1 Introduction

Beta titanium alloys have been available for the aerospace industry since the 1950s (for instance Ti–13V–11Cr–3Mo alloy), but the first significant example in a civil aircraft manufacturing is the usage of Ti–10V–2Fe–3Al high-strength forgings in Boeing 777.[1] This resulted in hundreds of kilograms of weight savings due to the replacement of high-strength low-alloy (HSLA) steel by lighter titanium alloy. Economic manufacturing of large components, typically landing gears, is achieved by hot near-net shape forging, which yields the desired mechanical properties. However, the processing window of Ti–10V–2Fe–3Al alloy is narrow and the eventual properties are sensitive to small fluctuations in both temperature and strain rate.[2] To circumvent this disadvantage, Ti–5Al–5Mo–5V–3Cr (abbreviated as Ti-5553) alloy was developed by Boeing on the basis of the Russian alloy VT22.[3] It offers a larger processing window around the beta-transus temperature (~ 845 °C) and can be also used at higher strengths than the Ti–10V–2Fe–3Al. Therefore, it has become the alloy of choice for airframe forgings.[4] The alloy is currently used in the Boeing 787 Dreamliner manufacturing.

Commercial processing includes either annealing just above beta-transus temperature to promote high toughness condition or processing just below beta-transus (in the α + β field) to achieve high-strength condition. In both cases, aging at 600 °C follows.[5] Even the lower strength condition (β-annealed and aged) achieves an impressive ultimate tensile strength of 1100 MPa. Therefore, the alloy was also considered for castings, resulting in similar properties as in the β-processed alloy.[6,7]

At least four different phases: β (stable high-temperature bcc phase), α (stable low-temperature hcp phase), α” (metastable orthorhombic phase), and ω (metastable hexagonal phase) occur in this alloy, depending on heat treatment scheme.[8,9] Out of these, formation of the stable α phase is most important for mechanical properties, particularly the volume fraction and the size of the α phase precipitates. During annealing at 570 °C, α laths nucleate after 90 s at the temperature and are as small as 150 nm in length and 10 nm in width; after 30 min of annealing they grow to the size of 1500 nm × 200 nm.[2] Higher annealing temperatures of 600–700 °C and longer annealing times lead to precipitation of larger α particles.[10,11] The best combination of mechanical properties, namely the strength above 1000 MPa and the ductility of 6–10 pct, is achieved by annealing at 600 °C for 1–4 h.[12,13]

A prospective area of applications of the Ti-5553 alloy are high-strength titanium fasteners that could replace the heavier ones made of steel. Contrary to large components like landing gear or engine mounting pylons, the fasteners can be produced by methods of powder metallurgy, for instance, by a field-assisted sintering technique (FAST).[14,15] This motivated us to explore the abilities of FAST not only for material compaction but also for an immediate direct (in situ) aging in the device to achieve the desired properties in a single processing run.

Field-assisted sintering technique (FAST), also known as spark plasma sintering (SPS) or pulsed electric current sintering (PECS) is a method of a material consolidation by electric current passing directly through the sintered specimen and its surroundings (typically a graphite die). The first patent concerning electric current sintering was already submitted in 1906,[16,17] but the method has been extensively exploited only in the last decades.[18] Its major advantages in comparison to other sintering methods are achieving very high heating rates (up to 1000 °C/min) and shortening the processing times by an order of magnitude.

Homogeneous temperature distribution within the die is required to achieve homogeneous final mechanical properties without additional treatment.[19,20] The use of a graphite felt as a thermal insulation was shown to greatly reduce the power consumption and improve the temperature homogeneity,[21] and was therefore also employed in this study. Thermal heterogeneity limits possible upscaling of the method for producing larger and near-net shape parts. However, a homogeneous rod with the length of 60 mm and the diameter of 10 mm was successfully sintered.[22] A potential for scalability was also shown on cylindrical samples with a diameter of 15–56 mm keeping aspect ratio constant.[23]

Near-net shape methods combined with FAST present a great potential for fast production of complex small parts. Different die arrangements can be designed to produce multiple specimens in a single sintering run.[18] Voisin et al.,[24] using a specially designed die and punches, were able to produce a net shape TiAl turbine blade. The material exhibited better creep resistance, but slightly worse RT ductility in comparison with the material manufactured by a conventional processing route. A so-called FAST-forge method consists of near-net shape sintering by FAST followed by a single-step precision forge.[25] Achieved properties of these Ti–6Al–4V alloy parts were comparable to the conventionally prepared material.

Ti-5553 alloy was also prepared by this two-step method, yielding the same microstructure as commonly observed in conventionally β-forged alloy,[26] but the initial powder was already alloyed. More recently, Liu et al.[27] prepared a Ti-5553-1Zr alloy using FAST and in situ aging from both milled and unmilled prealloyed powder, resulting in mechanical properties comparable to the commercial alloys. However, prealloyed powder was used in both cases, while we propose to employ blended elemental powder metallurgy (BEPM).

Advantage of BEPM methods is that prealloyed powder is not required, and composition of the alloy can be adjusted. This is of particular interest for development of new alloys which can be produced using a simple “mix and sinter” approach.[28,29,30] The sintering temperature and time can be easily adjusted to achieve either a heterogeneous material[31] or, more importantly, a homogeneous material.[28]

We propose the use of a single-run FAST procedure for all manufacturing steps: starting from the elemental powder blend, using the FAST for material compaction, homogenization, and for an immediate direct in situ aging in the device via adding an annealing step to the temperature program after the main sintering step.

2 Experimental Methods

The Ti–5Al–5V–5Mo–3Cr (commonly abbreviated as Ti-5553) alloy was prepared by field-assisted sintering technique (FAST) directly followed by aging using a prealloyed powder (hereinafter referred to as PA only) and a blend of elemental powders (referred to as BE only).

Basic information about the powders used for sintering is shown in the Table I. All powders were handled in air, except for the V powder (being most sensitive to severe oxidation at ambient conditions), which was handled in an Ar filled glovebox. Ti-15Mo master alloy was used for the BE sample, as the homogenization would be very slow if started from a pure Mo powder. A similar size distribution of the powders is desirable, as it facilitates homogeneous distributing of elements during mixing. The mixing of elemental powders was done using an in-house made rotational mixing device (a half-filled vial with the powder eccentrically placed within a rotating drum, 1 h duration) immediately before filling the sintering die. No ball milling was used.

Table I The Powders Used for Sintering
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The mixed BE powder was also characterized using SEM (Figure 1) to show the powder size and morphology and to check the blending quality. As shown in Figure 1(a), the BE powder consists mostly of the relatively large spherical Ti and Ti-15Mo powders, while the Al, V and Cr are irregular-shaped and smaller. The EDS mapping, shown in Figure 1(c), revealed good homogeneity of the blend without any agglomerates consisting of a single element only (which would be a sign of a poor mixing).

Fig. 1
figure 1

(a) The detailed BSE micrograph of the mixed BE powder used for sintering. (b) Low magnification overview of the mixed BE powder (BSE contrast) and (c) a corresponding EDS map overlay of the same area showing homogeneous distribution of elemental powder after mixing

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The sintering was performed in the SPS 10–4 furnace (Thermal Technology LLC, USA) in a vacuum of the order of 1 Pa. The sintering/aging programs are graphically presented in Figure 2. The process starts with applying the pressure of 20 MPa, subsequent heating up to 600 °C and a 15 min dwell to dissolve the Al particles and prevent their melting. The heating continues at a rate of 100 °C/min up to the sintering temperature (1400 or 1500 °C). After reaching the temperature, the pressure of 30 MPa is applied and the main sintering step, consisting of a 30 min dwell at the desired temperature, starts. Then, the electric current is turned off and the pressure is released to 5 MPa. The cooling rate is not actively controlled, the actual rate is comparable to air cooling. The sample is either cooled to the RT (black line in Figure 2), or aging steps are added to the program directly after the sintering (i.e., without cooling the sample to the RT). The first program includes two aging steps—30 min at 800 °C and subsequent 30 min at 600 °C (red line in Figure 2), the second program contains only one aging step—30 min at 600 °C (orange line in Figure 2). Both programs simulate the conventional thermal processing of the Ti-5553 alloy.[13] The temperature overshoot of the as-sintered sample (in black) at 5 min was caused by a malfunction of the temperature regulator, but it is considered to have no effect on the outcome of the sintering, as the main stage occurs at 1500 °C. A tungsten foil is used as a spacer between the powder and the graphite die to prevent carbon contamination at the high sintering temperatures used. The resulting samples have the diameter of 20 mm and the height of approximately 10 mm. The temperature is controlled using a C-type thermocouple placed in the bottom die, approximately 3 mm below the bottom of the sample.

Fig. 2
figure 2

The sintering process parameters during three different programs—sintering only (black line) and two different in situ aging schemes (red + orange lines). The applied piston pressure was identical for all samples and programs (blue line). The presented records relate to the BE samples sintered at 1500 °C. The temperature spike at the black line at 5 min is caused by a fault in the temperature regulator, but it is considered to have no effect on the outcome of the sintering and homogenization process (Color figure online)

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As the next step, a comparison is performed between a selected sintered sample and a conventionally cast and forged (CF) Ti-5553 alloy (produced by TIMET) using tensile tests. To overcome the problems associated with a tensile testing of miniaturized specimens, we have designed a method of preparation of large rod of material (having the diameter of 15 mm and the length of approximately 75 mm), allowing us to prepare tensile specimens of a standard size which can be directly compared to the conventionally prepared bulk alloy.

The experimental challenge lies in the application of high temperatures necessary to homogenize the alloy and the pressure distribution in the longer graphite die, which would normally lead to its failure. For this reason, a special housing was manufactured from a Ti–Zr–Mo (TZM) alloy (supplied by Plansee) having a good creep resistance and a high-temperature strength. The overall experimental assembly is shown in Figure 3.

Fig. 3
figure 3

The experimental setup for high-temperature sintering of large rods: (a) overall assembly within the sintering chamber (the counterweight is necessary to prevent sliding of the heavy TZM housing down), (b) a detail of the die design showing the inner graphite die in the TZM housing. The sintered specimen is shown in (c)

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It is not possible to perform the same in situ aging in the FAST furnace on the CF rod, which serves as a benchmark of mechanical properties. However, maintaining the same thermal history for both sintered and CF rods is crucial for the sake of comparison. Therefore, an alternative procedure was employed for the preparation of the mechanically tested rods, as illustrated in Figure 4: the sintered rod was cooled directly to the RT after the sintering stage and aged in conventional furnaces along with the CF rod. This heat treatment was designed to imitate the in situ aging in the FAST furnace as closely as possible. The heat treatment is based on the ‘sintered + 600 °C’ program in Figure 2, having the optimal microstructure (as shown in the Results section). However, heating to 600 °C from RT would not match the program, as the aging temperature during in situ sintering is achieved directly from higher temperatures (cf. Figure 2). Moreover, CF rod must be solution treated before final aging. Therefore, a two-step process was designed: solution treatment at 900 °C for 1 h, directly aged at 600 °C for 30 min and air cooled. Critical issue is the direct change from 900 °C to 600 °C without undercooling. Two preheated furnaces were used, and the samples, encapsulated in a fused quartz tube under high vacuum, were quickly transferred.

Fig. 4
figure 4

The heat treatment program for the samples used for measurements of mechanical properties. The green curve on the left, showing temperature evolution in the sintered rod, is the same one as that in Fig. 2. The heat treatment is schematically drawn on the right. The points in circles (crossing 900 °C during cooling from sintering temperature and the end of solution treatment) correspond to microstructurally equivalent conditions, allowing to match the in situ FAST aged samples with the conventionally heat-treated rods (Color figure online)

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A simple thermal simulation of the cooling after sample transfer was performed. The result is plotted in Figure 5, along with the experimental data of the temperature evolution of the cooling step during in situ aging in the FAST furnace (‘sintered + 600 °C’ program). It can be easily seen that while there are some differences, such as faster initial cooling of the samples transferred between furnaces, the steady aging temperature of 600 °C is reached after 10 min for both conditions. We may therefore conclude that the designed alternative heat treatment method is a satisfactory representation of the in situ aging in the FAST furnace. This is also further justified by a microstructural correspondence, shown later in the Results section.

Fig. 5
figure 5

Simulated cooling of a cylindrical rod (15 mm in diameter, cylinder bases are neglected). Initial temperature was 900 °C, and ambient temperature was set to 600 °C, only radiative heat transfer is permitted at the domain boundary (the rods are sealed in a fused quartz tube under high vacuum). The simulation data are plotted for the rod center (r = 0 mm) by a green/violet line, and the experimental data of the ‘sintered + 600 °C’ program (cf. Fig. 2) are plotted for comparison (orange line) (Color figure online)

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The oxygen, nitrogen, and hydrogen contents in all samples were measured by a carrier gas hot extraction (CGHE) elemental analyzer. A scanning electron microscopy (SEM) and an energy-dispersive X-ray spectroscopy (EDS) were performed using Zeiss Auriga Compact Crossbeam microscope. The microhardness was measured by a Qness 10A automatic testing machine using a Vickers indenter and 500 g of load (HV0.5 method). Tensile tests were done on an Instron 5882 universal testing machine using a specimen geometry illustrated in Figure 6, having a circular profile of 3 mm in diameter and a gauge length of 14 mm. The test was performed at a rate of 10–3 s−1.

Fig. 6
figure 6

A tensile specimen drawing. All dimensions are given in millimeters

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3 Results and Discussion

3.1 Process Optimization

The chemical analysis was performed on all powder samples and selected sintered samples. The results are summarized in the Table II. Hydrogen concentration was low in all powder samples, implying its low concentration also in the sintered samples, as there are no possible sources of H contamination during powder handling and sintering (done in air or vacuum). It is therefore not necessary to measure the content of H also in the sintered samples.

Table II The Results of a Chemical Analysis by a Carrier-Gas Hot Extraction (CGHE) Method
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The content of nitrogen is low in all sintered samples and comparable to the content expected from the known contamination of powders; no additional contamination occurred during sintering.

The oxygen content of the sintered BE samples is by 10–20 pct lower than the expected amount, calculated from the powder composition. This can be attributed to a fraction of oxygen-containing compounds (e.g., water vapor) adsorbed to the large surface of the powders (and thus included in the gas content measurements). However, these compounds desorbed and evaporated during the sintering in vacuum (and are not included in the gas content of the sintered samples). The oxygen content of the samples made of the prealloyed powder is much lower than that of the blended elemental powder (0.1 vs 0.2 wt pct, respectively), reflecting the difference in the composition of powders. While the oxygen content of the V, Cr, and Al powders is relatively high (cf. Table I), these powders are used in small fractions and the majority of the oxygen comes from the Ti and Ti–15Mo powders (0.16 and 0.20 wt pct, respectively), even though Ti complies with the Grade 1 ASTM technical standard (< 0.18 wt pct of O[5]). Oxygen is generally a strong α stabilizer, which could potentially change the behavior of the alloy during heat treatment. On the other hand, the alloys are typically used in the aged condition and the effect of oxygen on the final microstructure may not be that severe.

We have first characterized the microstructure in the as-sintered condition: the prealloyed (PA) powder sintered at 1000 °C, 1400 °C, and 1500 °C and the blended elemental (BE) powder sintered at 1500 °C. The lower temperatures of 1000 and 1400 °C were not used for the sintering the BE powder, because the material is not sufficiently homogenized, as recently shown for a Ti–Nb–Zr–O alloy.[28] The low magnification micrographs are shown in Figure 7. The grain size is gradually increasing as the sintering temperature increases—from grains below 100 μm at 1000 °C (Figure 7(a)) up to 1 mm at 1500 °C (Figure 7(c)). The BE sample contains grains of 500–800 μm in diameter after sintering at 1500 °C, comparable to the PA sample sintered at 1400 °C. This can be attributed to the increased oxygen content of the BE sample, impeding the grain growth in the β region.[32,33] No porosity was detected in any of the samples, as the temperatures used for sintering are sufficiently high to reach the full density The black dots in the micrographs are remnants of the polishing suspension as revealed by observations using secondary electrons (not shown). A fully dense CP Ti can be readily prepared by sintering at 800 °C[34] and a more demanding TNTZ alloy by sintering at 1100 °C.[35]

Fig. 7
figure 7

BSE micrographs comparing the Ti-5553 alloy sintered from (a) prealloyed (PA) powder at 1000 °C, (b) prealloyed powder at 1400 °C, (c) prealloyed powder at 1500 °C and (d) blended elemental (BE) powder at 1500 °C. The small black dots are remnants of the polishing suspension. The sintering compression direction is vertical in all micrographs

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The chemical homogeneity was measured using EDS line profiles over the length of 3.2 mm (10 μm step size), approximately along the diagonals of the Figure 7(c), (d)). The prealloyed sample sintered at 1500 °C was selected as a benchmark. The results are plotted in Figure 8. The overall homogeneity of the BE sample is comparable to the PA sample for all elements except for V, whose composition was significantly more heterogeneous in BE than in the PA benchmark sample (0.49 vs 0.22 wt pct, respectively) due to comparatively slower diffusion of V than that of other elements and consequently slightly slower dissolution and homogenization. Note that even this residual heterogeneity of 0.49 wt pct is low, considering the blended elemental approach.

Fig. 8
figure 8

A comparison of EDS line profiles and standard deviations of concentrations of four alloying elements. The length of the line is 3.2 mm with 10 μm step size. The smoothed line is for the sake of clarity only and was not used for calculation of σ

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The results of the microhardness measurements of the as-sintered samples are shown in Table III. The microhardness of the PA samples is significantly lower than that of the BE sample (290 HV vs 390 HV, respectively). The microhardness of the PA samples is the same within the measurement standard deviation. Moreover, the scatter of the individual values in the BE sample is significantly larger than in the PA samples. Both effects can be attributed to the higher oxygen content in the BE sample, influencing the microhardness both directly and indirectly (via influencing the phase composition).

Table III Comparison of a Vickers Microhardness of the Ti-5553 Alloy Prepared from Prealloyed and Blended Elemental Powders in the As-sintered Condition
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To explain the differences between PA and BE samples, the microhardness should be correlated with the detailed SEM micrographs shown in Figures 9(a) and (b). While the PA sample consists of mostly β phase with only individual α precipitates at grain boundaries (Figure 9(a)), the BE sample exhibits mostly α + β microstructure and a single-phase β is preserved only in limited amount (Figure 9(b)). This is caused by oxygen being a strong α stabilizer,[36] resulting in a widespread precipitation of α during cooling. A locally higher concentration of β stabilizing elements (caused by a residual heterogeneity, especially that of V (cf. Figure 8)) can overcome the effect of oxygen and shift the stability of the alloy back to the β region. At these spots, a single-phase β microstructure is retained. These variations in the phase composition and morphology, which can be homogenously found across the sample, strongly affect the microhardness, as two-phase α + β microstructure has generally significantly higher strength and microhardness than a single-phase β.[36] This consequently increases the standard deviation of microhardness measurements (7 HV vs 31 HV, respectively).

Fig. 9
figure 9

The detailed BSE micrographs of the prealloyed (PA) and blended elemental (BE) samples after various heating schemes. Note that while the microstructure is significantly different in the as-sintered condition (a, b), the heat treatments smooth the differences and similar microstructure is produced for both manufacturing approaches—two-step annealing at 800 °C and 600 °C (c, d) and a single-step annealing at 600 °C (e, f)

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The PA sample sintered at 1400 °C and BE sample sintered at 1500 °C (Figures 7(c) and (d), respectively) were selected for further direct in situ aging experiments because their grain size is similar, hence allowing a direct comparison.

Two aging schemes were employed, as described in the Experimental section: a two-step sequence (800 °C/30 min + 600 °C/30 min) and a single-step sequence (600 °C/30 min). The micrographs of both alloys after the treatments are presented in Figure 9 and compared with the as-sintered condition.

Despite the differences in the as-sintered condition (Figures 9(a) and (b)), all aged samples consist of a two-phase α + β microstructure. For a given aging scheme, the microstructure of the PA and BE samples is similar. The two-step aging resulted in an extensive precipitation of grain boundary α (Figures 9(c) and (d)), which occurred during the 800 °C aging step—the driving force for α precipitation inside the grains is low and only grain boundaries serve as preferential nucleation sites. The amount of the grain boundary α is lower in the PA sample due to the lower oxygen content. The fine α lamellae inside grains subsequently precipitated during the 600 °C step, when the precipitation driving force is sufficiently high. The lamellae in the two-step annealed BE sample (Figure 9(d)) are finer than in its PA counterpart (Figure 9(c)) due to the higher concentration of β-stabilizing elements, which were rejected from the grain boundary α particles, and lower concentration of O, which segregated into the grain boundary α from the matrix. This resulted in increased stability of the remaining β phase prior to the second aging step.

In contrast, during the single-step aging, α precipitates inside the grains directly and only a limited amount of grain boundary α is present (Figures 9(e) and (f)), resulting in similar size of α lamellae in the PA and BE samples. 600 °C is also a typical temperature used for aging of Ti alloys to produce a fine α + β microstructure, e.g. References 12 and 13.

Microhardness of the aged samples was subsequently measured, and the results are presented in Table IV. The microhardness of the as-sintered samples (from Table III) is also shown again for easy comparison. There is a significant increase in microhardness of the PA samples from 291 HV in the as-sintered condition to 370 and 380 HV in the aged condition due to the formation of a fine two-phase α + β microstructure. The single-step aged material is slightly harder than the two-step aged material due to the finer α lamellae size in the former material.

Table IV Vickers Microhardness of the Sintered Alloy After Three Different Aging Schemes
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Only a small increase from 391 HV to 421 and 429 HV is observed in the BE samples, because the as-sintered BE sample has already contained a significant amount of α phase (see Figure 9(b)), but the microhardness homogeneity of the BE sample improved dramatically, which is manifested by a decrease of a standard deviation from 31 to 10 HV. Slightly higher microhardness was measured in the two-step aging condition, again due to the finer α lamellae.

The PA and BE materials after single-step aging still exhibit a difference in microhardness despite the same microstructure with α lamellae of a similar size. This can be attributed to the increased oxygen content in the BE when compared to PA samples (0.19 vs 0.1 wt pct, respectively, cf. Table II), causing hardening of both α lamellae and β matrix.

As the presence of the grain boundary α particle is generally detrimental for mechanical properties, a single-step aging scheme was selected for manufacturing of the large specimen, which will be used to analyze the macroscopic mechanical properties.

3.2 Sintering of Large Samples

In the previous section, we have shown that the material prepared by FAST from elemental powder blend has the microstructure and microhardness which are comparable to the sintered prealloyed powder. However, for practical applications, the material strength and ductility must be determined from macroscopic tensile tests to allow direct comparison with the conventionally produced alloy. Therefore, we have produced a rod (15 mm in diameter, 75 mm in length) by FAST under the same conditions that were as close as possible to those used in production of the sample above. This allows us to prepare tensile specimens of standard size, which can be directly compared to a TIMET Ti-5553 cast and forged (CF) alloy. To ensure the completely same thermal processing, crucial for achieving the same microstructure, the sintered rod was subjected to the same aging process as the CF rod. The details are provided in the Experimental section of the paper.

After the heat treatment, oxygen and nitrogen levels were checked in both samples, the results are presented in Table V. The nitrogen level is very low in both samples. The oxygen of the sintered large sample is higher than that of the smaller sintered samples, but corresponds to the calculated oxygen level of the powder blend (cf. Table II). This may be caused by the sample geometry—i.e., it is not possible for gases to desorb from the powder surface and escape from a larger die before pores are closed during sintering.

Table V The Measured Oxygen and Nitrogen Content of the Rod Samples for Tensile Testing
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Microstructure was checked before the tensile tests; the corresponding micrographs are shown in Figure 10. Both materials exhibit a two-phase α + β lamellar microstructure of the same scale and morphology as the directly in situ aged sample shown above in Figure 9(f). Therefore, these three samples (Figures 5(f), 6(a), and (b)) represent the same microstructural state and can be directly compared to each other.

Fig. 10
figure 10

A backscatter SEM micrograph comparison of (a) the cast sample and (b) the sample sintered from the BE powders

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True stress–true strain tensile curves of the specimens are shown in Figure 11(a). Only one sintered specimen is shown, as the other one was damaged during preparation. Both materials, conventionally processed and sintered, exhibit a high yield strength of (1215 ± 4) MPa and 1183 MPa, respectively. A moderate ductility of 6 pct in tension was achieved in the sintered rod, which is fully comparable with (6.8 ± 2.3) pct of the CF rod. This relatively large scatter in ductility is caused by a combination of large grains and a relatively small specimen diameter, not having a sufficient number of randomly oriented grains on the cross-section. Nevertheless, it was clearly shown that by the means of FAST with BE powders as a precursor, we can produce material whose mechanical properties are fully comparable with the conventionally prepared alloys. Our values of ductility and strength are in a good agreement with the previously reported results, as shown in Figure 11(b). Similar results were achieved in an extensive study of influence of various aging schemes on mechanical properties and fracture behavior,[13] reporting optimal combination of yield strength of 1000 to 1100 MPa and ductility of 8–16 pct.

Fig. 11
figure 11

(a) Results of the tensile test: a comparison of the sintered and CF Ti-5553 rods. (b) Comparison of the measured yield stress and ductility with available experimental data after the same heat treatment (β solution treatment + 30 min, 600 °C aging). The mechanical properties of the sintered material are comparable to the CF alloy and to the properties measured on a similar Ti-5553-1Zr alloy

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4 Conclusions

The Ti-5553 alloy was successfully produced by a field-assisted sintering technique from blended elemental powders at 1500 °C. No additional contamination occurred during sintering. The alloy was directly aged in the sintering furnace after consolidation, without cooling to the room temperature, resulting in a typical α + β microstructure with good homogeneity.

A long Ti-5553 rod sintered from blended elemental powders exhibited the high strength of 1183 MPa and the ductility of 6 pct after aging imitating the in situ aging. These mechanical properties are fully comparable to the commercial cast and forged alloy prepared under the same conditions. The proposed sintering method integrates compaction and annealing into a single procedure. Moreover, FAST can be upscaled and adjusted for near-net shape processing. Therefore, the proposed manufacturing method can serve as an alternative to the conventional processing route for manufacturing of specialized parts.