Comparative Study of the Mechanical Properties and Fracture Mechanism of Ti-5111 Alloys with Three Typical Microstructures

Abstract

In this work, Ti-5111 alloys with equiaxed, bimodal and lamellar microstructures were prepared by various heat treatment processes. The room-temperature tensile properties, deformation microstructure and fracture mechanism of the alloys with different microstructures were investigated. Furthermore, the mechanism by which the microstructure affects the mechanical properties of Ti-5111 alloys with three typical microstructures was confirmed. The Ti-5111 alloy with a bimodal microstructure has minimum grain size and a large number of α_s/β phase boundaries, which are the primary reasons for its higher strength. Simultaneously, the excellent coordination in the deformation ability between the lamellar α_s and β phases is what enables the alloy with a bimodal microstructure to have the most outstanding mechanical properties. Additionally, the presence of a grain boundary α phase and the parallel arrangement of a coarse α_s phase are the main reasons for the inferior mechanical properties of the Ti-5111 alloy with a lamellar microstructure. The fracture mechanism of the alloy with an equiaxed microstructure is a mixed fracture mechanism including ductile fracture and destructive fracture. The fracture mechanisms of the Ti-5111 alloy with bimodal and lamellar microstructures are typical ductile fracture and cleavage fracture, respectively. These findings serve as a guide for the performance improvement and application of the Ti-5111 alloy.

1. Introduction

Titanium alloys have a wide range of applications in aerospace and marine engineering due to their high specific strength, good toughness and excellent corrosion resistance [1,2,3,4]. The Ti-5111 alloy is a typical near-α titanium alloy developed as a cost-effective alternative to Ti-100 (Ti-5Al-2Nb-1Ta-0.8Mo), offering excellent weldability, moderate strength and affordability. Furthermore, the fracture toughness and corrosion resistance of the Ti-5111 alloy are superior to those of the Ti-100 and Ti-6Al-4V alloys [5]. In addition, its excellent compatibility with various advanced composites makes it an ideal choice for composite assemblies. Currently, the Ti-5111 alloy is extensively utilized for submarine fasteners and ship hull materials [6,7].

In recent years, some scholars have conducted studies on the Ti-5111 alloy, recognizing its significant potential for marine engineering. Robinson et al. [8] investigated the relationship between the microstructure and ductility of the Ti-5111 alloy prepared by different casting processes. They found that the colonies comprising lamellar α and β phases were present in all castings, and the average grain size of the graphite mold casting sample was smaller than that of the investment casting sample. Stauffer et al. [5] studied the fatigue crack behavior of Ti-5111 alloy welds in various environments and found that the fatigue crack expansion rate is primarily influenced by the electrochemical corrosion environment. Kolli et al. [9] found that the addition of yttrium could refine grains and identified high-density nanoscale yttrium-rich precipitates in Ti-5111 alloy welds. Wang et al. [10] analyzed the effects of hydrides on the mechanical properties of the Ti-5111 alloy and considered that the β phase facilitated the rapid diffusion of hydrogen and weakening at the phase interface due to the increased hydrogen concentration, resulting in diminished mechanical properties. In addition, Wang et al. [11] investigated the crack extension behavior of twin-induced plasticity (TWIP) Ti-5111 alloys with three distinct microstructures and found that deformation twins significantly impeded crack extension.

The mechanical properties of titanium alloys are primarily governed by the α phase with a hexagonal close-packed (HCP) structure and the β phase with a body-centered cubic (BCC) structure [12]. The shape, size, volume fraction and distribution of α and β phases all have significant effects on the strength and plasticity of titanium alloys [13,14,15,16]. Titanium alloys can be classified into equiaxed, bimodal and lamellar alloys based on the microstructure, each exhibiting distinct mechanical properties and industrial applications [17,18,19]. It is of great significance to investigate the mechanical properties and deformation mechanism of Ti-5111 alloys with various microstructures for optimizing the properties of titanium alloy products. Currently, there is no systematic study on the mechanical properties and tensile deformation microstructure of the three Ti-5111 alloys. Especially, there is no report on the comparative study of the mechanical properties of the three Ti-5111 alloys by analyzing the microstructure evolution in different plastic deformation zones.

This paper aims to manipulate the three typical microstructures of the Ti-5111 alloy by different heat treatment processes and then establish the relationship between the three typical microstructures and mechanical properties, which is of great significance in improving the mechanical properties of the alloy. In this work, three types of Ti-5111 alloys with typical microstructures were prepared using different heat treatment processes. The tensile properties of the Ti-5111 alloy were tested, and the microstructure evolution at room temperature was analyzed by an optical microscope (OM) and transmission electron microscope (TEM). The mechanism of how the microstructure influences strength and plasticity was investigated. Furthermore, the fracture mechanism was analyzed by using a scanning electron microscope (SEM).

2. Experiment

The original material used in this study is the hot-rolled Ti-5111 alloy with a chemical composition of Ti-5.11Al-1.02Sn-0.98V-1.12Zr-0.84Mo. The β phase transition temperature of the alloy is 980 ± 5 °C, measured by the metallographic method. Three heat treatments were conducted on the hot-rolled Ti-5111 alloy for obtaining Ti-5111 alloys with an equiaxed microstructure (EM), bimodal microstructure (BM) and lamellar microstructure (LM), respectively. Figure 1 presents a schematic diagram illustrating the heat treatment processes to achieve the three typical microstructures. Heat treatments were performed with the following parameters: 900 °C × 0.5 h/AC, 980 °C × 0.5 h/AC and 1020 °C × 0.5 h/AC. The argon was used as a protective gas to prevent high-temperature oxidation during the heat treatment.

Ti-5111 alloys with different microstructures were processed into standard room temperature tensile specimens with the detailed parameters, as shown in Figure 2a. The tensile tests were conducted using a SHIMADZUAG-I250KN (Kyoto, Japan) precision universal testing machine at room temperature with a stretching rate of 1 mm/min for testing their yield strength, ultimate tensile strength, elongation and area reduction. Tensile tests of the alloys were carried out three times, and the final results were averaged. The metallographic observation was conducted by a Zeiss Axio Vert A1 microscope (Oberkochen, Germany). Before the above work, the specimens were ground successively on 180#, 240#, 600#, 800#, 1000#, 1500# and 2000# sandpaper, followed by mechanical polishing and etching with an HF:HNO₃:H₂O = 5:12:83 etching solution. The fracture morphology was observed using a JSM-IT100 SEM (JEOL, Tokyo, Japan) after the fractured part of the tensile specimen was soaked in anhydrous ethanol and cleaned in an ultrasonic cleaner for 10 min. The specimens near the fracture were cut into thin sheets with different sampling areas, as shown in Figure 2b. The microstructures of the near fracture region (NFR) and far fracture region (FFR) were observed using a metallographic microscope and a JEM-2100 high-resolution TEM (JEOL, Tokyo, Japan), respectively. The specimens were double-jet electrolytic-polished before TEM observation.

3. Results

3.1. Mechanical Properties

Figure 3a illustrates the engineering stress–strain curves of Ti-5111 alloys with three typical microstructures. Figure 3b,c show the strength and plasticity of three typical alloys, respectively. It is known that the EM Ti-5111 alloy has a moderate strength, with a yield strength of 719 MPa and an ultimate tensile strength of 787 MPa. The elongation and area reduction of the EM alloy are 16.5% and 34%, respectively. It is noteworthy that the BM alloy is not only superior to the EM alloy in terms of strength, with a yield strength and ultimate tensile strength of 722 MPa and 801 MPa, respectively, but also exhibits superior plasticity, with an elongation and area reduction of 17.5% and 43%. In addition, the LM alloy shows the lowest strength and plasticity, with a yield strength, ultimate tensile strength, elongation and area reduction of 713 MPa, 785 MPa, 11.5% and 19%, respectively. Overall, the BM alloy demonstrates excellent comprehensive mechanical properties, compared with the EM and LM alloys. The mechanical properties of Ti-5111 alloys with three different microstructures are analyzed in detail in Section 4.

3.2. Features of Three Typical Microstructures

Figure 4 shows the microstructure of three typical Ti-5111 alloys after different heat treatment processes. It can be seen in Figure 4(a₁,a₂) that the EM alloy mainly consists of a large number of homogeneously distributed equiaxed primary α phases (α_p), with a small number of β phases and diffusely precipitated short needle-like secondary α phases (α_s), indicating the occurrence of an α→β phase transition. Figure 4(b₁,b₂) shows the microstructure of the BM alloy heat-treated at 960 °C. It can be seen that the volume fraction and size of α_p decrease and become smaller, respectively. Additionally, the proportion of the phase boundary increases. At this time, the grain size of the BM alloy is the smallest under this heat treatment, and there is a high dislocation density between α_s and β phases, which is beneficial to improving strength [20]. Moreover, the increase in the volume fraction of the β phase is also beneficial for its plasticity because the β phase with a BCC structure has more slip systems compared to the α phase with an HCP structure [21,22]. Figure 4(c₁,c₂) show the microstructure of the LM alloy heat-treated at 1020 °C. With a further increase in the treatment temperature beyond the β transition temperature, α_s exists as lamellar grains, and the volume fraction of the β phase increases further. However, the grain size is larger due to the higher annealing temperature, which may reduce its strength. In addition, the grain boundaries α (α_GB) phase and some α colonies can be observed in the LM alloy, and the orientation of the lamellar α phase within a colony is similar.

3.3. Features of the Deformation Microstructure

Figure 5 shows the deformation microstructure of the NFR and FFR of Ti-5111 alloys with three different microstructures. It can be seen in Figure 5(a₁,a₂) that some equiaxed α grains are elongated, and some α_s grains show obvious bending deformation in the NFR of the EM alloy, while light deformation traces were observed in the FFR. It can also be found that small shear deformation occurs around the α/β grain boundary. Figure 5(b₁,b₂) show the deformation microstructure of the BM alloy. The severe plastic deformation of grains occurs in NFR, indicating that the BM alloy is more prone to dislocation slip, which is consistent with the excellent elongation found in the above tensile test results. In addition, it can be observed in Figure 5(b₂) that there are long deformation bands inside the FFR of the BM alloy, while other grains do not differ much from the grains in the original microstructure. Figure 5(c₁,c₂) show the optical microstructures of the NFR and FFR of the LM alloy. It can be seen that the thickness of the lamellar phase grains decreases in the NFR, and the lamellar α phase grains exhibit shear deformation, as indicated by the white circles. Previous investigations have found that the parallel arrangement of α/β phase boundaries is prone to be a source of crack initiation [23], and the crack would expand rapidly along the direction perpendicular to the lamellar α phase grains [24]. This is one of the key reasons for the poor plasticity of the LM alloy. No significant plastic deformation was found in FFR; however, traces of lamellar α phase grains starting to shear were also observed.

4. Discussion

4.1. Effect of the Microstructure on the Mechanical Properties

Figure 6 shows the TEM deformation microstructure in the NFR and FFR of the EM Ti-5111 alloy. It can be seen in Figure 6(a₁) that there are cracks and deformation twins in the near fracture area, which may be an important reason for the low dislocation density in this area. The crack deflects slightly when it passes through the twin boundary, indicating that the twin boundary is beneficial to the strength of the EM alloy by affecting crack growth. It can also be seen in Figure 6(a₂,a₃) that there is a high dislocation density and dislocation entanglement in the NFR, which is a significant reason for the moderate strength of the EM alloy. Figure 6(b₁–b₃) shows the deformation microstructure in the FFR of the EM Ti-5111 alloy. In Figure 6(b₁), it can be seen that the crack first expands along a β-sheeted grain, which should be due to the high dislocation density mentioned above in the β phase. Then, the crack is deflected at the α_s grain boundary during expansion, which is due to the additional energy consumed when deflection and bifurcation occur during crack growth [25,26]. At the same time, it can be seen in Figure 6(b₂) that the crack expands rapidly when it does not encounter the short needle-like α_s phase. The above phenomena indicate that the dispersed and precipitated short needle α_s have an effect of hindering crack growth, which is beneficial to the mechanical property of the EM alloy. Figure 6(b₃) shows the distribution of dislocations in the FFR of the EM alloy. The dislocation density in the FFR is far lower than that in the NFR. Moreover, unequally distributed dislocations and uneven grain boundaries indicate that the grain boundaries of the EM alloy are unstable, which is not conducive to the plasticity of the alloy. This may be a significant reason why the mechanical properties of the EM alloy are not optimal.

Figure 7 shows the TEM deformation microstructure in the NFR (a) and FFR (b) of the BM Ti-5111 alloy. First, from the entirety of Figure 7, it can be observed that masses of dislocations are accumulated in the BM Ti-5111 alloy, which explains why the BM alloy has the highest strength. Furthermore, as can be seen in Figure 7(a₁), a higher dislocation density around the interface between the α and /β phases indicates that the α/β phase boundary has the effect of impeding the dislocation movement, which is consistent with the findings of Tan et al. [27]. In general, the Ti alloy always precipitates parallel lamellar structures during heat treatment near the β transition point. Figure 7(a₂) shows that there are numerous α_s/β boundaries in the BM alloy, which is a key reason for the highest strength of the BM alloy. In addition, compared to the EM alloy, there were fewer cracks in the plastic deformation region of the BM alloy, and only a few short and discontinuous cracks were found, as shown in Figure 7(a₃). The formation of cracks soon led to the annihilation of nearby dislocations. Figure 7(b₁–b₃) shows the deformation microstructure in the FFR of the BM alloy. In Figure 7(b₁), it can be seen that a large number of dislocation lines and some dislocation entanglements exist inside α_p grains, possibly attributed to the variation in the distribution and density of dislocations within each grain with the local strain [28]. In contrast to equiaxed α_p, the dislocation distribution is more uniform. Additionally, it can also be observed that some shear slip lines with an angle of about 45° are formed inside the lamellar α and β phases, as this represents the direction of maximum shear stress [29]. Moreover, some of the shear slip lines were found to traverse the lamellar α and β phases, indicating that α and β phases are very well coordinated in deformation. This coordination is a significant factor contributing to the excellent plasticity of the BM alloy.

Figure 8 shows the TEM deformation microstructure in the NFR (a₁, a₂) and FFR (b₁, b₂) of the LM Ti-5111 alloy. It can be observed in Figure 8(a₁,a₂) that significant cracks are visible in the NFR of the LM alloy, indicating poor mechanical properties. As previously mentioned, dislocations tend to accumulate at the α_s/β phase boundary, resulting in a stress concentration. In LM alloys, a large number of α_s/β phase boundaries makes them more susceptible to cracking during the tensile process. Additionally, the grains are relatively coarse, causing the LM alloy to fracture. As can be seen in Figure 8(a₁), the crack rapidly propagates along the α_s/β phase boundary, leading to the rapid fracture of the LM alloy. Figure 8(a₂) shows the morphology of crack propagation in the LM alloy. The crack initially grows straight and then zigzags. Typically, cracks often propagate in a flat and straight manner in a specific direction. When a crack propagates without encountering grain boundaries with different phase correlation systems or without encountering dislocation entanglement, it tends to extend flat in a particular direction. However, as shown by the white circle in Figure 8(a₂), cracks demonstrate tortuous propagation influenced by dislocation cells. Only residual dislocation walls can be observed due to the significant number of dislocations consumed by the crack extension. Furthermore, it is observable that the crack ceases its development after traversing the dislocation cell, possibly indicating that fracture behavior has already taken place at different sites. Figure 8(b₁) shows the TEM morphology of the FFR of the LM Ti-5111 alloy, revealing the presence of needle-like α_s sheets that can be seen near the α_s/β phase boundary. This occurrence is attributed to the heat treatment temperature during the preparation of the LM alloy surpassing the β phase transition temperature [30,31]. Shear deformation occurs in some flaky grains mentioned above. Here, it can be seen in Figure 8(b₂) that the lamellar β phase accumulates numerous dislocations and undergoes shear fracture, which is due to the rapid accumulation of dislocations on the shear plane. The above phenomenon indicates that the limited coordination deformation ability of the α and β phases in the LM alloy leads to shear fracture in some grains. This is a key factor contributing to the low plasticity of the LM alloy.

4.2. Fracture Mechanism

Figure 9(a₁–a₃) show the macroscopic and microscopic fracture morphology of the EM Ti-5111 alloy. It can be seen in Figure 9(a₁) that there is slight necking near the fracture. Subsequently, microscopic fracture morphology was observed, as shown in Figure 9(a₂,a₃). The fracture morphology is characterized predominantly by numerous equiaxed dimples, indicating that ductile fracture is the primary fracture mechanism. However, some cleavage steps are found by observing the fracture morphology at high magnification, as shown by the white arrow in Figure 9(a₃), suggesting that cleavage fracture is also a contributing fracture mechanism. Substantially, the tensile fracture mechanism of the EM Ti-5111 alloy is a mixed fracture mechanism including ductile and cleavage fractures. In addition, some secondary cracks and pits were also found in the fracture morphology, due to the uneven plastic deformation and stress concentration [32,33]. Figure 9(b₁–b₃) show the macroscopic and microscopic fracture morphology of the BM Ti-5111 alloy. A notable necking phenomenon is observed near the fracture surface, distinguishing it from the other two types of alloys. It can be found in Figure 9(b₂,b₃) that the fracture exhibits numerous deep equiaxed dimples. Moreover, a notable phenomenon is observed where a large dimple encompasses a smaller one, as indicated by the white dotted circle. Furthermore, no cleavage steps or secondary cracks were identified. The above phenomena indicate that the fracture mechanism of the BM alloy is a typical ductile fracture, showcasing the alloy’s superior plasticity. The macroscopic fracture morphology of the alloy with a lamellar structure is shown in Figure 9(b₁). It can be found that the sample fractures along a 45° direction from the tensile axis, without significant necking, which correlates with its inferior properties. Figure 9(b₂,b₃) show the microscopic fracture morphology of the LM alloy. Unlike the EM and BM alloys, the fracture morphology of the LM alloy lacks dimples but exhibits distinct features such as cleavage steps, tongue shapes and tearing edges. The fracture morphology of the LM alloy presents typical cleavage fracture characteristics. These features collectively suggest that the LM alloy has limited plasticity [34], which is consistent with tensile test results.

5. Conclusions

This study comparatively investigated the tensile properties and deformed microstructure of the Ti-5111 alloy with equiaxed, bimodal and lamellar structures at room temperature. The influence mechanism of its microstructure on mechanical properties was analyzed, and the fracture mechanisms of the three Ti-5111 alloys were discussed. The key findings of the present study are summarized as follows:

(1)

The EM Ti-5111 alloy demonstrates moderate strength and plasticity, whereas the LM Ti-5111 alloy shows poor mechanical performance. Compared with the EM and LM alloys, the BM Ti-5111 alloy exhibits the most superior strength and plasticity. To achieve excellent mechanical properties in the Ti-5111 alloy, such as a fine balance of cold plasticity and strength, it should undergo a heat treatment process of 980 °C × 0.5 h/AC to obtain the BM Ti-5111 alloy.

(2)

The high strength of the BM alloy can be attributed to its small grain size and a large number of α_s/β phase boundaries. The exceptional coordinated deformation between the lamellar α_s and β phases plays a crucial role in the remarkable plasticity of BM alloys. In addition, the moderate strength of the EM alloy is influenced by the dispersed short needle-like α phase, which, although contributing to strength, hinders plasticity, resulting in suboptimal mechanical properties. The thicker lamellar α and β phases, the presence of the α_GB phase and the parallel arrangement of the α_s phase are the primary factors behind the inferior mechanical properties of the LM alloy.

(3)

The fracture mechanism of the EM alloy is a mixed fracture mechanism including ductile and cleavage. The fracture mechanism of the BM alloy is typical of ductile fractures. Different from the EM and BM Ti-5111 alloys, the fracture mechanism of the LM alloy is a typical cleavage fracture.