Investigation of High-Temperature Constitutive Behavior of Ti555211 Titanium Alloy Subjected to Plastic Deformation in the Different Phase Regions

Abstract

Ti555211 titanium alloy is subjected to plastic deformation in the dual-phase (α + β phase) zone and single-phase (β phase) zone at various deformation temperatures and strain rates. High-temperature constitutive equations of the alloy in the dual-phase zone and single-phase zone are established in order to describe deformation behavior of the alloy in the different phase zones. By comparing the constitutive equation of the alloy in the dual-phase zone with that of the alloy in the single-phase zone, the deformation activation energy of the former was found to be higher than that of the latter. It is obvious that the deformation activation energy of α phase is obviously greater than that of β phase. Furthermore, the microstructural evolution of the alloy is different in the dual-phase zone and single-phase zone. When the alloy was subjected to plastic deformation in the dual-phase zone, the size of the grains in the β phase increased with the decreasing strain rate. When the alloy was subjected to plastic deformation in the single-phase zone, the size of the grains in the β phase considerably increased with the increasing deformation temperature. In particular, in the microstructures of the alloy subjected to plastic deformation in the single-phase region, the elongated grains can be observed at higher strain rates. Furthermore, it is more difficult for the alloy to induce plastic deformation in the dual-phase region than in the single-phase region.

1. Introduction

As an excellent structural material, Titanium alloys are extensively applied in the fields of aerospace, oceans, petrochemistry, and so on because it possesses excellent high-temperature properties, outstanding resistance to fatigue, low density, high resistance to corrosion, and high specific strength [1,2,3,4,5]. To date, Titanium alloys that have been successfully applied mainly deal with metastable β type Ti-15-3 [6], β-21S [7], and near-β type Ti1023 [8], Ti5553 [9], Ti55531 [10], BT22 [11], TA10 [12], and so on. Many scholars have devoted themselves to investigating Titanium alloys in terms of properties [13], deformation behavior [14], and heat treatment regime [15]. Wu et al. [16] investigated the microstructural evolution and mechanical response of a new high-strength and high-ductility near-β titanium alloy (Ti6554) during an electropulsing-assisted microtension process. Shi et al. [17] investigated the deformation behavior and recrystallization mechanism of a near-β Ti-55511 titanium alloy in the β phase region at a high temperature, where they observed continuous dynamic recrystallization (CDRX) resulting from progressive sub-grain rotation, and discontinuous dynamic recrystallization (DDRX) resulting from grain boundary bulging, and the DRX mechanism was transformed from DDRX to CDRX with an increasing deformation temperature. Ratochka et al. [18] investigated the low-temperature superplasticity of ultrafine-grained Ti-5Al-5V-5Mo-1Cr-1Fe alloy at 550 ℃, which is obtained through multi-axial pressing. Yanchun Zhu et al. [19] investigated strain hardening exponents and strain rate sensitivity exponents of TA15, Ti40, and TC21 titanium alloys and comparatively investigated the hot processing property of the three typical titanium alloys.

A constitutive model of materials can be regarded as a mathematical model that can reflect the macroscale mechanical behavior of materials and, generally, it is used to describe the intrinsic nature of materials. A constitutive model of materials has become an effect tool for investigating the plastic deformation of metal materials and it has been widely used in the field of metal forming [20]. So far, the widely used constitutive models include the physical model and the phenomenological model [21]. The phenomenological constitutive model reflects plastic deformation behavior by means of numerical calculation and analysis and it is used for establishing the relationships between flow stress, strain rate, and deformation temperature in the course of thermal processing [22]. For example, Yuan et al. [23] investigated the constitutive behavior of Ti-6Al-4 alloy in the broader scope of temperature (800–959 °C), strain (0.1–0.5), and strain rate (0.0005–1 s⁻¹) based on stress-strain data from thermal compressive experiments. Zhao et al. [24] used constitutive equations and an artificial neural network to establish the constitutive model of the hot deformation behavior of Ti600 titanium alloy.

Ti555211 alloy is a kind of new, near-β type titanium alloy that is researched and developed by Western Superconducting Technologies Co., Ltd. The nominal chemical composition of Ti555211 titanium alloy is Ti-5.5Al-5Mo-5V-2Nb-1Fe-1Zr. Ti555211 titanium alloy possesses high hardenability and excellent comprehensive properties and thus it is suitable for manufacturing large-scale components, such as fuselage frames, landing gear, and frame rails. In general, high-temperature plastic forming is a necessary method to manufacture titanium alloy components and the involved process parameters have an important influence on the microstructures and properties of titanium alloy. High-temperature plastic forming of near-β type titanium alloy can be performed in the α + β phase, β phase, or cross-phase regions. Hence, it is significant to investigate the high-temperature plastic deformation behavior of Ti555211 titanium alloy in the various phase regions.

In the present work, high-temperature constitutive equations of Ti555211 titanium alloy in the dual-phase (α + β phase) zone and single-phase (β phase) zone were established in order to describe the deformation behavior of Ti555211 titanium alloy in different phase zones. Based on the different phase structures, Ti555211 titanium alloy presents different plastic deformation behavior and microstructural evolution.

2. Materials and Methods

In the present work, Ti555211 titanium alloy was used as a research object derived from Western Superconducting Technologies Co., Ltd. (Xi’an Shaanxi, China). The chemical composition of Ti555211 titanium alloy is illustrated in

Table 1. The chemical composition of Ti555211 titanium alloy (wt%).

Ti	Al	Mo	V	Nb	Fe	Zr
Bal	5.8	4.5	5.3	1.9	1.0	0.9

. In order to investigate the plastic deformation behavior of Ti555211 titanium alloys with different phase structures at high temperatures, thirty-two Ti555211 alloy samples with a height of 12 mm and a diameter of 8 mm were subjected to uniaxial compressive tests on the Gleeble3800 (Dynamic Systems Inc., Poestenkill, NY, USA) thermal-mechanical test machine using a 60% reduction in height, where the deformation rates were selected as 0.005, 0.05, 0.5, and 5 s⁻¹ and the deformation temperatures were chosen as 700, 750, 800, 850, 900, 950, 1000, and 1050 °C. Before compression, the compressive samples were annealed for 5 min at a given temperature in order to ensure the transformation from the room-temperature microstructure to the high-temperature equilibrium microstructure. During compression, the friction between the edges of the compressive samples and the dies was considered in order to avoid the influence of friction on the true stress–strain curves, in reference to past research [25,26]. In addition, the critical temperature of β-transus for Ti555211 titanium alloy was found to be about 875–880 °C by means of measuring the phase transformation point based on the metallographic method. In other words, Ti555211 titanium alloy belonged to the dual-phase structure at the range of 700–850 °C, which was composed of the α and β phases, whereas Ti555211 titanium alloy belonged to the single-phase structure at the range of 900–1050 °C, which consisted of the β phase. To analyze the microstructural evolution of Ti555211 alloys during plastic deformation at high temperatures, microstructures of the involved compressed Ti555211 alloy samples were characterized by means of an EBSD experiment. Before the EBSD experiment, the corresponding EBSD samples were subjected to electrolytic polishing, where the electrolyte was composed of 10% perchloric acid and 90% alcohol, and polishing temperature and polishing time were selected as 20 °C and 60 s, respectively.

3. Results and Discussion

3.1. Deformation Behavior of Ti555211 Alloy at High Temperatures

Stress–strain curves of Ti555211 alloy during plastic deformation at high temperatures can effectively reflect its deformation behavior [27]. Figure 1 shows the compressive stress–strain curves of dual-phase Ti555211 alloy at various strain rates and deformation temperatures. Figure 2 illustrates the compressive stress–strain curves of single-phase Ti555211 alloy at various strain rates and deformation temperatures.

As can be seen in Figure 1 and Figure 2, it was found that the stress–strain curves of dual- and single-phase Ti555211 alloys basically present similar variation trends, where the flow stresses decreased with the decreasing strain rate and they increased with decreasing deformation temperatures. In general, the smaller strain rate resulted in increasing the deformation time, which contributed to providing sufficient time for dislocation annihilation and grain boundary migration. Furthermore, increasing the deformation temperature contributed to the occurrence of dislocation annihilation, grain boundary migration, and dynamic recrystallization as well. These two reasons are responsible for the decrease in flow stresses. Simultaneously, it can be noted that Ti555211 alloy is sensitive to strain rates in single-phase and dual-phase zones.

3.2. Establishing Constitutive Equation of Ti555211 Alloy at High Temperatures

3.2.1. Constitutive Equation of Dual-Phase Ti555211 Alloy

Based on the Arrhenius equation, the constitutive equation of dual-phase Ti555211 alloy at high temperatures can be established. The Arrhenius equation is expressed as follows:

$$\dot{\epsilon }=\mathrm{A}{\left[\mathrm{s}\mathrm{i}\mathrm{n}\mathrm{h}\left(\mathsf{\alpha }\sigma \right)\right]}^{\mathrm{n}}\mathrm{e}\mathrm{x}\mathrm{p}\left(-\frac{\mathrm{Q}}{\mathrm{R}T}\right)$$

(1)

where $\dot{\epsilon }$ and $\sigma$ stand for strain rate (s⁻¹) and stress (MPa), respectively. T represents the absolute temperature (K). R and Q represent the gas constant 8.314 (J·mo⁻¹·K⁻¹) and activation energy, respectively. A, n, and α denote the material constants.

In order to solve the high-temperature constitutive equation of Ti555211 alloy, Equation (1) can be simplified as follows:

$$\dot{\epsilon }={\mathrm{A}}_1{\sigma }^{\mathrm{n}}\exp \left(-\frac{\mathrm{Q}}{\mathrm{R}T}\right),\mathsf{\alpha }\sigma \le 0.83$$

(2)

$$\dot{\epsilon }={\mathrm{A}}_2\exp (\mathsf{\beta }\sigma )\exp \left(-\frac{\mathrm{Q}}{\mathrm{R}T}\right),\mathsf{\alpha }\sigma \ge 1.2$$

(3)

where A₁, A₂, and β represent the material constants and A₁ = Aαⁿ, A₂ = A/2ⁿ, β = nα. Taking a natural logarithm in two sides of Equations (2) and (3) results in:

$$\ln \dot{\epsilon }=\ln {\mathrm{A}}_1+\mathrm{n}\ln \sigma -\frac{\mathrm{Q}}{\mathrm{R}T}$$

(4)

$$\ln \dot{\epsilon }=\ln {\mathrm{A}}_2+\mathsf{\beta }\sigma -\frac{\mathrm{Q}}{\mathrm{R}T}$$

(5)

According to Equations (4) and (5), material constants n and α for the high-temperature constitutive equation of Ti555211 alloy can be obtained. The peak stresses need to be derived from the corresponding stress–strain curves of Ti555211 alloy in order to solve the material constants n and α, and they are listed in

Table 2. The peak stresses of dual-phase Ti555211 alloy under various deformation conditions (MPa).

$\dot{\epsilon }$/s−1	$\mathit{T}$/°C
	700	750	800	850
0.005	253.41902	140.39049	86.42898	52.95557
0.05	311.06395	224.55433	137.64007	96.48557
0.5	429.17	302.23543	228.28947	156.84211
5	543.25127	380.24901	319.55999	245.62818

. Consequently, constants n and α can be solved by means of linear fitting, as shown in Figure 3, where n = 6.3133752 and β = 0.0288975, and thus α = β/n = 4.577 × 10⁻³ MPa⁻¹.

After soling the material constant α, it is necessary to optimize the material constraint n. Taking a natural logarithm in two sides of Equation (1) leads to:

$$\ln \dot{\epsilon }=\ln \mathrm{A}+\mathrm{n}\ln [\sinh \left(\mathsf{\alpha }\sigma \right)]-\frac{\mathrm{Q}}{\mathrm{R}T}$$

(6)

Based on Equation (6), linear fitting was implemented, as shown in Figure 4. As a result, the optimized n value can be solved and it is equal to 4.4677075.

Transforming Equation (6) can result in:

$$\ln [\sinh \left(\mathsf{\alpha }\sigma \right)]=\frac{\mathrm{Q}}{\mathrm{nR}}\times \frac{1}{T}+{\mathrm{C}}_1$$

(7)

where ${\mathrm{C}}_1=\frac{\ln \dot{\epsilon }-\ln \mathrm{A}}{\mathrm{n}}$. Linear fitting was performed based on Equation (7) so that the valve of activation energy Q is determined as 4.1517972 × 10⁵ J·mol⁻¹.

The Zener–Hollomon parameter (generally called parameter $Z$) can comprehensively reflect the impact of deformation temperature and strain rate on the constitutive behavior of Ti555211 alloy. Parameter $Z$ is expressed as follows:

$$Z=\dot{\epsilon }\exp \left(\frac{\mathrm{Q}}{\mathrm{R}T}\right)$$

(8)

Substituting parameter $Z$ into Equation (1) results in:

$$Z=\mathrm{A}{\left[\sinh \left(\mathsf{\alpha }\sigma \right)\right]}^{\mathrm{n}}$$

(9)

Taking a natural logarithm in two sides of Equation (9) results in:

$$\ln Z=\mathrm{lnA}+\mathrm{nln}\left[\sinh \left(\mathsf{\alpha }\sigma \right)\right]$$

(10)

According to Equation (10), linear fitting was implemented, as shown in Figure 5. As a result, the value of A can be determined as 5.48224 × 10¹³.

According to the aforementioned procedures, the involved material constants can be obtained so that the constitutive equation of dual-phase Ti555211 alloy can be established as follows:

$$\dot{\epsilon }=3.84962\times {10}^{19}{\left[\sinh \left(4.577187\times {10}^{-3}\sigma \right)\right]}^{4.4677075}\exp \left(\frac{-4.1517972\times {10}^5}{\mathrm{R}T}\right)$$

(11)

By introducing parameter $Z$, the flow stress σ of dual-phase Ti555211 alloy can be expressed as follows:

$$\sigma =218.474702\ln \left\{{\left(\frac{Z}{5.48224\times {10}^{13}}\right)}^{1/4.4677075}+\sqrt{{(\frac{Z}{5.48224\times {10}^{13}})}^{2/4.4677075}+1}\right\}$$

(12)

where parameter $Z=\dot{\epsilon }\cdot \exp \left(\frac{\mathrm{Q}}{\mathrm{R}T}\right)=\dot{\epsilon }\cdot \exp \left(\frac{4.1517972\times {10}^5}{\mathrm{R}T}\right)$.

3.2.2. Constitutive Equation of Single-Phase Ti555211 Alloy

Using the same manner to establish the constitutive equation of dual-phase Ti555211 alloy, it is necessary to solve the corresponding material constants in order to establish the constitutive equation of single-phase Ti555211 alloy. Similarly, the peak stresses need to be obtained from the corresponding stress–strain curves of single-phase Ti555211 alloy, as shown in

Table 3. The peak stresses of single-phase Ti555211 alloy under the various deformation conditions (MPa).

$\dot{\epsilon }$/s−1	T/°C
	900	950	1000	1050
0.005	36.59946	25.08206	21.83291	19.01952
0.05	63.52716	51.86333	42.47453	37.75538
0.5	118.73231	96.78551	91.57046	69.62649
5	196.46293	158.29796	119.8755	110.84324

.

In the same manner, the material constants n and β were solved based on linear fitting. Subsequently, the optimized n₁ was obtained. Finally, the values of A and Q were acquired. The specific fitting curves are shown in Figure 6. The corresponding fitting values are as follows: n = 3.860095, β = 0.057495, α = 1.4894711 × 10⁻², n₁ = 2.816993, A = 4.6059749 × 10⁷, and Q = 2.0817179 × 10⁵ J/mol.

Consequently, the constitutive equation of single-phase Ti555211 alloy can be established as follows:

$$\dot{\epsilon }=4.6059749\times {10}^7{\left[\sinh \left(1.48947\times {10}^{-2}\sigma \right)\right]}^{2.816993}\exp \left(\frac{-2.081717\times {10}^5}{\mathrm{R}T}\right)$$

(13)

$$\sigma =67.137925\ln \left\{{\left(\frac{Z}{4.6059749\times {10}^7}\right)}^{1/2.816993}+\sqrt{{(\frac{Z}{4.6059749\times {10}^7})}^{2/2.816993}+1}\right\}$$

(14)

where parameter $Z=\dot{\epsilon }\cdot \exp \left(\frac{\mathrm{Q}}{\mathrm{R}T}\right)=\dot{\epsilon }\cdot \exp \left(\frac{2.081717\times {10}^5}{\mathrm{R}T}\right)$.

By comparing the constitutive equation of dual-phase Ti555211 alloy with that of single-phase Ti555211 alloy, it can be found that the deformation activation energy of the former is greater than that of the latter. It is obvious that dual-phase Ti555211 alloy is composed of the α and β phases, whereas single-phase Ti555211 alloy consists of the β phase. The phenomenon indicates that the deformation activation energy of the α phase is obviously greater than that of the β phase.

3.3. Microstructural Evolution of Ti555211 Alloy

Figure 7 and Figure 8 illustrate the microstructures and grain orientation distributions of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 750 °C, respectively. In Figure 8, the grains with the same color possess the same orientation. It can be found that the microstructures of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 750 °C are composed of refine equiaxed grains, which show a certain preferential orientation. Based on EBSD data, in the case of 750 °C, the average grain sizes of dual-phase Ti555211 alloy samples subjected to plastic deformation decreased with the increasing strain rate. Figure 9 shows the phase distribution diagrams of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 750 °C, where red represents the α phase and the green represents the β phase. In the case of plastic deformation at 750 °C, strain rate did not have a considerable effect on the content of the α phase in dual-phase Ti555211 alloy. According to Figure 7 and Figure 9, at the higher strain rate, the size of the grains in the β phase is less than that in the α phase. When the strain rate was reduced to 0.005 s⁻¹, the grains in the α phase and β phase showed a slighter size difference, with the dual-phase Ti555211 alloy presenting more homogeneous microstructures. In other words, during the plastic deformation of dual-phase Ti555211 alloy at high temperatures, the size of the grains in the α phase was not sensitive to the strain rate, whereas the size of the grains in the β phase increased with the decreasing strain rate. According to Figure 8 and Figure 9, it can be found that in the microstructures of dual-phase Ti555211 alloy undergoing plastic deformation at high temperatures, the orientations of the grains in the α phase are dominated by $\left\{2\overline{1}\overline{1}0\right\}$ and $\left\{10\overline{1}0\right\}$, whereas the orientations of the grains in the β phase are dominated by $\left\{001\right\}$ and $\left\{111\right\}$.

Figure 10 and Figure 11 illustrate the microstructures and grain orientation distributions of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 850 °C, respectively. It can be observed that the microstructures of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 850 °C are composed of many refine equiaxed grains and a small number of long grains. Based on EBSD data, in the case of 850 °C, the average grain sizes of dual-phase Ti555211 alloy samples subjected to plastic deformation decreased with the increasing strain rate. Figure 12 shows the phase distribution diagrams of dual-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 850 °C. It can be observed that the α phase is distributed in the matrix of dual-phase Ti555211 alloy in a dispersive manner. It can be noted that the content of the α phase obviously decreased with the increasing deformation temperature. Unlike dual-phase Ti555211 alloy samples undergoing plastic deformation at 750 °C, the inhomogeneity of the involved microstructures increased with the decreasing strain rate in the dual-phase Ti555211 alloy samples subjected to plastic deformation at 850 °C. In addition, according to Figure 10, Figure 11 and Figure 12, it can be observed that in the microstructures of dual-phase Ti555211 alloy samples subjected to plastic deformation at 850 °C, the grain of the β phase possesses a larger size than that of the α phase, where the former is composed of dominant equiaxed grains and a small number of elongated grains, whereas the latter is almost composed of refined equiaxed grains. In addition, the size of the β grain increased with the decreasing strain rate and the size of the α grain slightly varied with the decreasing strain rate. Furthermore, in dual-phase Ti555211 alloy samples subjected to plastic deformation at 850 °C, the orientations of the α grain are dominated by $\left\{2\overline{1}\overline{1}0\right\}$ and $\left\{10\overline{1}0\right\}$, whereas the orientation of the β grain is related to the strain rate and its orientation is dominated by $\left\{112\right\}$ at the lower strain rate and by $\left\{111\right\}$ at the higher strain rate.

Figure 13 illustrates the grain orientation distributions of single-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 950 °C, respectively. It can be that, compared to dual-phase Ti555211 alloy, single-phase Ti555211 alloy possesses the larger grain size. In addition, the grains are gradually elongated with the increasing strain rate. This phenomenon indicates that the dynamic recrystallization of single-phase Ti555211 alloy is implemented more completely at the lower strain rate. In addition, it was found that, in the case of plastic deformation at 950 °C, the orientation of the grains in the single-phase Ti555211 alloy is dominated by $\left\{001\right\}$.

Figure 14 illustrates the grain orientation distributions of single-phase Ti555211 alloy undergoing plastic deformation at various strain rates at 1050 °C, respectively. It can be observed that the microstructures in the single-phase Ti555211 alloy, deformed at 1050 °C, are similar to those of the counterpart deformed at 950 °C, whereas the average grain size of the former is obviously greater than that of the latter. It is evident that the average grain size decreased with the increasing strain rate. In addition, it was noted that, during plastic deformation at 1050 °C, the orientation of the grains in the single-phase Ti555211 alloy is dominated by $\left\{001\right\}$ and $\left\{111\right\}$.

According to the aforementioned experimental results, it can be concluded that the constitutive behavior and microstructural evolution of Ti555211 titanium alloy in the dual-phase zone and single-phase zone exhibit obvious differences. As for the constitutive equations, in particular, the values of Q and n in the different phase zones present a very large distinction. The values of Q and n in the α + β phase zone approximate to 415 kJ/mol and 4.47, respectively, whereas the values of Q and n in the β phase zone are about 208 kJ/mol and 2.82, respectively. It is evident that the values of Q and n in the α + β phase zone are higher than the counterparts in the β phase zone. In fact, Saboori et al. investigated the hot deformation behavior of Ti–6Al–4V alloy and found that the deformation activation energy of the wrought Ti–6Al–4V alloy in the β phase zone was about 229.34 kJ/mol [28]. It is obvious that our result is very close to the work of Saboori et al. [28]. This phenomenon indicates that the deformation activation energy of titanium alloys in the β phase zone is very close to lattice self-diffusion activation energy [28]. Motallebi et al. claimed that the deformation temperature regime with respect to the β-transus temperature for titanium alloys is an important parameter in determining hot deformation behavior [29]. In fact, the magnitude of the n value is closely related to deformation mechanisms when metal materials are subjected to hot deformation [30,31,32]. Savaedi et al. claimed that one deformation mechanism refers to the glide and climb of dislocations in the climb-controlled regime and it corresponds to a value of 4.5 or 5; the other deformation mechanism is focused on the glide and climb of dislocations in the viscous glide regime and it corresponds to a value of 3 or so [33]. In the current investigation, it is evident that Ti555211 titanium alloy shows different plastic deformation mechanisms in the different phase regions. In particular, plastic deformation mechanisms of Ti555211 titanium alloy in the dual-phase zone are more complex, where the deformation compatibility of α + β phases in the dual-phase zones has a significant influence on the deformation of Ti555211 titanium alloy. Plastic deformation mechanisms of Ti555211 titanium alloy at high temperatures will be further investigated in the future.

4. Conclusions

• High-temperature constitutive equations of Ti555211 titanium alloy in the dual-phase zone and single-phase zone were established in order to describe the deformation behavior of Ti555211 titanium alloy in the different phase zones. By comparing the constitutive equation of Ti555211 alloy in the dual-phase zone with that of Ti555211 alloy in the single-phase zone, it was found that the deformation activation energy of the former is greater than that of the latter. It is obvious that the activation energy of the α phase is obviously greater than that of the β phase;
• It is evident that the microstructural evolution of Ti555211 alloy is different in the dual-phase zone and single-phase zone. When Ti555211 alloy was subjected to plastic deformation in the dual-phase zone, the size of the grains in the α phase was not sensitive to the strain rate, whereas the size of the grains in the β phase increased with the decreasing strain rate. When Ti555211 alloy was subjected to plastic deformation in the single-phase zone, the size of the grains in the β phase considerably increased with the increasing deformation temperature;
• It was found that, compared to dual-phase Ti555211 alloy, single-phase Ti555211 alloy possesses the larger grain size. In addition, the grains are gradually elongated with the increasing strain rate. This phenomenon indicates that dynamic recrystallization of single-phase Ti555211 alloy is implemented more completely at the lower strain rate. Furthermore, it can be concluded that it is more difficult for Ti555211 alloy to induce plastic deformation in the dual-phase region than in the single-phase region.