Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material

Abstract

Titanium alloys subjected to suitable thermomechanical processing (TMP) schedules can exhibit superplasticity. Most studies on superplasticity of titanium alloys are directed to sheet materials while studies on bulk materials are rather limited. Bulk Superplastic materials require lower load for forging aeroengine components. It further facilitates forming using non-conventional processes such as superplastic roll forming (SPRF). Multi axial forging (MAF), is employed here to achieve bulk superplasticity by imparting large strain without any concomitant change in external dimension. A comparison between uniaxial and MAF with respect to strain, strain path, initial microstructure and heat treatment was carried out to ascertain the microstructure refinement in Ti-6Al-4V alloy. A fine-grained structure was obtained after 3 cycles of MAF followed by static recrystallization at 850°C. Grain boundary sliding was observed in identified processing domain along with strain rate sensitivity (SRS) of 0.46 and maximum elongation of 815%. Validation of established ther¬momechanical sequence on a scaled-up work piece exhibited 640% elongation in domain (T = 820°C, ε ̇= 3 x 10-4/s) which indicated that the established TMP scheme can be used on a reliable and repeatable basis to achieve superplasticity in bulk material.

1. Introduction

Titanium alloys are used extensively in the aerospace industry owing to their high specific strength, good static and dynamic properties, corrosion resistance, etc. [1, 2]. As an engine stator and rotor, they are used extensively in the compressor section as rings, discs and blades. A typical aeroengine compressor (low and high pressure) disc varies from 15 to 100 cm in diameter. These compressor discs are generally manufactured using either conventional or advanced forging techniques such as isothermal or near isothermal forging. Near α and α+β titanium alloys that are used in the compressor region of aeroengine exhibit a martensitic structure with coarse grains in the as-cast and homogenised condition [2]. By subjecting the material to thermomechanical processing (TMP), the lamellar structure can be converted into a fine-grained equiaxed structure. Numerous studies have been carried out to identify TMP parameters to break lamellar structure, and various theories have been put forward to describe the continuous dynamic recrystallisation (CDRX) or globularisation of α lamellae [1, 2]. It has been suggested that the formation of α/α interface (due to sub grain or shear band formation) in contact with β platelet gives rise to surface-driven penetration of the β phase into the lamellae leading to the breaking of lamellae into equiaxed structure. Thin lamellae structure has been reported to show better globularisation kinetics [3] due to easy penetration of β phase into α lamellae. It has been reported that high strain (ε ≥ 3) is necessary to obtain a completely globularised structure [4] which could exhibit superplasticity.

Numerous studies have been carried out towards achieving superplasticity in various titanium alloy sheets [5, 6, 7]. However, work reported on realising superplasticity in the bulk material is very few. Multi-axial forging (MAF) was used to achieve sub-micron grain size in titanium alloy Ti-6Al-4V by Zherebstov et al. [8]. It was reported by Salishev et al. [4] that strain path change inherent in MAF process aids in increasing the globularisation kinetics or CDRX as multiple slip systems are activated during the strain path change. Poths et al. [9] subjected Ti-6Al-4V to monotonic and cyclic torsion in order to understand the effect of strain path on the globularisation kinetics of α lamellae. Based on the study and in contrast to the findings of Salishev et al. [4], Poths et al. [9] reported a decrease in kinetics of α globularisation and attributed the same to the change in strain path. Since contrasting results have been reported in literature and α lamellae globularisation is essential to achieve superplasticity in titanium alloys, it is imperative to carry out a systematic investigation on the influence of various factors that affect the globularisation kinetics of α lamellae. Further, as the objective is to achieve superplasticity in the bulk material, it is required to achieve this without modifying the external dimension of the material that facilitates subsequent secondary processing to produce the desired product/component.

A systematic study on the effect of strain, strain path, deformation temperature and starting or initial microstructure of Ti-6Al-4V on the globularisation or CRDX kinetics in a work horse titanium alloy Ti-6Al-4V is presented here. Further, a suitable thermomechanical process scheme that can maximise globularisation in the material on a reliable and repeatable basis is presented along with the temperature-strain rate regime under which the globularised material exhibits superplasticity.

2. Material

Ti-6Al-4V is an α+β alloy designed to provide moderately high strength, good fatigue strength and reasonable fracture toughness up to a temperature of 350°C. For aeroengine applications, the alloy is produced by vacuum arc remelting followed by thermomechanical processing in order to improve the structural integrity of the material. For the current study, triple vacuum arc remelted (VAR) titanium alloy Ti-6Al-4V ingot subjected to primary processing in the β and α+β field followed by mill annealing at 700°C for 1 h was procured from M/s M/s Mishra Dhatu Nigam, Hyderabad, India. The β transus of the 15 cm cylindrical mill annealed bars was reported to be 995 ± 5°C, and the same was reconfirmed through heat treatment experiments. Two billets of 15 cm diameter and 5 cm length were cut from the as-received mill annealed material. The billets extracted were coated with glass coating to prevent oxidation during high-temperature exposure and were subjected to heat treatment above the β transus at a temperature of 1015°C. The billets were held at this temperature for a period of 60 minutes to achieve thermal homogeneity. One billet was then removed from the furnace and quenched in water while the second billet was cooled in air to obtain martensitic and lamellar microstructures respectively. Typical optical microstructures obtained after water quenching and air cooling are shown in Figure 1a and b respectively. The prior β grain size was estimated to be 601 + 86 μm and 684 + 64 μm for the water quenched and air-cooled material, respectively.

Figure 1.

Typical microstructure of Ti-6Al-4V subjected to heat treatment at 1015°C for 60 min followed by (a) water quenching – martensitic structure and (b) air cooling – lamellar structure.

3. Effect of strain, strain path and microstructure

To evaluate effect of strain, strain path and initial microstructure on the globularisation kinetics of titanium alloy Ti-6Al-4V, both isothermal hot compression (uniaxial monotonic) and multi-axial (non-monotonic) compression experiments were carried out.

3.1 Isothermal hot compression

Cylindrical compression test samples with a constant height-to-diameter ratio of 1.5 were extracted from the water-quenched and air-cooled material using a wire-cut electro discharge machine (EDM) and lathe machine. The edges of the samples were chamfered to avoid fold formation during initial stages of deformation. The prepared samples were coated with glass (Deltaglaze 347) which acts as an oxidation resistor and lubricant. The coated samples were then heated to the desired test temperature and held at that temperature for 30 min in order to achieve thermal homogeneity. The cylindrical samples were then subjected to isothermal hot compression or uniaxial monotonic deformation (height reduction). A deformation of 25%, 40%, 58% and 78% was imparted to the samples at a constant true strain rate of 10⁻³/s using a computer-controlled 200kN formability and workability testing machine custom built by M/s BISS, Bangalore. The deformation imparted corresponds to an equivalent strain of 0.29, 0.58, 0.87 and 1.51, respectively. After deformation, all the samples were quenched in water and were cut parallel to the deformation direction, hot-mounted and subjected to metallographic investigation following standard procedures. For stereological measurements of α fraction globularised, automatic and semi-automatic procedures were used and the α lamellae with an aspect ratio of <2.0 was considered to be globularised as reported in various literatures [4, 9]. Standard statistical measures such as relative accuracy (RA) and 95% confidence level [10, 11] were used to ensure reliability of stereology measurements. The microstructure of Ti-6Al-4V with an initial lamellar structure subjected to increasing amount of strain through isothermal hot compression, i.e. uniaxial monotonic deformation is shown in Figure 2a–d. Deformation of α lamellae and increased globularisation of α lamellae with increasing strain can be readily observed from Figure 2. Maximum globularisation was observed in the sample subjected to 78% deformation (Figure 2d) that corresponds to a strain of 1.51.

Figure 2.

Ti-6Al-4V with martensitic microstructure subjected to (a) 25 (b) 40 (c) 58 and (d) 78 percentage reduction or deformation at 900°C with a strain rate of 10⁻³/s.

Conversion of lamellar structure of α phase into a globular morphology during deformation is considered to be a recrystallisation process, namely CDRX as against the discontinuous dynamic recrystallisation (DDRX) which has a distinctive nucleation and growth stage [12, 13, 14]. Globularisation of α lamellae present in the colony and at the grain boundary has been reported to take place by either sub grain or shear band formation. It was proposed by Margolin and Cohen [15] that subgrains form within the α lamellae during deformation followed by penetration of β phase into the α/α boundary with a simultaneous rotation of boundaries against each other resulting in coarsening of the recrystallised α when compared with the lamellae thickness from which it originated. Weiss et al. [3] reported formation of shear bands as main reason for CDRX of α lamellae. Irrespective of the mechanism, the formation of α/α interface in contact with β induces surface tension driven penetration of β phase resulting in globularisation. Balasundar [16] reported that both these mechanisms, namely sub grain and shear band formation operate in titanium alloys depending on the orientation of α lamellae and the processing conditions.

Based on the orientation of α lamellae with respect to the compression direction, bending and/or kinking of α lamellae can also be readily observed at different locations of the deformed sample. HCP α lamellae that have their c-axis aligned to the deformation direction require very high load or stress to deform [17] because such lamellae have a low Schmid factor and shear force along the slip direction. Though the globularisation fraction of α lamellae increases with increasing deformation, isolated regions where the α lamellae are still intact can be observed in the material even after imparting a stain of 151%. Though globularisation of grain boundary α lamellae could be observed after a stain of 58%, the prior β grain boundaries could be readily observed in the deformed material.

Microstructure of Ti-6Al-4V with an initial martensitic structure subjected to increasing amount of uniaxial monotonic deformation through isothermal hot compression is shown in Figure 3a–d. The influence of strain on the material with martensitic starting microstructure was observed to be similar to that in lamellar microstructure described above. However, the prior β grains were found to be destroyed completely by deformation, and it was not possible to identify them in the material. A quantitative discussion on the α fraction globularised is presented in Section 3.3.

Figure 3.

Ti-6Al-4V with lamellar microstructure subjected to (a) 25 (b) 40 (c) 58 and (d) 78 percentage reduction or deformation at 900°C with a strain rate of 10⁻³/s.

3.2 Multi-axial deformation

Conventional deformation process, such as forging, rolling, extrusion, etc., alters the dimension of the material that is being deformed. Severe plastic deformation is a metal working technique in which very high plastic strain can be imparted to the material without any concomitant change in the geometry or dimension. As the work piece material geometry is not altered, it provides an opportunity to deform the material repeatedly till the desired amount of strain is imparted. This large plastic strain results in the formation of ultrafine grain structure in the material. A large number of SPD processes are in vogue, MAF is one suchtechnique which is easier to implement as no special die or tooling is required for deforming the material. In MAF, the material is deformed in cyclic way along all the three orthogonal directions. Each MAF cycle consist of three deformation steps, i.e. imparting equal amount of deformation along the three orthogonal directions (X, Y, Z) as shown in Figure 4.

Figure 4.

Typical multi-axial forging cycle for imparting 40% deformation along all the three directions (all dimensions in mm).

The strain path is altered when the sample is rotated by 90° during MAF. This change in strain path is expected to assist in refining the grain size. After three processing or deformation steps, i.e. after a cycle of MAF, the work piece reverts back to its original dimension. Since the dimension of the work piece remains the same, it is possible to carry out MAF cycles multiple times and thereby impart a large amount of strain to the material as desired or till the failure of material. The amount of strain imparted for a given percentage of reduction and the number of MAF cycles can be calculated as per the following relation:

$${\epsilon }_{\mathit{eq}}=3\mathrm{N}ln\left(1-\mathrm{R}\right)$$E1

where R is the amount of reduction imparted per direction which is generally assumed to be the same along all the three orthogonal directions, and N is the number of MAF cycles.

Two numbers of cubic specimens with dimension of 2.5 x 2.5 x 2.5 cm³ were prepared from the water-quenched and air-cooled Ti-6Al-4V material respectively for carrying out MAF experiments at 900°C with a constant true strain rate of 10⁻³/s. The first sample from the water-quenched and air-cooled material was subjected to 25% reduction along all the three directions while the second sample was imparted with 40% reduction which corresponds to a cumulative strain of 87% for the sample subjected to 25% reduction and 153% for the 40% reduction sample. While carrying of MAF experiment, the sample was first deformed to the required amount in a particular direction (e.g. X direction), the furnace was opened, and then the sample was rotated and positioned for deforming along the next direction (e.g. Y direction). The furnace was closed and sample was reheated to the desired temperature and held at this temperature for 30 min before imparting the desired amount of deformation along this direction (e.g. Y direction). The procedure was repeated for deforming along the third direction (e.g. Z direction). After completion of desired number of MAF cycle, i.e. after imparting desired amount of reduction along all the three directions as shown in Figure 4, the samples were water-quenched to free the microstructure of the material.

Microstructure of Ti-6Al-4V with an initial lamellar structure subjected to a deformation of 25% and 40% along all the three orthogonal direction using MAF process is shown in Figure 5a and b respectively. No major microstructural change is observed in the material subject to 25% when compared with the initial starting microstructure expect for coarsening of α lamellae and few isolated regions of α globularisation within the grain and at grain boundaries. Partial globularisation of α can be observed at all the three faces of the same subjected to 40% deformation through MAF. Though isolated globularisation of grain boundary α was observed, the prior β grains were found to be intact and distinct.

Figure 5.

Microstructure of Ti-6Al-4V with lamellar microstructure subjected to (a) 25% and (b) 40% deformation along all the three orthogonal directions through MAF at 900°C with a strain rate of 10⁻³/s.

The microstructures of Ti-6Al-4V with martensitic structure subjected to 25% and 40% reduction using MAF process are shown in Figure 6a and b respectively. Similar to that of the air-cooled structure, martensitic structure also shows increasing globularisation with increasing amount of deformation. However, the breaking of the prior β grain boundaries and globularisation can be observed to be more in martensitic structure when compared with the lamellar structure.

Figure 6.

Microstructure of Ti-6Al-4V with martensitic microstructure subjected to (a) 25 and (b) 40% deformation per direction through MAF process at 900°C with a strain rate of 10⁻³/s.

3.3 Globularisation fraction

The globular α volume fraction was estimated in Ti-6Al-4V with martensitic and lamellar starting structure subjected to uniaxial (monotonic, i.e. no change in strain path) and multi axial (non-monotonic, i.e. changing strain path as the sample is rotated) deformation and is shown in Figure 7. It can be seen that, irrespective of the deformation process whether it is monotonic (uniaxial) or non-monotonic (multi-axial), the volume fraction of globularised α increases with increasing strain. The rate of increase in globularised α volume fraction is observed to be high for Ti-6Al-4V with martensitic structure when compared with the material with lamellar structure. This observation here concurs with the report of Shell et al. [18] where they have attributed the increased kinetics of globularisation to α lamellae thickness. When the α lamellae is thin, it is relatively easier for the β phase to penetrate the lamellae and transverse boundary when compared with thick lamellae. It can be noted that for an equivalent strain of 0.87 which corresponds to 58% reduction through uniaxial or monotonic deformation and 25% deformation along all the three directions through multi-axial or non-monotonic deformation, altering the strain path through multi-axial deformation results in reduced globularisation. However, with increasing equivalent strain from 0.87 to 1.51, i.e. increasing the deformation from 25 to 40% per direction in multi-axial deformation, the fraction of α phase globularised is quite comparable to that obtained through uniaxial deformation (with 78% reduction) as shown in Figure 7.

Figure 7.

Effect of strain and strain path on globularisation of α lamellae in Ti-6Al-4V.

It was reported by Banerjee et al. [19] that when Ti alloys are subjected to thermomechanical processing in the α/β regime, an equiaxed microstructure can be obtained only if the amount of deformation is greater than 30%. As stated earlier, the CRDX or globularisation of α lamellae depends on the formation of α/α interface by shearing, subgrain, buckling and kinking, etc. and the penetration of β phase. As the formation of α/α interface is a result of two contending processes of dislocation accumulation and annihilation, it is proposed that altering strain path before attaining a limiting critical strain leads to large annihilation of dislocations and a possible reduction or disappearance of the substructure. In the material subjected to 25% deformation through MAF, the strain path is changed before reaching the critical amount of deformation (~30%), and the fraction of globularisation is less due to large annihilation of dislocations. With increasing strain from 87 to 151%, adequate strain is available for accumulation of dislocations and formation of stable substructure. Hence, altering the strain path does not influence the fraction of globularisation. From Figure 7, it can be readily inferred that 40% deformation through non-monotonic MAF does not cause significant reduction in the globularisation kinetics when compared with uniaxial deformation. Concomitantly, no significant improvement has been observed as reported by Salischev et al. [4].

On the basis of above study, it can be readily inferred that a martensitic structure exhibits better globularisation kinetics in comparison to lamellar structure. Altering the strain path through non-monotonic multi-axial deformation below a critical limiting strain results in reduced globularisation, whereas beyond the critical strain, the results are quite comparable. The critical deformation limit has been identified to be ≥40% through the current investigation. Though required globularisation can be achieved by uniaxial deformation, it results in altering the dimension of the work piece, whereas multi-axial deformation process which does not cause any concomitant change in the external dimension [20]. Further, unlike uniaxial deformation which produces a dead metal zone at the fiction hill, no such un-deformed region is observed in multi-axial deformation; hence, multi-axial deformation process with 40% reduction per direction is best suited to globularise the α lamellae and refine the microstructure to obtain superplasticity.

4. Thermomechanical processing scheme to achieve ultrafine grain structure

To obtain grain refinement and achieve superplasticity, Ti-6Al-4V with a martensitic starting microstructure was subjected to three cycles of multi-axial deformation with 40% reduction per direction by progressively decreasing the deformation temperature at each cycle from 850, 800 and 750°C, i.e. first cycle of MAF is carried out at 850°C, second cycle of MAF on the sample is carried out at 800°C and so on. Microstructures of Ti-6Al-4V obtained after two and three cycles of multi-axial deformation are shown in Figure 8a and b respectively. It can be seen that three cycles of multi-axial deformation have resulted in complete globularisation of α lamellae.

Figure 8.

Microstructure of Ti-6Al-4V after (a) two cycles and (b) three cycles of MAF.

The mechanism of globularisation or CDRX of α lamellae depends on formation of α/α boundaries which is followed by penetration of β phase into α lamellae to separate the boundary [15]. Formation of stable substructure (α/α boundaries) or low-angle grain boundaries (LAGB) depends on interaction and multiplication of dislocation in the lamellar structure. When the HCP α lamellae c-axis is oriented parallel to the deformation direction, high shear stresses are required for the operation of basal and prism slip which is the major source of slip in HCP structure [17]. Such unfavourable orientated lamellae exhibit no major deformation leading to low dislocation density and less α/α boundaries which reduces the globularisation kinetics. However, rotation of specimen by 90^o during subsequent steps of deformation during multi-axial deformation leads to activation of prism slip system in such lamellae due to an increase in value of Schmid factor. Activation of prism slip system initiates the process of slip, and with increasing strain and grain rotation, the basal slip system also starts operating leading to formation of substructure, which eventually results in dynamic recrystallised or globularised equiaxed α grain. A comparison of the microstructures shown in Figure 6b, 8a and b reveals that with increasing the number of MAF cycle, globularisation increases. It can also be seen that the isolated α lamellae visible in Ti-6Al-4V specimen after two cycles of multi-axial deformation are converted into equiaxed grains after three cycles. Further, prior β grain boundaries were completely destroyed, and they are no longer distinguishable after three MAF cycles.

To gain further insight on the microstructure evolution, the material subjected to three cycles of multi-axial deformation was subjected to electron back scattered diffraction (EBSD) characterisation. The low-angle (LAGB; red line) and high-angle (HAGB; black line) boundary present in the material was estimated to be 24% and 76% respectively. The observed fraction of HAGB and LAGB is consistent with reported literature which states the presence of around 25–30% of LAGB in heavily deformed structures [20]. The average size of HAGB was estimated to be 1.65 ± 0.6 μm. Further, the inverse pole figure (IPF map) shown in Figure 9(b) reveals a random texture with no preferred orientation.

Figure 9.

Microstructure of Ti-6Al-4V after three cycles of MAF(a) band contrast image highlighting HAGB (black line) & LAGB (red line) (b)I PF colour map obtained through EBSD.

Superplasticity occurs by grain boundary sliding accommodated by diffusion at grain boundaries and lattice [21, 22]. A material with high fraction of HAGB will have higher degree of disorder. As diffusion occurs down the potential gradient, large fraction of HAGB will lead to higher gradient which increases the diffusion rate. It has been reported in the literature that the presence of high fraction of HAGB with maximum fraction lying in the range of 30–60° is necessary for ease of grain boundary sliding [4]. Further, a fairly uniform grain structure with grains of similar size is essential for superplastic forming. The presence of large grains and small grains (mixed or bimodal grains) in same microstructure has a negative effect on superplasticity. It has been reported that kinetics of diffusion is very slow around the large grain in comparison to the kinetics around smaller grain structure [23]. In the absence of slow diffusion around large grains, chances of cavitation are higher around the larger grains leading to early failure without appreciable tensile elongation. It is also well known that higher volume fraction of second phase improves superplastic property due to easier α/β grain boundary sliding [24]. Further, in titanium alloys, the diffusivity in BCC β grain is about of two orders of magnitude higher than α alloys [1, 2, 20]. Higher diffusivity of β phase improves superplastic property of titanium alloy. However, at the same time, β phase grows much faster than α phase, and hence, α phase helps in pinning the grain boundary of β and does not allow the grains to grow. Ideally, presence of both phases is essential for easier grain sliding and achieving optimal superplastic property [25].

As the material after three cycles of multi-axial deformation exhibits nearly 24% low-angle boundaries and predominantly α phase, it is essential to convert this microstructure into a one that can exhibit superplasticity. To improve microstructure and to obtain uniform grain size distribution with increased volume fraction of β phase and higher fraction of high-angle grain boundary, annealing of the deformed material was carried out at 850°C for 2 h followed by air cooling. A higher temperature was used for annealing to obtain the desired α and β proportions in the material. From the EBSD band contrast image shown in Figure 10a, the β phase present in the heat treated material was estimated to be ~15%. The fraction of HAGB in the heat-treated material was estimated to be 95% with an average grain size of 2.53 ± 0.65 μm. The IPF map shown in Figure 10b clearly indicates the random orientation of the grains. A random texture is important for a structure to exhibit superplastic behaviour since cavitation may occur along the transverse direction during deformation due to strain incompatibility [21]. The misorientation profile of the grains shows significant improvement in HAB with a high fraction of grains boundary in the range of 30–60°. The resultant microstructure satisfies the condition of superplasticity as a major mode of deformation.

Figure 10.

Microstructure of Ti-6Al-4V after three cycles of MAF and heat treatment at 850°C for 2 h followed by air cooling (a) band contrast image highlighting HAGB (black line) & LAGB (red line) (b) IPF colour map (c) misorientation plot obtained through EBSD.

5. Superplastic domain identification and validation

In order to identify the temperature-strain rate domain under which the material exhibits superplasticity, the SRS of the material subjected to three cycles of MAF and heat treatment was evaluated by carrying out isothermal hot compression tests over a range of temperature and constant true strain rate. Using the flow curves, the SRS of the material was estimated using standard relations [26] and plotted as a function of temperature and strain rate as shown in Figure 11. It can be seen that the material exhibits a maximum SRS of 0.46 between 810°C and 825°C for a constant strain rate of 3×10⁻⁴/s.

Figure 11.

(a) Strain rate sensitivity map at ε = 0.5 (b) initial tensile specimen (gage length = 6 mm, diameter = 4 mm) (c) sample after tensile deformation at 820°C with a strain rate of 10⁻³/s strain rate and (d) sample after tensile deformation at 820°C with a strain rate of 3×10⁻⁴/s strain rate.

In order to validate the domain identified, cylindrical tensile samples of 0.6 cm gage length and 0.4 cm diameter (Figure 11b) were prepared from the material subjected to three cycles of MAF and heat treatment. Tensile tests were carried out at 820°C with a strain rate of 10⁻³ and 3×10⁻⁴/s. A maximum elongation of 815% was obtained in the sample tested with a strain rate of 3×10⁻⁴/s (Figure 11d). As the strain rate is increased to 10⁻³/s, the % elongation obtained decreases from 815 to 447% as shown in Figure 11c. During tensile deformation, necking starts in the weak part of the structure and further deformation gets concentrated at this necked region. A tri-axial state-of-stress exists in the neck region and strain rate in this region does not follow the strain rate of the specimen. The strain rate in the region depends on the rate of decrease of the area which is given by [26]:

$$\frac{dA}{dt}={\left(\frac{P}{C}\right)}^{\frac{1}{m}}\left(\frac{1}{A^{\left(1-m\right)/m}}\right)$$E2

As the rate of decrease in area is inversely proportional to strain rate sensitivity (m), a higher m value leads to slower strain rate in the necked region. So a higher elongation is observed during tensile deformation.

In order to further substantiate the findings, EBSD characterisation of the material subjected to compression test (after three cycles of multi-axial deformation and heat treatment) at 800°C with a strain rate of 3×10⁻⁴/s was carried out to evaluate the grain size, boundary fractions, etc. From the band contrast image shown in Figure 12, the fraction of HAB and the average grain size were estimated to be 83% and 2.64 ± 0.89 μm respectively. Comparing the microstructure obtained in the material after deformation in the superplasticity domain and that obtained after three cycles of multi-axial deformation and heat treatment (Figure 10a) clearly indicates no major change in grain size. Therefore, from the observations on % elongation and microstructural features, it can be confirmed that the heat-treated material after subjecting to three cycles of multi-axial deformation exhibits superplasticity between 800 and 840°C when deformed with a strain rate of 3×10⁻⁴/s

Figure 12.

Band contrast image of fine grained Ti-6Al-4V after hot compression at 800°C and strain rate of 0.0003 S⁻¹ highlighted with HAGBs (black), LAGBs (red).

6. Scaling-up

In order to ensure viability of the identified thermomechanical scheme for industrial scale processing, a large-size billet of 15 cm × 12 cm × 9 cm was prepared and subjected to MAF using 2000MT hydraulic forge press. Figure 13 compares the size of scaled up Ti-6Al-4V billet with that of smaller-size specimen. The microstructure evolution in the large billet after each cycle of MAF is shown in Figure 14. It can be clearly seen that the globularisation fraction increases with increasing cycle and a completely globularised structure is achieved after three cycles of MAF. Post-deformation heat treatment for the material was carried out at 850 and 900°C for 2 h followed by air cooling, and the resulting microstructures are shown in Figure 15. Heat treatment led to increase in the volume fraction of beta phase with slight coarsening in the grain size.

Figure 13.

Typical microstructure of large size Ti-6Al-4V after (a,d) first MAF cycle (b,e) second MAF cycle (c, f) third MAF cycle.

Figure 14.

Microstructure of Ti-6Al-4V obtained after third MAF cycle followed by heat treatment at (a) 850°C/2 h/ac and (b) 900°C/2 h/ac.

Figure 15.

Comparison of scaled-up Ti-6Al-4V billet and small size specimen.

In order to validate super plasticity, standard cylindrical tensile samples of 2.0 cm gage length and 0.4 cm diameter were prepared and subjected to tensile testing. Tensile tests were carried out at 820°C with a Strain rate of 3 x 10⁻⁴ and 10⁻³/s. Tensile test was carried out for as deformed (three cycles of MAF) specimen also. It can be seen from Figure 16 that a maximum elongation of 640% has been obtained for the material.

Figure 16.

Tensile elongation of Ti-6Al-4V subjected to three cycles of MAF followed by heat treatment at 850°C/2 h/ac and (a) tested at έ = 0.0003/s, T = 820°C (b) tested at έ = 0.001/s, T = 820°C (c) Tensile elongation of Ti-6Al-4V subjected to three cycles of MAF tested at έ = 0.001/s, T = 820°C.

The result is comparable with the one achieved during tensile test of smaller-size MAF processed specimen. The obtained results have been also compared with available literature as shown in Table 1. It can be seen that obtained m values and elongation values are comparable with the data available in literature. Elongation obtained during tensile test of specimen extracted from smaller-size MAF specimen and large-size MAF billets clearly indicates superplastic behaviour of three cycles of MAF Ti-6Al-4V alloy. Hence, it can be concluded that the TMP scheme established using smaller sample is repeatable and is validated by repeating the process using 2000 T hydraulic Forge press.

Processing condition	Grain size (μm)	Temp (°C)	Strain rate (s^−1)	Elongation (%)	m	Ref
Rolled sheet	0.3	775	1 × 10^−4	800	0.45	[27]
		775	1 × 10^−3	600	0.45
		875	1 × 10^−4	720	0.41
		875	1 × 10^−3	380	0.41
	3	775	1 × 10^−4	200	0.32
		875	1 × 10^−4	420	0.33
		875	1 × 10^−3	390	0.33
SPD processed bulk samples	0.2	600	5 × 10^−4	500	0.34	[28]
	5	800	5 × 10^−4	600	0.4
Hot Rolled	3	700	1 × 10^−4	583	0.66	[29]
SPD processed (HPT)	0.2	650	1 × 10^−4	568	0.36	[30]
		750	1 × 10^−2	504	0.46
DMRL 3 cycle MAF + Heat treated	2.53	820	3 × 10^−4	640	0.46
		820	1 × 10^−3	525	0.38

Table 1.

Comparison of achieved superplastic property with available literature.

7. Conclusions

From the study it can be concluded that three cycles of MAF with 40% reduction per direction (cumulative effective strain of ~4.6) leads to complete globularisation of martensitic structure with an average α grain size of 1.65 ± 0.6 μm. Annealing of the deformed material at 850°C increases HAB fraction and β phase volume fraction but with a marginal increase in grain size (2.53 ± 0.65 μm). The MAF + annealed material exhibits a maximum SRS of 0.46 when deformed between 810°C and 825°C with constant strain rate of 3 x 10⁻⁴/s. A maximum tensile elongation of 815% was obtained with strain rate of 3 x 10⁻⁴/s at 820°C. TMP designed was implemented on a large-size work piece under near isothermal condition, and the process was found to be reliable and repeatable to obtain superplasticity in bulk Ti-6Al-4V in large-size specimen for aeroengine applications.

Acknowledgments

The funding provided by Defence Research and Development Organisation is acknowledged.

Conflict of interest

The authors declare no conflict of interest.