Effect of Secondary α Phase on Stress Corrosion Cracking of a Novel Metastable β Titanium Alloy in 3.5% NaCl Solution

Abstract

The effect of the secondary α phase on stress corrosion cracking of a novel metastable β titanium alloy, Ti-6Mo-5V-3Al-2Fe, in 3.5% NaCl solution was investigated by slow strain rate testing. Fine acicular secondary α phase was obtained by aging at the low temperature of 520 °C, and coarsened rod-like secondary α phase was obtained by aging at the high temperature of 680 °C. The electrochemical measurement results and slow strain rate testing results show that the microstructure contained with fine acicular secondary α phase exhibits better corrosion resistance and less stress corrosion cracking susceptibility. The fracture morphology exhibits a mixed fracture characteristic with shallow and small dimples, as well as tear ridges and flat facets with undulating surfaces. The combination of Absorption Induced Dislocation Emission and Hydrogen Enhanced Localized Plasticity is the main mechanism for stress corrosion cracking. Fine acicular secondary α phase with narrow spacing leads to less accumulated dislocations and smaller localized stress, so that has a beneficial effect on stress corrosion performance.

1. Introduction

Titanium alloys exhibit high specific strength, sufficient fatigue resistance, and good toughness and are widely used for structural applications in many engineering fields [1,2,3]. Particularly, the superior corrosion resistance and excellent strength of lead titanium alloys are ideal for marine engineering applications [4]. For instance, titanium alloys have been used to make oil country tubular goods in the offshore petroleum exploration industry and cooling systems in seawater-cooled plants [5,6]. Seawater contains various kinds of solutes and microorganisms and gives a corrosion environment for metals [7]. A protective oxide film forms on the surface of titanium alloys because of spontaneous passivation, resulting in the corrosion resistance of seawater. However, the oxide film can be ruptured under the ongoing mechanical load and lead to the unprotected metal matrix being corroded by seawater. Therefore, titanium alloys for marine engineering structural applications may be suffered from stress corrosion cracking (SCC).

Many efforts have been made to investigate the SCC of titanium alloys. Pazhanivel et al. [8] investigated the SCC of Ti-6Al-4V alloy in an aqueous 3.5% NaCl environment by performing a slow strain rate test. They demonstrated that the loss in elongation of the specimen in an aqueous 3.5% NaCl environment is attributed to corrosion susceptibility and hydride formation at the α/β interface. Leon et al. [9] proposed that the decreased amount of α phase and increased amount of β phase has a beneficial effect on the structural integrity of the passive protective film. This is one of the reasons why the stress corrosion performance of Ti-6Al-4V alloy in 3.5% NaCl aqueous solution is improved. Several findings show that corrosive attack mainly occurs at the α/β interface because of the micro-galvanic effect between α and β phases so that the SCC tends to take place at such place in a 3.5% NaCl environment [10,11,12]. Li et al. [13] found that the plastic deformation behaviors are different between α phase and β phase of TC2 alloy, as well as affect the hydrogen embrittlement. The SCC susceptibility of TC2 alloy is accelerated in simulated seawater due to the enhanced anodic dissolution and hydrogen embrittlement affected by deformations. Cao et al. [14] investigated the SCC properties of Ti-8Al-1Mo-1V alloy and proposed that α2 precipitates promote localized planar slip of SCC facets so that SCC accelerates. Dong et al. [15] investigated the stress corrosion cracking of TC4 ELI alloy with different microstructures in a 3.5% NaCl solution. The TC4 ELI alloy with lamellar microstructure exhibited significantly higher stress corrosion sensitivity than the alloy with equiaxed microstructure. The research work of Gao et al. [16] confirmed that the titanium alloy Ti6321 has stress corrosion sensitivity in a 3.5% NaCl solution. Junior et al. [17] investigated the SCC behavior of Ti-6Al-4V alloy at different aging stages and found that the SCC crack propagation was preferential through prior β grains. Ahn et al. [18] found that the change in cooling rate can lead to the random orientation of acicular α platelets, resulting in different SCC susceptibility.

The hexagonal close-packed structural α phase and body-centered cubic structural β phase exist together for titanium alloys. The contrasting crystal structure between α phase and β phase leads to different deformation behaviors under continuous stress loading [19]. Furthermore, it has been reported that the different volume fractions and morphology of α phase and β phase leads to different corrosion behaviors in titanium alloys [20]. Therefore, the α phase can greatly influence the SCC behaviors of titanium alloys. Some findings mentioned above [8,9,13,14,15,18] also indirectly confirmed this deduction. For metastable β titanium alloy, the precipitation of the secondary α phase is the main factor for alloy strengthening. The secondary α phase can be regulated by heat treatment and strongly affect the strength and ductility [21]. However, less attention has been paid to the effect of the secondary α phase on SCC behavior in metastable β titanium alloy. Ti-6Mo-5V-3Al-2Fe alloy is a novel metastable β titanium alloy, which shows potential for high strength [22]. However, limited available information on the relationship between the secondary α phase and SCC behavior exists at present. Therefore, this work is to investigate the effect of the secondary α phase on stress corrosion cracking of the Ti-6Mo-5V-3Al-2Fe alloy in a 3.5%NaCl solution. The objective of the present work is to improve the understanding of the SCC mechanism in the alloy, as well as to provide meaningful references for its marine engineering application.

2. Materials and Methods

The experimental materials are the Ti-6Mo-5V-3Al-2Fe (wt.%) alloy. The alloy ingot was produced by triple vacuum arc remelting. An alloy ingot with 120 mm in diameter and 200 mm in height was prepared. Then, β forging was carried out to obtain the alloy plates. The final forging temperature was 20 °C above β-transus (870 °C). The alloy plates with 90 mm in width and 32 mm in thickness were obtained. Subsequently, the as-forged alloys were heat treated to precipitate the secondary α phase. Two aging treatments were adopted to obtain the secondary α phase with large characteristic differences. According to the literature [22], the fine secondary α phase would precipitate during aging at the temperature range of 450–550 °C. When the aging temperature is larger than 600 °C, the secondary α phase would be significantly coarsened. Thus, the specimens were divided into two sets during heat treatment. One set was aged at a low temperature of 520 °C for 8 h, and another set was aged at a high temperature of 680 °C for 8 h. The purpose of adopting different aging temperatures was to obtain a secondary α phase with the difference in microstructural characteristics.

The chemical composition of the alloys before and after heat treatments were analyzed by ARL PERFORM X4200 X-ray fluorescence spectrometer. The microstructure was characterized by ZEISS GeminiSEM300 scanning electron microscope (SEM, Carl Zeiss AG, Oberkochen, Germany) and electron backscattered diffraction (EBSD, Oxford Instruments, Oxford, Britain). The specimens for SEM detection were cut from the aged alloy plates by electrical discharging machining and ground on metallographic sandpaper. Subsequently, the specimens were electropolished at~−15 °C in a solution of 6% perchloric acid, 35% n-butyl alcohol and 59% methanol.

The corrosion resistance of the aged alloy was gauged by potentiodynamic polarization and electrochemical impedance spectroscopy (EIS, Bio-Logic, Grenoble, France). Electrochemical measurements were performed at VSP-300 electrochemical experimental system in 3.5 wt.% NaCl aqueous solution. A conventional three-electrode cell was used in electrochemical tests. The specimen was used as the working electrode, a saturated calomel electrode was used as the reference electrode, and a carbon rod was used as the counter electrode. The open-circuit potential was measured at first. Potentiodynamic polarization curves were measured from the cathodic and anodic sides of corrosion potential, respectively. The scanning rate of the potentiodynamic polarization curves was 0.5 mV/s. EIS measurements were performed using a frequency range from 100 kHz and 10 mHz at a 10 mV amplitude over the open-circuit potential.

Slow strain rate testing (SSRT) of the aged alloys was carried out to investigate the SCC behaviors. SSRT was conducted at a WDML-10 slow strain rate stress corrosion testing machine with a strain rate of 1 × 10⁻⁶ s⁻¹ in the air, as well as in 3.5% NaCl aqueous solution, respectively. Before tests, the specimens were immersed for 12 h to ensure a stable surface condition. Except for the standard distance area, the rest of the area was sealed with epoxy resin, installed in the corrosion solution tank and further sealed with silicone rubber. The fracture morphology of the SSRT specimens was observed by using an S-3400N scanning electron microscope (Hitachi, Tokyo, Japan).

3. Results

3.1. Chemical Composition and Microstructure

The chemical composition of the alloys before and after heat treatments are shown in

Table 1. Chemical composition of the alloys before and after heat treatments.

Heat Treatment	Mo (wt.%)	V (wt.%)	Al (wt.%)	Fe (wt.%)	Ti (wt.%)
Before heat treatment	6.24	5.19	2.98	2.14	Bal.
aged at 520 °C for 8 h	6.13	5.12	2.99	2.22	Bal.
aged at 680 °C for 8 h	6.22	5.15	2.98	2.19	Bal.

. It can be seen that the chemical composition did not change significantly after the aging treatment. The microstructure of the aged alloy is shown in Figure 1. After aging treatment, a large amount of dispersed secondary α phase precipitate on the β matrix. When being aged at a relatively lower temperature of 520 °C, the precipitated secondary α phase is acicular with a high aspect ratio and has a large quantity (Figure 1a). When a relatively high aging temperature of 680 °C is adopted, the precipitated secondary α phase is coarsened and presents a rod-like shape with a low aspect ratio, and its quantity is reduced (Figure 1b). Similar aging response phenomena have been found in other metastable β titanium alloys, such as Ti-5Al-4Zr-8Mo-7V alloy [23], Ti-3.5Al-5Mo-6V-3Cr-2Sn-0.5Fe alloy [24], and Ti-5Al-3Mo-3V-2Cr-2Zr-1Nb-1Fe alloy [25]. The microstructures shown in Figure 1a and Figure 1b are referred to as “AM1” and “AM2” in the following parts, respectively. The morphology and volume fraction of the secondary α phase identified by the electron backscattered diffraction (EBSD) are listed in

Table 2. Microstructural characteristic of α phase in the aged alloys.

No.	Aging Treatment Process	Characteristic of α Phase	Volume Fraction of α Phase/%
AM1	520 °C for 8 h	Fine and narrow spaced	33.2
AM2	680 °C for 8 h	Coarsened and widely spaced	38.7

. It can be found that the secondary α phase volume fraction of the AM2 alloy is higher than that of the AM1 alloy. Additionally, the spacing of coarsened rod-like secondary α phase is significantly wider than that of the fine acicular secondary α phase.

3.2. Polarization Curves and EIS Analysis

The potentiodynamic polarization curves of the AM1 and AM2 alloys in a 3.5% NaCl aqueous solution are shown in Figure 2. According to Figure 2, the corrosion potential (E_corr) and corrosion current density (I_corr) are identified by Tafel extrapolation and are summarized in

Table 3. The corrosion potential, corrosion current density and passive current density.

Alloys	Ecorr (mV)	Icorr (μA/cm2)
AM1	−356	0.445
AM2	−412	1.730

. The E_corr of the AM1 alloys exhibits a higher value of −356 mV, as compared to the E_corr of −412 mV of the AM2 alloys. Moreover, the I_corr of the AM1 alloys is 0.445 μA/cm², which is lower than that of the AM2 alloys with 1.73 μA/cm². The relatively low current density values for corrosion demonstrate the relatively better corrosion resistance of the AM1 alloys compared to the AM2 alloy.

The Nyquist and Bode plots for the AM1 and AM2 alloys in a 3.5% NaCl aqueous solution are shown in Figure 3. According to Figure 3a, the capacitive loop radius for the AM2 alloys is smaller than that for the AM1 alloys. From Figure 3b, the impedance modulus |Z| at the frequency of 0.01 Hz for the AM2 alloys, which can reflect the characterization of the corrosion protection of the oxide film on the alloys, is much less than that for the AM1 alloys, although the two phase-angle plots are partially superimposed. All of these EIS results reflect that the compact property of the oxide film growing on the surface of the alloys is worse at AM2 alloys than that at AM1 alloys. According to the microstructural divergence between AM1 and AM2 alloys, the abovementioned difference in passivation and corrosion resistance may be related to the difference in microstructural characteristics of the secondary α phase.

3.3. SSRT Results

The stress–strain curves of SSRT, which were tested in air and in a 3.5% NaCl aqueous solution, are shown in Figure 4. It can be found that the strength and elongation of the specimens tested in a corrosive environment obviously decreased compared with the specimens tested without a corrosive environment.

In order to quantitatively evaluate the stress corrosive performance of the AM1 and AM2 alloys, an SCC susceptibility index (I_scc) is calculated according to the following formula [26].

I_scc = [(E_air − E_sol)/E_air] × 100%

(1)

where E_air is the elongation of SSRT in air, and E_sol is the elongation of SSRT in the corrosive environment. The SCC susceptibility index of the AM1 and AM2 alloys is 3.5% and 8.8%, respectively. The increased I_scc values clearly indicate that the SCC susceptibility of the AM2 alloys is more than that of the AM1 alloys. It can be speculated that the coarsened and widely spaced secondary α phase leads the alloy to be more susceptible to SCC.

3.4. Fracture Morphologies

The fracture morphologies of the AM1 alloys after SSRT are shown in Figure 5. The dimples on the fracture surface of the AM1 alloys tested in the air are deep and large, as shown in Figure 5a. For the AM1 alloys tested in 3.5% NaCl aqueous solution, the fracture morphology exhibits a mixed characteristic of ductile fracture and brittle fracture, as shown in Figure 5b. It is noteworthy that the dimples located in the ductile region present a shallow and small morphology under the influence of the corrosion environment, as shown in Figure 5c. The fracture morphologies of the AM2 alloys after SSRT tested in a 3.5% NaCl aqueous solution are shown in Figure 6. Similar to the AM1 alloys tested in a 3.5% NaCl aqueous solution, a mixed fracture characteristic of ductile and brittle can be observed in the AM2 alloys, as shown in Figure 6a. In addition to the shallow and small dimples, tear ridges (indicated by the white arrow) can be observed in the ductile region, as shown in Figure 6b. Moreover, a flat facet (circled by the dashed line) with undulating surfaces surrounded by dimples can be observed, as shown in Figure 6c. These special fracture morphologies provide threads for SCC mechanism analysis of the alloys.

4. Discussion

The precipitation behavior of the secondary α phase is dominated by the transformation driving force in β titanium alloy [27]. Lower aging temperature provided a higher undercooling degree as well as brought about a larger transformation driving force and higher nucleation rate. Moreover, lower aging temperature leads to a slow elements diffusion rate, resulting in a slow growth rate of the secondary α phase. Thus, the relatively lower aging temperature of 520 °C leads to fine acicular α precipitates. On the contrary, the larger driving force for the growth of the secondary α phase would be provided by a higher aging temperature. Meanwhile, the nucleation driving force could be reduced because of the lower undercooling. Then, coarsened secondary α phase precipitate under a relatively higher aging temperature of 680 °C. Furthermore, a region with enriched β-stabilizer forms in the vicinity accompanied by the precipitation of the secondary α phase due to the low tolerance of β-stabilizers. Such a region is so stable that hard to precipitate. Thus, coarsened precipitates lead to wider interspacing of the secondary α phase in AM2 alloys, compared with that in AM1 alloys.

Many research results [14,28] show that the SCC mechanism for titanium alloys is related to absorbed hydrogen. For the titanium alloys in 3.5% NaCl aqueous solution, two electrochemical actions take place simultaneously, i.e., the formation of an oxide film and the dissolution of the bare titanium alloy matrix [29]. The hydrogen for external absorption was produced by the chemical reaction that occurred on the alloy surface during the formation of the oxide film and the dissolution of bare metal. The occurred reactions during the repassivation of titanium alloys are shown in Equations (2)–(4), the reactions during the anodic dissolution are shown in Equations (5)–(7), and the production of hydrogen is shown in Equation (8) [30].

Ti → Ti³⁺ + 3e⁻

(2)

Ti³⁺ + 2H₂O → TiO²⁺ + 2H⁺ + e⁻

(3)

Ti + 2H₂O → TiO₂ + 4H⁺ + 4e⁻

(4)

Ti → Ti²⁺ + 2e⁻

(5)

Ti²⁺ + H⁺ → [Ti³⁺ H]_ads

(6)

[Ti³⁺ H]_ads → Ti³⁺ + H⁺ + e⁻

(7)

H⁺ + e⁻ → H

(8)

As shown in the area marked by the dotted box of polarization curves (Figure 2), the anodic current densities of both AM1 and AM2 alloys slowly increase with the increase in potential. It can be inferred that the slow dissolution of the metal occurs instead of maintaining passivation, even if the oxide films are grown on the metal surfaces. The same phenomenon also exists on the potentiodynamic polarization curves of Ti-6Al-4V alloy tested in 3.5 wt.% NaCl aqueous solution [30]. Therefore, titanium dissolution can be promoted. The anodic current density of the AM2 alloys increases faster than that of the AM1 alloys (the dotted box of Figure 2). Moreover, the corrosion potential of the AM2 alloys is more negative. Thus, the reactions during the anodic dissolution that occurred in the AM2 alloys would be more intense, resulting in a higher hydrogen concentration.

The fracture morphologies of the alloys tested in 3.5% NaCl aqueous solution present two typical characteristics:

(1) On the fracture surfaces of the alloys tested in 3.5% NaCl aqueous solution, shallow and small dimples could be observed (Figure 5c and Figure 6b). The formation of dimples with such characteristics has been clarified by an Absorption Induced Dislocation Emission (AIDE) mechanism [31]. It is proposed that intermetallic bond is weakened by absorbed hydrogen resulting in acceleration of dislocation emissions from crack tips [32]. The increasing dislocation density around the crack tip leads to a resharpening of the crack tip so that shallow and small dimples are formed.

(2) Tear ridges, as well as flat facets with undulating surfaces, could be observed on the fracture surfaces of the alloys tested in 3.5% NaCl aqueous solution (Figure 6b,c). It can be inferred that high-degree dislocation movement and localized plastic deformation occurred around cracks [14]. A Hydrogen Enhanced Localized Plasticity (HELP) mechanism can be used to explain such a phenomenon. The adsorbed hydrogen atom spread quickly to the front of crack tips through a continuous β phase, resulting in an increase in solute hydrogen concentration around crack tips [30]. The interaction between hydrogen and dislocation stimulates dislocation activation and then leads to localized plastic deformation [30].

In summary, the mixed fracture mode with shallow dimples and localized plasticity mainly contributed to the function of absorbed hydrogen, which is dominated by AIDE and HELP mechanisms. Moreover, these two mechanisms work together to trigger the SCC of the alloy in a 3.5% NaCl aqueous solution.

For metastable β titanium alloy during deformation, dislocations tend to activate within β matrix rather than within slender α precipitates. In addition, the interfaces between α phase and β matrix can act as an obstacle for dislocations slip because of a considerable lattice mismatch [33]. Then, A pile-up of dislocations can form in front of the interfaces between α phase and β matrix. The localized stress resulting from a pile-up of dislocations τ is as Equation (9) [34].

τ = Nτ_sb_T

(9)

where N is the number of accumulated dislocations, τ_s is the applied stress, and b_T is the Burgers vector of the alloy.

The number of accumulated dislocations N is as Equation (10) [23].

$$N=\pi {\tau }_s(1-v_T)/G_T{b}_T·\lambda /2$$

(10)

where v_T is Poisson’s ratio of the alloy, G_T is shear modulus of the alloy, λ/2 is the length of the pile-up zone of dislocations, and λ is the length of α phase spacing.

According to Equations (9) and (10), there is a positive correlation between τ, N, and λ. Thus, the increase in α phase spacing leads to the increase in accumulated dislocations quantity, which results in the increase in localized stress. Compared with the AM1 alloys, the AM2 alloys have wider secondary α phase spacing (Figure 1). The localized stress resulting from dislocations pile-up of the AM2 alloys is higher than that of the AM1 alloys. This stress concentration resulting from dislocation pile-up is also the initiation of cracks. Moreover, because of the action of absorbed hydrogen, the interatomic band is weakened, and the emission and motion of dislocations are promoted by AIDE and HELP mechanisms. Therefore, the AM2 alloys are more prone to crack initiation. In addition, worse corrosion resistance leads to a higher hydrogen concentration around the AM2 alloys. As a result, the AM2 alloys are more tend to occur SCC. Thus, the influence mechanism of the secondary α phase on SCC for the alloy in a 3.5% NaCl aqueous solution is proposed, shown in Figure 7. It can be found that compared with the coarsened rod-like secondary α with wide spacing, fine acicular secondary α with narrow spacing has a beneficial effect on the stress corrosion performance of the alloy.

5. Conclusions

In this work, the effect of the secondary α phase on stress corrosion cracking in 3.5% NaCl aqueous solution of a novel metastable β titanium alloy, Ti-6Mo-5V-3Al-2Fe, was investigated. The following conclusions can be drawn from this work;

(1) Compared with the alloy containing coarsened and widely spaced secondary α phase, the alloy containing fine and narrow spaced secondary α phase exhibits better corrosion resistance and less SCC susceptibility.

(2) The SCC mechanism of the alloy is a combination of Absorption Induced Dislocation Emission and Hydrogen Enhanced Localized Plasticity.

(3) Fine acicular secondary α phase with narrow spacing has a beneficial effect on stress corrosion performance. An influence mechanism of the secondary α phase on SCC for the alloy in a 3.5% NaCl aqueous solution is proposed.