids: (a) modified TNM alloy subjected to the first-stage heat treatment at 1177 °C and (b) modified TiAl4822 alloy with 2.14Cr.
Figure 11. Microstructures near ruptured position in creep specimens, showing the locations of creep voids: (a) modified TNM alloy subjected to the first-stage heat treatment at 1177 °C and (b) modified TiAl4822 alloy with 2.14Cr.
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Figure 12. Results obtained from tests conducted to analyze the microstructure, Vickers hardness (HV), and brittleness at room temperature for alloys aimed at a single phase corresponding to each phase in the Ti–Al–Cr ternary alloy system at 1200 °C.
Figure 12. Results obtained from tests conducted to analyze the microstructure, Vickers hardness (HV), and brittleness at room temperature for alloys aimed at a single phase corresponding to each phase in the Ti–Al–Cr ternary alloy system at 1200 °C.
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Table 1. Chemical analysis results of modified TiAl4822 alloy produced in this study.
Table 1. Chemical analysis results of modified TiAl4822 alloy produced in this study.
Composition (Mass%)
TiAlNbCrCuFeNiSiCONHCa
Bal.32.64.622.64<0.0050.0230.0080.014<0.0050.07<0.002<0.0020.044
Table 2. Charpy impact test results at room temperature (RT) and 700 °C for forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys.
Table 2. Charpy impact test results at room temperature (RT) and 700 °C for forged Ti–43.5Al–xCr and cast Ti–46.0Al–xCr alloys.
Composition (at. %)Production MethodHeat Treatment ProtocolHot Isostatic Pressing (HIP) ConditionsCharpy Impact Test Results at RTCharpy Impact Test Results at 700 °C
AlCrMean Absorbed Energy (J/cm2)Standard Deviation (SD)Mean Absorbed Energy (J/cm2)SD
43.52.0Cast → hot-
forging at
1330 °C		1280 °C/
5 h/furnace cooling (FC)		-5.850.5710.162.22
2.55.580.5711.152.60
3.05.040.5611.131.88
3.53.700.6810.351.27
4.03.470.297.711.06
46.02.0Cast-1200 °C/
4 h/
186 MPa		5.220.6111.371.57
2.54.730.6111.601.78
3.04.570.7511.462.24
3.53.480.6910.771.35
4.03.250.489.221.38
Table 3. Charpy impact test results at RT and 700 °C for modified TNM and TiAl4822 alloys.
Table 3. Charpy impact test results at RT and 700 °C for modified TNM and TiAl4822 alloys.
AlloysComposition (at. %)Production MethodHeat Treatment ProtocolHIP
Conditions		Charpy Impact Test Results at RTCharpy Impact Test Results at 700 °C
AlNbCrMoBMean Absorbed Energy (J/cm2)SDMean Absorbed Energy (J/cm2)SD
TNM alloy43.54.0-1.00.1Cast → hot-
forging at
1330 °C		1177 °C/3 h AC → 850 °C/6 h FC 5.86 1.41 6.07 0.91
1207 °C/3 h AC → 850 °C/6 h FC-6.66 1.93 7.77 0.78
1237 °C/3 h AC → 850 °C/6 h FC8.43 2.21 8.50 1.30
TiAl 482247.02.01.79--Casting-1200 °C/
4 h/186 MPa		6.08 1.08 15.54 3.12
1.916.06 1.32 15.30 3.32
2.144.59 0.98 12.34 2.05
Table 4. Summary of the effects of the α2- and βo-phases ratios on the practical properties of TiAl alloys.
Table 4. Summary of the effects of the α2- and βo-phases ratios on the practical properties of TiAl alloys.
Change in Phase RatioImpact ResistanceMachinabilityCreep StrengthHigh-Temperature Fatigue Strength
RT700 °C
Increase in α2-phase
(increase in lamellar structure)		ImprovedImprovedSmall effectImproved [29]Varies depending on conditions [30]
Increase in βo-phaseReducedReducedReducedReduced Reduced [19,20]
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Tetsui, T.; Mizuta, K. Detrimental Effects of βo-Phase on Practical Properties of TiAl Alloys. Metals 2024, 14, 908. https://doi.org/10.3390/met14080908

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Tetsui T, Mizuta K. Detrimental Effects of βo-Phase on Practical Properties of TiAl Alloys. Metals. 2024; 14(8):908. https://doi.org/10.3390/met14080908

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Tetsui, Toshimitsu, and Kazuhiro Mizuta. 2024. "Detrimental Effects of βo-Phase on Practical Properties of TiAl Alloys" Metals 14, no. 8: 908. https://doi.org/10.3390/met14080908

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