Tial alloy produced by direct energy deposition

Abstract

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Description

[0001] TIAL ALLOY PRODUCED BY DIRECT ENERGY DEPOSITION

[0002] FIELD OF INVENTION

[0003] THIS invention is in the field of alloys made by additive manufacturing, in particular direct energy deposition, more in particular Laser Engineering Net Shaping.

[0004] BACKGROUND OF INVENTION

[0005] Aircraft engine components are typically manufactured using cutting-edge and specialized materials, like ultrahigh-temperature ceramics, superior steels, nickel- based (Ni-based) superalloys or titanium-based (Ti-based) alloys, that are hard, expensive and problematic to fabricate using conventional methods. Moreover, the most extensively used Ti alloys for aero-engine parts are titanium aluminides (TiAl) and Ti6AI4V due to their outstanding properties. Weight reduction to save energy, enhance engine efficiency and reduce CO2 emissions is a major goal in all mobility sectors. The high demand for higher fuel efficiency, weight reduction and lower engine noise in aircraft engines drive the development of lightweight high- temperature structural materials which can partially replace the currently used heavier superalloys. The TiAI-based alloys have several advantages over the conventional titanium alloys, such as higher specific strength and modulus, lower density, high strength at elevated temperatures, and superior higher oxidation resistance due to the formation of a surface passivated alumina layer. Gamma-TiAl alloys have been widely used in various high-temperature applications such as aerospace engine blades and high-performance combustion engines due to their low-density levels, high specific strength retention at high temperature, excellent creep properties, and good oxidation resistance.

[0006] However, the major drawbacks for the structural applications of y-TiAl alloys using conventional fabrication techniques such as forging and extrusion is their limited workability and room temperature ductility. This inherent poor room -temperature ductility, insufficient fracture toughness and poor oxidation at temperatures above 800 °C limits their wide-scale applications. This impedes its further applicability for engineering parts with complicated geometry via traditional processing technologies, and in some cases up to 65 times the manufacturing cost of nickel superalloys. Also, the application of TiAl is restricted owing to technical hitches when it is being machined due to its inherent brittle nature that is attributed to the phases of the intermetallic present.

[0007] AM as an emerging technology enables the fabrication of near-net-shape components with complex geometries with relatively short lead times at higher efficiency and lower costs, compared with conventional manufacturing technologies. The main distinction for AM technologies is done according to the type of energy source utilised and the way the material is being processed. AM can be classified into two major groups; Powder Bed Fusion based technologies (PBF) and Directed Energy Deposition (DED). In the DED process which is also referred to as laser metal deposition (LMD), direct laser metal deposition (DLMD), direct metal deposition (DMD), direct laser deposition (DLD) and laser engineered net shaping (LENS), powder particles are sprayed through a nozzle onto a substrate where a high-power focused laser beam melts the powder along a predefined CAD path while in the PBF process, the powder to be melted are pre-placed on the platform before the laser or beam melts the powders based on the design paths. AM technologies and the ability to produce near-net-shape components are becoming increasingly important for TiAI-based alloys, which are difficult to process via conventional machining due to their pronounced brittleness. Metal additive manufacturing techniques such as PBF and DED are currently being explored to fabricate y-TiAl alloy components.

[0008] Since the curiosity in TiAI-based alloys has increased vastly but applications seem to be limited due to processing problems in manufacturing parts. The present invention is aimed at fabricating of Ti-AI alloy by additive manufacturing that will display improved engineering performances and processable so as to prepare parts of intricate geometries for aerospace applications.

[0009] SUMMARY OF THE INVENTION

[0010] IN ACCORDANCE WITH A FIRST ASPECT OF THE INVENTION, THERE IS PROVIDED an alloy produced by additive manufacturing, the alloy comprising: at least one or two metals and alloying elements; wherein the alloying elements comprising grain refining agents and grain stabilizing agents.

[0011] The at least one or two metals may preferably be two metals, and the two metals may be aluminium and titanium.

[0012] The grain refining agent may be selected from a group comprising of silicon, boron, carbon, titanium diboride and combinations thereof. The grain refining agent may be silicon.

[0013] The grain stabilizing agent may be selected from a group comprising molybdenum and vanadium and combinations thereof.

[0014] The additive manufacturing employed to fabricate / produce the alloy may be in the form of direct energy deposition, in particular LENS.

[0015] The titanium aluminide-based alloy may be deposited on a metal alloy substrate such as Ti6AI4V.

[0016] The titanium aluminide based alloy may be deposited on the metal alloy substrate either at ambient / room temperature (i.e. , about 23 degrees Celsius) or at an elevated temperature.

[0017] The elevated temperature may be between about 750 and 850 degrees Celsius, that is not less than 750 degrees Celsius and not more than 850 degrees Celsius, preferably between 800 and 850 degrees Celsius.

[0018] The metal alloy substrate may be connected to a heating means (e.g., a heating platform) arranged to heat and maintain the temperature of the metal alloy substrate at the elevated temperature.

[0019] The metal alloy substrate may also be rotated by a rotation means, for example by a rotation platform that may be connected to the heating means, during the fabrication / production / making of the titanium aluminide-based alloy.

[0020] The produced titanium aluminide-based alloy may comprise: from 44 to 48 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1.0 to 1.5 at.% molybdenum; and from 6 to 10 at.% vanadium; and balance is titanium. The titanium aluminide-based alloy may, comprise a duplex microstructure.

[0021] The duplex microstructure may comprise columnar grains and lamellar grains.

[0022] The columnar grains may be 02+y columnar grains.

[0023] Unmelted Al, p-phase and ^-TisSis-phase may be present at the grain boundaries and within the lamellar grains. The titanium aluminide based alloy may have a heterogenous microstructure having grain sizes ranging between 4 and 14 pm.

[0024] The produced titanium aluminide based alloy may be subjected to heat treatment at a temperature between about 1200 and about 1400 degrees Celsius for a predefined period and subsequently subjected to cooling, preferably furnace cooling. The heat treated titanium aluminide alloy may therefore comprise: from 46 to 50 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1.0 to 1.5% molybdenum; from 7 to 10 at.% vanadium; and balance is titanium.

[0025] The titanium aluminide based alloy may, after being subjected to heat treatment at a temperature of about 1200 degrees Celsius for a predefined period of about 60 minutes and subsequently subjected to cooling, preferably furnace cooling, comprise a duplex microstructure comprising of columnar and lamellar grains.

[0026] The duplex microstructure may have grain sizes of about 80 pm.

[0027] The titanium aluminide based alloy may, after being subjected to heat treatment at a temperature of about 1400 degrees Celsius for a predefined period of about 60 minutes and subsequently subjected to cooling, preferably furnace cooling, comprise a fully lamellar microstructure.

[0028] The fully lamellar microstructure may have grain sizes of about 100 pm.

[0029] The ultimate tensile strength of the titanium-aluminide based alloy, after being subjected to heat treatment and cooling, is between 500 and 650 MPa, preferably between 550 and 650 MPa, more preferably between 586 and 630 MPa. The yield strength of the titanium-aluminide based alloy, after being subjected to heat treatment and cooling, is between 300 and 500 MPa, preferably between 350 and 450 MPa, more preferably between 368 and 438 MPa.

[0030] The starting feedstock to produce the titanium aluminide-based alloy, by direct energy deposition, may be selected from powders and wires, preferably powders. The feedstock may comprise of individual, unmixed powders of essentially titanium powder, essentially aluminum powder, essentially vanadium powder, essentially molybdenum powder, and essentially silicon powder.

[0031] The powders may be individually and simultaneously co-deposited, at predefined flowrates, on the substrate and subjected to an energy source, preferably a laser energy beam.

[0032] The simultaneous co-depositing of the individual powders (i.e. , multi-deposition of the individual powders simultaneously or selectively on the substrate) may assist in limiting or eliminating powder, in particular aluminium powder, evaporisation during the making of the titanium aluminide-based alloy.

[0033] The flowrates of the titanium, aluminium, and silicon powders deposited on the substrate may be kept constant at between about 2.2 and 2.3g / min, 0.4 and 0.5 g / min, and 0.02 and 0.025g / min, respectively. The carrier gas for the titanium, aluminium and silicon powders may be argon.

[0034] The flowrates of the vanadium powder deposited on the substrate may range between 0.1 and 0.5g / min, preferably between 0.35 and 0.4g / min. The carrier gas for vanadium powder may be argon.

[0035] The flowrates of molybdenum deposited on the substrate may range between 0.05 and 0.16g / min, preferably kept constant at 0.05g / min. The carrier gas for the molybdenum powder may be argon.

[0036] The laser power may be kept constant at about 450 Watts. The titanium aluminide-based alloy may be fabricated in the form of a layer or multiple layers.

[0037] The titanium aluminide- based alloy produced by direct energy deposition may be in the form of a component or portion thereof.

[0038] ACCORDING TO A SECOND ASPECT OF THE INVENTION THERE IS PROVIDED AN ADDITIVE MANUFACTURING METHOD OF MAKING AN ALLOY, THE METHOD COMPRISING: depositing a feedstock of separate materials comprising one or more metals and alloying elements on a substrate, wherein the alloying elements are selected from a group comprising of at least one stabilizing agent and at least one grain refining agent; subjecting the deposited feedstock of separate materials comprising the one or more metals and alloying elements to an energy source (i.e., a focused energy source), so as to melt the materials to form a pool of molten metal alloy; cooling the molten metal alloy to form a first layer of cooled metal alloy; and optionally repeating the steps of depositing, subjecting the deposited feedstock of materials to an energy source, and cooling as many times as possible to prepare a desired number of layers of the cooled metal alloy.

[0039] The substrate may be a titanium-based substrate, such as a Ti6AI4V metal alloy plate.

[0040] The depositing of the materials may be by simultaneously or selectively codepositing the materials on the substrate and subjecting same to the energy source. The feedstock materials may each be in powder format.

[0041] The materials may comprise essentially titanium in powder format, essentially aluminium in powder format, essentially vanadium in powder format, essentially molybdenum in powder format and essentially silicon in powder format.

[0042] The feedstock may comprise of independent materials which are selectively or simultaneously co-deposited onto the substrate, in the building direction of the energy source, and are not mixed together.

[0043] In particular, the deposition may be by means of a multi-powder deposition method where each powder is deposited on the substrate from a dedicated container such as a hopper. The powders may be deposited simultaneously at different flowrates. In another version, the powders may be selectively deposited on the substrate at predetermined flowrates.

[0044] The orientation of the laser beam and individual powders being deposited onto the substrate may be concentric. The laser beam and the deposition of the individual powders may be controlled simultaneously during deposition of the metal powders.

[0045] The combined feedstock may comprise: from 44 to 48 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1 .0 to 1 .5 at.% molybdenum; from 6 to 10 at.% vanadium; and balance is titanium. The energy source may be a laser energy beam that is arranged to follow a predefined path according to a predefined algorithm.

[0046] The method may comprise making the alloy on the substrate, with the substrate not being externally heated other than by the energy obtained from the energy source.

[0047] Preferably, the method may comprise making the alloy on the substrate, with the substrate being externally heated by a secondary energy source other than the energy obtained from the energy source used to make the alloy.

[0048] The secondary energy source may be provided by a heating platform supporting the substrate.

[0049] The method may comprise rotating the substrate, about an axis of rotation, relative to the direction of the energy source used to make the alloy, during the making of the alloy.

[0050] The rotation may be provided by a rotation means which may be coupled to the heating platform.

[0051] The method may further comprise, after making the alloy of a desired number of layers, subjecting the alloy to heat treatment for a predefined period.

[0052] The heat treatment may comprise heating the alloy in a furnace at a temperature between about 1200 and about 1400 degrees Celsius. The temperature in the furnace may be increased incrementally for a predefined period of time until it reaches the temperature of between about 1200 and 1400 degrees Celsius, for example the temperature may be increased at a rate of 20 degrees Celsius per minute. The heat treatment may be conducted in an inert environment. The furnace may be purged with and constantly supplied with an inert gas, such as argon, so as to create the inert environment.

[0053] The alloy may be subjected to heat treatment for a period of about 60 minutes.

[0054] After heat treatment, the method may comprise cooling, preferably furnace cooling, at a cooling rate of between 20 and 25 degrees Celsius per minute.

[0055] ACCORDING TO YET ANOTHER ASPECT OF THE INVENTION, THERE IS PROVIDED AN ALLOY PRODUCED ACCORDING TO THE METHOD OF THE SECOND ASPECT OF THE INVENTION.

[0056] The alloy produced according to the second aspect of the invention may be a titanium aluminide-based alloy.

[0057] The titanium aluminide based alloy may, after being subjected to heat treatment at a temperature between 1200 and 1400 degrees Celsius for a predefined period and subsequently subjected to cooling, preferably furnace cooling, comprise: from 46 to 50 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1.0 to 1.5% molybdenum; from 7 to 10 at.% vanadium; and balance is titanium. The titanium aluminide based alloy may, after being subjected to heat treatment at a temperature of about 1200 degrees Celsius for a predefined period of about 60 minutes and subsequently subjected to cooling, preferably furnace cooling, comprise a duplex microstructure comprising of columnar and lamellar grains.

[0058] The duplex microstructure may have grains ranging between 4 and 14pm.

[0059] The titanium aluminide based alloy may, after being subjected to heat treatment at a temperature of about 1200 degrees Celsius for a predefined period of about 60 minutes and subsequently subjected to cooling, preferably furnace cooling, comprise a fully lamellar microstructure.

[0060] The fully lamellar microstructure may have grain sizes of about 100 pm.

[0061] BRIEF DESCRIPTION OF DRAWINGS

[0062] The invention will now be described in more detail by way of non-limiting example, with reference to worked experimental examples and to the accompanying drawings, in which:

[0063] Figure 1 shows SEM Images of the Powder Morphology: (a) Titanium, (b) Aluminium, (c) Silicon, (d) Molybdenum and (e) Vanadium;

[0064] Figure 2 shows schematic representation of LENS Process at the Deposition Head;

[0065] Figure 3 shows the schematic Design of the Modified Optomec LENS System;

[0066] Figure 4 shows a temperature History Profile of the Rectangular Block Build at

[0067] (a) Ambient Temperature; and (b) 800 °C; Figure s shows the Forming of the Rectangular Block Build at Ambient Temperature for (a) 49.6 seconds, (b) 109.1 seconds, (c) 168.6 seconds and (d) 218.2 seconds; the build at 800°C for (e) 49.6 seconds, (f) 109.1 seconds, (g) 168.6 seconds and (h) 218.3 seconds;

[0068] Figure 6 shows a complete Build of the Rectangular Block Build at (a) Ambient temperature and (b) 800°C;

[0069] Figure 7 shows a surface Temperature Profile of the Complete Rectangular Block Build at (a) Ambient Temperature; and (b) 800°C;

[0070] Figure 8 shows a Stress Deformation Contour Profile of the Rectangular Block (in 2D) Build at (a) Ambient Temperature; and (b) 800°C;

[0071] Figure 9 shows SEM Images of as-build Ti-AI-Si-x(Mo+V) Sample Processed at (c) 0.05 g / min Mo + 0.50 g / min V, and (d) 0.05 g / min Mo + 0.35 g / min

[0072] V Fabricated on the Heating Platform;

[0073] Figure 10 shows SEM images of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Sample Fabricated on the Heating Platform Heat Treated at (a) 1200°C / 60 mins / FC and (b) 1400°C / 60 mins / FC;

[0074] Figure 11 shows SEM Images of Ti-AI-S i-0.05 g / min Mo + 0.50 g / min V Samples Fabricated on the Heating Platform Heat Treated at (a) 1200°C / 60 mins / FC and (b) 1400°C / 60 mins / FC;

[0075] Figure 12 shows EBSD Analysis of As-Built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min

[0076] V Alloy Showing (a) FSD Image, (b) EDS Layered Image, and (c) Phase Mapping;

[0077] Figure 13 shows EBSD Analysis of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy Heat Treated at 1400 °C / 60 mins / FC Showing (a) FSD Image, (b) EDS Layered Image, and (c) Phase Mapping; Figure 14 shows Grain Boundary Area of Ti-AI-Si-0.08 g / min Mo + 0.35 g / min V Alloy Heat Treated at 1400°C / 60 mins / FC (a) SEM Image with Inset of FSD Image, and (b) Phase Mapping of the Selected Area;

[0078] Figure 15 shows XRD Analysis of Ti-AI-Si-0.08 g / min Mo + 0.35 g / min V Alloy;

[0079] Figure 16 shows microhardness of As-Build Ti-AI-Si-x(Mo+V) Sample Processed on the Heating Platform;

[0080] Figure 17 shows microhardness of Heat-Treated Ti-AI-0.025 g / min Si-x(Mo+V) Samples Fabricated on the Heating Platform;

[0081] Figure 18 shows nanoindentation Load-Displacement Curves of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy;

[0082] Figure 19 shows True Stress-True Strain Curve of As-Built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy Showing the Regions within the Curve;

[0083] Figure 20 shows True Stress-True Strain Curve of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Heat-Treated at 1200°C / 60 mins / FC Alloy Showing the Regions within the Curve;

[0084] Figure 21 shows True Stress-True Strain Curve of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min Heat-Treated at 1400°C / 60 mins / FC Alloy Showing the Regions within the Curve;

[0085] Figure 22 shows Coefficient of Friction and Wear Rate of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy; and

[0086] Figure 23 shows TGA Curves of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy. EXAMPLES

[0087] FEATURES OF THE INVENTION as described above will be apparent from the worked experimental examples that follow.

[0088] Raw materials characterization

[0089] Raw materials comprising of commercially Pure (CP) individual elemental powders of Ti, Al, Si, Mo and V were used with Ti6AI4V alloy metal plate as the substrate. As-received powders were all subjected to particle size analysis and examined with SEM-EDS. The SEM images of these metal powders are presented in Figure 1. All metal powders were in the required range of 45-90 pm needed for processing in the LENS machine.

[0090] The feedstock of the combined individual powders comprises about:

[0091] 44-48 at.% Al,

[0092] 0.9-1. O at.% Si,

[0093] 1.0-1.5 at.% Mo,

[0094] 6-10 at. % V, and

[0095] Balance is Ti.

[0096] Experimental procedure

[0097] Methodology

[0098] The Direct Energy Deposition (DED) LENS Optomec 850R machine (Albuquerque, NM, USA) that used a 1 kW IRE-Polus Group (IPG) fiber laser as a source of energy, was used for the printing of the TiAI-based composite alloys. Each of the samples were deposited onto metal plate substrates of Ti6AI4V alloy via laser in-situ alloy deposition (i.e., co-deposition or multi-powder deposition) of the individual powders. The LENS system deposition head had five feeding nozzles made of copper, with one of each nozzle being dedicated to one of each powder that is not mixed with any of the other powders, such that the individual powders are simultaneously or selectively co-deposited from their respective nozzles at predefined deposition rates.

[0099] The parameters used in the LENS machine to develop various quintenary alloy Ti-AI-

[0100] Si-x(Mo+V) samples are presented in Table 1 and Table 2. Table 1 : Parameters used for preparing Quintenary Ti-AI-Si-x(Mo+V) Alloy samples

[0101] Table 2: Laser Processing Parameters A heating platform, in the form of a ceramic heating platform that has a heating element and thermostat controller attached to it was introduced in the DED LENS machine to heat the samples as they were being prepared / fabricated in the LENS machine. The heating platform had the capability of rotating the substrate during the fabrication of the alloy. The heating platform was thus used to continuously keep the temperature of the substrate constant, at a temperature of about 800 degrees Celsius, to prevent cracking of the alloy.

[0102] Sample Fabrication

[0103] Figure 2 depicts the representation of typical DED LENS technique at the deposition head when the laser melts the powders on the substrate. The Al and Ti powders were supplied by the hoppers on the LENS machine, whereas other powders (Si, Mo and V) were externally fed into the LENS system through powder hoppers of GTV (Verschleiss-Schutz Gmbh, Luckenbach, Germany). Therefore, the design setup employed to fabricate the Ti-AI based alloy is a modified five hopper powder feeder LENS system as shown in Figure 3. Owing to the cracking susceptibility of TiAI- based alloys during DED processing, a heating platform was incorporated into the LENS machine. The heating platform was used to control and / or reduce thermal gradient during sample printing.

[0104] Heat Treatment

[0105] After the fabrication of the samples, heat treatment was performed in an inert atmosphere to avoid oxidation and other contaminations. A heat treatment furnace (Kejia Furnace, China) that has an inlet for the continuous supply of Argon to prevent contaminations and oxidation during heat treatment, was utilized for performing heat treatment. The samples were heat treated at various temperatures between 1200°C and 1400°C for 1 hr, followed by furnace cooling (FC) at a rate of 20-25°C per minute. Sample Preparations

[0106] To prepare the samples for further analysis, the fabricated samples were sectioned along the build directions using the Struers Discotom-5 (Advance Laboratory Solutions, South Africa) cutting Machine. The machine is controlled manually and highly suitable for sectioning an extensive range of materials. The sectioned samples were then mounted in phenolic resin using an automatic mounting press Press AMP 50 (Laryee Technology Co. LTD., China) for preparation before characterizations. This was followed by manual grinding using grit papers starting with P80, P320, P1200 and P4000 grit size followed by polishing with OP-S suspension fluid. The Struers TegraForce-5 (Struers Inc, USA) grinding / polishing machine was used in this work. Kroll’s reagent with 92 ml distilled water (H2O), 6 ml hydrochloric acid (HCI) and 2 ml hydrofluoric acid (HF) and was used as an etchant for making the microstructure visible for microstructural and phase analysis.

[0107] Microstructure and Phase Examination

[0108] The JEOL JSM-6010PLUS / LA (JEOL Ltd., Tokyo, Japan) analytical SEM equipped with EDS was used for the examination of the microstructure, phases, morphological variations, grain size and other microstructural observations. Elements present within the samples were identified by EDS. Composition and distribution analysis of elements, mainly Ti, Al, Si, Mo and V were done. The electron back-scattered diffraction (EBSD) analysis was performed using the ZEISS Crossbeam 540 (Zeiss International, Oberkochen Germany) FIB-SEM furnished with symmetry EBSD detector. Identification of phases and calculation of phase fractions in the alloy samples were done. The phases present in the alloys were characterized via XRD analysis using the PANalytical Empyrean (Malvern Panalytical Ltd, United Kingdom) machine. This approach is centered on studying the intensity of the scattered beam of X-ray as a function of incident and scattered angle of polarization, wavelength or energy on a sample. This employs a radiation source of Cu-Ka monochromator to study the crystallographic structure of samples. Commercial software X-Pert HighScore Plus was used for plotting and identification of the peaks in the graphs plotted using Origin software. The peak positions and background were ascertained based on the intensities and peak positions; a routine search match was conducted. Samples were X-rayed over 20 between 5° to 110° and 0.02° step size.

[0109] Microhardness Testing

[0110] Microhardness testing, also known as the Vickers hardness testing, was employed on the samples using Zwick / Roell, ZHVp (ZwickRoell GmbH, Germany) Vickers’ hardness tester. The microhardness testing machine makes use of an optical measuring technique of ASTM E-384 procedure which stipulates a range of light loads utilizing a diamond indenter. A load force of 500gf for 10 seconds dwelling time was used to make a total of thirty indentations (30) each on the samples; while average Hv values were reported to represent the microhardness values. A quick approximation of the YS was calculated from microhardness using the empirical relationship between microhardness and YS as shown in the below Equation: where, oyis the yield strength in MPa and Hvin MPa is Hvx 9.81 .

[0111] Nanoindentation Testing The nanoindentation testing was carried out using Anton-Paar TTX-NHT3 (Anton Paar GmbH, Austria) nanoindentation tester that has a Berkovich indenting tip with a radius of 20 nm. Before performing tests on all the samples, the machine was calibrated using fused silica, as a reference sample. The maximum loading force was set at 200 mN and was allowed to hold for 20 seconds before unloading for 20 seconds. The nanoindentation testing process records the loading against the displacement (depth of penetration). ear and TGA Tests

[0112] The Anton-Paar TRB3tribometer (Anton Paar GmbH, Austria), using a pin-on-disc based on ASTM G99 and ASTM G133, was used to evaluate the wear and tribology attributes of the samples. The tribometer systematically measures the results of the friction between the two surfaces with the aid of a friction force sensor and plots the resultant friction data in real-time. The wear tests were performed at 15 m / min sliding speed and 20N normal applied load at RT. This method is regarded as sufficiently trustworthy to estimate the functioning of the alloys in terms of tribology.

[0113] The DED fabricated alloys were studied with the aid of a thermogravimetric analyzer (TGA) using PerkinElmer thermogravimetric analyzer (TGA 4000; PerkinElmer Inc., USA) used to determine the change in the weight of the specimen relative to the unit time with temperature. An inlet gas of N2 and O2 with a composition of 79% and 21 %, respectively, was infused at 20 ml / min. At 50°C / min heating rate a maximum temperature of 900°C was attained for the test. A built-in accurate weigh balance with a pan inside the chamber during heating monitors the change in weight of the samples at various stages of the experiment up to the maximum temperature. TGA graphs were plotted from the results obtained; this consists of temperature (°C) on the x-axis and mass loss (or gain in mg) on the y-axis.

[0114] Results and Discussions

[0115] Temperature Profile during Deposition

[0116] Figures 4a and 4b presents the transient temperature profile of one rectangular block built on the substrate kept at ambient temperature (i.e., 23°C), and another rectangular block that is built on a substrate that is heated and kept at an elevated temperature of about 800°C. The acute rise in temperature depicts the laser interaction during the melting of the metal powder along the build direction; while the drop-in temperatures depict the laser idle time when the laser is off and the melt pool solidifies during cooling. The peak points show the completion of a particular layer of the build. However, in both cases the maximum temperatures of the succeeding layer are higher than the preceding layers, except at the time, 120 seconds where peak temperature drops slightly to less than the preceding layer. This was attributed to the prolonged idle time in switching from building the layers along the x-axis to building along the y-axis. Although, the cooling rate for both scenarios seems similar, the minimum temperature in each case are not similar. This is because both simulations were subjected to similar conditions but the pre-heating temperature (of the substrate plate) makes the build in Figure 4b experience higher temperature and stability during the DED processing of the TiAl alloy. The highest temperature recorded for the build process at ambient temperature was about 1100°C at the 20,hlayer when building along the y-axis. But the highest temperature recorded for the build process at 800°C was about 1900°C noticed at the 17thand 18thlayer when building along the y-axis.

[0117] Generally, when building along the y-axis, the temperature was very high regardless of the processing temperatures. This was ascribed to the shorter length of the build in the y-axis and the re-melting first set of layers build along the x-axis. It is understood that the DED process has a rapid heating and cooling regime which affects part quality and formation of cracks especially for intermetallic alloys like TiAI. This is influenced by the temperature variations due to inhomogeneous distribution and thermal gradient. It was observed that the average temperature gradually increases all through the build. This demonstrates that the temperature would continuously increase as more layers are added.

[0118] Melt Pool

[0119] The fabrication of the rectangular block in stages at different times with a focus on the stability of the melt pool region is presented in Figure 5. The yellowish-white colored portion denotes areas where the temperature is higher than TiAI alloy liquidus temperature. As observed from Figure 5a and 5e, at 49.6 seconds, the melt pool and the built for both cases looks very stable and somewhat similar. But at 109.1 seconds (Figure 5b and 5f), the build processed at ambient temperature (23°C) starts to get distorted at the edges while the build process at 800°C was still very stable without any evidence of defects. Furthermore, more distortions were more visible at 168.6 seconds for the build process at 23°C (Figure 5c) and the melt pool seems unstable at this point. It was inferred that the distortion would lead to crack initiation and propagation. However, the build process at 800°C (Figure 5g) still shows good melt pool stability without any apparent defects noticed. But at 218.2 seconds (Figure 5d and 5h), the build process at ambient temperature (23°C) exhibits lots of distortion all over the surface and at the edges while the build process at 800°C was showing very minimal defects on the surface of the build.

[0120] Figure 6 shows the complete build of the rectangular block build processed at 23°C and 800°C. The examination of the builds suggests that the processing at about 800°C would produce a part that has a more stable melt pool to achieve good component geometry.

[0121] Surface Temperature Field

[0122] The surface temperature profile of the rectangular block build at ambient temperature and 800°C is presented in Figures 7a and 7b, respectively. The thermal cycle and peak temperatures have undergone by each layer have a significant impact on the dimensional accuracy and final mechanical properties of the produced component. It was observed that the TiAl deposited at about 800°C had better dimensional accuracy than the build process at about 23°C. The highest temperatures were noticed at locations of the greatest heat flux applied. But peak temperature varied with the intensity at the surface of the TiAl deposited along longitudinal and width paths. The temperature gradient distribution was essentially noticed on the surface as the higher energy density. The material cooling rates during LENS processing were obtained from the thermal analysis transient temperature profile on the melt pool surface. The TiAl alloy liquidus temperature 1807°C and solidus temperature 1719°C has a close-range which is assumed to be linear in variation with time. Stress Deformation

[0123] Figures 8a and 8b shows the stress deformation contour profile of the rectangular block (in 2D) built at ambient temperature and at about 800°C, respectively. It was observed that the maximum Von Mises stress (0.52 GPa or 520 MPa) were similar irrespective of the processing temperatures. However, the lower limit of the deposit processed at about 800°C was at 1.69 x 10'4GPa while the deposit build at ambient temperature was 1.39 x 10'4GPa. This shows that an increase in temperature increases the residual stress within the deposit but does not exceed the UTS of TiAl alloy. Logically, the residual stress obtained should not be more than the YS of the material to avoid cracking. The preheating of the substrate decreases the probabilities of TiAl samples cracking during DED processing.

[0124] Microstructural evolution

[0125] As-build Ti-AI-Si-x(Mo+V) alloy samples

[0126] Figures 9 shows the SEM images of as-built Ti-AI-Si-x(Mo+V) alloy samples fabricated on the heating platform (i.e. , substrate heated to about 800 degrees Celsius) with the addition of 0.05 g / min Mo + 0.50 g / min V and 0.05 g / min Mo + 0.35 g / min V; while Table 3 presents the compositions of the as-built Ti-AI-Si-x(Mo-rV) alloy samples fabricated on the heating platform (i.e., substrate heated to about 800 degrees Celsius) with the addition of 0.05 g / min Mo + 0.50 g / min V and 0.05 g / min Mo + 0.35 g / min V. The as-built microstructures of Ti-AI-Si-x(Mo+V) fabricated on the heating platform contains largely O2 / y lamellae with relatively small amount of |3- phase and ^-TisSis. It was also suspected to contain unmelted Al within the matrix. The quintenary alloys of Ti-AI-Si-x(Mo+V) show lamellar microstructures promoted by the addition of 0.50 g / min and 0.35 g / min of V including the presence of high Al. Both Al and aid the formation of a2+y / y lamellae thereby reducing the quantity of p formed. As the Mo+V is increased from 0.05 g / min Mo + 0.35 g / min V to 0.05 g / min Mo + 0.5 g / min V, the Si remain virtually constant while the high quantity of V causes a reduction in Al from 47.95 at.% to 46.92 at.%. It was deduced that increasing the p- phase in the 0.05 g / min Mo + 0.50 g / min V alloy was caused by the rise in V from 6.13 at.% to 9.35 at.%. Moreover, the Al content is high enough to favor the production of y-phase needed to promote lamellar formation to achieve good mechanical properties. The small addition of Mo being a very strong p-stabilizer is expected to help improve the mechanical properties. The p-phase tends to cause shrinkage of the Q2 / y grains due to its higher specific strength, thus, at high quantities would be a detriment to the microstructure and mechanical properties of the alloys.

[0127] Table 3: Composition of As-Build Ti-AI-0.025g / min Si-x(Mo+V) Alloys Fabricated on the Heating Platform

[0128] Heat treated Ti-AI-Si-x(Mo+V) alloy samples

[0129] Figure 10 shows the SEM images of the heat-treated Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy fabricated on the heating platform; while Figure 11 is the SEM images of the heat-treated Ti-AI-Si-0.05 g / min Mo + 0.50 g / min V alloy fabricated on the heating platform. The phases present in the heat-treated Ti-AI-Si-x(Mo+V) alloy samples fabricated on the heating platform are a2+y / y, P and phases. The phases identified in all the alloys being developed were all identical but with varying proportions of those phases in their microstructures. It was noticed that the heat treatment technique adopted promotes microstructural refinement because of the presence of Si. This can be easily noticed for all heat-treated samples especially all the 1400°C / 60 mins / FC(furnace cooled) as compared to 1200°C / 60 mins / FC forming a lower quantity of p-phase.

[0130] Table 4 presents the compositions of heat-treated Ti-AI-Si-x(Mo+V) alloy samples fabricated on the heating platform. The heat-treated samples generally had high Al content than their as-built alloys. The composition of Mo in the heat-treated at 1200°C / 60 mins / FC 0.05 g / min Mo + 0.35 g / min V alloy slightly increases while others remain unchanged with the as-built sample content. Meanwhile, the Al content increases for all the heat-treated samples regardless of the mass flow rate of Mo+V heat treatment temperatures. The composition of V reduces with increased heat treatment temperatures but the Si was relative unchanged irrespective of the heat treatment and mass flow rate of Mo+V. The formation of y is promoted by the Al content thereby restricting p and phases. Likewise, V which is a weak [B-stabilizer also contributes 02 / 7 stability in the microstructure. Thus, the ^-TisSis phase would easily be able to dissolve the p-phase since Mo which is a strong p-stabilizer is present in minor quantities. Consequently, dissolution of the p was achieved in the presence of the ^-TisSis phase.

[0131] Table 4: Composition of Heat-Treated Ti-AI-0.025g / min Si-x(Mo+V) Alloys Fabricated on the Heating Platform It was observed that the as-built sample shows a2+y columnar grains of with suspected unmelted Al, [3-phase and ^-TisSis-phase noticed both at the grain boundaries and within the lamellar grains. Al dissolution was observed to result in the transformation and formation of large-grained lamellar microstructures after heat treatment. The as-built microstructure demonstrates grain disintegration with the lamellae colonies and the grain boundaries interfaces. The grain disintegration appears to reduce after heat treatment but still exist at the grain boundaries interfaces. The grain size of the as-built quintenary Ti-AI-Si-Mo-V alloy was between 13.467 pm to 4.056 pm while the heat-treated samples’ were much bigger with grain sizes of about 80 pm and about 100 pm for the 1200°C / 60 mins / FC and 1400°C / 60 mins / FC heat-treated samples, respectively. This is credited to the heat treatment temperature regime adopted leading to grain growth and lamellar coarsening. The 1200°C / 60 mins / FC heat-treated sample reveals a DP microstructure. But the 1400°C / 60 mins / FC heat-treated sample indicates a FL microstructure with a precipitate of suspected [3-phase and ^-TisSis particles that were anticipated but not visibly noticed. The as-built microstructure appears to be inhomogeneous with grains of both coarse and fine irregular lamellar structures.

[0132] Heat treating at 1200°C tremendously decreases the amount of [3-phase within the microstructure. However, heat treating at 1400°C led to lamellar coarsening and grain growth. The heat treatment temperature was noted would not affect the untransformed y-grains that tend to be spheriodized. These are positioned within the matrix of the lamellae grains and at the grain boundaries. Consequently, it was inferred that a-grain growth was not restricted by y-grains during heat treatments. Due to relatively heavy metal like Mo, p phase morphology was brighter as observed from the SEM images. The 1400°C / 60 mins / FC heat-treated sample showed very large lamellar colonies displaying characteristic orientations of lamellae. The grain boundary disintegration of the lamellar grains is believed to have occurred through twinning or recrystallisation of a-grains. Thus, it is understood that the 1400°C / 60 mins / FC heat-treated sample display ordered lamellae but the alloy appeared deformed with discontinuous coarsening close to the Q2-laths. The fine columnar microstructure observed in the as-built alloy is different from the NG (near gamma) equiaxed microstructure generally anticipated for most TiAl alloys produce via AM technologies. The suspected phases of p and ^-TisSis seem to be located at the lamellae interfaces which suggest the limitation of dislocation observed at the grain boundaries. The kinetics of transformation was influenced by the Si through the lamellar spacing refinement of the c^Zy colonies during heat treatment. Thus, ^-TisSis precipitation via Si addition leads to interface stability and reduction of dislocation movement. However, Si pileup would cause vacancy initiations within the lamellar interface thereby reversing its positive effects of microstructural stability.

[0133] The Ti-AI-Si-Mo-V alloy after heat treatment showed a decline in dislocation motion which is expected to result in microstructural stability against thermal degradation and creep improvement. Since Mo is a very strong p-stabilizer, the influence of Mo to stabilize the p-phase was much reduced for this alloy as Mo was about 1.0 at.% of the alloy composition. Thus, limiting the quantity of p / po phase in the overall alloy. However, the y+c / y lamellae colonies were stabilized by V which was about 6.0 at.% of the alloy composition. It was noticed that the y-phase grew steering coarse lamellar grains formation due to slow cooling during heat treatment and fast growth rate accompanying high-temperature solidification reactions. Moreover, the a-phase was nucleated via peritectic reaction which is expected to lead to the diverse crystallographic orientation of the a-phase. Since, the quintenary Ti-AI-Si-Mo-V alloy contains both p and y stabilizers, during the phase transformations the a-phase field is expected to have a a+p+y+ phase transformation system. However, due to the strong ^-stabilizing nature of Mo, the eutectoid line transforms to a Po+a+a2+y+^ phase region, thus producing two new eutectoid points. A minor but important phase that cannot be visualized via SEM is the coo-phase due to its extremely fine nature. This phase has been transformed from the |3o(co) phase and observed within the Po- phase of lamellar colonies. It is significant for this type of material because the alloys’ mechanical properties are exceptionally sensitive to the microstructure.

[0134] Phase analysis

[0135] The phase identification analysis through EBSD and XRD was only performed on the alloys of interest which were the as-built and 1400°C / 60 mins / FC heat-treated alloys. Figure 12 shows the EBSD analysis of as-built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy; while Table 5 and Table 6 presents the phase statistics and grain statistics of the as-built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy. As observed from the phase map analysis (Figure 12c), the EBSD shows the presence of 02-TisAI, y-TiAl, a-TiAh and confirms the presence of suspected phases of Po-TiAl and ^-TisSis. The phase was dominated by the y phase which is 79.97% volume fraction (Table 6) of the phase presence for the as-built quintenary alloy; while the a and 02 phases were 6.69% and 7.80% volume fraction of the phases present, respectively. The ordered o-phase was 1.83% volume fraction of the phase but the disordered BCC p phase was not detected. This further proves that Mo stabilizes the Po instead of the disordered p phase. Moreover, the amount of the (Bo formed was small which was attributed to about 1.0 at.% of Mo in the composition of the alloy. However, the volume fraction of ^-TisSis phase present is 3.34% which contributes to dissolving some amount of the Po-phase present in the alloy. The metastable a-TiA detected indicates the formation of supersaturated a owing to the y phase being the dominating phase in the microstructure.

[0136] Table 5: Phase Statistics of As-Built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy

[0137] Table 6: Grain Statistics of As-Built Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy

[0138] This was credited to the at.% of Al which favors y-phase formation and the characteristics of V to promote (a2+y) / y lamellae. Figure 13 shows the EBSD analysis of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy heat-treated at 1400 °C / 60 mins / FC; while Table 7 and Table 8 presents the phase statistics and grain statistics of the heat-treated Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy. As observed from the phase map analysis (Figure 13c), the EBSD shows the presence of c^-TisAI, y-TiAl, a-TiAh and confirms the presence of suspected phases of Po-TiAl and ^-TisSis. The dominate of the y-phase has been reduced due to heat treatment but the volume fraction is 53.22% (Table 7) of the phase presence after heat treatment. This was suggested to be attributed to phase transformation reactions occurring at a temperature above the a transus temperature and Al evaporation at 1400 °C for a prolonged amount of time. However, the a and 02 phase volume fractions increase to 23.45% and 15.97%, respectively. The amount of the ordered |3o phase volume fraction was also decreased to 1.21 % of the phase while the disordered BCC p phase could not still be detected. Thus, confirming Mo propensity to stabilize the |3o instead of the disordered p phase. Conversely, the ^-TisSis phase volume fraction increases to 6.14% which was almost double the amount present in the as-built sample. The increased amount of the ^-TisSis phase could be attributed to the precipitation of more ^-TisSis phase that dissolves some portion of the po-phase in the present in the a-phases that promote ^-TisSis phase formation. The metastable a-TiAh was still detected at the (a2+y) / y lamellae colonies.

[0139] Table 7: Phase Statistics of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy Heat Treated at 1400°C / 60 mins / FC

[0140] Table 8: Grain Statistics Table of Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V Alloy Heat Treated at 1400°C / 60 mins / FC

[0141] Based on the FSD image in Figure 13a, the microstructure after heat treatments could be said to be homogenous when compared to the as-built microstructure. Also, the heat-treated sample exhibited refined microstructural effects which display almost uniform orientation distribution of y lamellae and a / a2-phase that were nucleated from the y-phase. The refining of the lamellae was achieved through combined effects of dislocations motions, ^-TisSis particles and Al-supersaturated 02- TisAI grains. The volume fraction of p / po-phase was reduced via heat treatment because Mo strongly favors Po-phase formation which could hinder the mechanical properties of the alloy.

[0142] Figure 14 shows the EBSD analysis within a grain boundary region of the Ti-AI-Si- 0.05 g / min Mo + 0.35 g / min V alloy heat-treated at 1400 °C / 60 mins / FC while the phase statistics are presented in Table 9. The interface shows that the a and ^-TisSis were precipitated within the O2-phase sites but the Po was scarcely observed. This confirms that a certain amount of the p-phase was dissolved by -TisS is within the cophase which serves as a nucleation site for the precipitation phases in this alloy. The phase volume fraction of the Po phase at the grain boundary is lesser (1.19%) even though the volume fraction of the O2-phase was higher at this section (29.12%). The Po-phase in this situation was reduced through the growth of y grain from y-laths within colonies of the lamellae. The initiation of y coarsening within the lamellae without nucleation was propelled by structural and compositional stabilization of the metastable B2 p-phase structure. The substitution of Al by Mo causes a slight decrease in lattice parameter of po but this phase is not mechanically stable except Mo substitutes up to 4 at.% Al. When the p-phase becomes stable it limits the coarsening of a-grains by encircling a, therefore, restraining the grain boundaries mobility. From literatures, it is understood that the addition of alloying elements such as W, Ta, Nb and Mo moves the boundary of a2+y / y towards the Al area but Mn, Cr and V cause a shift to the Ti region.

[0143] The avoidance of y-grains recrystallization can be achieved with the a2-laths. In the as-built state, 02 volume fraction is 7.80% which led to substantial lamellae thinning. The increase in volume fraction to 15.97% leads to a corresponding decrease in the amount of y-y interfaces due to extensive dissolution and transformation of two adjacent y-laths. The presence of coarse ^-TisSis particles suggests liquid state transformation in the precipitation of the phase which would subsequently diminish the tensile properties of the alloy. Si promotes 02-laths dissolution invariably increasing y-phase recrystallization. This is attributed to the stronger ability of Si to create vacant sites in the alloy matrix. However, (j-TisSis nucleation is heterogeneous within the y-lath which characteristically ensues at dislocation sites. Thus, indicates that j-TisS is particles favorably precipitates at y-laths lamellae stacking faults.

[0144] Table 9: Phase Statistics of the Selected Grain Boundary Area of Ti-AI-Si-0.08 g / min Mo + 0.35 g / min V Alloy Heat Treated at 1400°C / 60 mins / FC

[0145] The XRD pattern in Figure 15 confirms the phases present in the microstructure. Figure 15 shows the XRD pattern of the Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V fabricated via the modified LENS machine. The XRD results show a pattern with several peaks a2-TisAI and y-TiAl with few minor peaks of Po-TiAl and -TisSis were detected. The Po-phase was not noticeable for the as-built sample but the 1400 °C / 60 mins / FC samples display more visibly and new XRD peaks which is an indication that phase transformation occurred during the heat treatment. Studies about interactions of elemental powder produce metastable phases such as TiAh and TiAl2 owing to an occurrence known as Kirkendall porosity. This could lead to p phase formation that is detrimental to the mechanical properties of TiAI-based alloys at high temperatures. The appearance of high peaks of a2+y at 20 between 40° to 45° shows the transformation of lamellar solid solution. The peaks of the a2+y for the heat-treated sample was higher indication more lamellar formation.

[0146] However, twinning occurred in both as-built and the heat-treated sample between 20 equals 40° to 55° but the twinning was less obvious for the heat-treated sample. This is indicative of stress relaxation in the heat-treated sample. XRD sensed a relational rise in y-phase peak intensities while 02 peaks were dissolved and the twin y peaks somewhat reduced. It is generally known that a rise in Al content decreases stacking fault energy. Since there was a reduction in Al content it can be interpreted that the stacking fault energy of the 1400°C / 60 mins / FC heat-treated sample increases. This further explains the dilation of the grain boundaries noticed in the SEM image of the 1400°C / 60 mins / FC heat-treated sample. It is worth mentioning that slip motions in TiAI-based alloys vary depending on Al content, thus, influencing the deformation and twinning. This shows that the heat-treated sample would exhibit little or no twinning because of reduced Al content. Moreover, materials that induce twinning possess improved resistance to impact and high formability.

[0147] Microhardness

[0148] As-build Ti-AI-Si-x(Mo+V) alloy samples

[0149] Figure 16 shows the microhardness of the as-built Ti-AI-Si-x(Mo+V) alloy samples fabricated on the heating platform with the addition of 0.05 g / min Mo + 0.50 g / min V and 0.05 g / min Mo + 0.35 g / min V. It is observed that additions of Mo+V do not show many changes in the microhardness of the TiAI-based alloys. This is evident in the average microhardness values of 514 Hvo.s and 526 Hvo.s for alloys produce with 0.05 g / min Mo + 0.35 g / min V and 0.05 g / min Mo + 0.50 g / min V. This gives a corresponding YS of 1681 MPa and 1720 MPa, respectively. This indicates that additions Mo+V could produce a TiAI-based alloy with better mechanical properties than separate additions of Mo and V. Nevertheless, it is understood that Si causes an increase in hardness by producing ^-TisSis. Thus, the need to balance of the amount phases formed in the microstructure and composition to obtain better TiAI- based alloys characteristics.

[0150] Heat treated Ti-AI-Si-x(Mo+V) alloy samples

[0151] Figure 17 presents the microhardness of heat-treated Ti-AI-Si-x(Mo+V) alloy samples fabricated on the heating platform with the addition of 0.05 g / min Mo + 0.35 g / min V and 0.05 g / min Mo + 0.50 g / min V. The microhardness values generally reduce for heat-treated sample regardless of the mass flow rate of Mo+V added. The microhardness values were generally slightly lesser than the other alloys fabricated on the heating platform. This was attributed to very low p and formation including high Al content which promotes the formation of more y phases. Since the y phase has the lowest hardness value among the phases present within the microstructure. Generally, the microhardness values of the heat-treated Ti-AI-Si-x(Mo+V) alloys fabricated on the heating platform shows very little disparity hardness values regardless of heat treatment temperatures. The average microhardness values for the 0.05 g / min Mo + 0.35 g / min V are 487 Hvo.s and 442 Hvo.s while the 0.05 g / min Mo + 0.50 g / min V gave 513 Hvo.s and 484 Hvo.s for heat treatment performed at 1200 °C / 60 mins / FC and 1400 °C / 60 mins / FC, respectively. This corresponds to YS of 1592 MPa and 1445 MPa for 0.05 g / min Mo + 0.35 g / min V and YS of 1678 MPa and 1583 MPa for 0.05 g / min Mo + 0.50 g / min V. This shows that Mo+V additions and heat treatment reduce the microhardness values of the TiAI-based alloys. Also, it was understood that improved mechanical properties are achievable through minor additions of V, Mo and Si.

[0152] Mechanical Properties from Nano-Indentation Results

[0153] The instrumented indentation testing technique also called nanoindentation testing or depth-sensing indentation was adopted. After analyzing the load-displacement curves, the mechanical properties were determined. The reported mechanical properties in this work are Young’s Modulus (E), yield strength (YS) and ultimate tensile strength (UTS). Figure 18 shows the load-displacement curves of the as-built and heat-treated Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy; while Figures 19, 20 and 21 shows the true stress-true strain curves of the as-built and heat-treated (1200 °C and 1400 °C, respectively) Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy. The loaddisplacement curve was analyzed using the Oliver and Pharr method and the mechanical properties were determined as discussed earlier. The load-displacement curves indicate materials behaviour. All the samples exhibited elastic-plastic behaviour. But the as-built sample demonstrated more plastic deformation than the heat-treated samples. The unloading section which describes the stiffness of the alloys suggests that the material stiffness was relatively the same. The as-built sample has a stiffness of 1.03809 x106N / m, the 1200°C / 60 mins / FC sample recorded a stiffness value of 0.93312 x 106N / m while the 1400°C / 60 mins / FC sample stiffness is 0.94432 x 106N / m.

[0154] Figures 19, 20 and 21 show the true stress-true strain curves with inserts of regions within the curve for the as-built, 1200°C / 60 mins / FC and 1400°C / 60 mins / FC Ti-AI- Si-0.05 g / min Mo + 0.35 g / min V alloy, respectively. The UTS and YS value of the as-built alloy is 450 MPa and 325 MPa, respectively, while the E is 155 GPa. The E value of the as-built sample was about 7.7% lesser than the E value of GE commercial alloy (I.E., GE-4822 alloy) with 168 ±2 GPa at RT (Room Temperature). But the YS in comparison to GE which is 326 MPa was practically the same (0.3% < GE). However, the GE alloy UTS of 422 MPa indicates that the as-built sample was about 6.6% greater, thus, withstanding higher stress at RT.

[0155] However, the UTS and YS value of the 1200°C / 60 mins / FC heat-treated sample is 586 MPa and 438 MPa, respectively, while the E is 191 GPa. The E value of the 1200°C / 60 mins / FC sample was about 13.7% greater than the E value of GE commercial alloy with 168 ±2 GPa at RT. Also, the YS in comparison to GE which is 326 MPa was about 34% greater, indicating the 1200°C / 60 mins / FC sample would withstand higher stress than GE before experiencing plastic deformation. Also, the GE UTS of 422 MPa indicates that the 1200°C / 60 mins / FC sample was about 38.9% greater. Therefore, the 1200°C / 60 mins / FC sample would be able to undergo more plastic deformation and would withstand higher stress at RT before fracture. The UTS and YS value of the 1400°C / 60 mins / FC heat-treated sample is 630 MPa and 368 MPa, respectively, while the E is 179 GPa.

[0156] Moreover, the E value of the 1400°C / 60 mins / FC sample was about 6.5% greater than the E value of GE commercial alloy with 168 ±2 GPa at RT. Likewise, the YS in comparison to GE which is 326 MPa was about 13% greater, indicating the 1400°C / 60 mins / FC sample would withstand higher stress than GE before experiencing plastic deformation. Also, the GE UTS of 422 MPa indicates that the 1400°C / 60 mins / FC sample was about 49.3% greater. Therefore, the 1400°C / 60 mins / FC sample would be able to undergo more plastic deformation than any of the other samples and would withstand the highest stress at RT before fracture. Although the 1400°C / 60 mins / FC sample would experience plastic deformation at lower stress compared to the 1200°C / 60 mins / FC sample, the UTS value of 630 MPa indicates that it has better toughness than the as-built sample, 1200°C / 60 mins / FC sample and commercially available GE alloy. Since the microstructure of the 1400°C / 60 mins / FC alloy sample displays FL microstructure, the indication of higher fracture toughness is in tandem with what is known in the art. Also, the DP microstructure of the 1200°C / 60 mins / FC explains high YS due to the presence of O2-phase that is difficult to deform and lesser y-phase with higher ductile performance. Generally, lamellar microstructures exhibit better reasonable mechanical properties but the mean c / y colonies spacing lamellar has a huge influence on the properties. Moreover, finer lamellae microstructures would give better strength, creep resistance and higher fracture toughness. In the design of microstructures, the heat treatment strategy could be modified to achieve a balance y:a2:Po / B2 phase in order to eliminate and / or suppress detrimental brittle phases. In TiAI-based alloys, brittle cracking originates from grain boundaries of lamellar colonies, thus, causing largegrained microstructures to be susceptible to cracking. The nonexistence of twinning can be connected to a lack of ductile properties. However, in the present invention, the heat-treated samples showing well-defined twinning in the y-phase experienced more plastic deformation with higher ductility than the ones with lower twinning.

[0157] Tribological and Thermogravimetric Behaviour

[0158] To ascertain the durability of TiAI-based alloy components in service, evaluating the tribological behaviour and high-temperature stability is very crucial. This represents a critical issue for aero-engine applications as the turbocharger operating temperature range is between 700-950°C. An important characteristic of high-temperature TiAI- based alloys is long-term stability in the presence of air where the microstructure and the mechanical properties need to be sustained throughout the component’s lifespan. Figure 22 presents the coefficient of friction (CoF) and wear rate of the as- built and heat-treated Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy. As observed from Figure 22, the CoF generally reduces from the as-built sample (0.586) to the 1200°C / 60 mins / FC sample (0.498) and the 1400°C / 60 mins / FC sample (0.492) having the lowest value of CoF. However, the wear rate does not follow the same pattern.

[0159] The sample heat-treated at 1200°C / 60 mins / FC showed the highest wear rate value of 1.645 x 10'4mm3 / N / m, followed by the as-built sample with a value of 4.266 x 10'5mm3 / N / m while the sample heat-treated at 1400°C / 60 mins / FC had the lowest wear rate of 2.88 x 10'6mm3 / N / m. The tribological behaviour of the quintenary Ti-AI-Si-Mo- V alloy is highly unusual because the high wear rate recorded by the 1200°C / 60 mins / FC sample was not expected. However, the high rate to wear is indicative of more material removal during the sliding test. Moreover, the as-built sample had a higher wear rate than the 1400°C / 60 mins / FC. This was expected due to the FL microstructure in the 1400°C / 60 mins / FC sample while the as-built sample exhibited more precipitates of hard phases like 02, po and ^-TisSis.

[0160] Figure 23 shows the TGA Curves of the as-built and heat-treated Ti-AI-Si-0.05 g / min Mo + 0.35 g / min V alloy. The results of the TGA presents the change in mass against temperature relative to the initial mass of the sample per unit area. Positive mass change means that the corrosion products formed are solid and remained attached to the surface of the samples. This mass gain depends on the chemical bonding formed with air (oxygen and nitrogen). In situations where spalling of the solid scale occurs or corrosion products are volatile, negative mass change is recorded which indicted that alloy components are lost. It was observed that mass loss for the as-built sample was in the negative direction with a slight reduction in mass at 130°C. The cumulative mass loss continued to reduce up to 800°C which results in a slight rise in mass in the positive direction. However, the mass gained at that temperature could not balance the initial material loss. The heat-treated sample demonstrated better high-temperature stability but most of the mass loss was also in the negative direction indicating that the formation of corrosion products leads to alloy components loss or the layers formed on the surface of the samples are volatile.

[0161] A slight drop in mass was noticed for the 1200°C / 60 mins / FC sample at 292°C while the 1400°C / 60 mins / FC sample also experienced a drastic decrease in mass at 600°C. Generally, all the samples displayed positive mass changes at temperatures above 800°C, denoting that the oxide scales are being attached to the samples. Thus, the alloys which demonstrate less spalling would be more suitable to turbocharger wheel applications because the operating temperature is between 850°C to 900°C. It can be problematic to predict the oxidation behaviour of TiAI- based alloys having varying microstructures and compositions. Also, oxidation responses are determined by the experimental set-up, temperature, type of atmosphere used and oxygen impurity initially within the alloy. The presence of Si is reported to minimize the mechanisms of oxidation due to the barrier effects induced via silica and silicide. But layer exfoliation by Si addition leads to a deteriorating effect on oxidation. Mo enriches the TiAl alloy layer to form a metals / oxide interface, thereby decreasing the oxygen diffusion and Ti into the bulk material. The benefits of Mo are ascribed to its capability to increase alumina at the outer section of the oxide layer and reducing simultaneously the solubility of oxygen into the whole material. Oxidation resistance still presents a strong restriction for extensive use of TiAl alloys in automotive and aerospace fields. The oxidation resistance is mainly dependent on the Al content because it influences the ratio of titanium oxides and alumina inside the oxide scales based on the concentration from the built material. The alumina layer demonstrates passivation behaviour and grows slowly compared to the formation of titanium oxides. The Al content also determines the ratio of the 02 and y phases within the microstructure which consequently influences the resistance to oxidation because Al distribution in y-TiAl is faster than a2-TisAI.

[0162] Conclusions

[0163] In the invention, TiAI-based alloys were successfully fabricated via in-situ alloying by means of DED processing with Si, Mo and V as alloying elements followed by heat treatment at between about 1200 °C to 1400 °C for 60 minutes then furnace cooling.

[0164] Based on the observed experiment results and corresponding analyses, several conclusions are summarized as follows that: The quintenary Ti-AI-Si-Mo-V alloy demonstrates 02+y columnar grains for as- built samples with p-phase and ^-TisSis-phase noticed both at the grain boundaries and within the lamellar grains. The EBSD analysis of the Ti-AI-Si-Mo-V identified the same phases as the quaternary alloys but the amount of the Po formed was small which was attributed to about 1 .0 at.% of Mo in the composition of the as-built alloy. In the XRD pattern of the as-built Ti-AI-Si-Mo-V alloy Po-phases was not noticeable but the heat-treated sample displays visibly new XRD peaks which is an indication that phase transformation occurred during the heat treatment. The quintenary Ti-AI-Si-Mo-V alloy UTS value of 630 MPa for the 1400°C sample indicates that it has better toughness than the as-built, 1200°C sample as well as the commercially available GE alloy. The tribological behaviour of the quintenary Ti-AI-Si-Mo-V alloy showed that the as-built sample had a higher wear rate than the 1400°C heat-treated sample. Generally, all the Ti-AI-Si-Mo-V alloy samples displayed positive mass changes at temperatures above 800°C, denoting that the oxide scales are being attached to the samples. Thus, the alloys which demonstrate less spalling would be more suitable for turbocharger wheel applications because the operating temperature is between 850°C to 900°C.

[0165] While the invention has been described in detail with respect to a specific embodiment and / or example thereof, it will be appreciated that those skilled in the art, upon attaining an understanding of the foregoing may readily conceive of alterations to, variations of and equivalents to these embodiments.

Claims

CLAIMS1. A method of producing a titanium aluminide-based alloy by in-situ direct energy deposition, the method comprising:(i) depositing on a substrate, separate feedstocks one each essentially consisting of aluminium, titanium, a grain refining agent, a grain stabilizing agent, and optionally another grain refining or stabilizing agent;(ii) using a focused energy source to melt the separate feedstocks as they are deposited on the substrate, thus melting the feedstocks and forming a melt pool of the alloy;(iii) cooling the melt pool to form a cooled layer of the alloy; and(iv) optionally repeating steps (i) to (iii) to build successive layers of the alloy.

2. The method according to claim 1 , comprising simultaneously depositing at least four separate feedstocks, preferably five separate feedstocks.

3. The method according to claim 1 , comprising selectively depositing at least four separate feedstocks, preferably at least five separate feedstocks.

4. The method according to any one of claims 1 to 3, wherein the feedstocks are selected from a group comprising of powders, wires, and a combination thereof.

5. The method according to claim 4, wherein the separate feedstocks comprising at least five separate powders one each essentially consisting of titanium, aluminium, the grain refining agent, the grain stabilizing agent, and another grain stabilizing agent.

6. The method according to claim 5, wherein the grain refining agent is silicon.

7. The method according to claim 5 or claim 6, wherein the grain stabilizing agent is molybdenum.

8. The method according to any one of claims 5 to 7, wherein the another grain stabilizing agent is vanadium.

9. The method according to any one of claims 1 to 8, wherein the substrate is heated and maintained substantially at a constant heated temperature by a second energy source during steps (i) to (iv), wherein the second energy source is an exogenous energy source other than the focused energy source.

10. The method according to claim 9, wherein the second energy source is a heating means for heating the substrate and maintaining the substrate at a constant temperature.11 . The method according to claim 9 or claim 10, wherein the substrate is heated and maintained at a temperature that is arranged to prevent cracking of the formed titanium-aluminide based alloy.

12. The method according to any one of claims 9 to 11 , wherein the substrate is heated and maintained at a constant temperature of about 800 degrees Celsius.

13. The method according to any one of claims 1 to 12, wherein the substrate is a titanium-based substrate.

14. The method according to claim 13, wherein the titanium-based substrate is a Ti6AIV4 alloy.

15. The method according to any one of claims 1 to 14, comprising subjecting the titanium-aluminide based alloy to heat treatment, and cooling the heat treated titanium-aluminide based alloy.

16. The method according to 15, wherein the heat treatment is effected in an inert atmosphere at a temperature of between about 1200 and about 1400 degrees Celsius.

17. The method according to claim 15 or claim 16, wherein the cooling of the heat treated titanium-aluminide based alloy, occurs after subjecting titanium-aluminide based alloy to heat treatment, and the cooling of the heat treated titanium-aluminide based alloy is effected at a rate of between 20 and 25 degrees Celsius / m inute.

18. The method according to any one of claims 15 to 17, wherein the heat treatment and cooling are effected in a furnace.

19. The method according to any one of claims 15 to 18, wherein the ultimate tensile strength of the titanium-aluminide based alloy, after being subjected to heat treatment and cooling, is between 500 and 650 MPa, preferably between 550 and 650 MPa, more preferably between 586 and 630 MPa.

20. The method according to any one of claims 15 to 19, wherein the yield strength of the titanium-aluminide based alloy, after being subjected to heat treatment and cooling, is between 300 and 500 MPa, preferably between 350 and 450 MPa, more preferably between 368 and 438 MPa.

21. The method according to any one of claims 15 to 20, wherein the titanium- aluminide based alloy, after being subjected to heat treatment and cooling, comprising:from 46 to 50 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1.0 to 1.5% molybdenum; from 7 to 10 at.% vanadium; and balance is titanium.

22. The method according to any one of claims 15 to 21 , wherein the titaniumaluminide based alloy, after being subjected to heat treatment and cooling, has a fully lamellar microstructure.

23. The method according to any one of claims 15 to 21 , wherein the titaniumaluminide based alloy, after being subjected to heat treatment and cooling, has a duplex microstructure comprising lamellar and columnar grains.

24. The method according to any one of claims 1 to 23, wherein the focused energy source is selected from a group comprising of a laser beam, electron beam, and plasma arc, in particular a laser beam.

25. A titanium-aluminide based alloy produced in accordance with the method of any one of claims 1 to 24.

26. A titanium-aluminide based alloy produced by in-situ direct energy deposition of distinct powders one each essentially consisting of titanium, aluminium, vanadium, silicon, and molybdenum, the titanium-aluminide based alloy comprising: from 44 to 48 at.% aluminium;from 0.9 to 1 .0 at.% silicon; from 1 .0 to 1 .5 at.% molybdenum; from 6 to 10 at.% vanadium; and balance is titanium.

27. A titanium-aluminide based alloy which has been heat treated and cooled after been produced by direct energy deposition of distinct powders one each essentially consisting of titanium, aluminium, vanadium, silicon, and molybdenum, the titanium-aluminide based alloy comprising: from 46 to 50 at.% aluminium; from 0.9 to 1 .0 at.% silicon; from 1.0 to 1.5% molybdenum; from 7 to 10 at.% vanadium; and balance is titanium.