1 Introduction

Titanium and its alloys have been widely applied to aerospace, orthopedic implant and the chemical industry due to their excellent properties such as low density, high specific strength, excellent biocompatibility and good corrosion resistance.[1,2,3,4] Among them, near-β titanium alloys has received widespread attention in recent years due to their high specific strength and ultimate tensile strength (UTS) after post heat-treatment.[4,5] Presently, the most widely studied near-β titanium alloy is Ti–5Al–5Mo–5V–3Cr (Ti-5553), which has also been successfully used in the Boeing Dreamliner 787 to fabricate truck beams in landing gear modules.[6,7] Most of the literature on near-β titanium alloys is mainly focused on additive manufacturing and thermal mechanical processing, resulting in all Ti-5553 alloy components being prepared by additive manufacturing and machining or thermal mechanical processing and machining methods.[8,9,10,11]

Recently, with the rapid development of military equipment, lightweight weapons and equipment is becoming an important issue, as they affect the improvement of an individual’s combat capabilities. Unlike the demand for the high-temperature mechanical properties of titanium alloys in the aerospace field, military equipment primarily requires the room-temperature mechanical properties of titanium alloys. Therefore, high-strength near-β titanium alloys are receiving widespread attention from the army. However, the high cost of both of these preparation processes limits their application in the army.[7,12,13] Thus, reducing the production cost of high-strength near-β titanium alloys is key to broadening their application field. Investment casting is a near-net-shape manufacturing method, improving material usage and reduce production costs.[13,14,15] In addition, the products produced using this method are widely used in the aviation, aerospace, and petrochemical fields, indicating that the products produced using this method have high quality and reliability.[16] Nevertheless, the cast coarse Widmanstätten structure significantly reduces the mechanical properties of investment casting near-β titanium alloys, thus making them unable to meet practical usage requirements. Therefore, regulating the microstructures and mechanical properties of investment casting near-β titanium alloys is a key scientific problem to their application. Several reports have described forged near-β titanium alloy microstructures and properties that are sensitive to hot working and post heat-treatment technologies.[17,18] However, no information is available on the regulation of the microstructure and mechanical properties of investment casting near-β titanium alloys. For near β titanium alloy products formed by hot processing, Ding et al. believe that near β titanium alloy should first experience the hot rolling or forging treatment to obtain discontinuous grain boundary α phase.[19] However, near-β titanium alloy products produced by investment casting cannot undergo grain hot deformation processing. Therefore, understanding the effect of post heat-treatment methods and process parameters on the microstructural evolution and mechanical properties of investment casting near-β titanium is essential for its application. Herein, we chose the novel near-β Ti–5Al–4Cr–4Mo–4V–3Zr (Ti-1300) titanium alloy developed by the Northwest Institute for Nonferrous Metal Research in China as the research object.[20,21,22] To understand the influence of post heat-treatment methods and process parameters on the microstructural evolution and machinal properties of investment casting Ti-1300 alloy, we designed two heat-treatment methods, β solution and single aging and β solution and double aging, and seven heat-treatment-process parameters. The results of this study contribute to understanding the effect of the morphology and volume fraction of the primary α (αP) and secondary α (αS) phases on the mechanical properties and to achieving strength–plasticity matching of investment casting Ti-1300 alloy.

2 Experimental Procedure

Samples of the investment casting Ti-1300 alloy were prepared by a vacuum shell furnace. The pouring temperature and cooling time are 1800 °C and 2 hours, respectively. The dimensions of the samples were Φ30 × 150 mm. The chemical composition of this as-cast Ti-1300 alloy was then determined using inductively coupled plasma analysis, as shown in Table I. To eliminate defects in as-cast samples, all samples were treated with hot isostatic pressing at 920 °C and 130 MPa for 2.5 h min. The β-transit temperature was metallographically measured to be approximately 830 °C. To investigate the influence of the morphology and volume fraction of the αP and αS phases on the machinal properties of investment casting Ti-1300 alloy, two post heat-treatment methods, β solution and single aging (named as SA) and β solution and double aging (named as SDA), and seven post-heat treatment process parameters were designed. The details of the two-post heat-treatment methods and the seven post-heat treatment process parameters are shown in Figure 1. The post heat-treatment process parameters were as follows: SA-750 (880 °C/1.5 hours + 750 °C/2 hours), SA-700 (880 °C/1.5 hours + 700 °C/2 hours), SA-650 (880 °C/1.5 hours + 650 °C/2 hours), SA-600 (880 °C/1.5 hours + 600 °C/2 hours), SDA-650 (880 °C/1.5 hours + 750 °C/2 hours + 650 °C/3 hours), SDA-620 (880 °C/1.5 hours + 750 °C/2 hours + 620 °C/3 hours), and SDA-570 (880 °C/1.5 hours + 750 °C/2 hours + 570 °C/3 hours).

Table I Chemical Composition of the Investment Casting Ti-1300 Alloy
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Fig. 1
figure 1

Investment casting Ti-1300 alloy post heat-treatment process parameters: (a) SA and (b) SDA. FC, furnace cooling; AC, air cooling

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Specimens for optical microscopy (OM; Axio Vert.A1; ZEISS) and scanning electron microscopy (SEM; JSM 7200F; JEOL) observations were metallographically ground, polished, and etched using Kroll’s reagent (5 vol pct HF + 15 vol pct HNO3 + 80 vol pct H2O) for 30 s. The microstructural evolution and fracture morphologies of the different post heat-treatment methods and process parameters in the samples were then analyzed. Quantitative analysis of the microstructure characteristics was performed using Imaging Pro software. Tensile specimens were processed according to GB/T 228.1–2021 of China. Tensile tests were conducted using an Instron-type testing machine at a cross speed of 8.3 × 10−6 m/s before the yield point and a cross speed of 8.3 × 10−5 m/s after the yield point.

3 Results and Discussion

3.1 As-Cast Ti-1300 Alloy Microstructure and Mechanical Properties

Figure 2 shows the microstructures, stress–strain curves, and fracture morphologies of the as-cast Ti-1300 alloy. Figure 2(a) shows that the as-cast microstructure of the Ti-1300 alloy comprises coarse β grains and a continuous α grain boundary (αGB) distributed on the β-grain boundaries, marked with yellow arrows in Figure 2(a). Figures 2(a) and (b) shows that the coarse β grains comprise variously sized and oriented α colonies (αC) composed of parallel distributions of the lamellar αP phase and β residual (βR) matrix, marked with yellow arrows in Figures 2(a) and (b). In addition, the αS phase is observed in the bulk βR matrix. To further analyze, the widths of the αP phase were quantitatively measured using Imaging Pro software, as shown in Figure 2(c). The width distribution of the αP phase with an 81.5 pct volume fraction ranged from 0.3 to 0.9 μm. This indicates that the microstructure of the as-cast Ti-1300 alloy is a typical coarse Widmanstätten structure, primarily owing to the low cooling rate of the investment casting process providing sufficient growth time for the αP phase. Figure 2(d) shows the stress–strain curve of the as-cast Ti-1300 alloy. The samples broke before reaching the specified yield strength of 2 pct. The ultimate tensile strength (UTS) is 1143.0 MPa, the elongation (EI) is 1.2 pct, and the calculated average elastic modulus (E) is 105.9 GPa. According to Wen et al.[23] report, the plasticity and strength of the TC21 alloy with a lamellar microstructure decreased as the thickness of the αP phase sheet decreased because thin α plates decrease the size of the plastic deformation zone of the crack tip. As shown in Figures 2(a) through (c), the low EI of the as-cast Ti-1300 alloy was owing to the precipitation of the thin-layer αP phase. To further investigate the fracture behavior, the macroscopic and microscopic fracture morphologies of the tensile sample were analyzed using SEM, as shown in Figures 2(d) through (f). The obvious tear ridges are observed in in Figure 2(d), indicating that intergranular crack propagation along the continuous α boundary distributed on β grain boundaries are the dominant failure mechanism in the as-cast Ti-1300 alloy. To further understand the fracture mechanism, the uneven and flat areas marked with red circles in Figure 2(d) were analyzed, as shown in Figures 2(e) and (f). Secondary cracks, microvoids and dimples are observed. However, the dimples are small and shallow, meaning that the tear area has low toughness. As shown in Figure 2(f), the intergranular fracture area displays shell-shaped dimples, indicating that the brittleness of the as-cast Ti-1300 alloy results in its inability to be applied in engineering practice. Therefore, the regulation of the microstructure and mechanical properties using post-heat treatment is a key issue for determining whether investment Ti-1300 alloy can be applied in engineering.

Fig. 2
figure 2

Microstructures, stress–strain curve and fracture morphologies of the investment casting Ti-1300 alloy. (a) OM and (b) SEM microstructures. (c) Width of the primary α plates. (d) Stress–strain curve. Inserted image in (d), (e) and (f) fracture morphologies. αGB, continuous α grain boundary distributed on the β-grain boundaries. αC, α colonies composed of parallel distributions of the lamellar primary α (αP) phase and β residual (βR) matrix. αS, secondary α

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3.2 Effect of SA Post heat-Treatment on the αP Phase and Mechanical Properties

According to published reports,[24,25,26] the β grains, αC, and αP and αS phases in titanium alloys with multilevel microstructures significantly influence their mechanical properties. For example, Terlinde et al.[26] believes that a decrease in the volume fraction of the αP phase leads to a significant decrease in plasticity in Ti–10V–2Fe–3Al titanium alloy owing to a soft αP phase and hard β phase. Because β grain size cannot be effectively regulated using post heat-treatment methods, we investigated the influence of post heat-treatment process parameters on the αC and αP phase and analyzed the effect of the αC and αP phase on the mechanical properties of investment casting Ti-1300 alloy. The detailed process parameters for samples SA-750, SA-700, SA-650, and SA-600 are shown in Figure 1(a).

Figure 3 shows the SEM images and widths of the αP plates of the investment casting Ti-1300 alloy after SA post heat-treatment method, and the detail parameters is shown in Figure 1(a). Figures 3(a), (d), (g), and (j) shows that the grain-boundary α phase transition thins and becomes discontinuous with decreasing annealing temperature, and no obvious αC was observed except in some grains of SA-700 sample. Only the αP phase precipitates in β grains after SA post heat-treatment, as shown in Figures 3(b), (e), (h), and (k). To further analyze the volume fractions and widths of the αP plates shown in Figures 3(b), (e), (h), and (k), they were measured using Imaging Pro software, as shown in Figures 3(c), (f), (i), and (l). The volume fractions of the αP phase in the samples were 36.9, 61.7, 63.9, and 66.3 pct, indicating that decreasing the aging temperature is beneficial for the precipitation of the αP phase. The β phase is a high-temperature metastable phase and the α phase is a room-temperature stable phase in titanium alloys. In addition, after the completion of the β-solution treatment at 880 °C for 1.5 hours, the furnace was cooled to the aging temperature at a cooling rate of approximately 1 °C/min, which means that a low aging temperature results in a long holding time. Thus, the volume fractions of the αP phases increased with decreasing aging temperature.

Fig. 3
figure 3

SEM images of the investment casting Ti-1300 alloy after SA post heat-treatment. (a) to (c) SA-750, (d) to (f) SA-700, (g) to (i) SA-650, and (j) to (l) SA-600. αGB, continuous α grain boundary distributed on the β-grain boundaries. αC, α colonies composed of parallel distributions of the lamellar primary α (αP) phase and β residual (βR) matrix. αS, secondary α

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The main widths of αP plates in these samples were between 1.0 and 2.5 μm, between 0.5 and 2.5 μm, between 0.9 and 1.8 μm, and between 0.8 and 2.0 μm. Compared with the as-cast state, not only the width of the αP plates after SA post heat-treatment increases, but also the distribution range becomes larger. Although the widths of the αP plates in these samples after SA post heat-treatment did not significantly change, the distribution interval frequency of the αP phases significantly change. According to Wang et al.,[27] as the temperature difference between the aging temperature and solution temperature increases, the driving force of the β → ɑ phase transformation increases during treatment, which is conducive to the formation of nuclei in random directions. However, a high nucleation rate limits the αP-phase growth of partial preferential nucleation. Therefore, the distribution interval frequency of the αP phase decreased with decreasing primary aging temperature, except for SA-700. In general, after the SA post heat-treatment, the morphology of the αP phase in the investment casting Ti-1300 alloy can be effectively regulated.

Figure 4 shows the variation in the tensile properties of the investment casting Ti-1300 alloy with annealing temperature. As shown in Figure 4(a), the engineering stress of the SA-750 sample decreases after the yield point as the engineering strain increases. Figure 3(b) shows that the microstructure of SA-750 sample is composed of 63.1 pct bcc-structure β phase and 36.9 pct hcp-structure α phase. According to Schwab,[28] this stress drop is common for materials with a bcc crystal structure at low deformation temperature. This effect can be explained by two theories. First, the static theory, describes discontinuous yielding as the dissociation of dislocations from a solute atomic environment.[29] Second, the dynamic theory, assumes that grain boundaries act as a source of new mobile dislocations.[30] Both theories are based on the low dislocation density of materials.[31] Herein, all Ti-1300 alloy samples were obtained by investment casting without any thermal mechanical processing. Thus, we speculated that all samples have a low dislocation density, which meets the basic requirements of these two theories. The engineering stresses of the other three samples remained unchanged after the yield point. According to published reports,[22,32] α/β interfaces can hinder the dislocation slip between these two phases and further act as stress concentrators. Based on the statistical results shown in Figures 3(f), (i), and (l), the hcp-structure primary α phase content in these three samples is greater than 60 pct, meaning that these three samples have a high density of α/β interfaces. A high-volume fraction of the αP phase is beneficial for preventing stress decreasing with strain in near-β titanium alloys.

Fig. 4
figure 4

Tensile properties of the Ti-1300 alloy at different annealing temperatures. (a) Stress–strain curves, (b) Ultimate Tensile strength (UTS) and yield strength (YS), and (c) Elongation (EI) and reduction area (RA)

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Figure 4(b) shows that the UTS and yield strength (YS) slightly increase with decreasing aging temperature from 750 °C to 600 °C. The UTS and YS of the alloy increased by approximately 6.7 and 4.7 pct, respectively. Based on the results shown in Figure 3, the increase in UTS and YS can be ascribed to the increasing volume fraction of the αP phase. This means that the αP phase is harder than the β phase. However, Terlinde et al.[26] show that the αP phase is softer than the β phase in Ti–10 V–2Fe–3Al titanium alloy. This indicates that the degree of softness or hardness between the αP and β phases vary in different alloys. Figure 4(c) shows that the EI and reduction area (RA) remained unchanged, except for SA-700. The decreasing EI and RA of SA-700 was owing to the formation of coarse Widmanstätten structures in some β grains, as shown in Figure 3(d). Briefly, compared with the as-cast Ti-1300 alloy mechanical properties, after SA post heat-treatment, the strength of the investment casting Ti-1300 alloy decreased, whereas the plasticity significantly increased. Wang et al.[27] show that a small αP phase width increases dislocation motion resistance. In addition, by comparing Figures 2(b) and 3(b), (e), (h), and (h), the orientation relationship of the αP phase is no longer apparent, which is conducive to hindering crack diffusion.

To further investigate the influence of microstructure after SA post heat-treatment on fracture behavior, the tensile fracture morphologies of the four samples were investigated using SEM, as shown in Figure 5. Figures 5(a), (c), (e), and (g) show that the macroscopic fracture morphologies of these samples are uneven and have obvious tearing marks. Moreover, very few intergranular fractures were observed in SA-600, indicating that the fracture mechanism of these four samples is transgranular fracture. For further study, the typical regions shown in Figures 5(a), (c), (e), and (g) (marked with red circles) were analyzed, as shown in Figures 5(b), (d), (f), and (h), respectively. The fracture morphology of these samples exhibits fractured features, including dimples, cleavage planes, and secondary cracks. However, it is mainly dominated by dimples, which are large and deep. Large and deep dimples can extend crack propagation paths and propagation time.[26,27,28] Therefore, these samples exhibit high EI and RA. By contrast, the mechanical properties of the investment casting Ti-1300 alloy can be effectively improved by adjusting the morphology and volume fraction of the αP phase.

Fig. 5
figure 5

Fracture morphologies of the investment casting Ti-1300 alloy after SA post heat-treatment. (a) and (b) SA-750, (c) and (d) SA-700, (e) and (f) SA-650, and (g) and (h) SA-600

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3.3 Effect of SDA Post heat-Treatment on the αS Phase and Mechanical Properties

Based on the above analysis, it was concluded that SA-750 sample has a low volume fraction, optimal morphology of the αP phase, and a good matching relationship between strength and plasticity. It also left enough room to study the relationship between the αS phase and the machinal properties. We designed the SDA post heat-treatment method to investigate the influence of morphology and lamellar width of the αS phase on mechanical properties, and the detailed post heat-treatment process parameters for samples SDA-650, SDA-620, and SDA-570 are shown in Figure 1(b). Figure 6 shows the SEM images of the investment casting Ti-1300 alloy after applying the SDA post heat-treatment. Figures 6(a) through (c) shows that the microstructures after SDA post heat-treatment are composed of the αP phase and βR matrix. As shown in Figure 3, the αP phase is precipitated during the first aging at 750 °C and the morphology and volume fraction of the αP phase is unchanged. This means that secondary aging will not affect the morphology or volume fraction of the αP phase. To analyze the effect of the αS phase on the properties, a high magnification SEM was used, and the results are shown in Figures 6(e), (f). A high density of the acicular αS phases is precipitated in βR matrix by the secondary aging treatment. The widths of the acicular αS phases decrease with decreasing secondary aging temperature. The widths of the acicular αS phases under these three-post heat-treatment processes corresponded to approximately 0.2, 0.13, and 0.06 μm. The nucleation and growth of the acicular αS phase have a competitive relationship. Generally, a high aging temperature can provide a high nucleation activation energy. This means that SDA-650 sample has the maximum nucleation activation energy, which can effectively hinder the growth of the acicular αS phase. In addition, some αS phases are spheroidized when the secondary aging temperature is increased to 650 °C. By contrast, a high secondary aging temperature is beneficial for obtaining αS phases with small aspect ratios.

Fig. 6
figure 6

SEM images of the investment casting Ti-1300 alloy after SDA post heat-treatment. (a) and (d) SDA-650, (b) and (e) SDA-620, and (c) and (f) SDA-570. αP, primary α. βR, β residual. αS, secondary α

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Figure 7 shows variation in the tensile properties of the investment casting Ti-1300 alloy after SDA post heat-treatment. The UTS and YS increase with decreasing secondary aging temperatures; however, the EI and RA decrease with decreasing secondary aging temperatures. Deng et al.[31] report, the soft coarse αP phase mainly provides ductility in the bimodal structural sample, while the fine αS phase determine its strength. Ding et al.[19] think that high aging temperature can promote the growth of αs phase in βR matrix and the decrease in the α/β interface area in the alloy, which result in the weakening of the second phase strengthening effect. Wu et al.[31] believe that α/β interfaces can hinder dislocation slip between these two phases. As shown in Figure 6, we can find that as the aging temperature increases, the αS phase volume fraction and α/β interface area decrease. Therefore, we can think that the increase of UTS and YS with decreasing secondary aging temperatures can be ascribed to the weakening of interface strengthening effect. Although high volume of α/β interfaces can enhance the interface strengthening effect, it will significantly hinder dislocation slip between these two phases and further act as stress concentrators. Additionally, thin α plates decrease the size of the plastic deformation zone of a crack tip. Therefore, The EI and RA show a decreasing trend as the aging temperature decreased. The inserted picture in Figure 7(d) shows engineering stress vs engineering strain of this work in comparison with typical cast titanium alloy. The investment casting Ti-1300 alloy after SA and SDA process post-treatment shows a high strength–plasticity matching relationship compared with the typical cast titanium alloy, and overcoming the strength–plastic trade-off.

Fig. 7
figure 7

Tensile properties of the Ti-1300 alloy at different secondary aging temperatures. (a) Stress–strain curves. (b) Ultimate tensile strength (UTS) and yield strength (YS). (c) Elongation (EI) and reduction area (RA). (d) Stress–strain curves of as-cast, SA-650 and SDA-620, and inserted picture in (d) shows engineering stress vs engineering strain of this work in comparison with typical cast titanium alloy

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To further investigate the influence of the SDA post heat-treatment on fraction behavior, the tensile fracture morphologies of these samples were investigated, as shown in Figure 8. The macroscopic morphologies of the fracture surfaces in the samples are composed of rough regions with some flat regions, marked with yellow arrows in Figures 8(a), (c), and (e). This indicates that the fracture mechanism of the investment casting Ti-1300 alloy after SDA post heat-treatment is a mixed mode of transgranular and intergranular fracturing.[31] The volume fraction of the flat regions, however, increases with decreasing secondary aging temperature, indicating that the main fracture mechanism changes from transgranular fracture to intergranular fracture. For further analysis, the rough regions in Figures 8(a), (c), and (e) (marked with red circles) were analyzed, as shown in Figures 8(b), (d), and (f), respectively. SDA-650 and SDA-620 exhibit fractured features, including dimples, microvoids, and secondary cracks; however, only dimples are observed in SDA-570. The size and depth of the dimples become smaller and shallower with decreasing secondary aging temperature. Therefore, the EI and RA of the investment casting Ti-1300 alloy decrease with decreasing secondary aging temperature. By contrast, the widths of the αS phases decrease with decreasing secondary aging temperature, which results in an easy concentration of stress at the tip of the acicular αS phase. After heat treatment by the SDA-620 process, the UTS, YS, EI, and RA of the investment casting Ti-1300 alloy were 1094 MPa, 1017 MPa, 7.0, and 12 pct, respectively.

Fig. 8
figure 8

Fracture morphologies of investment casting Ti-1300 alloy after SDA post heat-treatment. (a) and (b) SDA-650, (c) and (d) SDA-620, and (e) and (f) SA-570

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4 Conclusions

In this work, the influence of post heat-treatment methods and process parameters on the morphology and volume fraction of the αP and αs phases and the corresponding mechanical properties in the near-β titanium alloy Ti-1300 alloy fabricated by investment casting were studied systematically. The main conclusions are drawn as:

  1. (1)

    The microstructure of as-cast Ti-1300 alloy prepared by investment casting is a typical coarse Widmanstätten structure, exhibiting high strength of 1143.0 MPa but almost zero elongation.

  2. (2)

    After SA post heat-treatment, the EI and RA of investment casting Ti-1300 alloy was effectively improved. The EI increased from 1.2 pct in the as-cast sample to 13.5 pct in the SA-600 post heat-treated sample, while also achieving a UTS of 949 MPa. These superior properties stem from the combined effects of microstructures, including grain boundary α phase, primary α phase, and α colonies.

  3. (3)

    After SDA post heat-treatment of the SA-750 sample, the UTS and YS were increased further and the EI and RA decreased with decreasing secondary aging temperatures. The UTS increased from 889 MPa in SA-750 sample to 1205 MPa in the SDA-570 sample, while also achieving an EI of 4 pct.

The increase of UTS and YS with decreasing secondary aging temperatures can be ascribed to the weakening of interface strengthening effect. The EI and RA decrease with decrease of aging temperature can be ascribed to high volume of α/β interfaces can significantly hinder dislocation slip between these two phases and further act as stress concentrators. Additionally, thin α plates decrease the size of the plastic deformation zone of a crack tip.