Effect of Aging Treatment on the Microstructure and Mechanical Properties of Ti-3Al-8V-6Cr-4Mo-4Zr Alloy

Abstract

The mechanical properties of beta titanium alloys can be improved by precipitating the α phase in the β-phase matrix and controlling the microstructure via appropriate aging treatments. In this study, heat treatment in the range of 400 to 550 °C is performed to optimize the aging of Ti-3Al-8V-6Cr-4Mo-4Zr alloys. The increase in the aging temperature and holding time increases the hardness and compressive yield strength owing to the precipitation of the secondary α phase in the β matrix. The precipitation driving force at 400 °C is low because of the slow diffusion rate, and therefore the improvements in the hardness and strength are small. At temperatures above 500 °C, phase separation occurs rapidly (β → β + β′), and the β′ phase acts as a nucleation site for the secondary α phase. The phase transformation from the β′ to the secondary α phase is promoted at 500 °C, resulting in the highest hardness (406.3 HV) and compressive yield strength (1433.8 MPa) at 24 h. At 550 °C, the secondary α phase grows and the hardness and compressive yield strength degrade. These results can be effectively applied to manufacture springs with excellent formability and mechanical properties.

1. Introduction

The increasing environmental regulations on fuel efficiency and greenhouse gasses have led to an increased demand for ecofriendly vehicles. It is essential to reduce the weights of such vehicles to improve fuel efficiency [1,2]. Therefore, titanium alloys with higher specific strength (strength/density) and specific fatigue strength (fatigue strength/density) than SPS 6 materials, which are currently being used in automotive springs, are being actively researched for application in automotive parts such as valves and springs [3,4,5,6]. C. L. Li et al. [7] investigated the 400 °C to 600 °C range, for application in automobile springs. K. Wang et al. [8] used the metastable β titanium alloy Ti-5Mo-3Cr-Fe-3Zr and observed a higher maximum tensile strength (1235 MPa) with air cooling by varying the cooling rate after solution treatment, compared to quenching.

In the automotive industry, replacing conventional steel springs with titanium alloy springs offers several advantages in terms of weight reduction, corrosion resistance, long-term cost-effectiveness, and safety considerations. Firstly, titanium has a density that is only 56% of that of steel, and its shear modulus of elasticity is approximately 50% of that of steel. Therefore, theoretically, a weight reduction of up to 72% is achievable for springs with equivalent performance. Additionally, titanium alloy exhibits excellent corrosion resistance, eliminating the need for complex anti-corrosion treatments during manufacturing [9]. In other words, titanium springs remain stable in harsh environments over the long term, reducing maintenance costs. Finally, titanium springs possess excellent vibration damping characteristics, requiring less energy for acceleration and deceleration, and they are easier to control, significantly enhancing passenger safety.

The stable phase of titanium alloys at room temperature varies depending on the type and amount of alloying elements added. For example, β alloys with a BCC structure at room temperature exhibit better fatigue resistance and lower elastic modulus than α alloys. They have excellent formability and are easy to process for parts with complex shapes [10,11]. In addition, they generally have a lower elastic modulus (55–91 GPa) than α and α + β alloys [12]. Owing to these advantages, their application in automotive springs is expanding. Automotive suspension springs made of β-titanium alloys are approximately 53% lighter than those made of steel materials with the same fatigue life, and are being proposed as a way to improve the fuel efficiency in the automotive industry [13]. However, research on the interrelationship between the mechanical properties and microstructure of β-titanium alloys is scarce, compared to that on α and α + β alloys [14,15].

The hardness and tensile strength of β-titanium alloys can be improved by precipitating the α phase in the β phase matrix through aging treatment and by controlling the fraction and size [16,17,18]. When the alloying elements added are lean, the ω phase precipitates, and when they are rich, phase separation occurs as β + β′. This shows the precipitation strengthening effect [19]. However, the excessive precipitation of the α phase in the β phase matrix can increase the elastic modulus; therefore, proper microstructure control is required to balance the strength and elastic modulus [20].

It is possible to increase the strength of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy, a metastable β-titanium alloy, by controlling the grain size and α precipitation phase through heat treatment. In addition, it is known to have excellent processability due to its low Cr content and low possibility of segregation [21]. Therefore, this alloy is being studied for application in aircraft structural materials and automobile engine springs [14,15]. However, most of the previous studies on Ti-3Al-8V-6Cr-4Mo-4Zr alloys have focused on heat treatment to improve the mechanical properties. Research on the interrelationship between the microstructure and mechanical properties after heat treatment is lacking. B. H. Lee et al. [22] showed that the increase in density of the α/β interfacial area after the aging of the metastable β titanium alloy Ti-3Al-8V-6Cr-4Mo-4Zr effectively enhances hardness and tensile strength. J. W. Lee et al. [15] showed that the tensile strength of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy slightly increases with the aging temperature. In the case of a tension coil spring, it is subjected to tensile or compressive stress simultaneously with torsion. However, research on the compressive properties of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy with respect to aging treatment is insufficient. Therefore, in this study, the microstructure changes associated with aging treatment conditions of a metastable Ti-3Al-8V-6Cr-4Mo-4Zr β-titanium alloy were observed. In addition, the correlation between hardness and room temperature compression properties was investigated.

2. Experimental Section

The alloy used in this study was manufactured into Φ 15 bars through a drawing process at ‘Daewon Kang-Up Co.’ (Cheonan, Republic of Korea). The chemical compositions of the alloys are presented in

Table 1. Chemical compositions of Ti-3Al-8V-6Cr-4Mo-4Zr alloy.

Elements	Al	V	Cr	Mo	Zr	Si	Fe	O	C	N	Ti
wt.%	3.38	8.3	5.82	4.13	3.92	0.12	0.06	0.08	0.007	0.008	Bal.

. The Ti-3Al-8V-6Cr-4Mo-4Zr alloy is a metastable β alloy ([Mo]eq = 10% to 25%) with a large amount of added β stabilizing elements such as V, Cr, and Mo; the Mo equivalent ([Mo]eq) is 16.8% [12]. In addition, the β-transition temperature is determined to be approximately 709 °C [23].

To relieve the stresses in the alloy produced by the drawing process, it was subjected to a stress-relief heat treatment at 950 °C (above the β-transition temperature) for 25 s, followed by air cooling. The strength of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy can be improved through alpha phase precipitation when aged between 400 and 600 °C, but it is known that there is little difference in hardness when aged for more than 24 h [24,25,26]. The alloy was coated with delta glaze to prevent surface corrosion during aging. It was then charged at aging temperatures of 400 °C, 450 °C, 500 °C, and 550 °C and held for 2, 8, 16, and 24 h, respectively. It was then water-cooled to comparatively analyze the microstructure and mechanical properties under different aging conditions. A total of 16 aging conditions were tested. The detailed heat-treatment process is presented in Figure 1.

To observe changes in microstructure according to aging conditions, the specimen was cross-sectioned, hot mounted, and polished to a mirror surface using #220–#2000 sandpaper and 6 μm, 3 μm, 1 μm, and 0.04 μm abrasives. After etching with 96 mL H₂O, 2 mL HNO₃, and 2 mL HF solutions for a few seconds, the microstructure changes were observed using an optical microscope (OM, BX53M, Olympus, Tokyo, Japan). An X-ray diffraction (XRD) analyzer (XRD-7000, Bruker D8, Karlsruhe, Germany) was used to analyze the phase changes due to the aging, with 2θ = 30–90° and step size = 0.067°. The phase evolution with temperature and time was analyzed using an energy-dispersive X-ray spectrometer (EDS, JSM-7610F Plus, JEOL, Tokyo, Japan).

Hardness tests and room-temperature compression tests were performed to evaluate the mechanical properties under each aging condition. A 1 kgf load was maintained for 15 s using a Vickers hardness tester (HM-200, Mitutoyo, Kawasaki, Japan), and the average value was calculated after measurements at 14 points, excluding the maximum and minimum values. The aged bar specimens were machined into Φ 6 × 9 mm specimens according to ASTM-E9 [27] for room-temperature compression testing. The dynamic universal materials testing machine (BESTUM-10MD, Ssaul Bestech, Seoul, Republic of Korea) was used to perform three compression tests at a strain rate of 1 × 10⁻³/s until the strain reached 70%. The elongation and average compressive yield strength values were calculated.

3. Results and Discussion

3.1. Microstructure Changes for Different Aging Conditions

The XRD patterns, optical micrographs, and EDS spectra of specimens subjected to stress-relief heat treatment at 950 °C for 25 s and air cooled are presented in Figure 2, accompanied by the corresponding data in

Table 2. EDS analysis results of Ti_xZr_ySi_z precipitation.

Element	wt.%	at.%
Al	3.09	5.53
Si	1.70	2.91
Ti	74.09	74.64
V	8.60	8.15
Cr	5.61	5.20
Zr	3.16	1.67
Mo	3.75	1.89

and

Table 3. EDS analysis results of β matrix.

Element	wt.%	at.%
Al	3.06	5.53
Si	0.06	0.10
Ti	75.53	76.92
V	8.84	8.47
Cr	5.91	5.55
Zr	3.09	1.65
Mo	3.51	1.78

. The XRD results contained α-phase and β-phase peaks, and the microstructure observations confirmed the formation of fine precipitates in the equiaxed crystal structure and β matrix. EDS analysis was performed to clearly identify the phases of the precipitate through composition analysis, and the EDS spectrum of the white precipitate showed a relatively high content of Si and Zr (Figure 2c). The EDS spectrum of the β base indicated a lower content of Si, compared to that in the white precipitate (Figure 2d), thus confirming that the white precipitate present in the β matrix was the Ti_xZr_ySi_z phase. These precipitates appear when Zr is added to the titanium alloy; Zr reduces the solubility of Si, which promotes the formation of precipitates [21,28]. It has also been reported that the Ti_xZr_ySi_z phase does not precipitate when Zr is not added [29,30]. The Ti_xZr_ySi_z phase precipitates in the matrix because of the rapid cooling rate after heat treatment, and the size and fraction of the phase increases with increasing heat-treatment temperature [29,30]. By controlling the aging conditions to properly distribute the Ti_xZr_ySi_z phase within the β matrix, the grain growth can be hindered, resulting in grain refinement [21,31].

To analyze the influence of the aging-treatment conditions on the secondary α phase, optical micrographs of the microstructures of specimens subjected to 2 h to 24 h of aging treatment at 400 °C to 550 °C are presented in Figure 3. The microstructure change with the increase in aging time at a temperature of 400 °C was insignificant. This was because of the low driving force for the secondary α-phase precipitation due to the relatively slow diffusion rate at low aging temperatures [32]. During 2 h of aging at 450 °C, 500 °C, and 550 °C temperatures, αGB was obtained because of the preferential precipitation of the secondary α phase on the β grain boundary. αGB was formed by the migration of the α stabilizing elements (O and Al) present inside the β phase toward the β grain boundary by diffusion during aging below the β-transition temperature [33]. Subsequently, as the aging time increased, the secondary α phase precipitated inside the β matrix owing to the relatively high temperature, increasing the amount of precipitation.

The precipitation mechanism of the secondary α phase according to aging treatment temperature is shown in Figure 4. As mentioned earlier, at 400 °C, the diffusion rate of alloying elements slows and the growth of precipitates decreases, so a very long aging treatment time is required to complete the phase transformation to the secondary α phase [34]. Therefore, the microstructure change with increasing aging time is minimal. In addition, the diffusion rate of atoms increased with the increase in aging temperature. The time for precipitation and the growth of the secondary α-phase occurred rapidly as the aging temperature increased from 450 °C to 550 °C, confirming the coarsening of the secondary α phase. This is because higher aging temperatures result in a higher driving force for growth but a reduced driving force for secondary α nucleation [35].

To analyze the phase changes with aging conditions, the XRD patterns of specimens subjected to 2 h to 24 h of aging at 400 °C to 550 °C are presented in Figure 5. For all aging conditions, α and β-phase peaks were observed, and no additional precipitated phase peaks were observed. To analyze the changes in the phase-volume fractions of the α and β phases with aging conditions using XRD patterns, Equation (1) was used to calculate the phase volume fraction, and the obtained results are presented in

Table 4. Volume fractions of the aged Ti-3Al-8V-6Cr-4Mo-4Zr alloy.

Temp. (°C)	Time (h)	Vf,α(%)	Vf,β (%)
400	2	1.45	98.55
	8	2.30	97.70
	16	0.90	99.10
	24	1.17	98.83
450	2	1.33	98.67
	8	1.73	98.27
	16	1.76	98.24
	24	7.53	92.47
500	2	0.99	99.01
	8	5.41	94.59
	16	6.79	93.21
	24	23.62	76.38
550	2	1.25	98.75
	8	8.14	91.86
	16	20.08	79.92
	24	27.13	72.87

.

$$V_{f,\alpha }=\frac{A_{\alpha }}{A_{\alpha }+A_{\beta }}{V}_{f,\beta }=\frac{A_{\beta }}{A_{\alpha }+A_{\beta }}$$

(1)

At aging temperatures above 450 °C, the volume fraction of the secondary α phase increased with the increasing aging time and showed a trend similar to that of the microstructure. At 400 °C, the volume-fraction change in the secondary α phase was insignificant because of the slow diffusion rate and lack of driving force for precipitation at a relatively low temperature. At the aging temperature of 500 °C, the volume fraction of the secondary α phase increased from 6.79% to 23.62%, with the increase in aging time, and at the aging temperature of 550 °C, the volume fraction of the secondary α phase increased from 8.14% to 27.13%, with the highest secondary α-phase volume fraction (27.13%) obtained at 550 °C/24 h. As shown, with the increase in aging time at aging temperatures above 500 °C, additional α-phase peaks were observed around 2θ = 39 ° and 41°, and the volume fraction of the secondary α-phase increased sharply. This can be seen in Figure 6, which shows the time–temperature transformation (TTT) diagram of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy, where the β′ phase, which acts as a nucleation site for the secondary α phase at aging temperatures above 500 °C, undergoes a rapid phase separation from β → β + β′ because of the short aging time [30,36]. Therefore, the phase transformation from the β′ to the secondary α phase is accelerated, which explains the high secondary α-phase fraction [19].

3.2. Evaluation of Mechanical Properties Changes

To observe the changes in the mechanical properties due to the secondary α phase precipitated during aging, Vickers hardness measurement and room temperature compression tests of the aged specimens were performed and the results are presented in Figure 7. To minimize the influence of the barreling phenomenon during the room-temperature compression test, the average values of the compressive strength and compressive yield strength at 30% compressive strain were calculated and are shown in

Table 5. Compressive yield strength and compressive strength of Ti-3Al-8V-6Cr-4Mo-4Zr alloy at strain 30%.

Temp. (°C)	Time (h)	Compressive Yield Strength (MPa)	Compressive Strength (MPa) (at Strain 30%)
-	-	908.1 ± 4.1	1378.3 ± 10.2
400	2	928.6 ± 5.4	1442.7 ± 6.9
	8	935.0 ± 4.1	1425.5 ± 21.2
	16	947.7 ± 6.2	1492.4 ± 23.7
	24	942.5 ± 13.6	1425.0 ± 10.1
450	2	935.6 ± 7.4	1386.2 ± 4.5
	8	937.8 ± 12.9	1403.9 ± 47.0
	16	1000.3 ± 4.9	1476.8 ± 12.8
	24	1202.1 ± 5.7	1683.0 ± 20.6
500	2	930.1 ± 9.6	1389.0 ± 67.4
	8	1064.0 ± 29.5	1525.3 ± 73.1
	16	1349.8 ± 18.9	1822.0 ± 43.3
	24	1433.8 ± 21.5	1859.9 ± 11.7
550	2	915.0 ± 12.3	1396.8 ± 54.4
	8	981.5 ± 5.8	1488.4 ± 59.6
	16	1110.5 ± 2.6	1651.6 ± 29.3
	24	1143.5 ± 6.1	1648.6 ± 26.2

.

For all aging temperature conditions, except 400 °C, the hardness and compressive yield strength significantly improved. This was due to the increased precipitation of the fine secondary α-phase with longer aging time, which increased the α/β interfacial area. The α/β interface effectively impedes the movement of dislocations during plastic deformation [37]. Therefore, increasing the α/β interfacial area significantly contributes to improving the strength of the alloy. However, the 400 °C condition did not have enough driving force for the secondary α phase to precipitate because of the slow diffusion rate caused by the relatively low aging temperature. Therefore, the amount of precipitation of the secondary α phase was the lowest and the α/β interfacial area was the smallest (α volume fraction: 0.90% to 2.30%; β volume fraction: 97.70% to 99.10%), and the strengthening effect with increasing aging time was insignificant [32].

For the 450 °C condition, the change in hardness and compressive yield strength was minimal because of the small amount of secondary α-phase precipitation until 8 h. Then, as the aging time increased, the amount of fine secondary α-phase precipitation within the β grain boundary and β matrix increased, resulting in a sharp increase in hardness and compressive yield strength (311 HV to 372.5 HV; 937.8 MPa to 1202.1 MPa), with each value increasing by 19.8% and 28.2%, respectively. For the 500 °C condition, the hardness and compressive yield strength values increased rapidly (306.1 HV → 406.3 HV; 930.1 MPa → 1433.8 MPa) after 2 h due to the increase in the amount of fine secondary α-phase precipitation caused by the fast diffusion rate at the relatively high temperature, with the highest hardness and compressive yield strength obtained at 24 h, showing increases of 32.7% and 54.1%, respectively. For the 550 °C case, the hardness and compressive yield strength increased rapidly (304.2 HV → 349.6 HV; 915.0 MPa → 1110.5 MPa) for aging times from 2 h to 16 h and gradually increased to 24 h, with increases of 14.9% and 21.4%, respectively. However, this condition showed lower hardness and compressive yield strength compared to the 500 °C condition. This was due to the rapid phase separation of β → β + β′ and β′ → secondary α phase during the short aging time at 500 °C and above, as mentioned earlier. Moreover, the higher temperature caused the secondary α phase to grow, resulting in a softening effect [19]. As a result, as the aging temperature and time increased, the amount of precipitation of the secondary α phase increased, which improved the mechanical properties; however, in the case of 550 °C, the secondary α phase coarsened because of the relatively high temperature, which caused a degradation in the mechanical properties.

To analyze the changes in compressive elongation with aging, the compressive stress–strain curves of the aged specimens were obtained and are presented in Figure 8. In the 500 °C and 550 °C cases, the elongation decreased somewhat (500 °C: 36.0% → 31.9%; 550 °C: 44.0% → 39.5%) with the increase in aging time from 16 h to 24 h. This was due to the precipitation of the secondary α phase into the β matrix, which reduced the ductility with increasing strength. Under the same aging time of 24 h, the elongation decreased (450 °C: 39.1% → 500 °C: 31.9%) because of the increase in the amount of secondary α precipitation as the aging temperature increased from 450 °C to 500 °C. On the other hand, as mentioned earlier, the elongation at 550 °C actually increased to 41.6% owing to the softening effect caused by the higher temperature.

4. Conclusions

In this study, a Ti-3Al-8V-6Cr-4Mo-4Zr alloy was aged in the range of 400 °C to 550 °C, and the microstructure changes and mechanical properties were analyzed. The following conclusions were obtained:

(1)

The metastable β-titanium Ti-3Al-8V-6Cr-4Mo-4Zr alloy was subjected to stress-relief heat treatment at 950 °C for 25 s, and a fine Ti_xZr_ySi_z phase was precipitated in the β-base owing to the rapid cooling rate. Subsequent aging treatments at 400 °C to 550 °C for 2 h, 8 h, 16 h, and 24 h resulted in the generation of a secondary α phase, preferentially on the β grain boundary, in all specimens except that treated at 400 °C, where the driving force for precipitation was low. The secondary α phase precipitated inside the β matrix as the aging time increased.

(2)

At aging temperatures below 500 °C, the aging time significantly contributed to the amount of secondary α phase precipitated, which increased with increasing aging time. This improved the hardness and compressive yield strength. At 550 °C, the β′ phase, which acted as a nucleation site for the secondary α phase, rapidly underwent phase separation from β → β + β′, which accelerated the phase transformation to the secondary α phase, resulting in the coarsening of the secondary α phase. The softening phenomenon reduced the hardness and compressive yield strength and increased the compressive elongation.

(3)

The best mechanical properties (406.3 HV and 1433.2 MPa) were obtained at 500 °C/24 h because of the increase in the α/β interfacial area as a large amount of fine secondary α phase was generated. This condition was confirmed to be the optimal process condition.

(4)

The correlation between the distribution of precipitation phases and mechanical properties during the aging treatment of the Ti-3Al-8V-6Cr-4Mo-4Zr alloy was confirmed, and optimal aging process conditions were derived to secure the mechanical properties required for springs and aircraft structural materials.