2020 Volume 61 Issue 10 Pages 2017-2024
Titanium alloys have applications in air frames for commercial aircraft, and jet engine components such as fans and compressor disks, which function at low temperatures (up to 673 K). Near β-type Ti–5Al–2Sn–2Zr–4C–4Mo (Ti-17) exhibits greater strength, crack propagation resistance, and creep resistance at intermediate temperatures compared to the (α + β)-type Ti–6Al–4V. It is important to estimate the fatigue life of engine components made of Ti-17. This requires problem quantitative relationship between the fatigue properties and microstructural factors of Ti-17. Therefore, the fatigue properties including tensile properties and microstructures of Ti-17 samples fabricated by hot-forging at various temperatures, followed by high- and low-temperature solution treatment (ST), and same aging treatment were investigated to define a quantitative relationship between the fatigue properties and the microstructures.
The microstructures of all forged Ti-17 samples exhibit elongated prior β-grains composed of two microstructural feature regions: acicular α and fine equiaxed α-phase regions. The volume fraction of the acicular α region decreases with increasing ST temperature. The Vickers hardness, 0.2% proof stress and tensile strength increases with increasing ST temperature. However, the elongation and reduction of area exhibit a reverse trend. The Ti-17 samples forged at 1173 K followed by solution treatment at 1073 K and aging treatment exhibits the highest fatigue limit of around 975 MPa. The fatigue strength of the forged Ti-17 samples is strongly related to the microstructural factor such as the volume fraction of the equiaxed α-phase region, which is one of the crack initiation sites in the forged Ti-17 samples subjected to low temperature ST and aging, and the strength difference between the acicular α-phase and the fine (α + β)-phase, which leads to the crack initiation in the forged Ti-17 sample subjected to high temperature ST and aging.
This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 84 (2020) 200–207.
In recent years, the air frame weight reduction and improvements in engine thermal efficiency have been the important issues in aerospace industry. Till the end of the 20th century, steel and aluminum alloys were used for constructing airframes. However, titanium alloys used in the construction of air frames and aircraft engines is on the rise owing to their high specific strength (strength/density). Recently, the consumption of titanium alloys has been rapidly increasing due to carbon fiber-reinforced plastic (CFRP), which exhibit excellent compatibility with titanium alloys from the viewpoint of thermal expansion coefficient.1)
Titanium alloys are expected have wide range of applications as structural materials in the current aerospace industry. Titanium alloys occupy an important position in the air frame and aircraft engine. Reduction of air frame weight and cost and to improvements in thermal efficiency and power of the aircraft engine are still necessary. Therefore, light weight, high heat resistance, and enhanced mechanical properties are necessary for the titanium alloys used in aircraft. The aircraft engine is composed of a fan, compressor, combustor, and turbine from the entrance side. The gas temperature is approximately 775 K at most, around the fan or compressor near the combustor, and approximately 1875 K at most around the entrance of the turbine located at the rear of the combustor.2) Hence, Ni-based super alloys with high heat and creep resistance are used for engine parts exhibiting high temperatures, but their high density is an issue. Excellent corrosion resistance, toughness, fatigue, creep, and impact characteristics are required for the fan and compressor blades. Therefore, research and development of titanium alloys with high mechanical performance at relatively high temperatures has progressed for applications in fans and compressors.2)
Ti–6Al–4V (Ti-64) is a general titanium alloy used in aircraft structures. Ti-64 is the most commonly used (α + β)-type titanium alloy and has balanced characteristics. At present, plenty of data and operating experience exist for Ti-64. However, since the upper temperature limit for its use is approximately 773 K, the application of Ti-64 around the aircraft engine is limited, and it is mainly used in the fans. To solve this problem, Ti–5Al–2Sn–2Zr–4Mo–4Cr (Ti-17), which is a near β-type titanium alloy, has been developed.3–5) Ti-17 exhibits superior heat resistance, tensile strength, fracture toughness, and creep characteristics compared to those of Ti-64 and is expected to be used in the fan and shaft around the aircraft engine. However, at present, the relationship between the microstructure and the mechanical properties is not fully understood. Therefore, the relationship between the microstructure and the mechanical properties, focusing on the fatigue properties of the forged Ti-17 at various temperatures subjected to solution treatment at various temperatures, followed by the same aging treatment is investigated in this study.
The forged Ti–5Al–2Sn–2Zr–4Mo–4Cr (Ti-17) fabricated by Kobe Steel Co. Ltd., whose chemical composition is shown in Table 1, was used in this study.
The schematic representation of the homogenization and hot forging process is given in Fig. 1. In this process, the ingot of Ti-17 was first subjected to homogenization at 1203 K for 2 h under vacuum and then to hot forging (HF) at 1023, 1073 and 1173 K followed by air cooling (AC), respectively. Hot forging was carried out with a working ratio of 75% and a strain rate of 0.033 s−1.6) Ti-17 forged at 1023, 1073, and 1173 K are referred to as HF/1023K, HF/1073K, and HF/1173K, respectively. Figure 2 shows an overview of a hot-forged Ti-17 with a pancake-like shape.
Schematic drawing of homogenization treatment and hot forging (HF).
Overview of pancake shaped Ti-17 after hot forging.
The schematic representation of the solution treatment (ST) and aging treatment (AT) subjected to the forging of Ti-17 is shown in Fig. 3. In this process, the forged Ti-17 was subjected to solution treatment at a relatively lower temperature (L) of 1073 K, which is lower than the β-transus temperature of 1163 K and a relatively higher temperature (H) of 1143 K for 4 h followed by water quenching (WQ). Finally, the same aging treatment (AT) at 893 K for 8 h followed by AC was conducted. The forged Ti-17 subjected to AT after ST at low and high temperatures followed by AT will be referred to as L/STA and H/STA, respectively, herein. HF/1023K, HF/1073K and HF/1173K subjected to solution treatment at low and high temperature followed by AT will be referred to as HF1023K·L/STA, HF1023K·H/STA, HF1073K·L/STA, HF1073K·H/STA, HF1173K·L/STA and HF1173K·H/STA, respectively.
Schematic drawing of solution treatment (ST) and aging treatment (AT).
Small pieces of specimen for microstructural observations were cut from Ti-17 subjected to each HF and STA and embedded in polyester resin. Specimens were wet-polished using a #4000 grid SiC emery paper and finished by buff polishing using a silicon (SiO2) suspension. The microstructural observations were carried out using an optical microscopy (OM). The phase constitution was evaluated using an X-ray diffractometer (XRD) with a voltage of 30 kV, a current of 10 mA, a diffraction angle 2θ between 30 and 90 degree and a scanning speed of 2θ = 10 degree/min at room temperature.
2.4 Evaluation of mechanical propertiesThe small piece specimens used for the microstructural observations were wet-polished using emery papers up to #4000, and their Vickers hardness was measured from near the surface to the center of the specimens using a micro-Vickers hardness tester with an indention load of 1.96 N and a holding time of 10 s.
Dog-bone type specimens for tensile tests (Fig. 4(a)), and dog-bone type specimens for fatigue tests (Fig. 4(b)), were machined from Ti-17 subjected to each HF and STA.
Geometries of (a) tensile and (b) fatigue specimens.
The surfaces of the specimens for tensile tests were wet-polished using emery papers up to #1500. Tensile tests were carried out at a strain rate of 0.5%/min up to a strain of 5%, after which, at a strain rate of 1.5%/min using an Instron-type testing machine of 10 kN capacity at room temperature (285 K).
The surfaces of the specimens for the fatigue tests were wet-polished by #4000 grid SiC emery paper and finished by buff polishing using a SiO2 suspension. Fatigue tests under load-controlled conditions were carried out at a stress ratio of 0.1, and a frequency of 10 Hz using an electro-hydraulic servo fatigue-testing machine of 50 kN capacity, at the room temperature of 285 K. The fractured surfaces of the specimens after the fatigue tests were observed using a scanning electron microscope (SEM).
Low-magnification optical micrographs of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are shown in Fig. 5. Every microstructure shows an elongated prior β-grain due to HF. Needle-like α-phases (acicular α-phases) and grain boundary α-phase are observed in prior β-grains. According to the XRD analysis, the diffraction profiles were nearly the same in both L/STA and H/STA, but the diffraction peaks of the β-phase were higher in H/STA than L/STA.
Optical micrographs of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA (×100).
Figure 6 shows the relationship between the aspect ratio of the prior β-grains HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA. The aspect ratio in this case is defined as the length of the short axis divided by the length of the long axis. The number of prior β-grains showing a high aspect ratio, with a maximum aspect ratio of 0.18, tends to increase with increasing HF temperature.
Relationship between aspect ratio of prior β grain, ARpβ, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
High-magnification optical micrographs of HF/1173K subjected to L/STA and H/STA are shown in Fig. 7 as representative examples showing each phase of the acicular α-phase, equiaxed α-phase, or grain boundary α-phase. In addition to the acicular α-phase, which is the main constituent phase (Fig. 7(a)), a small portion of equiaxed (α + β)-phase area (Fig. 7(b)), and the grain boundary α-phase at the prior β grain boundary (Fig. 7(c)), are observed in the microstructure. The equiaxed α-phase region tends to exist near the prior β-grain boundary, and the aspect ratio of the equiaxed α-phase is remarkably large.
Optical micrographs of (a) acicular α-phase, (b) equiaxed α-phase and, (c) grain boundary α-phase of HF/1173K subjected to L/STA and H/STA (×1000).
The volume fractions of the acicular α-phases of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are listed in Table 2. The volume fraction of the acicular α-phase is significantly larger in L/STA than in H/STA, which is smaller with increasing ST temperature, because the amount of the transformation of the α-phase to the β-phase is greater when the ST temperature approaches the β-transus temperature.7) However, the aspect ratio of the acicular α-phases of H/STA is nearly the same as that of L/STA with increasing ST temperature.
The number ratio and aspect ratio of the acicular α-phase of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA are shown in Fig. 8. The number ratio of the acicular α-phase with a small aspect ratio is proportional to the HF temperature. In other words, the acicular α-phase elongates with increasing HF temperature. This is because the amount of induced strain in the microstructure by forging decreases with increasing HF temperature, leading to the formation of the acicular α-phase with relatively small aspect ratio. Furthermore, the grain growth occurs by heat treatment after HF allows the aspect ratio of the acicular α-phase to be smaller. The trend in the number ratio and aspect ratio of the acicular α-phase of HF/1023K, HF/1073K and HF/1173K subjected to H/STA were similar.
Number ratio, Naα, and aspect ratio, ARaα, of acicular α-phase of HF/1023K, HF/1073K and HF/1173K subjected to L/STA.
The volume fractions of the equiaxed α-phase of HF/1023K, HF/1073K and HF/1173K subjected to L/STA are shown in Fig. 9. The volume fraction of the equiaxed α-phase in L/STA tends to decrease with increasing HF temperature. The volume fraction of the equiaxed α-phase in H/STA is remarkably low, below 1%.
Volume fraction of equiaxed α-phase, Veα, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA.
The continuous ratio of the grain boundary α-phase of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA is shown in Fig. 10. The continuous ratio of the grain boundary α-phase is the ratio of the total length segment of the grain boundary α-phase to the total length of the grain boundary. The continuous ratio of the grain boundary α-phase of HF/1023K is relatively high, approximately 30% in both L/STA and H/STA. The continuous ratio of the grain boundary α-phase increases with increasing HF temperature in both cases, and reaches approximately 80% at maximum in HF/1173. The reason for this phenomenon is that the splitting of the grain boundary α-phase is inhibited because the amount of strain induced by working decreases with increasing HF temperature and thus, the continuity of the grain boundary α-phase increases.
Continuous ratio of grain boundary α-phase, Gα, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
The Vickers hardness of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are shown in Fig. 11. The Vickers hardness value of each HF subjected to L/STA ranges from 397 HV to 407 HV, while that of each HF subjected to H/STA is from 429 HV to 444 HV. The maximum Vickers hardness value of 446 HV is obtained for HF/1073 subjected to H/STA. The Vickers hardness value increases by approximately 30 HV by H/STA as compared with the case of L/STA owing to the increment in the age-hardening ability.
Vickers hardness, HV, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
Tensile properties such as tensile strength, 0.2% proof stress, elongation, and reduction of area of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are shown in Fig. 12. The tensile strength and 0.2% proof stress of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA are from 1160 to 1210 MPa and to H/STA from 1060 to 1120 MPa, respectively. On the other hand, the tensile strength and 0.2% proof stress of HF/1023K, HF/1073K, and HF/1173K subjected to H/STA are from 1330 to 1390 MPa and from 1260 to 1320 MPa, respectively, which are approximately 100 MPa and 200 MPa, respectively, greater than those of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA. These trends of increment in the tensile strength and 0.2% proof stress are similar to those in the Vickers hardness.
Tensile properties of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
Elongation and reduction of area of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA ranges from 13.2% to 14.5% and 21.5% to 25.0%, respectively. However, the elongation and reduction of the HF/1023K, HF/1073K, and HF/1173K subjected to H/STA are from 5.5% to 8.4% and 9.8% to 11.5%, respectively, and they are smaller than those of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA. Therefore, this trend is opposite to that observed in the tensile strength and 0.2% proof stress.
A higher elongation in L/STA than in H/STA is due to the lower age hardening in L/STA than in H/STA. Remarkably high strength and low ductility are obtained in H/STA with a high age hardening. Although improving ductility has been reported owing to the texture formation of the precipitated α-phase according to the Burgers relation $(\{ 1\bar{1}0\} _{\beta }//(0002)_{\alpha })$ or splitting the grain boundary α-phase and ambiguous grain boundary α-phase,8) the clear relationship between the continuity of grain boundary α-phase and improved ductility is not recognized in this study.
The relationship between the tensile strength and elongation of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are shown in Fig. 13. Since HF/1023K, HF/1073K, and HF/1173K subjected to L/STA exhibited elongation over 10% and strength over 1 GPa, the balance of strength and elongation is judged to be better in L/STA than in H/STA.
Relationship between tensile strength, σB, and elongation, EL, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
The Vickers hardness, tensile strength and 0.2% proof stress are strongly related to the volume fraction of the α-phase (acicular α-phase). In general (α + β)-type titanium alloys, the β-phase and fine precipitated α-phase exist around the primary α-phase. In such microstructures, the strength depends mainly on the fine precipitated α-phase.5) Therefore, the amount of the β-phase increases with increasing ST temperature, and the amount of the fine precipitated α-phase increases remarkably with increasing ST temperature, leading to high strength.9)
3.3 Fatigue characteristicsThe S-N curves of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA are shown in Fig. 14. The relationship between the maximum cyclic stress and the cycles to failure, namely the S-N curves of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are studied. The data surrounded by open circles indicate the specimens with the fatigue crack initiation from the sub-surface. The regression equation10) for each S-N curve is as follows:
| \begin{equation} \sigma_{\text{max}} = -152.04\times \log (\text{Nf}) + 1876.0 \end{equation} | (1) |
| \begin{equation} \sigma_{\text{max}} = -110.88\times \log (\text{Nf}) + 1569.7 \end{equation} | (2) |
| \begin{equation} \sigma_{\text{max}} = -165.68\times \log (\text{Nf}) + 1932.6 \end{equation} | (3) |
| \begin{equation} \sigma_{\text{max}} = -149.92\times \log (\text{Nf}) + 1696.6 \end{equation} | (4) |
| \begin{equation} \sigma_{\text{max}} = -101.82\times \log (\text{Nf}) + 1568.9 \end{equation} | (5) |
| \begin{equation} \sigma_{\text{max}} = -142.07\times \log (\text{Nf}) + 1663.5 \end{equation} | (6) |
S-N curves of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
For HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA subjected to L/STA and H/STA, their fatigue limits exhibit the highest values of around 900 and 975 MPa, respectively, at HF/1073K and their fatigue ratios (fatigue limit/tensile strength) are 0.68 and 0.87, respectively. The latter fatigue ratio is greater than that of the high-fatigue strength Ti-64 with an equiaxed α structure, which exhibits a fatigue ratio of 0.79.11)
The relationship between the fatigue limit and the aspect ratio of the prior β-grains of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA are shown Fig. 15. The maximum aspect ratio of the prior β-grain is 0.18. The fatigue limit increases with increasing aspect ratio of the prior β-grain. Furthermore, for L/STA, the fatigue limit increases with decreasing volume fraction of the equiaxed α-phase area as shown in Fig. 16. Hence, for L/STA, although two microstructural factors (aspect ratio of the prior β-grain and the volume fraction of equiaxed α-phase) affect the fatigue limit, decreasing the volume fraction of the equiaxed α-phase near the grain boundary, (a fatigue crack initiation site), mainly improves the fatigue limit.
Relationship between fatigue limit, σfl, and aspect ratio of prior β grain, ARpβ, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
Relationship between fatigue limit, σfl, and volume fraction of equiaxed α-phase in prior β grain, Veα, of HF/1023K, HF/1073K and HF/1173K subjected to L/STA.
The standard deviations obtained from the S-N curves of HF/1023K, HF/1073K, and HF/1173K subjected to L/STA and H/STA are shown in Fig. 17. For HF/1023K, HF/1073K and HF/1173K subjected to both treatments, the scatter of the standard deviation tends to decrease with increasing HF temperature. In L/STA, the ratio of the sub-surface fatigue crack initiation tends to decrease with increasing HF temperature, but it does not change with HF temperature in H/STA.
Standard deviations, SD, obtained from S-N curves of HF/1023K, HF/1073K and HF/1173K subjected to L/STA and H/STA.
The macroscopic fractographs and fractographs near the crack initiation sites of HF/1173K subjected to L/STA and H/STA obtained from fatigue tests in the high-cycle fatigue life region are shown in Fig. 18 as a representative example. The fatigue crack is initiated from the surface or subsurface of the specimen and propagates radially towards its center. In some cases, striations are observed in the stable fatigue crack propagation area, and dimples are observed in the fast fatigue crack propagation area. Therefore, the fatigue fracture morphology is similar to that of a ductile metallic material. This type of fatigue fracture was similar in both low- and high-cycle fatigue life regions. A large number of secondary cracks were observed in L/STA as compared with the case in H/STA.
Fractographs of HF/1173K subjected to L/STA and H/STA obtained from fatigue tests at high fatigue life region.
Cross sections of the near crack initiation site of HF/1173K subjected to both treatments are shown in Fig. 19 as a representative example. In L/STA, the fatigue crack initiates near the interface between the equiaxed α-phase and the acicular α-phase, while, in H/STA, it initiates near the coarse acicular α-phase or the split grain boundary α-phase.
Cross section of near crack initiation site of HF/1173K subjected to L/STA and H/STA (×1000).
In HF/1023K, HF/1073K and HF/1173K subjected to L/STA, the rate of the subsurface fatigue crack initiation decreases with increasing HF temperature because of the decrease in the volume fraction of the equiaxed α-phase, which will increasingly become a fatigue crack initiation site (Fig. 16). Decreased rate of the subsurface crack initiation leads to improved fatigue limit. On the other hand, in HF/1023K, HF/1073K, and HF/1173K subjected to H/STA, the rate of the subsurface crack initiation is almost constant with increasing HF temperature. As the coarse acicular α-phase scatters in the fine (α + β)-phase regions (Fig. 7(a)), which becomes a fatigue crack initiation site (Fig. 19), even in the HF/1173K, the fatigue crack initiates due to the strength difference in between the coarse acicular α-phase and the fine (α + β)-phase.
The relationships between the microstructure and the mechanical properties, with a focus on the fatigue properties of Ti–5Al–2Sn–2Zr–4Mo–4Cr (Ti-17), developed for use in the fans and compressors of aircraft engines, forged at various temperatures subjected to solution treatment at various temperatures followed by the same aging treatment were investigated in this study. The following results were obtained.
The authors would like to express great thanks to Dr. Y. Itsumi at Kobe Steel Co., Ltd. Kobe, Japan and Dr. Y. Yamabe-Mitarai at National Institute for Materials Science, Tsukuba, Japan for fabricating and supplying forged Ti-17. This work was partly supported by the Structural Materials for Innovation, Cross-ministerial Strategic Innovation Promotion Program (SIP), Cabinet Office, Government of Japan.
