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Illarionov, A.G.; Stepanov, S.I.; Naschetnikova, I.A.; Popov, A.A.; Soundappan, P.; Raman, K.H.T.; Suwas, S. Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb. Encyclopedia. Available online: https://encyclopedia.pub/entry/41251 (accessed on 19 May 2024).
Illarionov AG, Stepanov SI, Naschetnikova IA, Popov AA, Soundappan P, Raman KHT, et al. Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb. Encyclopedia. Available at: https://encyclopedia.pub/entry/41251. Accessed May 19, 2024.
Illarionov, Anatoliy G., Stepan I. Stepanov, Inna A. Naschetnikova, Artemiy A. Popov, Prasanth Soundappan, K. H. Thulasi Raman, Satyam Suwas. "Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb" Encyclopedia, https://encyclopedia.pub/entry/41251 (accessed May 19, 2024).
Illarionov, A.G., Stepanov, S.I., Naschetnikova, I.A., Popov, A.A., Soundappan, P., Raman, K.H.T., & Suwas, S. (2023, February 15). Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb. In Encyclopedia. https://encyclopedia.pub/entry/41251
Illarionov, Anatoliy G., et al. "Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb." Encyclopedia. Web. 15 February, 2023.
Additive Manufacturing of Intermetallic Alloys Based on Ti2AlNb
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This entry is adapted from the peer-reviewed paper 10.3390/ma16030991

Titanium alloys based on orthorhombic titanium aluminide Ti2AlNb are promising refractory materials for aircraft engine parts in the operating temperature range from 600–700 °C. Parts made of Ti2AlNb-based alloys by traditional technologies, such as casting and metal forming, have not yet found wide application due to the sensitivity of processability and mechanical properties in chemical composition and microstructure compared with commercial solid-solution-based titanium alloys. Metal additive manufacturing (MAM) has attracted the attention of scientists and engineers for the production of intermetallic alloys based on Ti2AlNb. 

Ti2AlNb-based alloy metal additive manufacturing microstructure

1. General Characteristics of Alloys Based on Ti2AlNb: Phases, Phase Diagrams, Alloying Elements, Microstructure and Properties

Titanium alloys containing the O-phase developed based on the Ti-Al-Nb system usually contain 18–30% of Al and 12.5–30% of Nb [1][2]. The equilibrium phase composition of O-phase alloys is based on the isopleth sections of the Ti-Al-Nb phase diagram, which corresponds to constant Al content (22, 23, 25, 27.5 at%) and varied Nb content [3][4][5][6].
According to the diagram, depending on the content of Al and Nb, different phases get stabilized at different temperature regimes. Phases available in different regimes include:
  • − Ordered (B2) and disordered (β) solid solution based on the BBC lattice;
  • − α2-phase based on intermetallic Ti3Al with ordered HPC lattice;
  • − Ordered O-phase-based on Ti2AlNb with orthorhombic lattice, which is divided into two types in [6]: The 1st type (high-temperature O-phase)—Al ordered and Ti and Nb disordered—O1-phase, the 2nd type (low-temperature)—Ti, Al and Nb ordered—O2-phase. Most diagrams have no such division and use the generalized name O-phase.
By analogy with alloys based on titanium solid solutions, alloying elements in alloys based on orthorhombic titanium aluminide can be divided into α-stabilizing elements (Al), β-stabilizing elements (Nb, Mo, V, Ta, Fe, Si and H) and neutral elements (Zr, Hf, B and Y). To estimate the composition of the multicomponent O-phase alloy in the Ti-Al-Nb phase diagrams, the so-called structural equivalents are utilized by analogy with titanium alloys based on solid solutions. The structural Al equivalent ([Al]eq) allows people to estimate the influence of the alloying elements and impurities on the stability of the α(α2) phase relative to the corresponding influence of Al. This equivalent of α-stabilizing elements and neutral elements, introduced by Rosenberg [7], is described by the following ratio:
[Al]eq = %Al + %Sn/3 + %Zr/6 + 10%O (mass.%).
Since these alloys can be doped with other β-stabilizers, the concept of the Nb equivalent ([Nb]eq) can be employed. [Nb]eq can be estimated by two methods. The first method [1] proceeds from the fact that the stabilizing elements in the O-phase occupy the Nb lattice position, and according to this method, 1% of any β-stabilizing element is equivalent to 1% of Nb. The second method [4] considers the effect of β-stabilizers on the stability of the high-temperature β-solid solution in O-phase alloys by analogy with their effect on the stability of the matrix β-phase in solid solution-based titanium alloys during quenching. It takes into account the critical concentrations for binary alloying systems (by analogy with [Mo]eq [8]), above which only the β-solid solution is present and the martensitic transformation in the alloy is suppressed. In [4], the following critical concentrations (wt.%) were considered: 36 Nb; 11 Mo; 15 V; 45 Ta; 22 W; 4.5 Fe; 6.5 Cr; 6.5 Mn; 9.5 Co; and 8.5 Ni. Based on these values for O-alloys, the structural [Nb]eq, which evaluates the effect of each β-stabilizing element in the alloy on the stability of the β-solid solution relative to Nb, can be calculated by the formula [4]:
[Nb]eq = %Nb + %Mo/0.31 + %V/0.42 + %Ta/1.25 + %W/0.61 + %Fe/0.13 + %Cr/0.18 + %Mn/0.18 + %Co/0.26 + %Ni/0.24
The mechanisms and stages of phase transformations (β(B2)-O, α2-O, β(B2)-α2 and β(B2) + α2-O in O-phase alloys at constant temperature (during aging or holding after quenching at annealing temperatures) are addressed in quite a few studies and their detailed analysis is given in reviews [9][10][11][12].
The type of structure formed in O-phase alloys during deformation and heat treatment is determined by the temperature ranges and cooling rates used during processing and can result in single-phase β(B2), two-phase (β(B2) + O) or three-phase (β(B2) + α2 + O) microstructures with different morphologies of O (α2)-precipitates, namely lamellar, equiaxed and globular-lamellar (duplex) structures and also different levels of grain refinement. Obtaining different structural phase states is possible by varying the temperature rate parameters of heat treatment and deformation processing.
According to [3][13][14][15][16][17][18][19][20][21][22][23][24], a favorable effect on ductility and impact toughness of O-phase alloys results from an increase in the volume fraction and refinement of β(B2)-grain size, the absence of grain boundary α2-precipitates and thickening of O-phase plates. An increase of secondary O-phase volume fraction along with a reduction of α2-phase and the formation of refined O-phase platelets enhance strength characteristics.
The influence of chemical composition, deformation processing of O-phase alloys with various types of structure and range of mechanical properties is discussed in detail in [25]. It was shown that in O-phase alloys, such as Ti-22Al-25Nb, after isothermal forging and subsequent strengthening heat treatment (quenching and aging), the following range of properties can be achieved: UTS = 1040–1172 MPa, YS = 875–1065 MPa; EL = 6.5–14.8; RA = 11–15.3%.

2. Classification of the MAM Processes Used for Ti2AlNb Based Alloys

2.1. Powder Bed Fusion (PBF)

In powder bed fusion additive manufacturing, powder is applied with a recoater to the platform layer by layer. A thermal energy source, such as a laser or electron beam, fuses the powder before the next layer is applied. Laser powder bed fusion (L-PBF) manufacturing occurs in controlled environment atmospheres of argon or nitrogen with a deposition rate in the range from 2–20 sm3/h. The surface roughness of L-PBF parts lies in the range from 10–20 μm [26]. Optionally, preheating the base platform up to 1000 °C can be applied [27].
Electron beam powder bed fusion (EB-PBF) takes place under controlled vacuum conditions. The electron beam is focused and deflected by means of an electromagnetic lenses, instead of mirrors as in L-PBF. The deposition rates typical of EB-PBF correspond to the range from 55–80 sm3/h and provide surface roughness up to the values in the range from 15–30 μm [26]. Pre-heating each layer by scattered electron beam to temperatures about 0.6 of the melting temperature is required to avoid so-called ‘smoking’ [28].
There are some disadvantages to the powder bed fusion method related to the efficiency of the process both for the transformation power from electricity to laser and for the production output rate. The main advantages are the good resolution, reduced material wastage and efficient recycling of the un-melted powder.

2.2. Direct Energy Deposition (DED)

In direct energy deposition or direct metal deposition (DMD), a four- or five-axis arm moves around, depositing melted material around a fixed object. An electron beam, wire arc or laser melts the powder or wire feedstock adding the material to the substrate. Simultaneous deposition and melting during the construction of the parts should be provided. In the case of metal, a powder is coaxially fed through the laser head and provides a much better finish than wire. In contrast, an independent feeding system in the case of wire is employed. The DED deposition rate lies in the range from 20–160 sm3/h, and surface roughness is typically above 30 μm [26]. The desired effect with the wire can be achieved through post-processing. The problem of significant thermal stresses arises due to the high deposition rates and large melt pools.
In laser direct energy deposition (L-DED), the shielding gas’ flow head protects the melt pool from oxidation and carries the powder to the melt pool. The electron beam direct energy deposition (EB-DED) works exclusively with wire as a feedstock due to poor powder flow in a vacuum.
WAAM (wire arc additive manufacturing) combines principles similar to arc welding processes and the digital control of electric arc displacement. WAAM-based processes are more efficient than L-DED due to the larger energy requirements. Machining to obtain the final parts produced by EB-DED and WAAM is assumed; therefore, the surface roughness is irrelevant.
GMA-DED is the simplest, cheapest and most widely spread process to implement due to its direct wire-feeding, which is coaxial with the nozzle of the welding torch. In gas tungsten arc direct energy deposition (GTA-DED) [29] and plasma arc direct energy deposition (PA-DED) [30], non-consumable electrodes generally made of tungsten are used. The feedstock wire is supplied through an additional wire-feeding unit. PA-DED is the only WAAM process that can use powder along with the wire as feedstock. It is characterized by higher deposition rates, greater energy concentration, better stability and less thermal distortion than in GTA-DED.

2.3. Binder Jetting (BJ)

The powder-based material is applied to the build platform with a recoater and then the print head selectively deposits the liquid adhesive on top, which both adheres the powder particles in the same layer and between the adjacent layers together. Following a layer, the product is lowered on the platform repeatedly to create more layers until the ‘green part’ with low strength and approximately 60% relative density is finished. Various materials can be utilized including polymers, ceramics and metals. The green part is then heated in a controlled atmosphere to remove the adhesive and to sinter the individual particles into a fully dense metal part.
The main drawbacks of binder jetting are the increase in post-processing time and it may not be the best choice for creating structural parts. The increase in density during sintering is obtained via shrinkage and loss of dimensional precision [31].
The primary advantage of binder jetting is that metal powder can be easily reinforced with ceramics as the process occurs at room temperature, which provides minimal distortion of parts associated with thermal effects.

2.4. Sheet Lamination (SL)

Sheet lamination with respect to metals is a process that binds layers using adhesive bonding, ultrasonic welding or friction stir welding [32][33] to form a single piece that is subsequently machined into the required part. However, sheet lamination has not found a broad application in the production of structural parts made of titanium alloys based on Ti2AlNb.

2.5. Hybrid Additive Manufacturing

In contrast, hybrid additive manufacturing (HAM) gains increasing interest in the production of metal parts with improved capabilities compared with standard DED and PBF processes. In [34], HAM is defined as a combination of two or more established manufacturing processes into a new, combined set-up. In [26], the HAM is divided into two groups.
The first group contains processes where two or more combined energy sources/tools result in a synergetic effect in the processing zone. The second group accounts for the processes, where the synergetic effects are obtained by a controlled combination of processes acting separately in order to fabricate parts in a more efficient and productive way, e.g., combination of DED and point forging.
First, relatively high energy is required for melting. Secondly, the low ductility of these materials can lead to the appearance of cracks caused by thermal stresses at high cooling rates [35]. Moreover, it leads to a low processability of the alloys at the ambient temperatures. Thirdly, Ti alloys are susceptible to contamination, when heated in oxygen atmosphere. Thus, the use of a protective atmosphere or vacuum is required for fusion.
DED and PBF are mainly used as they provide an optimal combination of mechanical properties with high manufacturability, i.e., capacity and minimal tolerance for machining. However, the process parameters vary in a wide range. Different supply schemes, rates and geometries of feedstock materials are utilized in MAM processes as discussed in the previous section. Finally, the substrate materials and the pre-heat conditions also vary widely, from pure Ti at room temperature to Mo heated at 1000 °C.

3. The Influence of MAM Process Parameters on Structure, Phase Composition and Mechanical Properties of the Alloys Based on Ti2AlNb

3.1. The Influence of MAM Process Parameters on Density and Microhardness

Comprehensive research data showed the influence of AM processing parameters on the properties of as-built parts made of alloys based on orthorhombic titanium aluminides: density, hardness, elastic modulus, mechanical properties at room and elevated temperatures (650, 750 °C) and, in particular, the effect of volumetric energy density—VED = P/(S⋅H⋅D⋅L) [J/mm3], where laser power (P), scanning speed (S), hatch distance (HD) and layer thickness (L). Hatch distance and platform pre-heating temperature on density of the AM parts was evaluated for L-PBF.
According to previous research, the dependence of density on VED [36][37], the maximum density (above 99.5%) is typical of alloys produced with VED of about 60 ± 10 J/mm3. This is explained in [37] by the fact that at a lower VED, the decrease in density and increase in porosity is associated with incomplete melting of the Nb powder due to insufficient energy input in L-PBF of a mixture of elemental Ti, Al and Nb powders. A similar effect was observed earlier in [38] during AlSi10Mg alloy manufacturing. A high VED resulted in overheating of the melt pool and evaporation of Al, which is in accordance with the previous studies of the L-PBF of TiAl [39]. When using the powder based on orthorhombic titanium aluminides [36], the density decrease with decrease in VED is associated with an increase in the viscosity of the melt pool, which hinders the escape of bubbles from the alloy in the liquid state. Increased VED leads to spattering and partial evaporation of the melt pool, which results in an increase in the fraction of pores [40][41]. Note that heating the platform to 200 °C [36][42] and 600–700 °C [43] provides the highest relative density of the fused material.
Based on the analyzed data, it is still difficult to determine the optimal pre-heating temperature. However, the cracks can occur at pre-heating temperatures of 200 (MAPS), 200 and 500 °C (GA) in alloys manufactured from GA and MAPS powders. Moreover, the density was lower at pre-heating temperatures of 900 (MAPS), 980 °C (GA) [43]. Cracks in this case are associated with residual stresses that occur due to fast heating and cooling rates in MAM processes. The porosity vs. laser power dependence passes through a maximum at a certain energy and is similar to the effect of VED.
Increased microhardness of O-phase alloys is associated with the precipitation of second phases (O, α2, TiC) in the B2/β matrix. This occurs as a result of cyclic thermal [44][45][46][47] or thermomechanical [48] effects during MAM as well as during pre-heating of the platform at 500 °C and higher [43], subsequent heat treatment [37][49][50][51] and addition of SiC [52]. Note that the refinement of the precipitation of the second phases results in higher hardness [37][43].
The maximum microhardness of 570 ± 5 HV was obtained for a sample produced from Ti-24Al-25Nb-1Zr-1.4V-0.6Mo-0.3Si GA powder by L-PBF (at a platform temperature of 600 °C). This ensured 100% decomposition of the B2/β matrix with the formation of refined particles of the O-phase. In [53], the microhardness and elastic modulus of individual phases (O, B2, α2) were determined by nanoindentation and compared with the previous research [54][55][56]. The difference in the values of the microhardness in different studies can be associated with the different measurement conditions (load, holding time), in which different works vary within a wide range—from 0.3 N and 10 s [47] to 300 N and 30 s [48].

3.2. The Influence of MAM Process Parameters and Heat Treatment on Mechanical Properties at Room Temperature

Compressive mechanical properties of the of O-phase alloys were investigated in [45][52]. According to [45], the compressive properties of the Ti-22Al-25Nb alloy samples manufactured by TWPF method depend on the sample height due to differences in the phase-structural state. The lower regions of the sample have lower strength with ductility compared to the upper regions of the sample. This can be explained by the formation of thinner and longer precipitates of O- and α2- phases in the upper regions compared to the bottom [45]. The compressive strength values of the composite obtained from Ti-24Al-25Nb-1Zr-1.4V-0.6Mo-0.3Si alloy powder with 0–15% SiC whiskers lie in the range of 2300 (0% SiC)—2550 MPa (5% SiC) [52]. Further increase in SiC above 5% resulted in the reduction of strength. According to [52], the increase in strength with the addition of SiC whiskers is associated with a refinement of the grain structure of the B2-matrix, due to the inhibition of grain growth by the precipitation of dispersed TiC particles. A further strength decrease is associated with the higher fraction of the B2/β matrix formed as a result of its stabilization by Si due to the dissolution of SiC-whiskers in the B2/β solid solution during melting.
Numerous results on tensile mechanical properties at room temperature have been obtained [2][27][36][42][43][44][45][46][48][50][57][58][59][60][61][62][63] and less research is devoted to elevated temperatures of 650 °C [2][27][60][62][64] and 750 °C [59]. A wide range of values characterizes the strength and ductility of MAM O-phase alloys—from ultimate tensile strength (UTS) and elongation (EL) of 220 MPa and 0% after L-PBF [43] to 1258 MPa after heat treatment [64] and 25.7% after DED accompanied by point forging [48].
The best combination of mechanical properties (UTS = 879 ± 107 MPa [45], EL = 19.3% [63]) of O-phase alloys manufactured by the WAAM method are lower than for L-PBF samples (UTS = 1090 ± 9.27 MPa, EL = 22 ± 0.48% [36]), UTS = 1144 MPa, EL = 24.25% [42], PF-LD (UTS = 1169 ± 10 MPa, EL = 25.7 ± 0.6% [48]. The main reason for the lower properties of WAAM alloys is the coarser-grained structure with the average B2/β-grain size exceeding 300 µm [45][46]. The average grain after L-PBF and DED-PF size did not exceed 40 μm [36][42] and 55 μm [48], respectively. Point forging provides the highest combination of properties with a relatively small dispersion [63], apparently due to the formation of a more uniform melt pool during manufacturing.
Previous research devoted to the mechanical properties of O-phase alloys manufactured using various types of powders showed that the use of a mixture of elemental metal powders does not provide a required combination of properties [65] due to embrittlement of the alloy as a result of oxygen contamination. An increased set of properties is provided by hot isostatic pressing (HIP) (1160 °C, 3 h) after MAM [66].
The best combination of tensile properties UTS = 1169 ± 10 MPa, EL = 25.7 ± 0.6% was obtained for Ti-21.18Al-25.3Nb-0.12O-0.1N alloy, manufactured by DED + PF. Thus, recrystallization after point forging results in the fine-grained B2/β structure [57]. A similar level of properties was obtained in [36][42] on the L-PBF Ti-22Al-25Nb alloy manufactured according to the parameters given which provides the maximum relative density and the formation of a fine-grained structure with an average grain size of 38 μm. At the same time, such a combination of properties is obtained in the considered alloys immediately after manufacturing in a metastable state with high content of the B2/β matrix (95% or more) that is thermodynamically unstable at elevated operating temperatures.
Annealing [36][62][65][66] or hardening heat treatment including quenching and aging [27][37][47][58][59][60][64] stabilizes the microstructure of the alloy due to the decomposition of the metastable B2/β-solid solution. However, the effect of hardening heat treatment on the mechanical properties of MAM alloys is uncertain.
The effect of hardening heat treatment on the structure, phase composition and mechanical properties of the L-PBF Ti-22Al-25Nb alloy is considered in [64]. An increase in quenching temperature from 950 to 1050 and 1100 °C leads to a decrease in the YS from 949 to 900 MPa and EL from 24.9% to 12.1 and 14.3% with a maximum tensile strength of 981 MPa after quenching at 950 °C. The observed ductility drop is associated with the precipitation of O, α2 -phases along the B2/β-grain boundaries, as well as the growth of the B2/β-grain after heating at 1100 °C. At the same time, the volume fraction of the second phases after quenching is insignificant (no higher than 7%), but the precipitations are quite coarse. This does not allow obtaining higher strength in hardened alloys in comparison with the as-built almost single-phase B2/β state with a fine-grained structure with maximum solid solution strengthening. As shown in [64], during subsequent aging at 700 °C fine O-phase particles precipitate both in the body and along the B2/β-grain boundaries. This is accompanied by a decrease in EL down to 0.6–1.4% and contributes to the brittle fracture (after preliminary quenching from 1050, 1100 °C). At the same time, aging at 700 °C does not provide significant hardening, except for an alloy quenched from 950 °C. An increase in the aging temperature up to 830 °C after quenching from 950 °C leads to the formation of coarser O, α2-precipitates, which provide a higher ductility of 6.1%. However, the strength properties correspond to that of as-built samples. A similar tendency in mechanical properties of O-phase based alloys is also observed in other studies [27][63], which reveals that, at the moment, there are no heat treatment routes available that provide a strength-ductility combination higher than in an as-built state.
The authors of [42] considered the effects of the hatching distance (HD), inhomogeneity of the properties throughout the part height [44], the orientation of the sample on the building platform [58] and doping [2]. The superior properties for L-PBF were obtained at HD 0.12 and 0.16 mm in [42]. The authors of [44] revealed an increase in strength and a decrease in ductility from the bottom to the top of the part due to a decrease in the volume fraction of embrittling O + α2-phases during L-PBF of Ti-22Al-25Nb alloy. A study of the mechanical properties of anisotropy of the MAM Ti-22Al-25Nb [58] after hardening heat treatment (quenching from 960 °C followed by aging 850 °C for 24 h) showed that the strength and ductility are slightly higher in the horizontal direction compared to the vertical one. However, considering the dispersion of properties, the values overlap in both directions. In [2], the effect of the TiB2 addition to Ti2AlNb powder on the properties of welded joints of L-DED Ti-22Al-25Nb alloy was investigated. The highest strength and ductility of welded joints were achieved with a minimum fraction of TiB2 of 3.1%. An increase in the TiB2 content and coarsening of its particle size led to a drop in mechanical properties due to the negative effect of borides on the ductility of O-phase alloys.

3.3. The Influence of MAM Process Parameters and Heat Treatment on Mechanical Properties at Elevated Temperatures

Previous studies have shown that at a test temperature of 650 °C, the UTS and EL lie in the range of 365–833 MPa and 0–10%, respectively. The lowest values of UTS (365 MPa and 375 ± 32.6 MPa) at an EL of 0.4% and 0.76 ± 0.05% were obtained in Ti-22Al-25Nb, Ti-24.8Al-22.3Nb alloys produced by L-PBF [64] and DWAAM [46] methods. In [64], the reason for the low mechanical properties may be due to residual stresses caused by the lack of platform pre-heating. The alloy in the as-built state consisted of a single-phase B2/β-solid solution in which decomposition took place with the formation of a dispersed O-phase, during heating to the testing temperature. A similar embrittlement effect (UTS = 400 MPa, EL = 0%) of an alloy with a B2 structure in an as-built state during tests at 650 °C was observed in [27]. On the one hand, the low combination of mechanical properties in [46] is probably associated with the large B2/β-grain size formed during the fusion of the alloy (the average size is about 350 μm, individual grains are up to 1.2 mm). On the other hand, it can be due to chemical inhomogeneity and Nb-rich zones, which stabilize the matrix B2/β phase. The highest combination of mechanical properties at 650 °C (UTS = 820–833 MPa, EL = 2.6–3%) was obtained at Ti–22Al–24Nb-0.5Mo [27] and a welded joint of the Ti-22Al-25Nb alloys [62] manufactured by L-PBF and L-DED, respectively. Similar results are characteristic of a structure consisting mainly of O-phase platelets and a minor fraction of α2 and B2 phases of the samples that were additionally heat treated (quenched from 1000 °C, aged at 800 °C, 24 h [27] and annealed at 850 °C, 2 h [62]).
High-temperature tests at 750 °C of Ti-22Al-25Nb alloy (B2 + α2 + O-phase structure) in [59] manufactured by LMD method and subjected to stress relief annealing at 550 °C, 2 h, followed by hardening heat treatment (960 °C, 3 h + 800 °C, 24 h) revealed a considerable spread of UTS in the range of 645–745 MPa, and EL of 1.5–3%. An increase in the test temperature by 100 °C from 650 to 750 °C had not resulted in a significant decrease in strength properties. This is an encouraging result in terms of the use of MAM for O-phase alloys operating at temperatures in the range of 650–750 °C. Production of welded joints of O-phase alloys by the DED method using TiB2 additives [2] showed that the highest set of properties at a test temperature of 650 °C is obtained with a relatively small (3.1%) TiB2 addition. This provides a tensile strength of 638–640 MPa and elongation of 11.4–12.6% in the non-heat-treated state. At higher contents of TiB2, both strength and ductility decrease due to an increase in the embrittlement effect of TiB2, especially with the introduction of a coarse fraction.

3.4. The Influence of MAM Process Parameters and Heat Treatment on Structure and Phase Composition

The typical microstructure of O-phase alloys obtained by MAM consists of the B2/β-grains with a different fraction of grain boundary O- or α2-precipitates and intragranular precipitates of lamellar O-phase. Plate-like colonies grow mainly from the boundaries or appear in the form of individual small O-platelets between the colonies. The closer the pre-heating or processing temperature to the B2/β region, the lower the volume fraction of the second phases in the structure due to their dissolution and coagulation. Generally, the O and α2 precipitates at the grain boundary dissolve in the last turn. The B2/β-grains grow when the fraction of the O and α2—phases decrease upon transition to a single-phase B2/β—region during post-processing. Chemical inhomogeneity of alloying elements in as-built state can take place if a mixture of elemental powders as well as Al and Ti-Nb alloy wires are used as feedstock materials. Thus, regions enriched with refractory elements primarily Nb on the one hand and with a more fusible Al on the other hand are formed.
The phase composition of the parts or regions of inhomogeneous elementals distribution in as-built state and after post-processing is determined by four main factors: (1) chemical composition of the analyzed alloy (region); (2) heating temperature; (3) holding time; and (4) cooling rate. In the 1st factor, the chemical composition allows for choosing the binary (Ti-Al, Ti-Nb) or ternary Ti-Al-Nb diagram to evaluate possible changes in the phase composition depending on the temperature of the final heat treatment (the 2nd factor). For more complex alloy compositions of Al, the Nb equivalents mentioned in Section 1 can be employed. The 3rd and 4th factors allow for the evaluation of the completeness of phase transformations at the heating temperature in accordance with the phase diagram. Lower heating temperatures and holding times with higher cooling rate result in incomplete phase transformation.
Generally, the solidification texture <001> B2/β typical of BCC metals is formed along building direction of MAM products [67]. The platform/substrate, powder pre-heating temperature, hatch distance and other printing parameters affect the crystallization texture. The simulation of time–temperature dependence during L-DED at different points throughout the sample height was conducted in [44]. The temperature change at each point of the synthesized O-phase alloy is characterized by a self-extinguishing oscillating amplitude of temperature. In a high temperature thermal cycle (HTTC), the heating temperature exceeds βtr (1060 °C). For the low temperature thermal cycle (LTTC) the heating temperature is below βtr. When βtr belongs to the temperature range between the highest and the lowest, the medium temperature thermal cycle (MTTC) takes place. The experiment in [44] proved that the thermal “history” in each region of the fused sample determine the final phase composition of this region.

References

  1. Kumpfert, J. Alloys Based on Orthorhombic Titanium Aluminide. Adv. Eng. Mater. 2001, 11, 851–864.
  2. Zhang, K.; Lei, Z.; Ni, L.; Zhou, H.; Chen, Y. Improvement in Microstructure and High-Temperature Brittleness of Laser-Welded Ti-22Al-25Nb Joints with the Addition of TiB2 Powder. J. Mater. Process. Technol. 2021, 288, 116848.
  3. Boehlert, C.J.; Majumdar, B.S.; Seetharaman, V.; Miracle, D.B. Part I. The Microstructural Evolution in Ti-AI-Nb O + Bcc Orthorhombic Alloys. Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 1999, 30, 2305–2323.
  4. Novak, A.V.; Alexeev, E.B.; Ivanov, V.I.; Dzunovich, D.A. The Study of the Quenching Parameters Influence on Structure and Hardness of Orthorhombic Titanium Aluminide Alloy VTI-4. Proc. VIAM 2018, 2, 5.
  5. Sagar, P.K.; Banerjee, D.; Muraleedharan, K.; Prasad, Y.V.R.K. High-Temperature Deformation Processing of Ti-24Al-20Nb. Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 1996, 27, 2593–2604.
  6. Muraleedharan, K.; Nandy, T.K.; Banerjee, D.; Lele, S. Phase Stability and Ordering Behaviour of the O Phase in TiAlNb Alloys. Intermetallics 1995, 3, 187–199.
  7. Rosenberg, H.W. Titanium Alloying in Theory and Practice. In The Science, Technology and Application of Titanium; Pergamon Press: Oxford, UK, 1970; pp. 851–859.
  8. Ilyin, A.A.; Kolachev, B.A.; Polkin, I.S. Titanium Alloys. Composition, Structure, Properties; Russian State Technological University Moscow (MATI): Moscow, Russia, 2009.
  9. Goyal, K.; Sardana, N. Phase Stability and Microstructural Evolution of Ti2AlNb Alloys-a Review. Mater. Today Proc. 2021, 41, 951–968.
  10. Banerjee, D. The Intermetallic Ti2AlNb. Prog. Mater. Sci. 1997, 42, 135–158.
  11. Zhang, H.; Yan, N.; Liang, H.; Liu, Y. Phase Transformation and Microstructure Control of Ti2AlNb-Based Alloys: A Review. J. Mater. Sci. Technol. 2021, 80, 203–216.
  12. Liu, S.S.; Cao, J.X.; Zhou, Y.; Dai, S.L.; Huang, X.; Cao, C.X. Research and Prospect of Ti2AlNb Alloy. Zhongguo Youse Jinshu Xuebao/Chin. J. Nonferrous Met. 2021, 31, 3106–3126.
  13. Banerjee, D.; Gogia, A.K.; Nandi, T.K.; Joshi, V.A. A New Ordered Orthorhombic Phase in a Ti3AlNb Alloy. Acta Metall. 1988, 36, 871–882.
  14. Chen, X.; Weidong, Z.; Wei, W.; Xiaobo, L.; Jianwei, Z. Coarsening Behavior of Lamellar Orthorhombic Phase and Its Effect on Tensile Properties for the Ti-22Al-25Nb Alloy. Mater. Sci. Eng. A 2014, 611, 320–325.
  15. Wei, W.; Weidong, Z.; Chen, X.; Xiaobo, L.; Jianwei, Z. Designed Bimodal Size Lamellar O Microstructures in Ti2AlNb Based Alloy: Microstructural Evolution, Tensile and Creep Properties. Mater. Sci. Eng. A 2014, 618, 288–294.
  16. Wang, S.; Xu, W.; Sun, W.; Zong, Y.; Chen, Y.; Shan, D. Study on Microstructure Evolution and Mechanical Properties of Ti2AlNb-Based Alloy under Canning Compression and Annealing. Metals 2019, 9, 980.
  17. Cowen, C.J.; Boehlert, C.J. Microstructure, Creep, and Tensile Behavior of a Ti-21Al-29Nb(at.%) Orthorhombic+B2 Alloy. Intermetallics 2006, 14, 412–422.
  18. Chen, X.; Weidong, Z.; Wei, W.; Xiaobo, L.; Jianwei, Z. The Enhanced Tensile Property by Introducing Bimodal Size Distribution of Lamellar O for O+B2 Ti2AlNb Based Alloy. Mater. Sci. Eng. A 2013, 587, 54–60.
  19. Wang, W.; Zeng, W.; Xue, C.; Liang, X.; Zhang, J. Microstructure Control and Mechanical Properties from Isothermal Forging and Heat Treatment of Ti-22Al-25Nb (at.%) Orthorhombic Alloy. Intermetallics 2015, 56, 79–86.
  20. Wang, W.; Zeng, W.; Liu, Y.; Xie, G.; Liang, X. Microstructural Evolution and Mechanical Properties of Ti-22Al-25Nb (At.%) Orthorhombic Alloy with Three Typical Microstructures. J. Mater. Eng. Perform. 2018, 27, 293–303.
  21. Wei, Z.Y.; Hu, K.M.; Sa, B.S.; Wu, B. Pressure-Induced Structure, Electronic, Thermodynamic and Mechanical Properties of Ti2AlNb Orthorhombic Phase by First-Principles Calculations. Rare Met. 2017, 40, 1–11.
  22. Boehlert, C.J. Part III. The Tensile Behavior of Ti-Al-Nb O+Bcc Orthorhombic Alloys. Metall. Mater. Trans. A 2001, 32, 1977–1988.
  23. Zhou, Y.; Cao, J.X.; Huang, X.; Wang, B.; Mi, G.B. Effects of Aging Temperature on the Microstructure and Mechanical Properties of a Ti2AlNb Based Alloy. Mater. Sci. Forum 2016, 849, 368–375.
  24. Wang, W.; Zeng, W.; Li, D.; Zhu, B.; Zheng, Y.; Liang, X. Microstructural Evolution and Tensile Behavior of Ti2AlNb Alloys Based A2-Phase Decomposition. Mater. Sci. Eng. A 2016, 662, 120–128.
  25. Goyal, K.; Sardana, N. Mechanical Properties of the Ti2AlNb Intermetallic: A Review. Trans. Indian Inst. Met. 2021, 74, 1839–1853.
  26. Pragana, J.P.M.; Sampaio, R.F.V.; Bragança, I.M.F.; Silva, C.M.A.; Martins, P.A.F. Hybrid Metal Additive Manufacturing: A State–of–the-Art Review. Adv. Ind. Manuf. Eng. 2021, 2, 100032.
  27. Yang, X.; Zhang, B.; Bai, Q.; Xie, G. Correlation of Microstructure and Mechanical Properties of Ti2AlNb Manufactured by SLM and Heat Treatment. Intermetallics 2021, 139, 107367.
  28. Murr, L.E.; Martinez, E.; Amato, K.N.; Gaytan, S.M.; Hernandez, J.; Ramirez, D.A.; Shindo, P.W.; Medina, F.; Wicker, R.B. Fabrication of Metal and Alloy Components by Additive Manufacturing: Examples of 3D Materials Science. J. Mater. Res. Technol. 2012, 1, 42–54.
  29. Baufeld, B.; Van der Biest, O.; Gault, R. Additive Manufacturing of Ti–6Al–4V Components by Shaped Metal Deposition: Microstructure and Mechanical Properties. Mater. Des. 2010, 31, S106–S111.
  30. Martina, F.; Mehnen, J.; Williams, S.W.; Colegrove, P.; Wang, F. Investigation of the Benefits of Plasma Deposition for the Additive Layer Manufacture of Ti–6Al–4V. J. Mater. Process. Technol. 2012, 212, 1377–1386.
  31. Oh, J.-W.; Nahm, S.; Kim, B.; Choi, H. Anisotropy in Green Body Bending Strength Due to Additive Direction in the Binder-Jetting Additive Manufacturing Process. Korean J. Met. Mater. 2019, 57, 227–235.
  32. Brandt, J. A Review of Ultrasonic Additive Manufacturing for Particle Accelerator Applications. In Proceedings of the 11th International Conference on Mechanical Engineering Design of Synchrotron Radiation Equipment and Instrumentation, Chicago, IL, USA, 26–29 July 2021; pp. 185–188.
  33. Derazkola, H.A.; Khodabakhshi, F.; Simchi, A. Evaluation of a Polymer-Steel Laminated Sheet Composite Structure Produced by Friction Stir Additive Manufacturing (FSAM) Technology. Polym. Test. 2020, 90, 106690.
  34. Zhu, Z.; Dhokia, V.G.; Nassehi, A.; Newman, S.T. A Review of Hybrid Manufacturing Processes—State of the Art and Future Perspectives. Int. J. Comput. Integr. Manuf. 2013, 26, 596–615.
  35. Bidare, P.; Jiménez, A.; Hassanin, H.; Essa, K. Porosity, Cracks, and Mechanical Properties of Additively Manufactured Tooling Alloys: A Review. Adv. Manuf. 2022, 10, 175–204.
  36. Zhou, Y.H.; Li, W.P.; Wang, D.W.; Zhang, L.; Ohara, K.; Shen, J.; Ebel, T.; Yan, M. Selective Laser Melting Enabled Additive Manufacturing of Ti–22Al–25Nb Intermetallic: Excellent Combination of Strength and Ductility, and Unique Microstructural Features Associated. Acta Mater. 2019, 173, 117–129.
  37. Polozov, I.; Sufiiarov, V.; Kantyukov, A.; Popovich, A. Selective Laser Melting of Ti2AlNb-Based Intermetallic Alloy Using Elemental Powders: Effect of Process Parameters and Post-Treatment on Microstructure, Composition, and Properties. Intermetallics 2019, 112, 106554.
  38. Brandl, E.; Heckenberger, U.; Holzinger, V.; Buchbinder, D. Additive Manufactured AlSi10Mg Samples Using Selective Laser Melting (SLM): Microstructure, High Cycle Fatigue, and Fracture Behavior. Mater. Des. 2012, 34, 159–169.
  39. Kenel, C.; Dasargyri, G.; Bauer, T.; Colella, A.; Spierings, A.B.; Leinenbach, C.; Wegener, K. Selective Laser Melting of an Oxide Dispersion Strengthened (ODS) γ-TiAl Alloy towards Production of Complex Structures. Mater. Des. 2017, 134, 81–90.
  40. Li, R.; Niu, P.; Yuan, T.; Cao, P.; Chen, C.; Zhou, K. Selective Laser Melting of an Equiatomic CoCrFeMnNi High-Entropy Alloy: Processability, Non-Equilibrium Microstructure and Mechanical Property. J. Alloys Compd. 2018, 746, 125–134.
  41. Qiu, C.; Panwisawas, C.; Ward, M.; Basoalto, H.C.; Brooks, J.W.; Attallah, M.M. On the Role of Melt Flow into the Surface Structure and Porosity Development during Selective Laser Melting. Acta Mater. 2015, 96, 72–79.
  42. Zhou, Y.H.; Li, W.P.; Zhang, L.; Zhou, S.Y.; Jia, X.; Wang, D.W.; Yan, M. Selective Laser Melting of Ti–22Al–25Nb Intermetallic: Significant Effects of Hatch Distance on Microstructural Features and Mechanical Properties. J. Mater. Process. Technol. 2020, 276, 116398.
  43. Polozov, I.; Sufiiarov, V.; Kantyukov, A.; Razumov, N.; Goncharov, I.; Makhmutov, T.; Silin, A.; Kim, A.; Starikov, K.; Shamshurin, A.; et al. Microstructure, Densification, and Mechanical Properties of Titanium Intermetallic Alloy Manufactured by Laser Powder Bed Fusion Additive Manufacturing with High-Temperature Preheating Using Gas Atomized and Mechanically Alloyed Plasma Spheroidized Powder. Addit. Manuf. 2020, 34, 101374.
  44. Tang, Y.J.; Zhang, Y.Z.; Liu, Y.T. Numerical and Experimental Investigation of Laser Additive Manufactured Ti2AlNb-Based Alloy. J. Alloys Compd. 2017, 727, 196–204.
  45. Li, Z.; Cui, Y.; Wang, L.; Zhang, H.; Liang, Z.; Liu, C.; Du, D. An Investigation into Ti-22Al-25Nb in-Situ Fabricated by Electron Beam Freeform Fabrication with an Innovative Twin-Wire Parallel Feeding Method. Addit. Manuf. 2022, 50, 102552.
  46. Li, Z.; Cui, Y.; Yu, Z.Y.; Liu, C. In-Situ Fabrication of Ti2AlNb-Based Alloy through Double-Wire Arc Additive Manufacturing. J. Alloys Compd. 2021, 876, 160021.
  47. Valkov, S.; Bezdushnyi, R.; Petrov, P. Synthesis, Structure and Mechanical Properties of Ti-Al-Nb Coatings Formed by Electron Beam Additive Technique. Vacuum 2018, 156, 140–145.
  48. Zhang, S.; Xi, M.; Liu, R.; Li, M.; Guo, X.; Gui, Y.; Wu, J. Fabricating Ti–22Al–25Nb Intermetallic with Ductility Higher than 25% by Advanced Printing Technique: Point-Forging and Laser-Deposition. Mater. Sci. Eng. A 2022, 850, 143520.
  49. Grigoriev, A.; Polozov, I.; Sufiiarov, V.; Popovich, A. In-Situ Synthesis of Ti2AlNb-Based Intermetallic Alloy by Selective Laser Melting. J. Alloys Compd. 2017, 704, 434–442.
  50. Polozov, I.; Starikov, K.; Popovich, A.; Sufiiarov, V. Mitigating Inhomogeneity and Tailoring the Microstructure of Selective Laser Melted Titanium Orthorhombic Alloy by Heat Treatment, Hot Isostatic Pressing, and Multiple Laser Exposures. Materials 2021, 14, 4946.
  51. Zhao, X.H.; Liu, Y.T.; Ye, C.; Zhang, Y.Z. Effects of Heat Treatment on Microstructure and Microhardness of Laser Melting Deposited Ti-22Al-25Nb Alloy. Mater. Res. Innov. 2015, 19, 823–830.
  52. Polozov, I.; Sufiiarov, V.; Starikov, K.; Popovich, A. In Situ Synthesized Ti2AlNb-Based Composites Produced by Selective Laser Melting by Addition of SiC-Whiskers. Mater. Lett. 2021, 297, 129956.
  53. Chen, H.; Liu, Z.; Cheng, X.; Zou, Y. Laser Deposition of Graded γ-TiAl/Ti2AlNb Alloys: Microstructure and Nanomechanical Characterization of the Transition Zone. J. Alloys Compd. 2021, 875, 159946.
  54. Cai, X.Q.; Wang, Y.; Yang, Z.W.; Wang, D.P.; Liu, Y.C. Transient Liquid Phase (TLP) Bonding of Ti2AlNb Alloy Using Ti/Ni Interlayer: Microstructure Characterization and Mechanical Properties. J. Alloys Compd. 2016, 679, 9–17.
  55. Schloffer, M.; Iqbal, F.; Gabrisch, H.; Schwaighofer, E.; Schimansky, F.P.; Mayer, S.; Stark, A.; Lippmann, T.; Göken, M.; Pyczak, F.; et al. Microstructure Development and Hardness of a Powder Metallurgical Multi Phase γ-TiAl Based Alloy. Intermetallics 2012, 22, 231–240.
  56. Göken, M.; Kempf, M.; Nix, W.D. Hardness and Modulus of the Lamellar Microstructure in PST-TiAl Studied by Nanoindentations and AFM. Acta Mater. 2001, 49, 903–911.
  57. Liu, Y.; Zhang, Y.; Chen, Y.; Tang, Y. Microstructure and Mechanical Properties of Laser Melting Deposited TA15+Ti2AlNb Alloys. J. Aeronaut. Mater. 2017, 37, 61–67.
  58. Tang, Y.J.; Zhang, Y.Z.; Liu, Y.T.; Pan, W. Characteristics and Evolution Mechanism of Solidification Microstructure for Laser Additive Manufactured Ti2 Alnb-Based Alloy. Mater. Sci. Forum 2018, 913, 118–125.
  59. Liu, Y.; Gong, X.; Liu, M.; Yongzhong, Z. Microstructure and Tensile Properties of Laser Melting Deposited Ti2AlNb-Based Alloy. Chin. J. Lasers 2014, 41, 0103005.
  60. Zhang, Y.Z.; Liu, Y.T.; Zhao, X.H.; Tang, Y.J. The Interface Microstructure and Tensile Properties of Direct Energy Deposited TC11/Ti2AlNb Dual Alloy. Mater. Des. 2016, 110, 571–580.
  61. Lei, Z.; Zhang, K.; Zhou, H.; Ni, L.; Chen, Y. A Comparative Study of Microstructure and Tensile Properties of Ti2AlNb Joints Prepared by Laser Welding and Laser-Additive Welding with the Addition of Filler Powder. J. Mater. Process. Technol. 2018, 255, 477–487.
  62. Zhang, K.; Lei, Z.; Chen, Y.; Yang, K.; Bao, Y. Heat Treatment of Laser-Additive Welded Ti2AlNb Joints: Microstructure and Tensile Properties. Mater. Sci. Eng. A 2019, 744, 436–444.
  63. Li, Z.; Chang, B.; Cui, Y.; Zhang, H.; Liang, Z.; Liu, C.; Wang, L.; Du, D.; Chang, S. Effect of Twin-Wire Feeding Methods on the in-Situ Synthesis of Electron Beam Fabricated Ti-Al-Nb Intermetallics. Mater. Des. 2022, 215, 110509.
  64. Zhou, Y.H.; Wang, D.W.; Song, L.J.; Mukhtar, A.; Huang, D.N.; Yang, C.; Yan, M. Effect of Heat Treatments on the Microstructure and Mechanical Properties of Ti2AlNb Intermetallic Fabricated by Selective Laser Melting. Mater. Sci. Eng. A 2021, 817, 141352.
  65. Polozov, I.; Sufiiarov, V.; Popovich, A.; Masaylo, D.; Grigoriev, A. Synthesis of Ti-5Al, Ti-6Al-7Nb, and Ti-22Al-25Nb Alloys from Elemental Powders Using Powder-Bed Fusion Additive Manufacturing. J. Alloys Compd. 2018, 763, 436–445.
  66. Polozov, I.; Popovich, A.; Sufiiarov, V. Effects of Heat Treatment and Hot Isostatic Pressing on Microstructure and Mechanical Properties of Ti2AlNb-Based Alloy Fabricated by SLM. In Proceedings of the METAL 2021—30th Anniversary International Conference on Metallurgy and Materials, Brno, Czech Republic, 26–28 May 2021; pp. 948–952.
  67. Popov, V.V.; Lobanov, M.L.; Stepanov, S.I.; Qi, Y.; Muller-Kamskii, G.; Popova, E.N.; Katz-Demyanetz, A.; Popov, A.A. Texturing and Phase Evolution in Ti-6Al-4V: Effect of Electron Beam Melting Process, Powder Re-Using, and HIP Treatment. Materials 2021, 14, 4473.
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Anatoliy G. Illarionov
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Stepan I. Stepanov
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Inna A. Naschetnikova
,
Artemiy A. Popov
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Prasanth Soundappan
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K. H. Thulasi Raman
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Satyam Suwas
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