FAST-Forge of Diffusion Bonded Dissimilar Titanium Alloys

metals

Article FAST-forge of Diﬀusion Bonded Dissimilar Titanium Alloys: A Novel Hybrid Processing Approach for Next Generation Near-Net Shape Components

Jacob Pope and Martin Jackson *

The University of Sheffield, Department of Material Science and Engineering, Sir Robert Hadfield Building, Sheffield S1 3JD, UK; jpope1@sheffield.ac.uk * Correspondence: martin.jackson@sheffield.ac.uk; Tel.: +44-114-2225474

Received: 19 May 2019; Accepted: 3 June 2019; Published: 4 June 2019

Abstract: Material reductions, weight savings, design optimisation, and a reduction in the environmental impact can be achieved by improving the performance of near-net shape (NNS) titanium alloy components. The method demonstrated in this paper is to use a solid-state approach, which includes diffusion bonding discrete layers of dissimilar titanium alloy powders (CP-Ti, Ti-6Al-4V and Ti-5Al-5Mo-5V-3Cr) using field-assisted sintering technology (FAST), followed by subsequent forging steps. This article demonstrates the hybrid process route, firstly through small-scale uni-axial compression tests and secondly through closed-die forging of dissimilar titanium alloy FAST preforms into an NNS (near-net shape) component. In order to characterise and simulate the underlying forging behaviour of dissimilar alloy combinations, uni-axial compression tests of FAST cylindrical samples provided flow stress behaviour and the effect of differing alloy volume fractions on the resistance to deformation and hot working behaviour. Despite the mismatch in the magnitude of flow stress between alloys, excellent structural bond integrity is maintained throughout. This is also reflected in the comparatively uncontrolled closed-die forging of the NNS demonstrator components. Microstructural analysis across the dissimilar diffusion bond line was undertaken in the components and finite element modelling software reliably predicts the strain distribution and bond line flow behaviour during the multi-step forging process.

Keywords: spark plasma sintering; thermomechanical processing; closed-die forging; process modelling; powder consolidation; downstream processing; titanium powder processing

1. Introduction Titanium alloys are desirable to use in high-performance applications due to a combination of a high strength-to-weight ratio, excellent corrosion resistance, and fatigue performance. Consequently, titanium alloys are a key material for the aerospace industry in components such as gas turbine compressor parts, engine pylons, and landing gear forgings. Today, small to medium aerospace high-strength components, such as hook, pylon, and strut profiles are machined from a single titanium alloy forging. Optimising the mechanical properties of metallic components benefits performance, by reducing cost and weight through optimal design and minimising the environmental impact. a conventional approach for material property optimisation is to produce tailored microstructures, for example, the dual grain size distribution in a nickel superalloy aero-engine gas turbine disk. Hot isostatic pressing (HIPing), extrusion, and forging processes of nickel superalloy powder feedstock can produce fine-grain lamellar microstructures in the disk bore for low-cycle fatigue (LCF) resistance and coarse-grained microstructures in the disk rim for optimised creep resistance.

Metals 2019, 9, 654; doi:10.3390/met9060654 www.mdpi.com/journal/metals Metals 2019, 9, 654 2 of 19

Although dissimilar titanium alloys have been demonstrated that they can be joined through linear friction welding in blisk (bladed rotor disc) prototype parts [1,2], such solid-state joining of dissimilar titanium alloys has yet to be implemented into production parts. Yet in the future, optimised performance could be achieved in titanium components if different alloys and tailored microstructures were targeted for different sub-component regions. In the case of a gas turbine engine blisk for example, circumferential blades produced from a creep resistant, near α titanium alloy could be bonded to a rotor disk produced from a high strength, low-cycle fatigue resistant near β or α+β titanium alloy. A new hybrid powder metallurgy approach combining accelerated sintering with hot forging, could provide designers with the option to engineer future components with sections containing dissimilar titanium alloys, thus integrating controlled changes in properties across the component. a solid-state powder processing method known as field assisted sintering technology (FAST) or spark plasma sintering (SPS) is emerging as a flexible powder consolidated route for titanium alloys [3–15]. When compared to conventional powder processing and sintering methods, FAST is advantageous due to an electrical current, which has been shown to provide faster heating rates and consolidation compared to conventional sintering [16,17]. FAST has been demonstrated to be an effective sintering route for complex shapes, most notably a novel deformable interface approach produced fully dense CoNiCrAlY turbine blades [18]. Furthermore, FAST-forge [5] and FAST-DB [19] processing routes have been developed, primarily to provide an alternative route to produce titanium components from particulate feedstocks. The first step of FAST-forge is the consolidation of the titanium feedstock into an optimised forging preform using FAST. The subsequent step produces a near-net shape (NNS) component by closed-die forging to refine the as-FAST microstructure. Hot forging is beneficial in titanium alloy components not only to provide a NNS, but to induce flow lines into the product to enhance strength and breakdown the α + β structural unit size. In terms of bonding dissimilar alloys, solid-state methods are preferred over fusion joining methods, which are subject to solidification issues such as chemical segregation. Recently, FAST-DB has demonstrated the ability to diffusion bond dissimilar titanium alloy gas atomised powders in a simple cylindrical mould using the FAST process [19]. The bonds exhibited excellent structural integrity during tensile testing with no cracks or porosity evident at the bond line interface. Furthermore, the prior β grains at the bond interface had functionally graded chemistries with respect to the dissimilar alloys diffusion bonded by FAST. In this paper, dissimilar titanium alloys bonded through FAST were characterised and simulated through subsequent hot closed-die forging stages, in order to demonstrate a robust hybrid processing approach for next generation titanium components. Firstly, small-scale FAST-DB compression samples consisting of Ti-6Al-4V (Ti-6-4) and commercially pure titanium (CP-Ti) alloys were deformed using controlled uni-axial compression testing to understand the deformation characteristics. Additionally, for the first time, finite element (FE) software DEFORM 3D has been validated as a method for modelling the controlled deformation of FAST-DB samples to determine the bond interface evolution and strain profiles. Secondly, a semi-complex NNS eye-bolt component was produced using the FAST-DB processing route with subsequent closed-die forging steps. The microstructural evolution of two FAST-DB alloy combinations containing Ti-6-4 bonded with commercially pure titanium and Ti-5Al-5Mo-5V-3Cr (Ti-5553) bonded with Ti-6-4 were characterised post forging. The alloys Ti-6-4 and Ti-5553 were selected for the large-scale NNS forgings due to their increasing usage in the aerospace sector, making them ideal alloys to demonstrate the FAST-forge technology on dissimilar alloys.

2. Materials and Methods

2.1. Materials The Ti-6-4 and CP-Ti powders were both supplied by Phelly Materials Inc., Upper Saddle River, NJ, USA. They were produced via the hydride-dehydride (HDH) method and consequently have an Metals 2019, 9, 654 3 of 19 angularMetals 2019 morphology., 9, x REVISED 31 The May Ti-6-4 2019 powder chemistry achieved ASTM B348-11 Grade 5 standard3 of and 19 CP-Ti powder chemistry achieved Grade 1 standard. The Ti-5553 powder was gas atomised (GA) by CP-Ti powder chemistry achieved Grade 1 standard. The Ti-5553 powder was gas atomised (GA) by TLS Technik Spezialpulver from a forged billet and, therefore, the powder had a spherical morphology. TLS Technik Spezialpulver from a forged billet and, therefore, the powder had a spherical The particle size distribution (PSD) of both powders are shown in Figure1a, and the morphology of morphology. The particle size distribution (PSD) of both powders are shown in Figure 1a, and the CP-Ti, Ti-6-4, and Ti-5553 are shown in Figure1b–d, respectively. morphology of CP-Ti, Ti-6-4, and Ti-5553 are shown in Figure 1b–d, respectively.

Figure 1. (a) Graph showing the PSD of each powder; (b) corresponding secondary electron (SE) micrographFigure 1. (a) of Graph the CP-Ti showing HDH the powder; PSD of(c) each corresponding powder; (b SE) c micrographorresponding of thesecondary Ti-6-4 HDH electron powder; (SE) andmicrograph (d) the corresponding of the CP-Ti HDH SE micrograph powder; (c) of corresponding the Ti-5553 GA SE powder. micrograph of the Ti-6-4 HDH powder; and (d) the corresponding SE micrograph of the Ti-5553 GA powder. 2.2. Field Assisted Sintering 2.2. Field Assisted Sintering The large-scale FAST billet was processed at Kennametal Ltd. in Newport, Ghent, UK using an FCTThe Systeme large GmbH-scale FAST Type billet HP D was 250 processed spark plasma at Kennametal sintering furnaceLtd. in Newport, with a 250 Ghent, mm graphite UK using ring an mould.FCT Systeme Details GmbH of the fullType assembly HP D 250 are spark shown plasma in detail sintering in Figure furnace2. with a 250 mm graphite ring mould.The Details photographs of the full in Figureassembly3 show are shown the stages in detail of the in manual Figure lay-up2. of dissimilar titanium alloy powders prior to large-scale FAST processing. To reduce contamination, Ti-6-4 sheet acts as a temporary divider between dissimilar alloy powders during the manual lay-up. a dwell temperature of 1200 ◦C was used, with a dwell time of 60 mins and a pressure of 35 MPa (1720 kN). The chamber was then evacuated to approximately 2 10 3 bar and pulsed current was applied through the electrodes to × − begin heating. Due to the size of the specimen, a slow heating rate of 25 ◦C/min was required, and the average cooling rate in the water-cooled vacuum chamber was 4 ◦C/min. Graphite foil was lined on the inside of the mould, as well as on the punch surfaces in contact with the powder, to prevent powder from consolidating to the graphite tooling and to aid sample extraction. The processing chamber was held under vacuum for the entire duration of the process to reduce oxygen pickup. The temperature control was regulated using a vertical pyrometer, which measured the temperature 21 mm above the powder surface within the graphite tooling. To reduce heat loss via radiation, a woven carbon ﬁbre jacket housed the assembly.

Figure 2. FAST-DB processing at the Kennametal Ltd. UK facility showing photographs of; (a) the FCT Systeme FAST processing machine; (b) the 250 mm inner diameter assembly consisting of graphite tooling and an outer woven carbon fibre jacket; and (c) the sliced CAD model of the

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CP-Ti powder chemistry achieved Grade 1 standard. The Ti-5553 powder was gas atomised (GA) by TLS Technik Spezialpulver from a forged billet and, therefore, the powder had a spherical morphology. The particle size distribution (PSD) of both powders are shown in Figure 1a, and the morphology of CP-Ti, Ti-6-4, and Ti-5553 are shown in Figure 1b–d, respectively.

Figure 1. (a) Graph showing the PSD of each powder; (b) corresponding secondary electron (SE) micrograph of the CP-Ti HDH powder; (c) corresponding SE micrograph of the Ti-6-4 HDH powder; and (d) the corresponding SE micrograph of the Ti-5553 GA powder.

2.2. Field Assisted Sintering The large-scale FAST billet was processed at Kennametal Ltd. in Newport, Ghent, UK using an Metals 2019, 9, 654 4 of 19 FCT Systeme GmbH Type HP D 250 spark plasma sintering furnace with a 250 mm graphite ring mould. Details of the full assembly are shown in detail in Figure 2.

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aforementioned assembly showing the division of powders and the pyrometer hole passing through the tooling.

The photographs in Figure 3 show the stages of the manual lay-up of dissimilar titanium alloy powders prior to large-scale FAST processing. To reduce contamination, Ti-6-4 sheet acts as a temporary divider between dissimilar alloy powders during the manual lay-up. A dwell temperature of 1200 °C was used, with a dwell time of 60 mins and a pressure of 35 MPa (1720 kN). The chamber was then evacuated to approximately 2 x 10−3 bar and pulsed current was applied through the electrodes to begin heating. Due to the size of the specimen, a slow heating rate of 25 °C/min was required, and the average cooling rate in the water-cooled vacuum chamber was 4 °C/min. Graphite foil was lined on the inside of the mould, as well as on the punch surfaces in contact with the powder, to prevent powder from consolidating to the graphite tooling and to aid sample extraction. The Figure 2. FASTFAST-DB-DB processingprocessing atat thethe KennametalKennametal Ltd. Ltd. UK UK facility facility showing showing photographs photographs of; of; (a )(a the) the FCT processing chamber was held under vacuum for the entire duration of the process to reduce oxygen FCTSysteme Systeme FAST FAST processing processing machine; machine; (b) the (b 250) the mm 250 inner mm diameter inner diameter assembly assembly consisting consisting of graphite of pickup. The temperature control was regulated using a vertical pyrometer, which measured the graphitetooling and tooling an outer and woven an outer carbon woven ﬁbre carbon jacket; fibre and (jackec) thet; slicedand (c CAD) the modelsliced of CAD the aforementioned model of the temperatureassembly 21 showing mm above the division the powder of powders surface and thewithin pyrometer the graphite hole passing tooling. through To reduce the tooling. heat loss via radiation, a woven carbon fibre jacket housed the assembly.

Figure 3. Photographs of the large-scale manual powder layup at Kennametal Ltd. UK, showing; Figure 3. Photographs of the large-scale manual powder layup at Kennametal Ltd. UK, showing; (a) (a) titanium sheet and Perspex dividers in the 250 mm diameter mould; (b) dissimilar powders filled in titanium sheet and Perspex dividers in the 250 mm diameter mould; (b) dissimilar powders filled in the separate sections with agitation of powder using a brush; and (c) flattening of powder after the the separate sections with agitation of powder using a brush; and (c) flattening of powder after the dividers have been removed. dividers have been removed. 2.3. Hot Compression Testing and Determination of Flow Stress 2.3. Hot Compression Testing and Determination of Flow Stress To develop an understanding of the upset hot forging of FAST-DB dissimilar titanium alloy samples,To develop the rheological an understanding behaviour of of the the single upset alloy hot constituentsforging of FAST were- firstDB dissimilar measured totitanium provide alloy flow samples,stress data the for rheological the FE model. behaviour CP-Ti samplesof the single were alloy tested constituents at temperatures were offirst 800, measured 900, and 950to provide◦C and 1 flowat constant stress data strain for rates the FE of 0.1,mode 1,l. and CP- 10Ti ssamples− . Ti-6-4 were samples tested were at temperatures tested at temperatures of 800, 900, ofand 850, 950 950, °C −1 1 and at 1050 constant◦C and strain at constant rates of strain 0.1, 1, rates and 10 of 0.1,s . Ti 1,- and6-4 samples 10 s− . were tested at temperatures of 850, 950, and 1050The large°C and FAST-DB at constant billet strain was usedrates toof machine0.1, 1, and nine 10 s compression−1. cylinders at the bond interfaces whichThe contained large FAST a varying-DB billet volume was used fraction to machine of Ti-6-4 nine and compression CP-Ti. The cylinders cylinders were at the 15 bond mm height interfaces and which10 mm contained diameter providinga varying avolume height tofraction diameter of Ti ratio-6-4 of and 1.5:1, CP as-Ti. recommended The cylinders by were Roebuck 15 mm et al.height [20]. andBefore 10 mm deformation, diameter theproviding bond line a height at the to upper diameter and ratio lower of surface 1.5:1, as of recommended each cylinder by was Roebuck determined. et al. [20It was]. Before assumed deformation, the bond line the followed bond line a linear at the path upper between and lower the two surface surfaces. of each a 1.1 mmcylinder diameter was determined.thermocouple It was hole assumed was drilled the intobond the line centre followed of each a linear sample path from between the side, the wheretwo surfaces. a 1 mm A diameter 1.1 mm diameterN-type thermocouple thermocouple was hole inserted was drilled to achieve into the close centre temperature of each sample control from during the hot side, compressions. where a 1 mm diameterA Servotest N-type Testing thermocouple Systems was Ltd inserted(Egham, toUK) achieve thermomechanical close temperature compression control machine during was hot compressions.used to test the samples using strain rate controlled compression. FAST-DB samples were tested at 1 temperaturesA Servotest of 850,Testing 900, Systems and 950 ◦ LtdC at (Egham, constant UK strain) thermomechanical rates of 0.1, 1, and compression 10 s− . Each machine sample was used to test the samples using strain rate controlled compression. FAST-DB samples were tested at temperatures of 850, 900, and 950 °C at constant strain rates of 0.1, 1, and 10 s−1. Each sample was sprayed with boron nitride lubricant on the upper and lower surfaces contacting the dies to stop adhesion and reduce friction. The test regime included heating to the test temperature at 4 °C/s, using a fast-thermal treatment unit (FTTU), followed by a dwell period of 2 min to ensure temperature homogeneity throughout the sample. The sample was moved into the furnace using robot controlled

Metals 2019, 9, x REVISED 31 May 2019 5 of 19 Metals 2019, 9, 654 5 of 19 gripper arms, compressed and quenched with water to suppress any further changes in microstructure. During testing a data logger was used to record time, temperature, load, sprayed with boron nitride lubricant on the upper and lower surfaces contacting the dies to stop displacement, and platen velocity. The samples were compressed to half height to achieve a nominal adhesion and reduce friction. The test regime included heating to the test temperature at 4 C/s, strain of 0.69. ◦ using a fast-thermal treatment unit (FTTU), followed by a dwell period of 2 min to ensure temperature The true stress data was friction corrected, by applying a shear friction factor m̅ to account for homogeneity throughout the sample. The sample was moved into the furnace using robot controlled the characteristic barreling of the TMC sample, and adjusted for machine compliance during gripper arms, compressed and quenched with water to suppress any further changes in microstructure. compression. An origin correction was also applied to account for elastic deformation when During testing a data logger was used to record time, temperature, load, displacement, and platen necessary. The rheology data obtained from constituent Ti-6-4 and CP-Ti compression tests were used velocity. The samples were compressed to half height to achieve a nominal strain of ~0.69. in the FE material models. The data was inputted using a tabular format at three temperatures and The true stress data was friction corrected, by applying a shear friction factor m to account for the strain rates, whereby the FE software interpolates between the data points. characteristic barreling of the TMC sample, and adjusted for machine compliance during compression. An origin correction was also applied to account for elastic deformation when necessary. The rheology 2.4. Near-Net Shape, Closed-Die Forging data obtained from constituent Ti-6-4 and CP-Ti compression tests were used in the FE material models. The dataForging was inputtedpreforms using of dimensions a tabular format 22 mm at diameter three temperatures by 160 mm and length strain were rates, electro whereby-discharge the FE softwaremachined interpolates from the 250 between mm diameter, the data 35 points. mm thick FAST billet. The forging preforms were extracted from the FAST billet with the diffusion bond running along the preform length at the centreline. The 2.4.eye- Near-Netbolt demonstrator Shape, Closed-Die components Forging were closed-die forged at W.H. Tildesley Ltd., Wolverhampton, UK usingForging a Massey preforms 1.1 ofMSC dimensions drop hammer 22 mm forge diameter with foot by pedal 160 mm control length and were 11 kJ electro-discharge blow energy, as machinedshown in Figure from the 4a. 250 The mm semi diameter,-complex 35 eye mm-bolt thick geometry FAST billet. was Theselected forging in order preforms to generate were extracted a large fromrange theof strains FAST billetacross with the thebond di ffinterface.usion bond Figure running 4b shows along a photograph the preform of length a machined at the centreline.FAST-DB Thepreform eye-bolt laid demonstrator on top of the componentseye-bolt lower were die closed-die impression. forged During at W.H. the forging Tildesley operatio Ltd., Wolverhampton,ns, the stamper UKwas usinginstructed a Massey to use 1.1 as MSCmany drophammer hammer blows forge as required with foot to fill pedal the controldies. The and No. 11 5 kJElectem, blow energy,Ni-Cr- asMo shown die steel in Figuredies are4a. constantly The semi-complex gas heated eye-bolt between geometry 80 and was90 °C selected throughout in order the to forging generate process a large to rangereduce of the strains die chilling across effect. the bond interface. Figure4b shows a photograph of a machined FAST-DB preformA single laid onwo toprkpiece of the of eye-bolt identical lower preform die impression.geometry was During used as the a forgingsacrificial operations, part, which the remained stamper wasin the instructed gas furnace to use with as manytwo thermocouples hammer blows attached as required to control to fill the the dies. temperature The No. 5of Electem, the furnace Ni-Cr-Mo prior to forging. One thermocouple was inserted 10 mm into the end of the sample and the second in the die steel dies are constantly gas heated between 80 and 90 ◦C throughout the forging process to reduce themiddle die chillingof the workpiece effect. length. A 20 min soak time allowed the preform to achieve a homogeneous temperature.

Figure 4.4. (a) Photograph of the Massey 1.1 MSC drop hammer forge forge used used in in forging forging trials trials at at W.H. W.H. TildesleyTildesley Ltd;Ltd; ((bb)) PhotographPhotograph ofof thethe eye-bolteye-bolt geometrygeometry withwith aa representativerepresentative preformpreform resting resting on on top. top. A single workpiece of identical preform geometry was used as a sacriﬁcial part, which remained 2.5. Analysis Techniques in the gas furnace with two thermocouples attached to control the temperature of the furnace prior to forging.Each One compressed thermocouple sample was was inserted sectioned 10 mm in intohalf, the perpendicular end of the sample to the andbond the interface. second in The the samples middle ofwere the workpiecemounted in length. conductive a 20 min Bakelite soak time and allowed a typical the titanium preform metallographic to achieve a homogeneous preparation temperature. procedure was used with a 20-minute colloidal silica polishing stage due to the presence of the soft CP-Ti alloy. Prior to imaging with a camera, the macrographs were etched with Kroll’s reagent consisting of 2% HF, 6% HNO3, and 92% H2O for between 10 and 15 s. Macrographs of the uni-axially forged thermomechanial compression (TMC) samples were captured using a Nikon D7000 camera (Tokyo,

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2.5. Analysis Techniques Each compressed sample was sectioned in half, perpendicular to the bond interface. The samples were mounted in conductive Bakelite and a typical titanium metallographic preparation procedure was used with a 20-minute colloidal silica polishing stage due to the presence of the soft CP-Ti alloy. Prior to imaging with a camera, the macrographs were etched with Kroll’s reagent consisting of 2% HF, 6% HNO3, and 92% H2O for between 10 and 15 s. Macrographs of the uni-axially forged thermomechanial compression (TMC) samples were captured using a Nikon D7000 camera (Tokyo, Japan)attached with a 105 mm Sigma DG Macro HyperSonic Motor lens (Kawasaki, Japan). Optical light microscopy was carried out using a Nikon Eclipse LV150 microscope with cross-polarised light to assess the bond microstructure.

2.6. Finite Element Modelling Validation DEFORM 3D v11.2 [21] FE software was used to model the hot compression samples and the closed-die forging process. Isothermal compression simulations of FAST-DB Ti-6-4 with CP-Ti samples were used to predict the alloy distribution and strain profiles throughout the sample. Each sample was unique, containing varying fraction volumes of Ti-6-4 and CP-Ti alloys. The compression samples were modelled using two workpieces, designed in SolidWorks [22], and imported into DEFORM as an STL file. Each workpiece was meshed separately with a minimum element size of 0.3 mm, approximately 88,000 elements for both workpieces combined, and a fine internal mesh. Mesh windows were defined at the interface between the workpieces to capture the local deformation. Each workpiece was assigned the corresponding material data, and a sticking contact with a FEM DEFORM shear friction factor, m , of 20 was applied between the workpieces replicating monolithic workpiece behaviour. The friction factor of 20 is a unique number within DEFORM which eliminates any element sliding occurring. During the simulation, forced remeshing steps at linear displacement intervals were used to prevent severe element deformation from occurring, which would decrease the simulation accuracy. The upper platen movement was set up to provide a constant global strain rate and was stopped at the final sample height. The bottom platen remained stationary throughout the simulation. For consistency and repeatability, an identical variable friction factor, was applied between the workpieces and the platens. The variable friction factor started at 0.2, with a linear increase up to 0.5 at 0.55 of the total test time, and a final further linear increase to 0.9 until the end of the test. Compared to the compression tests, the closed-die forging is a relatively uncontrolled multi-step process, which is a challenge to capture using FE models. The multiple operations (MO) processor was used to model the process, consisting of; the workpieces cooling in air, cooling on dies, and multiple forging operations in the starting and finishing die impressions. Using the MO processor allows the transfer of the workpiece and dies from the previous operation. This is a more efficient alternative to multiple separate operations, which increases the complexity and time required for simulating the multi-step forging process. Again, as with the FE modelling of the compression testing, two separate workpieces were defined in the workpiece. a mesh was applied of approximately 250,000 elements, and mesh windows were applied at the interface between the two workpieces. a sticking condition was defined at the bond interface, so the two alloys deformed as a singular workpiece and did not separate during deformation. a friction factor of 0.3 was applied as the contact condition between the workpieces at the top and bottom dies. To prevent excessive element deformation, a forced remeshing step was set after every 10 steps to improve the reliability of the simulation. Additionally, to prevent high local temperatures a truncation temperature of 1500 ◦C for each workpiece was applied, which limits individual element temperature to 1500 ◦C. The hammer efficiency was set to a constant 90% for all hammer steps, therefore 9.9 kJ of energy is transformed into plastic deformation. The workpiece absorbs the energy of each forging step until the input energy becomes 0 kJ, at which point the next step proceeds after a 1 s gap between the previous and consecutive step. The simulation was complete when the final step energy was expended and the height of the simulated workpiece was within 5% of the actual workpiece height. Metals 2019, 9, 654 7 of 19

3. Results and Discussion

3.1. Determination of Flow Curves The two alloys exhibit differing flow stress behaviour at the temperature range between 800 and 1050 ◦C. This is the hot working range that the FAST-DB dissimilar titanium alloy preforms were closed-die forged at W.H. Tildesley, so characterising the individual alloys is key for the FE simulation Metals 2019, 9, x REVISED 31 May 2019 7 of 19 of the near-net shape component. TheThe lower strength strength CP CP-Ti-Ti exhibits exhibits work work hardening hardening up to up approximately to approximately 0.2 strain 0.2 prior strain to steady prior- to steady-statestate flow, seen flow, in seen Figure in Figure 5a–c. 5Thea–c. proportion The proportion of work of work hardening hardening to steady to steady-state-state flow flow is greatest is greatest at atthe the lowest lowest temperature temperature and and fastest fastest strain strain rat rate.e. Conversely, Conversely, the the flow flow stress stress at at 950 950 °C◦C and and above above the the ββ transustransus forfor CP-TiCP-Ti (which(which isis approx.approx. 910 °C)◦C) exhibits exhibits steady steady-state-state flow flow with with no no work work hardening. hardening.

Figure 5. Friction corrected flow stress curves for the full matrix of temperatures of CP-Ti TMC tests at; Figure 5. Friction corrected flow stress curves for the full matrix of temperatures of CP-Ti TMC tests (a) 0.1 s 1;(b) 1 s 1;(c) 10 s 1. Ti-6-4 TMC tests at; (d) 0.1 s 1;(e) 1 s 1; and (f) 10 s 1. at; (a) 0.1− s−1; (b) 1− s−1; (c) 10− s−1. Ti-6-4 TMC tests at; (d) 0.1 s−1; (e) 1 s−1−; and (f) 10 s−1.− The higher strength Ti-6-4 alloy exhibits flow stress behaviour which is more than twice the The higher strength Ti-6-4 alloy exhibits flow stress behaviour which is more than twice the magnitude of the CP-Ti at the same temperature range, see Figure5d–f. However, the low-strain magnitude of the CP-Ti at the same temperature range, see Figure 5d–f. However, the low-strain flow flow behaviour is different for Ti-6-4: After rapid work hardening up to a yield point flow softening, behaviour is different for Ti-6-4: After rapid work hardening up to a yield point flow softening, due due to the break-up of fine-scale plate-like secondary alpha, is observed particularly at the lowest test to the break-up of fine-scale plate-like secondary alpha, is observed particularly at the lowest test temperature of 850 C. The degree of flow softening reduces with increasing temperature and at the temperature of 850 ◦°C. The degree of flow softening reduces with increasing temperature and at the super transus test temperature of 1050 C, steady-state flow is exhibited. super transus test temperature of 1050 ◦°C, steady-state flow is exhibited. Therefore, in summary, the flow behaviour of the two alloys are similar at higher temperatures, Therefore, in summary, the flow behaviour of the two alloys are similar at higher temperatures, as both alloys exhibit steady-state flow. However, there is a mismatch in the magnitude of resistance to as both alloys exhibit steady-state flow. However, there is a mismatch in the magnitude of resistance deformation: At a test temperature of 950 C and strain rate of 1.0 s 1, the steady-state flow stress in to deformation: At a test temperature of 950◦ °C and strain rate of 1.0− s−1, the steady-state flow stress CP-Ti and Ti-6-4 is 20 MPa and 50 MPa, respectively. At the lower end of the hot working range, there is in CP-Ti and Ti-6-4 is 20 MPa and 50 MPa, respectively. At the lower end of the hot working range, greaterthere is contrast greater incontrast the flow in behaviour,the flow behaviour, particularly particularly at low strains: at low The strains: CP-Ti The exhibits CP-Ti work exhibits hardening, work whereashardening, the whereas Ti-6-4 displays the Ti-6 flow-4 displays softening. flow Forgers softening. prefer Forgers to work prefer with to materialwork with that material has steady-state that has flow,steady i.e.-state equivalent flow, i.e. resistance equivalent to deformationresistance to atdeformation low and high at low strains, and inhigh order strains, to control in order die to filling control and ensuredie filling homogeneous and ensure strainhomogeneous distributions. strain On distributions. the basis of On the the flow basis data of in the Figure flow5 ,data there in potentially Figure 5, couldthere potenti be a largeally mismatch could be a inlarge flow mismatch stress behaviour, in flow stress particularly behaviour, at lowparticularly strains, whenat low couplingstrains, when CP-Ti andcoupling Ti-6-4 CP in- aTi dissimilar and Ti-6-4 FAST in a dissimilar preform. FAST preform.

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3.2. Load-Displacement Curve Validation Selected friction corrected load-displacement data representing the full range of loads comparing the FE simulation and hot compressed FAST CP-Ti and Ti-6-4 constituent samples are shown in Figure6a,b, respectively. All tests showed an initial increase in load, followed by a linear increase until a point where the friction has more prominent influence on the load due to an increase in surface area, and a breakdown of the boron nitride lubricant between the platens and the workpiece. 1 This frictional effect was most prominent in the 10 s− tests where after approximately 6 mm of displacement,Metals 2019, 9, x REVISED there is 31 an May exponential 2019 increase in load, suggesting the friction is generally higher8 of in19 Metals 2019, 9, x REVISED 31 May 2019 8 of 19 these tests. There is excellent agreement between the experimental and simulation load-displacement 1 −1 data,displacementdisplacement although data, data, a slight although although underestimation a a sl slightight underestimation underestimation was observed was was for observed testsobserved at 10 for for s− tests testsstrain at at 10 10 rates. s s− 1strain strain This rates. providesrates. This This confidenceprovidesprovides confidence confidencein the FE model in in the the as FE aFE reliable model model method as as a a reliable reliable for modelling method method simple for for modelling modelling uni-axial simple hot simple compressions uni uni-axial-axial hot ofhot FASTcompressionscompressions material of derivedof FAST FAST frommaterial material powder. derived derived from from powder. powder.

FigureFigureFigure 6. 6. 6. FrictionFrictionFriction corrected corrected corrected load load dispdisplacement displacementlacement curves curves comparing comparing experimental experimental TMC TMC and and and FE FE FE simulationssimulationssimulations for; for; ( (a(aa)) )CP CP-TiCP-Ti-Ti alloy alloy material material material post post post FAST FAST FAST consolidation; consolidation; consolidation; and and and ( b(b) ) (Ti bTi)-6- Ti-6-46-4-4 alloy alloy alloy material material material post post postFASTFAST FAST consolidation. consolidation. consolidation.

Figure7 demonstrates the friction corrected load-displacement responses of FAST-DB samples, FigureFigure 7 7 demonstrates demonstrates the the friction friction corrected corrected load load-displacement-displacement responses responses of of FAST FAST-DB-DB samples, samples, asasas compared compared compared to to theirto their their corresponding corresponding corresponding single single single constituent constituent constituent FAST FAST processedFAST processed processed CP-Ti and CP CP- Ti-6-4Ti-Ti and and alloy Ti Ti-6- samples6-4-4 alloy alloy with a nominal test temperature of 950 C. Both tests with strain rates of 0.1 and 10 s 1 show−1 that samplessamples with with a a nominal nominal test test temperature temperature◦ of of 950 950 °C. °C. Both Both tests tests with with strain strain rates rates of of 0.1 0.1 and and− 10 10 s s− 1show show Ti-6-4thatthat Ti hasTi-6-6-4 the-4 has has higher the the higher ﬂowhigher stress flow flow compared stress stress compared compared to the CP-Ti to to t het alloy.he CP CP The-Ti-Ti alloy. FAST-DBalloy. The The dissimilarFAST FAST-DB-DB alloydissimilar dissimilar samples alloy alloy lie betweensamplessamples thelie lie between single between alloy the the Ti-6-4 single single and alloy alloy CP-Ti Ti Ti-6 load-6-4-4 and and responses, CP CP-Ti-Ti load load suggesting responses, responses, the loadssuggesting suggesting required the the for loads loads deformation required required couldforfor deformation deformation be estimated could could to determine be be estimated estimated the to necessary to determine determine machine the the necessary necessary capacities. machine machine capacities. capacities.

FigureFigure 7. 7. Friction corrected corrected load-displacement load-displacement graphs graphs comparing comparing single single alloy alloy compressions compressions with with Figure 7. Friction corrected load-displacement graphs comparing single alloy compressions with1 FAST-DBFAST-DB dissimilar dissimilar alloy alloy samples samples at at a testa test temperature temperature of 950of 950◦C °C and and a constant a constant strain strain rate rate of; ( aof;) 0.1 (a) s −0.1; FAST-DB dissimilar1 alloy samples at a test temperature of 950 °C and a constant strain rate of; (a) 0.1 and−1 (b) 10 s− . −1 ss−;1 ;and and ( b(b) )10 10 s s−.1 .

3.3.3.3. Small Small-Scale-Scale Hot Hot Compression Compression Forgings Forgings PriorPrior to to simulating simulating FAST FAST-DB-DB uni uni-axial-axial hot hot compressions, compressions, initial initial simulations simulations with with varying varying frictionfriction conditions conditions and and alloy alloy volume volume fraction fraction ratios ratios in in the the cylindrical cylindrical TMC TMC geometry geometry were were modelled, modelled, asas shown shown in in Figure Figure 8. 8. WhenWhen the the alloy alloy volume volume fraction fraction is is 1:1 1:1 with with z zeroero friction friction the the geometry geometry remains remains rectangular; rectangular; no no barrellingbarrelling is is observed observed and and the the bond bond interface interface remains remains straight, straight, see see Figure Figure 8a. 8a. As As no no friction friction is is present present thethe sample sample deforms deforms uniformly uniformly resulting resulting in in a a constant constant strain strain profile profile across across the the sample. sample. WhenWhen the the samp samplele is is a a singular singular alloy alloy and and with with the the addition addition of of a a constant constant friction friction factor, factor, barrelling barrelling occurs,occurs, as as shown shown in in Figure Figure 8b. 8b. At At the the top top and and bottom bottom of of the the compression compression sample, sample, low low strain strain regions regions existexist that that have have little little or or no no deformation. deformation. At At the the corners corners are are local local hig highh-strain-strain regions, regions, while while the the curved curved

Metals 2019, 9, 654 9 of 19

3.3. Small-Scale Hot Compression Forgings Prior to simulating FAST-DB uni-axial hot compressions, initial simulations with varying friction conditions and alloy volume fraction ratios in the cylindrical TMC geometry were modelled, as shown in FigureMetals8. 2019, 9, x REVISED 31 May 2019 10 of 19

Figure 8.FigureUni-axial 8. Uni-axial compression compression FE FE simulations simulations of of FAST FAST-DB-DB samples, samples, showing showing the bond the interface bond interface −1 1 locationlocation with corresponding with corresponding strain strain distribution. distribution. Simulation Simulation conditions conditions are are a strain a strain rate of rate 1 s of, 850 1 s °C,− , 850 ◦C, and a progressiveand a progressive friction friction factor factor of of 0.2-0.5-0.9. 0.2-0.5-0.9. ((a) 1:1 alloy alloy distribution distribution with with zero zerofriction; friction; (b) single (b ) single alloy with friction; (c) 1:1 alloy distribution with friction; (d) 1:1 alloy distribution with friction at alloy with friction; (c) 1:1 alloy distribution with friction; (d) 1:1 alloy distribution with friction at quarter deformation height; (e) dominant CP-Ti with friction; and (f) dominant Ti-6-4 with friction. quarter deformation height; (e) dominant CP-Ti with friction; and (f) dominant Ti-6-4 with friction. Figure 9 shows a comparison between the FE simulations and macrographs of the sectioned Whencompressed the alloy cylinders, volume deformed fraction at representative is 1:1 with nominal zero friction test conditions. the geometry The macrographs remains of these rectangular; no barrellingcylinders is demonstrate observedand the effect thebond of differing interface friction remains factors and straight, the influence see Figure of the8 a.bond As interface no friction is presentmorphology the sample. deforms Post uni-axial uniformly hot compression resulting and in a after constant etching strain with profile Kroll’s acrossreagent the the sample. visually Whenbrighter the alloy sample is the Ti is-6 a-4, singularwhile the brown alloy tinted and alloy with is thethe CP addition-Ti. With the of exception a constant of the friction test at factor, barrelling occurs, as shown in Figure8b. At the top and bottom of the compression sample, low strain Metals 2019, 9, 654 10 of 19 regions exist that have little or no deformation. At the corners are local high-strain regions, while the curved edges remain low-strain. Towards the centre of the sample the strain increases where the maximum strain is located. a symmetrical strain profile is observed and strain continuation across the interface is excellent despite the deformation complexity. The strain profile becomes more complex and loses symmetry with the addition of a second alloy and with friction still present, as shown in Figure8c. The bond interface bulges into the CP-Ti alloy due to the comparably lower flow stresses at the hot compression temperature compared to the Ti-6-4 alloy. The highest strain region seen in the centre shifts slightly to the right into the higher flow stress alloy distorting the strain profile. Within the CP-Ti alloy the highest strain does not exceed 1.3, as a consequence of the Ti-6-4 displacing the CP-Ti alloy. When the deformation is increased further to quarter height, the strain profile complexity increases significantly, see Figure8d. The strain profiles of the two alloys begin to behave as individual constituent samples and the FE simulated bond line is under high strain. The Ti-6-4 alloy behaves similarly to Figure8c, however, the maximum strain increases from 1 to 2.5. On the other hand, the CP-Ti alloy behaves unexpectedly with the highest strain occurring at the top and bottom surfaces, while in the centre of the sample the strain is only in the region of 1.6 to 1.9. Along the horizontal centreline is a high strain region between 1.9 and 2.5, where it is expected that dynamic recrystallisation processes are most likely to occur. If the Ti-6-4 to CP-Ti alloy ratio is increased to 1:3, see Figure8e, the bulging direction occurs in the Ti-6-4 alloy but to a greater extent compared to the 1:1 alloy volume fraction distribution condition. The highest strain remains in the centre of the sample but the continuation of constant strain across the interface does not occur. Instead there is a strain mismatch between the two alloys where the lower strength CP-Ti alloy accommodates a higher strain compared to the Ti-6-4 alloy. As the Ti-6-4 to CP-Ti alloy ratio reverses to 3:1, see Figure8f, the bulging becomes more dominant in the CP-Ti alloy, which has a lower resistance to deformation at the forging temperature compared to Ti-6-4. As a result, the bond interface bulging becomes broader into the CP-Ti during the hot compression. The highest strain remains in the centre of the sample and covers a slightly larger area (compared to Figure8e), demonstrating the higher flow stress of the Ti-6-4 alloy. Again, at the interface there is a significant strain mismatch: The CP-Ti alloy has a strain between 0.75 and 0.94 along the majority of the interface, while the Ti-6-4 remains a similar strain only in the centre of the sample at the interface. The bond interface evolution demonstrated in these models in Figure8 are also observed in the simulations with a diagonal starting bond interface morphology and the corresponding TMC samples. Figure9 shows a comparison between the FE simulations and macrographs of the sectioned compressed cylinders, deformed at representative nominal test conditions. The macrographs of these cylinders demonstrate the effect of differing friction factors and the influence of the bond interface morphology. Post uni-axial hot compression and after etching with Kroll’s reagent the visually brighter alloy is the Ti-6-4, while the brown tinted alloy is the CP-Ti. With the exception of the test at 850 ◦C 1 with a 0.1 s− strain rate, the CP-Ti volume percentage decreased from the top to the bottom surface. There is excellent agreement of the bond interface morphology when comparing the simulations to the compression cylinders. Particularly considering the assumption of a linear interface between the top and bottom surfaces prior to the hot deformation. The simulations also assume isothermal conditions and only contained friction corrected data, so the flow of the simulation is likely to deviate from the compression cylinder and, consequently, the bulging in the simulations varies slightly. Metals 2019, 9, x REVISED 31 May 2019 11 of 19

850 °C with a 0.1 s−1 strain rate, the CP-Ti volume percentage decreased from the top to the bottom surface. There is excellent agreement of the bond interface morphology when comparing the simulations to the compression cylinders. Particularly considering the assumption of a linear interface between the top and bottom surfaces prior to the hot deformation. The simulations also assume isothermal conditions and only contained friction corrected data, so the flow of the simulation Metals 2019, 9, 654 11 of 19 is likely to deviate from the compression cylinder and, consequently, the bulging in the simulations varies slightly.

Figure 9. Selected Ti-6-4 with CP-Ti FAST-DB TMC samples sectioned perpendicular to the bond Figure 9. Selected Ti-6-4 with CP-Ti FAST-DB TMC samples sectioned perpendicular to the bond interface after deformation at nominal test conditions as stated including effective strain profile, alloy interface after deformation at nominal test conditions as stated including effective strain profile, alloy profile, and macrograph of the sample. a progressive 0.2-0.5-0.9 progressive friction factor was used. profile, and macrograph of the sample. A progressive 0.2-0.5-0.9 progressive friction factor was used. 1 1 Cylinders tested under conditions of 850◦C at 0.1 s− and 900 ◦C at 10 s− demonstrate the Cylinders tested under conditions of 850°C at 0.1 s−1 and 900 °C at 10 s−1 demonstrate the dependence of the alloy distribution. Although Ti-6-4 has a higher flow stress at the test temperature, dependence of the alloy distribution. Although Ti-6-4 has a higher flow stress at the test temperature, the CP-Ti displaces the Ti-6-4 due to the dominance of the CP-Ti alloy in the sample. The cylinder tested the CP-Ti displaces the Ti-6-4 due1 to the dominance of the CP-Ti alloy in the sample. The cylinder under conditions of 950 ◦C at 1 s− has a greater percentage of Ti-6-4 and, consequently, displaces the tested under conditions of 950 °C at 1 s−1 has a greater percentage of Ti-6-4 and, consequently, CP-Ti alloy during the upset forging. The post deformation bond interface morphologies are likely displaces the CP-Ti alloy during the upset forging. The post deformation bond interface formed due to a combination of prior bond interface morphologies, alloy distribution from top to morphologies are likely formed due to a combination of prior bond interface morphologies, alloy bottom, and the hot deformation conditions. These findings demonstrate the importance of controlling distribution from top to bottom, and the hot deformation conditions. These findings demonstrate the the bond interface prior to hot working. It is envisaged that there will be greater control through importance of controlling the bond interface prior to hot working. It is envisaged that there will be process development which will occur as FAST-DB development elevates through the technology greater control through process development which will occur as FAST-DB development elevates readiness levels. through the technology readiness levels. 3.4. Large-Scale, Near-Net Shape Forging 3.4. Large-Scale, Near-Net Shape Forging Two eye-bolts were sectioned and analysed during this study: The first being eye-bolt 1 (EB1) from a FAST-DBTwo eye preform-bolts were of Ti-6-4 sectioned and CP-Ti. and analysed Eye bolt during 2 (EB2) this was study: processed The fromfirst being a FAST-DB eye-bolt preform 1 (EB1) of fromTi-6-4 a and FAST Ti-5553,-DB preform shown in of Figures Ti-6-4 10 anda and CP 11-Ti.a, Eye respectively. bolt 2 (EB2) Visually, was both processed eye-bolts from show a FAST excellent-DB preformbond integrity of Ti-6 with-4 and no Ti cracking-5553, shown at the bondin Figure interface. 10a and In EB1,Figure the 11a, two respectively. alloys can be Visually, distinguished both fromeye- boltseach show other excellent by the di ffbonderences integrity in amount with ofno flashcracking at the at topthe andbond bottom interface. of the In EB1, forging, the two see Figurealloys can10a. beAt distinguished the forging temperature, from each other the CP-Tiby the alloydifferences has a lowerin amount flow of stress flash comparedat the top and to the bottom Ti-6-4 of alloy the forging,and thus, see the Figure former 10a. alloy At the creates forging more temperature flash for, giventhe CP height-Ti alloy reduction. has a lower The flow die stress fill was compared mostly tocomplete; the Ti-6 however,-4 alloy and the thus, material the didformer not completelyalloy creates fill more the Ti-6-4 flash areafor given on the height inside reduction. of the thinner The “eye” die fillsection. was mostly In addition, complete; the bondhowever, line hasthe material been revealed did not due comp to theletely diff fillerent the oxide Ti-6-4 scale area colourson the inside of the two alloys on the surface of the component. It provides an initial estimation of the bond interface region, which is particularly valuable as it saves time and resources which would be beneficial when evaluating new components. Metals 2019, 9, x REVISED 31 May 2019 12 of 19 Metals 2019, 9, x REVISED 31 May 2019 12 of 19 of the thinner “eye” section. In addition, the bond line has been revealed due to the different oxide scaleof the colours thinner of “eye” the two section. alloys In oaddition,n the surface the bond of the line component. has been revealedIt provides due an to initial the different estimation oxide of thescale bond colours interface of the region, two alloys which o nis theparticularly surface of valuable the component. as it saves It timeprovides and resourcesan initial whichestimation would of bethe beneficial bond interface when region,evaluating which new is components.particularly valuable as it saves time and resources which would be beneficialIn EB2, thewhen alloys evaluating can be determined new components. by the differences in die fill. The Ti-6-4 alloy filled the die similarlyIn EB2, to thatthe alloysobserved can inbe EB1. determined Conversely, by the the differences Ti-5553 alloy in die demonstrated fill. The Ti- 6limited-4 alloy die filled fill theon thedie insidesimilarly of theto that “eye” observed section, in due EB1. to theConversely, higher flow the stressTi-5553 of alloythe high demonstrated strength metastable limited die beta fill alloy on the at theinside forging of the temperature. “eye” section, due to the higher flow stress of the high strength metastable beta alloy at the forgingThe sectioning temperature. strategy for the two eye bolts was selected in order to validate the FE model bond interfaceThe sectioningpredictions strategy and corresponding for the two eye strain bolts distribution was selected in in a orderhigh and to validate low strain the FEregion. model In bondEB1, theinterface low predictions and high strainand corresponding cross-sections strain were distribution located approximately in a high and 22.7low strain mm andregion. 75.0 In EB1, mm, respectivelythe low and, from high the strain reference cross end-sections of the eye were-bolt located and are approximately discussed further 22.7 in mmSections and 3.5 75.0 and mm, 3.6, respectively, from the reference end of the eye-bolt and are discussed further in Sections 3.5 and 3.6, seeMetals Figure2019, 9 10b., 654 In EB2, the high strain cross section was located at approximately 56.4 mm from12 ofthe 19 referencesee Figure end, 10b. see In EB2,Figure the 11b. high strain cross section was located at approximately 56.4 mm from the reference end, see Figure 11b.

Figure 10. (a) Photograph of the Ti-6-4 with CP-Ti near-net shape forging, eye-bolt 1; (b) Figure 10. (a) Photograph of the Ti-6-4 with CP-Ti near-net shape forging, eye-bolt 1; (b) corresponding cFigureorresponding 10. (a) FE Photograph model showing of the the alloy Ti-6- 4distribution; with CP-Ti and near (c) -thenet corresponding shape forging, FE eyemodel-bolt showing 1; (b) FE model showing the alloy distribution; and (c) the corresponding FE model showing the eﬀective thecorresponding effective strain FE model distribution. showing Approximate the alloy distribution; forging temperature and (c) the c orrespondingbetween 860 and FE model 878 °C, showing with a strain distribution. Approximate forging temperature between 860 and 878 C, with a total of nine totalthe effective of nine hammerstrain distribution. steps. Approximate forging temperature between ◦860 and 878 °C, with a hammertotal of nine steps. hammer steps.

Figure 11. 11.( a()a Photograph) Photograph of the of Ti-6-4 the Ti with-6-4 Ti-5553 with near-net Ti-5553 shape near-net forging, shape eye-bolt forging, 2; (b ) eye corresponding-bolt 2; (b) c cFEFigureorresponding model 11. showing (a) FEPhotograph model the alloy showing distribution; of the the Tialloy-6-4 and distribution; with ( ) the Ti-5553 corresponding and near (c) the-net c orresponding FEshape model forging, showing FE model eye the-bolt showing eﬀ ective 2; (b) straincorresponding distribution. FE model Approximate showing the forging alloy temperature distribution; betweenand (c) the 860 corresponding and 878 ◦C, with FE model a total showing of nine hammer steps.

In EB2, the alloys can be determined by the differences in die fill. The Ti-6-4 alloy filled the die similarly to that observed in EB1. Conversely, the Ti-5553 alloy demonstrated limited die fill on the inside of the “eye” section, due to the higher flow stress of the high strength metastable beta alloy at the forging temperature. The sectioning strategy for the two eye bolts was selected in order to validate the FE model bond interface predictions and corresponding strain distribution in a high and low strain region. In EB1, the low and high strain cross-sections were located approximately 22.7 mm and 75.0 mm, respectively, from the reference end of the eye-bolt and are discussed further in Sections 3.5 and 3.6, see Figure 10b. In EB2, the high strain cross section was located at approximately 56.4 mm from the reference end, see Figure 11b. Metals 2019, 9, x REVISED 31 May 2019 13 of 19

the effective strain distribution. Approximate forging temperature between 860 and 878 °C, with a total of nine hammer steps.

3.5. Comparison of FE Model with Near Net Shape Forging Metals 2019, 9, x REVISED 31 May 2019 13 of 19 Figure 12 shows the evolution of the FE model and the shape of EB2 after each individual closed- Metals 2019, 9, 654 13 of 19 die forgingthe effective step. Itst rainprovides distribution. a demonstration Approximate of forging the robust temperature FE modelling between software 860 and for878 the °C, simulationwith a of a relativelytotal of nine uncontrolled hammer steps. forging route for FAST-DB preforms. 3.5. Comparison of FE Model with Near Net Shape Forging 3.5. Comparison of FE Model with Near Net Shape Forging Figure 12 shows the evolution of the FE model and the shape of EB2 after each individual closed-dieFigure forging 12 shows step. the evolution It provides of athe demonstration FE model and ofthe the shape robust of EB2 FE after modelling each individual software forclosed the- simulationdie forging of step. a relatively It provides uncontrolled a demonstration forging of route the robust for FAST-DB FE modelling preforms. software for the simulation of a relatively uncontrolled forging route for FAST-DB preforms.

Figure 12. FE model and corresponding photographs of the eye-bolt 2 (Ti-5553 with Ti-6-4) shape evolution after each forging step.

Figure 13 compares the FE model to a low strain region location shown in Figure 10b of EB1. The first observation that can be made is that the prediction of the bond interface morphology seen in Figure 13b is in good agreement with the macrographs, considering the complex, multi-step forging. FigureFigure 12.12. FE modelmodel andand correspondingcorresponding photographs of the eye-bolteye-bolt 2 (Ti-5553(Ti-5553 withwith Ti-6-4)Ti-6-4) shapeshape There is slightly more curvature of the bond interface in the macro image and the CP-Ti material has evolutionevolution afterafter eacheach forgingforging step.step. spread further along the upper and lower contact surfaces during the hot working process. This is likelyFigureFigure to be a 13 result compares of the thecontinually the FE FE model model varying to to a alow low friction strain strain conditionsregion region location location between shown shown the in workpiece Figure in Figure 10b and 10of bEB1.the of dies, EB1.The Thecomparedfirstﬁrst observation observation to a constant that that can friction can be bemade factor made is used isthat that thein the the prediction prediction simulation. of of the the bond bond interface interface morphology morphology seen inin FigureFigureFigure 1313bb is is13a inin goodgoodshows agreementagreement the macrostructures withwith the macrographs, evolution across considering the low the strain complex, region multi multi-step cross-step-section forging. in ThereEB1.There Across isis slightlyslightly the morewholemore curvaturecurvature bond interface, ofof thethe bondbondsimilar interfaceinterface macrostructural inin thethe macromacro flow imageimage characteristics andand thethe CP-TiCP are-Ti materialobservedmaterial hashas in bothspreadspread alloys. furtherfurther In general alongalong thethe, the upperupper CP-Ti andand alloy lowerlower microstructure contactcontact surfacessurface has beens duringduring deformed thethe hothot significantly, workingworking process.process. while ThistheThis Ti isis- 6likelylikely-4 alloy toto beisbe only aa resultresult slightly ofof thethe deformed. continually This varying is confirmed friction in conditions the predicted betweenbetween strain thethe profile workpiece of the cro andandss the-section dies, showncomparedcompared in Figure toto aa constantconstant 13c. This frictionfriction is an factor factorexample usedused where inin thethe there simulation.simulation. is a clear mismatch in flow behaviour at the interface,Figure yet 13a there shows is no theevidence macrostructures of delamination evolutio or cran acrosscking atthe the low bond strain line. region cross-section in EB1. Across the whole bond interface, similar macrostructural flow characteristics are observed in both alloys. In general, the CP-Ti alloy microstructure has been deformed significantly, while the Ti- 6-4 alloy is only slightly deformed. This is confirmed in the predicted strain profile of the cross-section shown in Figure 13c. This is an example where there is a clear mismatch in flow behaviour at the interface, yet there is no evidence of delamination or cracking at the bond line.

Figure 13. (a) Photograph of the low strain cross cross-section-section of of the the eye eye-bolt-bolt 1 1 forging forging (Ti (Ti-6-4-6-4 and CP CP-Ti)-Ti) from Figure Figure 10 b; 10b; (b ) corresponding(b) corresponding simulation simulation showing showing the alloy thedistribution; alloy distribution; and (c) the corresponding and (c) the simulation showing the eﬀective strain distribution. Approximate forging temperature between 860 and

878 ◦C, with a total of nine hammer steps.

Figure 13a shows the macrostructures evolution across the low strain region cross-section in EB1. AcrossFigure the whole13. (a) Photograph bond interface, of the similarlow strain macrostructural cross-section of ﬂowthe eye characteristics-bolt 1 forging are(Ti-6 observed-4 and CP- inTi) both alloys.from In general, Figure the 10b; CP-Ti (b) c alloyorresponding microstructure simulation has been showing deformed the alloy signiﬁcantly, distribution; while and the ( Ti-6-4c) the alloy

Metals 2019, 9, x REVISED 31 May 2019 14 of 19

corresponding simulation showing the effective strain distribution. Approximate forging temperature between 860 and 878 °C, with a total of nine hammer steps.

The high strain cross section region of EB1 is shown in Figure 14a, consisting of the Ti-6-4 alloy on the left hand side and the CP-Ti alloy on the right. The central “eye” region of the workpiece is quite complex including long banded regions of recrystallised Ti-6-4, approximately 25–50 µm wide. The Ti-6-4 alloy has flowed from the centre of the section and bulged into the CP-Ti side, once again, demonstrating the mismatch in flow stresses between the two alloys at the approximate forging temperature range of 860–878 °C. Upon further inspection, it appears that the CP-Ti alloy extends along the bottom surface to the centre of the forged shape. However, there is no evidence of the CP- TiMetals alloy2019 along, 9, 654 the top surface suggesting the Ti-6-4 alloy has displaced the alloy completely into14 ofthe 19 CP-Ti side. In the bulk Ti-6-4 and CP-Ti alloys, the grain structure is only slightly deformed due to limited is only slightly deformed. This is confirmed in the predicted strain profile of the cross-section shown material flow within the component. Towards the edges of these sections, and as the width of the in Figure 13c. This is an example where there is a clear mismatch in flow behaviour at the interface, component reduces, extremely fine grains are observed as a result of the high strains produced in yet there is no evidence of delamination or cracking at the bond line. these regions. This continues into the thin section, as can be seen in Figure 14a. The corresponding The high strain cross section region of EB1 is shown in Figure 14a, consisting of the Ti-6-4 alloy simulation alloy distribution, seen in Figure 14b, has very good agreement considering the on the left hand side and the CP-Ti alloy on the right. The central “eye” region of the workpiece is magnitude of the displaced CP-Ti alloy. Notably the simulation has captured the displacement of the quite complex including long banded regions of recrystallised Ti-6-4, approximately 25–50 µm wide. CP-Ti alloy and the displacement of the Ti-6-4 into the CP-Ti side. The predicted bulging form and The Ti-6-4 alloy has flowed from the centre of the section and bulged into the CP-Ti side, once again, curvature of the bond line deviates from the experimental forging and the thickness of the CP-Ti alloy demonstrating the mismatch in flow stresses between the two alloys at the approximate forging in the thin “eye” region is slightly overestimated. However, the model could be developed further to temperature range of 860–878 ◦C. Upon further inspection, it appears that the CP-Ti alloy extends along improve the accuracy by increasing the number of elements while also decreasing the element size. the bottom surface to the centre of the forged shape. However, there is no evidence of the CP-Ti alloy The varying friction between the workpiece and the dies is difficult to determine, so a constant along the top surface suggesting the Ti-6-4 alloy has displaced the alloy completely into the CP-Ti side. friction factor was used.

Figure 14. ((aa)) Photograph of of the the high high strain strain cross cross-section-section of of the the eye eye-bolt-bolt 1 1 forging forging from from Figure Figure 10b;10b; (b) ccorrespondingccorresponding simulation showing the alloy distribution; and (c) the correspondingcorresponding simulation showing the effective effective strain distribution. Approximate Approximate forging forging temperature temperature between between 8 86060 and 878 °C,◦C, with a total of nine hammer steps. with a total of nine hammer steps. In the bulk Ti-6-4 and CP-Ti alloys, the grain structure is only slightly deformed due to limited material flow within the component. Towards the edges of these sections, and as the width of the component reduces, extremely fine grains are observed as a result of the high strains produced in these regions. This continues into the thin section, as can be seen in Figure 14a. The corresponding simulation alloy distribution, seen in Figure 14b, has very good agreement considering the magnitude of the displaced CP-Ti alloy. Notably the simulation has captured the displacement of the CP-Ti alloy and the displacement of the Ti-6-4 into the CP-Ti side. The predicted bulging form and curvature of the bond line deviates from the experimental forging and the thickness of the CP-Ti alloy in the thin “eye” region is slightly overestimated. However, the model could be developed further to improve the accuracy by increasing the number of elements while also decreasing the element size. The varying friction between the workpiece and the dies is difficult to determine, so a constant friction factor was used. Metals 2019, 9, 654 15 of 19

3.6. High Strain Microstructural Analysis Microstructural analysis of high strain regions of EB1, which is located at 52.3 mm from the reference end of the forging, as shown in Figure 10b, was carried out to assess the microstructural evolution adjacent to the diffusion bond interface (see Figure 15). Site “a” is in the centre of the Ti-6-4 side and shows a typical low strained, lamellar Ti-6-4 microstructure, consisting of slightly elongated grains perpendicular to the compression axis compared to in the as-FAST condition. Additionally, within the Ti-6-4 side is site “b”, which is situated on the edge of the sample showing a crack propagation from the surface into the sample in a recrystallised region. The crack does not appear to be a geometrical defect, but rather a crack that may have formed during the high strain rate, multi-step forging. Site ‘c’ is also located in the Ti-6-4 side, but in the region transitioning into the thin “eye” cross-section of the forging. The recrystallised band region shown at site “b” also passes through site “c” and continues into the thin “eye” cross-section. Under cross-polarised light, the recrystallised region is highlighted as a lighter shade of grey compared to the region that starts from the bottom of the component, suggesting the grains may have been orientated in a preferential crystallographic direction during the forging process. Site “d” is a continuation of the recrystallised band which is observed throughout the thin “eye” cross-section. At site ‘e’, the first observation of the bond interface between Ti-6-4 and CP-Ti is shown at the bottom of the micrograph. The interface is difficult to distinguish due to the deformed microstructures, however, the Ti-6-4 remains as a typical low strain transformed β microstructure, while the CP-Ti is high strained. Site ‘f’ lies in the thin “eye” cross-section, as clearly shown by the severely deformed microstructures in both the Ti-6-4 and CP-Ti alloys. a fine-grain recrystallised region passes through the Ti-6-4 alloy to eventually make direct contact with the CP-Ti. At site “g” this contact is observed. The CP-Ti shows quite a fine grain structure in contact with the fine Ti-6-4 microstructure at the bond interface. Slightly further from the bond interface at approximately 200 µm into the Ti-6-4, the grain size becomes coarser and shows a highly deformed grain structure instead. Site “h”, at the end of the bulging Ti-6-4 region, shows a return to a constant moderate strain Ti-6-4 microstructure. The CP-Ti alloy grain size also increases slightly. Site “i” provides another demonstration of how the grain size can vary towards the bond interface. In this micrograph, the CP-Ti gradually decreases in grain size towards the bond interface while the Ti-6-4 remains fairly constant. This behaviour is similar to the suggested local recrystallisation that occurs at the highly-strained diffusion bond interface, when CP-Ti with Ti-6-4 compression samples are plastically deformed. Finally, site “j” shows the CP-Ti microstructure in the alloy section, consisting of an acicular microstructure. A high-strain region in EB2 was also microstructurally analysed at 56.4 mm from the reference end, as defined in Figure 11b. The varying microstructures across the eye-bolt, both in the side sections and also in the thin “eye” section, are shown in Figure 16. Metals 2019, 9, 654 16 of 19 Metals 2019, 9, x REVISED 31 May 2019 16 of 19

FigureFigure 15. 15. LightLight micrographs micrographs of ofregions regions of interest of interest in the in high the highstrain strain cross- cross-sectionsection of eye of-bolt eye-bolt 1 (Ti-6- 1 4(Ti-6-4 with CP with-Ti): CP-Ti): (a) Bulk (a) BulkTi-6-4 Ti-6-4 microstructure; microstructure; (b) e (xteriorb) exterior crack crack initiated initiated due dueto excessive to excessive strain; strain; (c) r(ecrystallisedc) recrystallised Ti-6 Ti-6-4-4 region; region; (d) ( da) comparison a comparison between between a medium a medium strain strain Ti Ti-6-4-6-4 and and high high strain strain at at the the contactcontact surface; surface; ( (ee)) Ti Ti-6-4-6-4 with with CP CP-Ti-Ti interface at intermediate strains; strains; ( (ff)) h highigh strain bond bond interface interface forfor both both alloys; alloys; ( (gg)) b bondond interface interface consisting consisting of of fine ﬁne grained grained Ti Ti-6-4-6-4 and and small small grained grained CP CP-Ti;-Ti; ( (hh)) b bondond interfaceinterface consisting ofof lowlow strain strain Ti-6-4 Ti-6- and4 and high high strain strain CP-Ti; CP- (Ti;i) bond(i) bond interface interface consisting consisting of low of strain low strainTi-6-4 Ti and-6- high4 and strain high strain CP-Ti; CP and-Ti; (j )and bulk (j) CP-Ti bulk inCP an-Ti medium in an medium strained strained region. region.

A high-strain region in EB2 was also microstructurally analysed at 56.4 mm from the reference end, as defined in Figure 11b. The varying microstructures across the eye-bolt, both in the side sections and also in the thin “eye” section, are shown in Figure 16. Site “a” shows the microstructure towards the middle of the Ti-6-4 region and consists of a typical Widmanstätten microstructure, which has been slightly deformed by a small increase in strain in this region. Site 'b' is located on the bond interface within the Ti-6-4 alloy region and is comparable

Metals 2019, 9, x REVISED 31 May 2019 17 of 19 to an as-FAST microstructure [19]. The Ti-6-4 has a similar microstructure to that observed at site “a”, suggesting similar strain histories. The microstructure changes quite abruptly from the Ti-6-4 into the Ti-5553, as evidenced by the disappearance of the α laths, shown as the darker bands in Figure 16b. The Ti-5553 comprises of a fine sub-structure, which is absent of any prominent microstructural features, as a result of the fully-retained β microstructure produced from soaking the workpiece slightly above the β transus and subsequently quenched to below the β transus. At high magnifications using SEM analysis no evidence of the α precipitates were observed, as the α phase does not have sufficient time to precipitate out at the fast cooling rates. Some lighter and darker areas exist in the Ti-5553 microstructure, which are likely to be deformed grain boundaries and grain sub- structures. Site “c” shows an excellent example of how robust the bond interface can be when severely deformed with a wavy prior bond interface morphology; no defects or evidence of cracking are evident at the interface. Site “d” and “e” represent the Ti-5553 with Ti-6-4 bond interface at high strains, located within the thin section of the workpiece. Each micrograph can be divided into three distinct regions. The first is the bulk Ti-5553 microstructure, followed by a bonding region and, finally, the Ti-6-4 bulk microstructure. The bulk alloys are severely deformed, and there is clear evidence of the Ti-6-4 alloy flow at site “d”. The bonding region at site “d” has a darker appearance than the bulk Ti-5553 and contains fine α phase, which gradually evolves into the bulk Ti-6-4. The α phase in the bonding region follows the grain boundaries, which are more favourable paths for diffusion. Site “f” shows a high magnification micrograph of the interface where the two alloys meet eachMetals other2019, 9 , in 654 the centre of workpiece. There is evidence of diffusion at the interface by a 17 grain of 19 containing the α phase on one side and an absence on the other, demonstrated by [19].

Figure 16. High-resolution backscatter micrographs of regions of interest in the high strain cross section Figure 16. High-resolution backscatter micrographs of regions of interest in the high strain cross of eye-bolt 2 (Ti-5553 with Ti-6-4); (a) bulk Ti-6-4 microstructure in the side section; (b) bond interface section of eye-bolt 2 (Ti-5553 with Ti-6-4); (a) bulk Ti-6-4 microstructure in the side section; (b) bond in a relatively low strain location; (c) and an example of a wavy interface; (d) high-strain bond interface interface in a relatively low strain location; (c) and an example of a wavy interface; (d) high-strain on the upper side of the thin section; (e) the high-strain bond interface on the lower side of the thin bond interface on the upper side of the thin section; (e) the high-strain bond interface on the lower section; and (f) the bond interface at the middle of the eye-bolt. side of the thin section; and (f) the bond interface at the middle of the eye-bolt. Site “a” shows the microstructure towards the middle of the Ti-6-4 region and consists of a typical Widmanstätten microstructure, which has been slightly deformed by a small increase in strain in this region. Site ‘b’ is located on the bond interface within the Ti-6-4 alloy region and is comparable to an as-FAST microstructure [19]. The Ti-6-4 has a similar microstructure to that observed at site “a”, suggesting similar strain histories. The microstructure changes quite abruptly from the Ti-6-4 into the Ti-5553, as evidenced by the disappearance of the α laths, shown as the darker bands in Figure 16b. The Ti-5553 comprises of a fine sub-structure, which is absent of any prominent microstructural features, as a result of the fully-retained β microstructure produced from soaking the workpiece slightly above the β transus and subsequently quenched to below the β transus. At high magnifications using SEM analysis no evidence of the α precipitates were observed, as the α phase does not have sufficient time to precipitate out at the fast cooling rates. Some lighter and darker areas exist in the Ti-5553 microstructure, which are likely to be deformed grain boundaries and grain sub-structures. Site “c” shows an excellent example of how robust the bond interface can be when severely deformed with a wavy prior bond interface morphology; no defects or evidence of cracking are evident at the interface. Site “d” and “e” represent the Ti-5553 with Ti-6-4 bond interface at high strains, located within the thin section of the workpiece. Each micrograph can be divided into three distinct regions. The first is the bulk Ti-5553 microstructure, followed by a bonding region and, finally, the Ti-6-4 bulk microstructure. The bulk alloys are severely deformed, and there is clear evidence of the Ti-6-4 alloy flow at site “d”. The bonding region at site “d” has a darker appearance than the bulk Ti-5553 and contains fine α phase, which gradually evolves into the bulk Ti-6-4. The α phase in the bonding region follows the Metals 2019, 9, 654 18 of 19 grain boundaries, which are more favourable paths for diffusion. Site “f” shows a high magnification micrograph of the interface where the two alloys meet each other in the centre of workpiece. There is evidence of diffusion at the interface by a grain containing the α phase on one side and an absence on the other, demonstrated by [19].

4. Conclusions For the first time, the FAST process has been industrial-scaled using angular and spherical morphology powders to produce near-net shape forgings containing dissimilar titanium alloys. The material flow of FAST-DB compression samples is dependent on the flow stresses of the titanium alloys used, the processing conditions, and also the volume fraction of each alloy. Even though there is a mismatch in the magnitude of flow stress between alloys and flow characteristics, for example, Ti-6-4 exhibits flow softening and CP-Ti work hardening, excellent structural bond integrity is maintained throughout. In high strain regions in eye-bolts 1 and 2, the bond interface has been shown to retain high integrity despite the relatively uncontrolled drop forging processing. There is an absence of processing defects at the interface providing confidence that the bonds could endure typical titanium thermomechanical processing routes in aerospace and automotive applications. In Ti-6-4 with CP-Ti and Ti-6-4 with Ti-5553 bond combinations, the lower flow stress alloy at the forging temperature was displaced by the higher flow stress alloy, in both high strain and low strain regions. In the future, alloy compatibility of each bond combination needs to be considered, specifically differences in flow stresses. In addition, measurements of residual stress profiles across forged dissimilar bonds to ensure long term durability. There was excellent agreement between the FE simulations of the closed-die forging process of Ti-6-4 with CP-Ti bond combinations in the high strain region. The ability to predict the thermomechanical responses of a FAST-DB component, is a key milestone in the development of processing titanium alloy components by this approach.

Author Contributions: J.P. conducted the experimental work as part of his Ph.D study. M.J. was the academic supervisor. J.P. and M.J. conceived the idea and designed the experiments. Funding: This research was funded by Engineering and Physical Sciences Research Council’s Advanced Metallic Systems Centre for Doctoral Training (grant number EP/G036950/1). Acknowledgments: The authors acknowledge Adam Tudball from Kennametal Ltd. UK for support with processing the large FAST-DB billets. We gratefully acknowledge David Lunn and the staff at W.H. Tildesley Ltd. for the drop forging trials and useful discussions. Finally, thanks to Dean Haylock from The University of Sheffield for technical assistance, and Nicholas Weston, Ben Thomas, and Bhushan Rakshe for useful discussions. Conflicts of Interest: The authors declare no conflict of interest.

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