Boron modified titanium alloys

Journal Pre-proofs Boron Modified Titanium Alloys Gaurav Singh, Upadrasta Ramamurty PII: DOI: Reference:

S0079-6425(20)30017-7 https://doi.org/10.1016/j.pmatsci.2020.100653 JPMS 100653

To appear in:

Progress in Materials Science

Received Date: Revised Date: Accepted Date:

3 March 2019 3 July 2019 31 January 2020

Please cite this article as: Singh, G., Ramamurty, U., Boron Modified Titanium Alloys, Progress in Materials Science (2020), doi: https://doi.org/10.1016/j.pmatsci.2020.100653

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Boron Modified Titanium Alloys Gaurav Singha1 and Upadrasta Ramamurtyb aSchool bSchool

of Materials, University of Manchester, Manchester M13 9PL, UK of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Republic of Singapore

Abstract Titanium and its alloys are extensively used in a variety of high performance applications with the α+β alloy, Ti-6Al-4V, being the most popular. Conventionally Ti alloys in the as-cast condition possess highly coarse prior β grains, whose size is in the order of ~mm. Such large grain sizes not only reduce the strength of the alloy, but also impair its workability. The addition of about 0.1% wt.% B can result marked reduction in the grain size. The advantages offered by microstructural refinements, thus induced by trace addition of B in Ti alloys, are reviewed in this paper. Processing response of the as-cast alloys improved as a result of grain refinement due to B addition, leading to the possibility of removal or minimization of primary ingot breakdown steps. This can significantly bring down the cost of the finished Ti alloy components. Microstructural refinements with B addition on the mechanical performance of the alloys both at room and elevated temperatures are reviewed with emphasis on the microstructure-property correlations. Keywords: Constitutional supercooling; Phase diagram; processing; Powder Metallurgy; strain-hardening; Wear.

thermo-mechanical

Corresponding author: Gaurav Singh. Email: [email protected]; Tel: +91-80-6147-9478; Fax: +91-80-4663 3991 1

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Table of Contents 1. Introduction and background .....................................................................................8 2. Ti-B alloy system .......................................................................................................9 2.1 Phase diagram ......................................................................................................9 2.2 Grain size refinement via trace addition of B ....................................................10 2.3 Microstructures ..................................................................................................12 2.3.1 Microstructural parameters quantification with B content .........................13 2.3.2 Influence of TiB needles on the morphology of the α phase ......................13 3. Production methods..................................................................................................14 3.1 Ingot metallurgy.................................................................................................14 3.2 Powder metallurgy .............................................................................................15 3.3 Additive Manufacturing.....................................................................................16 4. Advantages of B addition in Ti alloys .....................................................................17 4.1 Grain size stability due to TiB needles ..............................................................18 4.2 Improvement in the cohesive strength of grain boundaries ...............................18 5. Mechanical properties ..............................................................................................19 5.1 Elastic modulus ..................................................................................................19 5.2 Tensile properties ...............................................................................................20 5.2.1 Heat-treatment effects .................................................................................22 5.2.2 Strengthening mechanisms .........................................................................23 5.2.3 Effect of thermo-mechanical processing.....................................................24 5.2.4 High temperature tensile properties ............................................................26 5.3 Hardness and compressive properties ................................................................28 5.4 High-cycle fatigue..............................................................................................28 5.5 Low-cycle fatigue ..............................................................................................30 5.6 Creep resistance .................................................................................................31 5.7 Damage tolerance...............................................................................................34 5.7.1 Fatigue crack growth...................................................................................34 5.7.2 Fracture toughness ......................................................................................35 5.8. Dynamic mechanical behavior..........................................................................36 5.9. Mechanical properties of AM Ti-B products ....................................................37 7. Effect of B addition on Tβ ........................................................................................38 8. Thermo-mechanical processing conditions..............................................................40 9. Joining......................................................................................................................43 10. Wear .......................................................................................................................45 11. Oxidation resistance ...............................................................................................46 12. Machining ..............................................................................................................47 13. Hydrogen embrittlement ........................................................................................48 14. Tensile properties at cryogenic temperatures.........................................................49 15. Biomedical properties ............................................................................................50 16. Summary and outlook ............................................................................................51 Acknowledgements ......................................................................................................52

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References ....................................................................................................................53 List of Tables ...............................................................................................................66 List of Figures ..............................................................................................................75

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Nomenclature ASB AM AC AES BM BA BCC BD BE BF BOR CAD CFCs CTE CS CSSR DA DED-L DRV DRX EDM EBSD EPMA ELI FCG FC FZ GTAW GND GB GBS HAZ HBA HCF HCP HRTEM HIP IPF ISM IM LCF L-R L-T MRR MMC MA NSR

Adiabatic shear band Additive manufacturing Air cooled Auger electron spectroscopy Base metal Beta anneal Body centered cubic Build direction Blended elemental Bright field Burger’s orientation relationship Computer aided design Carbon fiber-reinforced composites Coefficient of thermal expansion Constitutional supercooling Cyclic stress-strain response Duplex anneal Direct energy deposited – Laser Dynamic recovery Dynamic recrystallization Electric discharge machining Electron back scatter diffraction Electron probe micro-analyzer Extra low interstitial Fatigue crack growth Furnace cooled Fusion zone Gas tungsten arc welding Geometrically necessary dislocation Grain boundary Grain boundary sliding Heat affected zone High beta anneal High-cycle fatigue Hexagonal closed packed High-resolution transmission electron microscopy Hot isotactic pressing Inverse pole figure Induction skull melting Ingot metallurgy Low-cycle fatigue Longitudinal-Radial Longitudinal-Transverse Material removal rate Metal matrix composites Mill anneal Notch strength ratio (=σn/σu) 4

OQ OM OR PAM PPM P/M PA RKR HRC RD RT ROM SAED SEM SBF STA ST SRS SFT TMC TWR TEM VAR VHN WAAM XRD A α b β c C Cv da/dN DL dα dβ E ETiB Eα f G Hv KIC Kt L LW

Oil quenched Orientation microscopy Orientation relationship Plasma arc melting Parts per million Powder metallurgy Pre-alloyed Ritchie, Knott and Rice Rockwell hardness scale Rolling direction Room temperature Rule of mixture Selective area electron diffraction Scanning electron microscope Simulated body fluid Solution treat and aged Solution treatment Strain rate sensitivity Superplastic forming technique Titanium matrix composite Tool wear rate Transmission electron microscopy Vacuum arc remelting Vickers hardness number Wire arc additive manufacturing X-ray diffraction Creep constant HCP phase in Ti Burger’s vector of mobile dislocations BCC phase in Ti α colony size Paris coefficient Constraint factor Fatigue crack growth rate Self diffusion coefficient Primary or equiaxed α grain size Prior β grain size Elastic modulus Elastic modulus of TiB needle Elastic modulus of α phase Frequency of the fatigue cycle in Hz Shear modulus Hardness of the material Mode-I plain strain fracture toughness Stress concentration factor Liquid phase Average whisker length

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𝑙∗ m mW n N nc Nf Q R rW ro Smax STiB Sα T t TEutectic Tliquidus tr Tβ Vα VTiB Vβ y δc ΔK ΔKo ΔT ΔεT/2 ε εf εu ζ η ρ σ σc σFS σm σn 𝛴𝑠 σu σutc σy

Key microstructural length scale for slip in the HallPetch equation Paris exponent Weibull modulus Strain hardening exponent Number of fatigue loading cycles Creep exponent Number of cycles to fatigue failure Growth restriction factor Ratio of the minimum to maximum stress/strain of the fatigue cycle Radius of the whisker Inner cut-off radius of the dislocation Maximum stress of the fatigue cycle Aspect ratio of TiB needles Aspect ratio of α laths Temperature Time Eutectic temperature Liquidus temperature Time-to-rupture β-transus temperature Volume fraction of α phase Volume fraction of TiB needles Volume fraction of β phase Separation between TiB needles during Orowan strengthening mechanism Critical whisker length Stress intensity factor range Fatigue threshold (da/dN < 10-9 m/cycle) Degree of undercooling Total strain amplitude Strain Strain-to-failure (%) Uniform strain Instability parameter Power dissipation efficiency Density Dimensionless initial whisker length Stress Characteristic TiB strength Fatigue strength (at 106 cycles) Matrix strength Net section strength Extent of cyclic softening (𝜎𝑦 ― 𝜎′𝑦) Ultimate tensile strength Ultimate tensile strength of composite 0.2% offset yield strength

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τ 𝜎𝑜 𝑘𝐻 ― 𝑃 ∆𝜎𝐿 ― 𝑆 𝜎𝑦𝑚 ∆𝜎𝐻 ― 𝑃 ∆𝜎𝑂𝑅 ∆𝜎𝑦 Δσ Δσu 𝜎′𝑦 𝜀100ℎ 𝜀𝑚𝑖𝑛 𝛿𝐶𝑂𝐷

Matrix shear strength Hall-Petch coefficient Hall-Petch slope Enhancement in strength due to load-sharing mechanism Yield strength of the Ti matrix Enhancement in strength due to Hall-Petch mechanism Strengthening due to Orowan mechanism Enhancement in yield strength due to various strengthening mechanism Extent of work hardening (= 𝜎𝑢 ― 𝜎𝑦) Relative reduction strength after hydrogen charging at 700°C for 2 h in comparison to unhydrogenated alloys 0.2% offset cyclic yield strength Creep strain after 100 h exposure Minimum creep rate Crack opening displacement

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1. Introduction and background Ti and its alloys are extensively used in aerospace, chemical, marine and bio-medical industries. This is due to their attractive combinations of high specific strength, endurance to fatigue, good resistance to corrosion and oxidation, and excellent damage tolerant properties [1–4]. The aerospace sector accounts for more than 70% of Ti consumption worldwide [5]. Continuous demand for increasing fuel efficiency makes the usage of Ti alloys more attractive. This is illustrated in Fig. 1, which shows a steady rise in the usage of Ti alloys, especially in new airframe structures [6–8]. The galvanic compatibility between Ti and CFCs, which is better than that between Al and CFCs, is also a reason for the observed increase. The usage of Ti alloys in other sectors is relatively limited due to their higher cost. Fig. 2(a) shows the cost breakdown of Ti from extraction to final fabrication [9,10]. It is clear from it that approximately 60% of the total cost arises from the energy intensive melting and fabrication processes that are required for producing the finished components from the raw material. Thus, improving the processibility of Ti alloys can significantly reduce the cost of the finished product and, in turn, broaden their usage in different industrial sectors. In general, the microstructure of as-cast Ti alloys consists of highly coarse prior β grains, whose size can be in the order of few mm [11]. This necessitates several thermo-mechanical processing steps to break the as-cast structure down to sub-mm or micron length scales. Recently, it has been shown that the addition of trace amounts of B (up to 0.1 wt.%2) reduces the as-cast grain size of Ti alloys by an order of magnitude; we will discuss them in detail later. This can, in principle, reduce--if not completely eliminate--the number of high temperature processing steps (Fig. 2b), which can offer substantial cost savings in manufacturing Ti products. The microstructural refinement that occurs with trace addition of B is also beneficial in improving the mechanical properties. Together, these benefits have the potential to make Ti alloys much more affordable. The grain size refinement through B as an inoculant is widely and long used industrial practice for Al and W alloys [12,13]. In the context of Ti alloys, B addition was initially considered towards the possibility of it forming TiB reinforcement phase

2

All the compositions are in wt.%, unless otherwise explicitly stated.

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for improving the elastic modulus and high temperature properties via the composite route in early 1990s [14]. Towards this end, discontinuously reinforced (TiB and/or TiC) TMCs were processed via in-situ formation techniques [15]. Major obstacles like high processing cost and low fracture toughness and ductility at high volume fractions of the reinforcement phases limited the scope of such TMCs [14]. In trying to figure the optimum B addition, Zhu et al. [16] examined the effect of trace additions of B (< 0.57 wt.%) on the microstructure and tensile properties of cast Ti alloys for dental application. They observed B addition significantly refines the cast structure of Ti alloys. The alloys studied by Zhu et al. were melted on a lab scale (45 x 20 x 12 mm3) using a dental casting machine with significant oxygen contamination. After this first report, considerable amount of research work was performed to examine the effect of B on the microstructures and mechanical behavior of Ti alloys. Tamirisakandala and Miracle [17] provided a brief review of the state of the art in 2012. A large volume of literature on these alloys is published since, mandating a comprehensive overview of it, which is the aim of the present paper. 2. Ti-B alloy system 2.1 Phase diagram The equilibrium Ti-B phase diagram for B content up to 50 at.% is shown in Fig. 3 [18]. As seen, B is completely soluble in the liquid phase of Ti. But, its solubility in solid β or α phases is negligible (< 0.02 at.% [18]). Hence, all the B added (below 50 at.%) to Ti precipitates in the form of intermetallic TiB, whose crystal structure is B27 orthorhombic [19,20]. TiB grows in the form of short whiskers with needle-like morphology [19,20]. A HRTEM study of the interfacial structure between the TiB and the matrix by Feng et al. [19,21] indicates clean interface formation (Fig. 4), thereby promoting strong bonding between them. The transverse cross-sectional view of TiB needle shows that they are often hexagonal in shape, as shown in the inset of Fig. 4, consisting of (100), (101) and (101) planes. The longitudinal growth rate of TiB along its [010] direction is much faster in comparison to transverse direction [19], and hence TiB occurs in the form of needles. The binary Ti-B alloy system can be classified into hypoeutectic, eutectic, or hypereutectic depending upon the B content [22]. When the B content is below the eutectic limit (< 1.55 wt.% for Ti64 [23]), Ti-B alloys are considered as B-modified Ti

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alloys. The microstructure, processing response and mechanical properties of these alloys are similar to those without B. Above the eutectic limit, the volume fraction of TiB needles in the microstructure is sufficiently large and hence the alloys are considered as discontinuously reinforced TMCs. Here, TiB can exist in two forms: primary TiB needles, which are generally coarse in size due to unconstrained growth in the L + TiB phase field, and secondary or eutectic TiB needles that form as a result of eutectic reaction with fine size distribution [22,24]. Increasing VTiB improves strength, stiffness and wear resistance of the base Ti alloy [14,25]. However, premature cracking of the large primary TiB needles (100-200 μm) results in poor ductility and fracture toughness [14,25–27]. Therefore, hypoeutectic alloys with low B content (< 1 wt.%) are technologically relevant for fracture critical applications, as the reductions in ductility and toughness with B addition in these alloys is smaller. We will limit our discussion only to Ti alloys with B content in the hypoeutectic regime in this review. 2.2 Grain size refinement via trace addition of B Coarse prior β grain size in the cast Ti alloys poses challenges in terms of secondary working of the ingot. Additionally, Ti alloys in the cast condition exhibit inhomogeneous solidification textures: a <100> fibre texture in the columnar β grain regions and a weak or random texture in the interior of the ingot [11]. Consequently, Ti alloys are subjected to several thermo-mechanical processing steps to break the as-cast structure, randomize the texture, and to produce alloys with fine microstructures. In the last two decades, several researchers have shown that trace additions of B in a variety of Ti alloys, which include near-α, α+β and β alloys, results in significant microstructural refinement in the as-cast condition [28–34]. Macrographs obtained from the transverse section of VAR Ti-6Al-2.75Sn-4Zr-0.4Mo-0.45Si (Ti-1100) ingots with and without B are shown in Fig. 5 [32]. Two distinct regions are apparent from the macrostructure of B-free Ti-1100 ingot: while columnar grains nucleate at the periphery and grow towards the center, equiaxed grains can be seen at the center. Dramatic grain refinement with trace addition of B is directly evident from the macrographs Figs. 5(a) and (b). The columnar grain structure observed in the as-cast Ti-1100 ingot is absent in the B-containing alloy. It should be noted that grain refinement due to B addition is not sensitive to section size or melting routes such as VAR, ISM and PAM [29–35]. The variations of the as-cast prior β grain sizes in various commercial Ti alloys such as CP-Ti, Ti-6Al-4V (Ti64), Ti-6Al-2Sn-4Zr-2Mo-0.1Si 10

(Ti6242S),

Ti-6Al-5Zr-0.5Mo-0.25Si

(Ti685),

Ti-5.8Al-4Sn-3.8Zr-0.7Mo-0.3Si

(Ti8343), Ti-15Mo-2.6Nb-3Al-0.2Si (Beta 21S) and Ti-5Al-5V-5Mo-3Cr (Ti5553) are plotted against the B content in Fig. 6 [17,30,31,36]. From it, we see that the addition of 0.1 wt.% B refines prior β grain size in all of them significantly. For example, addition of 0.1 wt.% B in Ti64 reduces prior β grain size by an order of magnitude, from 1.7 mm to 200 μm. Close examination of Fig. 6 reveals grain size vs. B content plot possess a knee in the range ~0.04-0.1 wt.% B, indicating that there exists a critical level of B content to achieve the most effective grain refinement [17]. Further increase in B content beyond such a critical level (~0.1 wt.%) results in only marginal reduction in grain size [28]. The mechanism of microstructural refinement with B addition in cast Ti alloys is investigated in detail by Tamirisakandala et al. [28]. The solidification pathway taken by Ti64 alloy with B content in the hypoeutectic regime is illustrated in Fig. 7 [17]. During cooling from the T < Tliquidus, primary β-Ti nuclei form in the melt and grow upon cooling between Tliquidus and TEutectic. Owing to negligible solubility of B in β-Ti [37], the nuclei rejects excess B into the melt, which leads to the formation of a B rich layer around β-Ti nuclei. This in turn, results in CS due to the gradient in composition ahead of the L/β-Ti interface [28]. As the solidification progresses, CS provides additional driving force for the nucleation of new β-Ti nuclei ahead of the L/solid β-Ti interface and slows the growth kinetics of existing L/solid β-Ti interface [17,28]. At the eutectic temperature, the remaining liquid transforms to eutectic mixture β + TiB, with TiB needles present at the prior β grain boundaries. Further, β transforms to lamellar α and retained β for T < Tβ according to BOR. The possibility of TiB needles acting as heterogeneous nucleation sites for β-Ti nuclei can be ruled out, as TiB needles are formed at temperatures below T < TEutectic, while grain refinement occurs in the L + β phase itself [28]. The extent of grain refinement through B addition depends upon the dynamic balance between latent heat of production and heat extraction from the solid/L interfaces [28]. One would expect grain refinement to increase with increasing B content due to higher CS. However, experimental observations indicate saturation in grain size beyond 0.1 wt.% B content; see Fig. 6. Tamirisakandala et al. [28] explained

3

Proprietary material of TIMET®, commonly referred as Timetal 834

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this behavior on the basis of growth restriction factor, Q, which describes the relative contribution of solute to the development of CS. In the Ti alloy systems, B is one of the most powerful segregating solutes with a high Q in comparison to common alloying elements like Al and V [38,39]. According to the growth restriction theory, new β-Ti grains will continue to nucleate in the undercooled melt and refine the grain size until latent heat released is sufficient to suppress undercooling, ΔT. This hypothesis was supported by the observation of a good correlation between estimated values of Q for various B-modified Ti alloys such as Ti64, Ti6242S and grain sizes in them [28,38]. 2.3 Microstructures Depending upon the α (Al, O, N, C, B) and β (Mo, V, Ta, Nb) stabilizer contents, Ti alloys are generally classified into α, α+β and β alloys [40]. B is an α stabilizer with limited α stability [25]. CP-Ti in α Ti alloy class exists in single α phase whereas α+β Ti alloys annealed above Tβ contain transformed β microstructure. In contrast to α and α+β Ti alloys, β Ti alloys do not transform martensitically upon quenching and retain β phase in metastable form at RT [40]. Another important class of Ti alloys are those based on the TiAl intermetallic compound. However, we will not include them in our review here. Microstructures of as-cast α+β alloy Ti64 and β-Ti alloy Beta 21S with 0 and 0.1 wt.% B contents are shown in Fig. 8 [36,41]. Lamellar microstructure with alternate layers of hcp-α and bcc-β phases within a colony is seen for Ti64 with or without B. On the other hand, equiaxed grain structure with α platelets within the retained β matrix is observed in Beta 21S alloy. Discontinuous network of TiB needles is present at the prior β grain boundaries in Ti64-0.1B and Beta 21S-0.1B alloys with a necklace like arrangement. As evident from Fig. 8(a), a GB-α phase with an average thickness of ~10 μm appears at the prior β grain boundaries in B-free Ti64 alloy. Increasing the B content to 0.1 wt.% in Ti64 reduces and/or eliminates GB-α phase (Fig. 8b). In Beta 21S alloy too, β grain boundaries are decorated with a continuous layer of GB-α phase (Fig. 8c). However, the thickness of GB-α phase in Beta 21S is much finer (~0.5 μm) in comparison to Ti64 (~10 μm). It appears that the network of GB-α phase in Beta 21S0.1B alloy is completely broken (Fig. 8d). Similar observation was noted by Sarkar et al. [42] in another β–Ti alloy, Ti-15V-3Al-3Cr-3Sn (Ti-15-3), with 0.2 wt.% B content, where the GB-α phase network was reported to be absent. Thus, B addition could be beneficial in removal or disruption of GB-α phase network. Work presented in [43] 12

shows workability of Ti alloys improves as a result of reduction in thickness of GB-α phase layer. 2.3.1 Microstructural parameters quantification with B content The development of microstructural features, which span across different length scales, in α, α+β and β Ti alloys primarily depends upon a number of parameters such as the alloy type, the thermo-mechanical processing route employed, degree of reduction, cooling rate, and the annealing temperature [40]. The average values of various microstructural parameters such as dα, c, dβ, VTiB, STiB in various α, α+β and β Ti alloy with varying B contents are summarized in Table 1. More details can be found in the respective references provided in Table 1. It is clear from it that the minor additions of B not only refines the as-cast β grain size but also reduces the other microstructural length scales. The variations of dβ and c in as-cast Ti64 alloy is plotted as a function of B content in Fig. 9 [44]. Both dβ and c exhibit a similar trend with B; this is because α colonies nucleate and grow from the vicinity of these prior β grain boundaries [45]. In contrast, the variation in dα with B content in Ti alloys seems to be more controlled by the cooling rate during solidification [38]. Interestingly, the aspect ratio of α laths Sα appears to be affected by the B addition. In the B-free alloy, the α laths are long and slender, whereas in the B-containing alloys they are short and thick, as shown in Fig. 8(b) [41]. Close scrutiny of the microstructural parameters of the wrought alloys listed in Table 1 indicates that dβ or dα obtained after thermo-mechanical processing and heattreatment are not significantly affected by B addition as they are similar to those of the B-free alloy, see for e.g. Refs. [30,32,46–48]. This is because a substantial grain refinement has already taken place during the casting stage of the B containing alloy. Hence, subsequent processing steps result in only marginal refinement. 2.3.2 Influence of TiB needles on the morphology of the α phase It has been widely reported that TiB needles provide additional heterogeneous nucleation sites for α phase and, in the process, alter the morphology of α--from conventional lath-like to more equiaxed form during slow cooling from the β→α, in a variety of Ti alloys [49–53]. The low lattice mismatch between TiB and α phases (< 3.6%) aids in this heterogeneous nucleation of α [49]. Hill et al. [52] first reported a change in the morphology of α to coarse equiaxed rather than lamellar in FC (~5 °C/min) Ti64 alloy with VTiB = 4%. Nandwana et al. [50] studied in detail the role 13

of crystallographic ORs between the TiB, α, and β phases on the morphology of α nucleating at the TiB needles in Ti5553-0.5B alloy via EBSD OM technique. Figs. 10(a) and (b) shows pseudo-colored OM maps of two different regions with α precipitates in equiaxed and lath like morphologies. Here, α1, α2, and α3 are the variants of α precipitates associated with the TiB needles. α precipitates with equiaxed morphology are seen in Fig. 10(a); here TiB needle exhibit distinct ORs with all the three α precipitates. They are OR1- (0001)α || (001)TiB and  1120  α || [010]TiB; OR2(0001)α || (101)TiB and  1120  α || [010]TiB; and OR3- (0111) α || (001)TiB and  1120  α

|| [010]TiB for α1, α2, and α3, respectively [50]. However, OR between TiB/β and BOR

between the three α precipitates and β are not observed [54]. Similar analysis performed on α precipitates (α1 and α2) in lath like morphology in Fig. 10(b) reveals the existence of ORs between TiB/α and TiB/β hold, (011)β || (001)TiB and <111>β || [010]TiB [50]. Consequently, BOR between β and α phases is obeyed. This is in contrast to equiaxed morphology of α observed in Fig. 10(a) where β does not exhibit OR either with the TiB needle or the α precipitates. In summary, presence or absence of ORs between TiB/β plays an important role in governing the morphology of α. When TiB needles do not hold specific ORs with adjoining β, equiaxed morphology of α results. Otherwise, lath-like morphology forms, provided BOR is maintained between α precipitates and β [50]. This is interesting, especially since equiaxed α forms near TiB needles in Ti55530.5B without the necessity of any mechanical working and/or recrystallization steps in the α+β field. 3. Production methods Methods for the production of B-modified Ti alloys include conventional ingot metallurgy, P/M and more recently AM routes. An overview of the various processing routes and product forms are shown in Fig. 11 [17,39,55,56]. Secondary forming techniques include rolling, forging and extrusion. 3.1 Ingot metallurgy Both VAR and ISM are attractive ingot metallurgy routes to produce Ti-B alloys [29– 32,44] because of the low operating cost and time associated with them. In both these melting procedures, B source in the form of TiB2 or elemental B is used. In VAR, Ti sponge with B source is blended with a known composition of master alloy and then

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mechanically compacted to form blocks that are welded together and then melted in a VAR furnace. It was reported that Ti-B ingots having diameter ~ 140 mm weighing about ~20 kg are melted on a lab scale through this route [31]. Due to limited melt capacity of ISM (maximum ~40 kg), melting via VAR route offers potential for scale up. In ISM, Cu crucibles with external cooling system are employed for melting purpose [57]. Around 1700 °C, molten Ti-B melt forms and is allowed to cool in a controlled manner. Induction stirring effect arises due to the generation of an electromagnetic field, which minimizes segregation of solute atoms and facilitates uniform distribution of temperature throughout the melt. Ti alloy in contact with Cu crucible melts first and forms a skull which contains molten Ti-B inside it. Molten TiB in contact with Ti inhibits Cu contamination of the melt. Formation of Ti-B ingots through ISM technique provides high degree of supersaturation, which results in good surface finish along with the flexibility to melt any composition. In certain cases, solidified Ti-B ingots are subjected to standard HIP operation to heal any as-cast porosity present. 3.2 Powder metallurgy Near net shape processing techniques like PA and BE methods have been used in the past to prepare powder metallurgical Ti-B products [17,55,58,59]. In PA method, gas atomized Ti-B powders (having the same chemistry as the desired alloy) is outgassed and compacted using techniques like cold pressing, HIP, or direct powder extrusion. The compact is subsequently sintered, generally above Tβ in an Ar atmosphere under pressure, which results in simultaneous densification and in-situ formation of TiB needles with uniform distribution. It should be noted that no refinement in grain size occurs during sintering, as refinement in grain size due to B addition in Ti occurs only in the liquid state. A shorter growth time during rapid solidification of the starting powder in the PA method results in much refined microstructural features after sintering [55]. Studies reported in [60–62] observed nanometer sized TiB precipitates in addition to sub-micron scale TiB needles during solid state exposure of the Ti matrix with supersaturated B. These former are likely to provide dispersion strengthening. BE powder processing method is completely performed in the solid state [55]. Ti powder is either wet or dry blended with the powders of other alloying elements and B source. Subsequent steps are similar to those of the PA powder processing method. 15

Prolonged exposure (typically 6 to 8 h) during heat-treatment at high temperatures (above > 1200 °C), required for the conversion of B source into TiB, results in undesirable microstructural coarsening [55]. As a consequence of this, BE P/M products have inferior mechanical properties. PA and BE billet preforms can be subjected to secondary working to produce mill products. 3.3 Additive Manufacturing Recently, there has been considerable interest in AM of metal alloys for near-net shape and one-step fabrication of components via selective melting of powder beds with either laser or electron beams [63,64]. In addition to flexibility, on-demand manufacturing of components with nealy-100% fly-to-buy ratio, the line-by-line, layer-by-layer building method that is intrinsic to these processes imparts a mesoscopic structure to the built components that could be exploited for tailoring properties [65–67]. Amongst all the alloys that were explored for their suitability for AM hitherto, Ti64 is the most widely studied [68–70]. Production of B-modified Ti alloy components through AM processes like DED-L [56,71,72] and WAAM [39] was also explored by using Ti alloy feedstock in the form of powder or wire, blended with B powder. Due to rapid heating-cooling thermal cycles that are inherent in AM process, small melt pool size and directional heat extraction along the BD, the microstructures of Ti alloys produced through AM process are characterized by coarse epitaxial columnar β dendritic grains with mixtures of very fine α and martensitic 𝛼′ laths within β grains [39,56,63,64,69–72]. Because of the prevalence of the non-equilibrium solidification conditions, the microstructures of AM Ti alloys are similar to those in the fusion zone of welds [73]. The presence of coarse columnar β grains and GB-α phase detrimental to ductility and fatigue properties of Ti alloys fabricated through AM processes [69,74–76]. Moreover, the columnar β grain structures can induce significant mechanical anisotropy [77]. In this context, it was shown that the addition of trace amount of B can break the columnar grain morphology into an equiaxed one in Ti64, Ti-20V and Ti-12Mo alloys fabricated through DED-L and WAAM techniques [39,56,71,72]. It was also reported that the GB-α phase network at the β grain boundaries were broken or completely eliminated [39,56]. This is illustrated in Fig. 12 for DED-L Ti64-B alloy [56]. Increasing B content in DED-L Ti64 refines length and width of β grains and changes its morphology from columnar to short columnar or equiaxed, due to the competition between epitaxial columnar growth of β due to higher 16

cooling rate and the nucleation event in the solidifying melt pool due to ΔT [39,56,72]. The length and Sα of α laths decrease, while their width increases marginally with increasing B content. Unlike α+β alloys like Ti64, addition of 0.5 wt.% B in β-Ti alloys, Ti-20V and Ti-12Mo [72], fabricated through similar DED-L process as in Ti64, result in complete elimination of columnar β grain morphology. The observed contrasting columnar to equiaxed transition behavior in α+β and β Ti alloys is explained on the basis of two factors [56,72]. Firstly, freezing range of β-Ti alloys (~15 K) is much larger in comparison to Ti64 (~5 K). Addition of B was found to increase the freezing range ΔT. The amount of B added (0.5 wt.%) in Ti-20V and Ti-12Mo alloys was much larger in comparison to Ti64 (maximum 0.25 wt.%). Secondly, higher alloying elements present in β-Ti alloys in comparison to Ti64 will result in very high overall Q, which is sum of Q values of individual solutes. These two factors will increase ΔT and promote more equiaxed dendritic nucleation in the melt with increasing B addition. 4. Advantages of B addition in Ti alloys Apart from microstructural refinement, there are several unique advantages of TiB reinforcement in Ti alloys which are briefly discussed in [17] and are summarized below: 1. B does not dissolve in both α and β lattices; all the B added is utilized for TiB needle formation. This eliminates lattice embrittlement issue unlike H, O, N and C interstitial elements when present above their allowable limit. 2. TiB needles grow in the form of short whiskers (Fig. 4) and are efficient strengtheners owing to their high hardness. 3. The CTE of TiB is similar to Ti (CTETi = 9 x 10-6 K-1, CTETiB = 8.6 x 10-6 K-1 [78]), which, in turn, minimizes thermal residual stresses generated during cooling from the β/α+β phase. 4. TiB needles present at the grain boundaries restrict their mobility at high temperatures due to Zener drag mechanism (discussed in detail in section 4.1) [79,80]. 5. TiB reinforcement is both mechanically and chemically stable. The interface between Ti-TiB needles is clean [19,81] (Fig. 4), with no intermediate phase formation like in the case of SiC, Al2O3, Ti3N4 reinforcements [82–85].

17

6. Increase in the specific strength (σy/ρ) and specific stiffness (E/ρ) of Ti alloys without any weight penalty, as the densities of TiB and Ti are almost equal (ρTi = 4.57 g/cm3 and ρTiB = 4.56 g/cm3 [78]). 4.1 Grain size stability due to TiB needles Ti alloys are susceptible to rapid grain coarsening at elevated temperatures especially in the β phase [86]. The influence of TiB needles on the grain growth kinetics of Beta 21S-0.1B alloy was investigated by Cherukuri et al. [80]. For this, they annealed this alloy with 0 and 0.1 wt.% B for times varying between 30 min to 10 h at temperatures above their Tβ. Fig. 13 shows the variation of dβ as a function of annealing time at different temperatures. In the B-free alloy, grain size increases considerably due to annealing. In contrast, dβ of Beta 21S-0.1B alloy remains unchanged at all annealing T-t combinations. Similar observation was also made on P/M processed Ti64-1B alloy, where average dβ remains unchanged after thermal exposure at 1200 °C for 1 h [17,58]. Cherukuri et al. ascribed this to Zener pinning effect of TiB needles present at the β grain boundaries and modeled it using a modified Zener drag model that takes various factors such as VTiB, STiB, orientation of TiB needles with respect to grain boundary into account [80]. The necklace like network of TiB needles at the β grain boundaries may also provide additional resistance to grain growth. 4.2 Improvement in the cohesive strength of grain boundaries Prior works on various Ni and Ti-based intermetallic alloys have shown that B segregates to the grain boundaries in these alloys and improves their cohesive strength, thus suppressing intergranular failure [87–90]. For example, Tang et al. [89] found that the addition of 0.2 at.% B to Ti-22Al-20Nb-2W (at.%), an O+B2 based Ti aluminide alloy, increases its ductility from 0.8 to 3.4%. At such low B concentration, segregation of B atoms at the B2 grain boundaries was reported. Does segregation of B atoms to the prior β grain boundaries occurs in Ti alloys too? To answer this question, Luan et al. [91] examined grain boundary chemistry of the tensile fracture surfaces of as-cast Ti64 alloy with 0, 0.02 and 0.05 wt.% B contents using AES. Detection of B peaks in the Auger spectra in the grain boundary regions of 0.02 and 0.05B alloys, see Fig. 14(a), indicates that indeed B segregates to the grain-boundaries of cast Ti64 alloys and improves its cohesive strength [91]. As a result of this, fracture mode changed

18

completely from intergranular in B-free alloy (Fig. 14b) to predominately transgranular in B-modified alloys (Fig. 14c). B peaks in the transgranular regions were mostly from the TiB needles. Quantitative estimates further indicate that the relative concentration of B in the transgranular region (mostly from TiB) was much higher in comparison to grain boundary region [91]. Thus, the observed ductility enhancement in as-cast Ti alloys with trace B additions cannot be attributed solely to grain-size refinements. Improvement in the cohesive strength of β grain boundaries due to segregation of B appears to play an important role as well [91]. 5. Mechanical properties B-modified Ti alloys are contenders to replace existing Ti alloys for various applications. Therefore, a detailed understanding of their mechanical behavior is essential, which is reviewed below. 5.1 Elastic modulus Table 1 summarizes the reported values of E of Ti alloys with varying B contents in different processing conditions. As seen, E increases with the B content. For instance, addition of 0.55 wt.% B in as-cast Ti64 [92] and 0.2 wt.% B in Ti-1100 alloys [32] enhances their E by approximately ~11% over B-free counterparts. Insignificant amount of stiffening is noted in β-Ti alloys like Beta 21S and Ti5553 with 0.1 wt.% B addition [36]; E increases merely by only 3.2 and 6%, respectively. This could be due to lower E of the bcc-β Ti matrix [93], which is the dominating phase in these alloys. The E of Ti alloys depends upon various factors like moduli of individual phases and their volume fractions, interstitial content (especially O and N) and orientation of the loading direction with respect to the axis [94]. Sen and Ramamurty [92] investigated E of as-cast Ti64 alloy with varying B contents both experimentally and theoretically. Variation of E with B content in Ti64 is shown in Fig. 15 [92]. Except 0.04B alloy, E increase linearly with B content (or TiB needles), which was attributed to much stiffer TiB needles (~4 times that of α phase, ETiB = 384.5 ± 40.2 GPa and Eα = 132.2 ± 12.2 GPa) even though it is present in much smaller volume fractions. Analysis using isostrain ROM, which provides the upper bound, that takes into account the variations of Vα, Vβ, VTiB with the B content, as well as texture and orientation effects, show that the predicted and experimental E are in good agreement, except for Ti640.55B where the estimated value is significantly lower than the measured value. This

19

was attributed to necklace type arrangement of TiB needles (inset of Fig. 15) at the prior β grain boundaries which provides extra stiffening to the matrix [92]. The high temperature stiffness of B-modified Ti alloys in the temperature range 25-550 °C have been reported to improve with B addition [95–97], see Table 2. The enhancement in stiffness in extruded alloys was found to be higher in comparison to as-cast alloys [95]. Development of strong (0002)α basal texture during extrusion which aligns axis with the loading direction was suggested to be the reason behind [95]. Further, effective load-transfer to TiB needles occurs as they preferentially orient themselves parallel to the loading direction [95]. 5.2 Tensile properties RT tensile properties of near-α, α+β, and β Ti alloys with varying B contents are summarized in Table 1. As evident from it, addition of B to Ti alloys enhances their strength. Variations of σu and εf in a variety of Ti alloys like Ti64 [44], Ti685 [31], Beta 21S [36], Ti5553 [36] in the as-cast condition are plotted as a function of B content in Fig. 16. Following are some key observations. Addition of 0.1 wt.% B in Ti64 increases its σu by ~10%. However, the extent of strengthening with B addition in Ti685, Ti5553 and Beta 21S alloys was found to be negligible. It seems that dramatic grain size refinement observed in as-cast Ti alloys with trace addition of B does not lead to significant improvement in strength. Form Fig. 16(b), it is clear that addition of small amount of B ≤ 0.1 wt.% in Ti64 and Beta 21S alloys improves its ductility. This initial enhancement was attributed to beneficial effect of microstructural refinement due to B addition [44]. However, further increase in B content beyond 0.1 wt.% leads to a loss in ductility. In Ti685 and Ti5553 alloys, εf decreases monotonically with increasing B content. It appears that the increase in εf due to microstructural refinement is offset by the increased volume fractions of hard and brittle TiB needles at higher B concentrations (> 0.1 wt.%) [44]. Fractographic analyses performed on tensile tested specimens indicate to ductile fracture appearance with the observation of dimples in samples with low B content. In contrast, quasi-cleavage fracture features appear in those with higher B content. Ductility of β-Ti alloys is relatively low due to the presence of a GB-α at the prior β grain boundaries. This phase is soft in comparison to the β matrix, and early crack nucleation occurs at these locations due to strain localization within this GB phase [98,99]. Tamirisakandala et al. [36] noted that the continuous layer of GB-α (~0.5 μm) 20

in Beta 21S is replaced by thin and small α platelets at the prior β grain boundaries in Beta 21S-0.1B alloy, see Fig. 8(d). Consequently, fracture path changes from straight in Beta 21S to more tortuous in Beta 21S-0.1B (Fig. 17), which explains enhanced ductility in the latter. To summarize, increasing the B content beyond 0.1 wt.% leads to significant drop in ductility, and in the process, more than offsets the beneficial effect of B. Key microstructural length scale for slip, 𝑙 ∗ , (for Ti alloys 𝑙 ∗ could be dα, c or dβ) controlling the strength of B-modified Ti alloys is also investigated by several authors [31,44,100]. The relationship between σy and 𝑙 ∗ is often described by the HallPetch relationship [101]: 𝜎𝑦 = 𝜎𝑜+ 𝑘𝐻 ― 𝑃(𝑙 ∗ )

―1/2

……………………………….

(1)

This possibility was explored in as-cast Ti64 [44] and Ti685 [31] alloys. Good correlation between σy and c-1/2 in Ti64 and dα-1/2 in Ti685 was obtained. This shows that enhancement in strength with B addition in Ti alloys is primarily due to microstructural refinement. Due to low VTiB in alloys with B content ≤ 0.1 wt.%, extent of strengthening due to load-sharing mechanism by the TiB whiskers would be negligible. Strengthening due to load-sharing mechanism would be feasible to some extent, but only in higher B-containing alloys (≥ 0.3 wt.%); this is discussed in section 5.2.2. Several studies have shown that strain-hardening behavior of Ti64 alloy at RT increases marginally with B addition with n ~ 0.1 [31,44,102]. This is expected as B does not go into solid solution of α, which is the major constituent phase in α and α+β Ti alloys. Then, the observed increase in n with the B addition could be due to increased dislocation interaction with α-β interfaces, as their density is expected to increase with B addition. To obtain further insight into it, detailed dislocation and slip trace analysis in the as-cast Ti64-B alloys deformed in compression was performed using optical profilometry and TEM in [103]. Planar slip at lower B contents and homogeneous deformation at higher B contents was observed. It was suggested that much finer colonies in the latter inhibits easy slip transfer across α colony boundaries [103]. Activation of multiple slip systems at these boundaries results in enhanced dislocation interaction and thus strain-hardening [103]. 5.2.1 Heat-treatment effects

21

B-modified Ti alloys can exhibit differences in microstructural response to heattreatment and subsequent mechanical properties relative to B-free alloys due to differences in the Tβ, α/β phase fraction, and the influence of TiB needles on the β→α phase transformation kinetics [52,58]. McEldowney et al. [58] studied the microstructure evolution and tensile properties of P/M processed and extruded Ti64 and Ti64-1B alloys subjected to various standard heat-treatments like MA, DA, STA, BA and HBA. See Table 1 for abbreviation used and heat-treatment details. Representative backscattered SEM images of Ti64 and Ti64-1B alloys after various heat-treatments are shown in Fig. 18 [58]. An extensive reduction during extrusion resulted in complete recrystallization of primary α. MA condition microstructure showed equiaxed α with broken β network at the grain boundaries or triple points (Figs. 18 a and b) while equiaxed α + transformed β in DA (Fig. 18 c and d). Grain coarsening due to thermal exposure at 940 °C was observed in Ti64 (Fig.18c), which otherwise was absent in the Ti64-1B alloy (Fig. 18d). The STA produced globular primary α plus tempered martensitic α' and retained β in Ti64 and Ti64-1B alloys (Fig. 18 e and f). Heat-treatment above the Tβ in BA and HBA conditions leads to the formation of completely lamellar structure in Ti64 and Ti64-1B alloys (Fig. 18 g to j). Tensile properties of Ti64-1B alloy in different heat-treated conditions are summarized in Table 1 and strength of Ti64-1B alloy is compared with Ti64 in Fig. 19 [58]. Both σy and σu of Ti64-1B alloy are significantly higher than the respective values of Ti64 in all the conditions examined. Except STA condition, strength of Ti64-1B alloy does not show much dependence on the heat-treatment. STA condition produced maximum strengthening in Ti64 and Ti64-1B alloys due to formation of martensitic α', which is more resistant to plastic deformation [40,104]. The strengths of Ti64-1B alloy in MA and DA conditions are nearly-identical Surprisingly, a drop in σy by 10% over as-received condition was noted by McEldowney et al. [58] in BA and HBA conditions. The morphology of TiB in P/M processed Ti64-1B alloy was circular/equiaxed in contrast to needle type in IM processed alloys. The change in morphology of TiB from needle type to equiaxed may have lesser grain boundary pinning effect due to lower STiB when the alloys where annealed above Tβ. The resulting coarsening of lamellar structure leads to lower σy in both BA and HBA conditions. The εf for P/M Ti64-1B alloy in different heat-treated conditions was in the range of 10-13%, see Table 1.

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5.2.2 Strengthening mechanisms B-modified Ti alloys can be viewed as particle or discontinuously reinforced composites reinforced with TiB whiskers. Consequently, the following different types of strengthening mechanisms likely to operate in B-modified Ti alloys [58,105]. 1. Load-sharing mechanism: load transfer from the matrix to TiB reinforcement through interfacial shear stresses. Enhancement in σy of the composite with cylindrical or platelet type reinforcement through load-sharing mechanism can be expressed as [106], ∆𝜎𝐿 ― 𝑆 =

𝜎𝑦𝑚𝑉𝑇𝑖𝐵𝑆𝑇𝑖𝐵 4

…..…………..……………..

(2)

Load-transfer predominately depends upon the bond integrity and nature of the interface. 2. Hall-Petch strengthening: as already mentioned, the increment in yield strength associated with a reduction in grain size due to B addition can be estimated using the classical Hall-Petch relationship: ∆𝜎𝐻 ― 𝑃≅𝑘𝐻 ― 𝑃(𝑙 ∗ )

―1/2

…………………………………

(3)

3. Increase in the dislocation density due to difference in the CTE between the Ti matrix and the TiB whiskers [107]. 4. Enhanced strength due to generation of GNDs at the Ti-TiB interfaces to accommodate strain incompatibility [108]. 5. Orowan strengthening due to bowing of dislocations around TiB needles separated by distance y [109]. Increase in σy due to Orowan mechanism is given by: ∆𝜎𝑂𝑅 =

0.13𝐺𝑏 𝑟𝑜 ln 𝑏 𝑦

()

……….……………..

(4)

The extent of strengthening from mechanisms (3) to (5), listed above, can be neglected due to the fact that CTEs of TiB and Ti are similar and the former forms a coherent interface with the latter, see Fig. 4 [19,21]. TEM observations on the failed tensile specimens in [30], which does not show higher dislocation density adjacent to TiB whisker, supports this. Therefore, strengthening due to GNDs or CTE mismatch is ruled out [30,31]. Orowan strengthening via bowing of the dislocations around TiB needles is only feasible when TiB's size is within the submicron range (~100-300 nm). Such

23

length scales are achieved only when B is present in supersaturated form [58,60–62]. For Ti64 alloy with B content in the hypoeutectic regime processed through conventional IM or P/M routes, the size of TiB generally varies from ~10-100 μm. As the size and inter-particle spacing of TiB needles are considerably larger in comparison to dominant scale for dislocation motion, strengthening from the Orowan mechanism would be insignificant. The enhancement in yield strength (Δσy) due to load-sharing and Hall-Petch mechanisms is quantified in as-cast Ti64 alloys with B content up to 0.55 wt.% using an average STiB value of 10, kH-P ≈ 0.424 MPa 𝑚 (obtained by fitting least square regression fit between σy and c-1/2) and VTiB data from [44]. A good agreement between theoretical estimates and experimental values of Δσy was found (Table 3), in spite of the variations in the shape, size and orientation of TiB needles, and not considering the strengthening of the matrix in terms of σym due to increased density of interfaces with B addition in Eq. (2) [58]. 5.2.3 Effect of thermo-mechanical processing While a considerable research was performed on the tensile properties of as-cast Bmodified Ti alloys, work in the wrought condition is limited. Cast Ti alloys are inherently associated with porosity and display substantial scatter in their properties, especially fatigue [98,110]. Therefore, thermo-mechanical processing of as-cast Ti alloys is essential to heal any porosity present and convert them into other product forms with tailored microstructures. This is achieved through processes like rolling or forging followed by annealing. The tensile properties of Ti64, Ti-1100, Ti834 and VT8 (Ti-6.5Al-3.3Mo-0.3Si) alloys with B-content up to 0.2 wt.% are examined in wrought condition and compared with B-free counterparts [30,46–48,111,112]. Results of these studies are listed in Table 1, which clearly show that the strength of B-modified alloys in thermo-mechanically processed and heat-treated either in the (α+β) or β regimes are superior to B-free alloy in identical condition. For B content ≤ 0.1 wt.%, enhancement in ductility was observed in the alloys heat-treated in the (α+β) or β regime. This trend is similar to as-cast alloys although the extent of enhancement was less with increasing B content. On a contrary, a slight reduction in ductility was observed at higher B content ~0.2 wt.% in Ti834 alloy [30]. This was ascribed to cracking at the brittle TiB needles either during processing or tensile testing.

24

The microstructural features of B-modified alloys are similar to B-free alloys, for comparable thermo-mechanical processing and heat-treatment conditions, with dβ or dα being not significantly affected by B addition, see for e.g. Refs. [30,32,46–48]. Increase in the strength of wrought Ti-1100 and Ti834 alloys followed by solution heattreatment in the α+β/β phase have also been reported in [30,32,47]. For all the solution heat-treatment temperatures examined [30,32,47], the B-modified alloys show significantly higher σy and σu; see Fig. 20 for wrought Ti834 alloy. Change in the morphology of α from lath-type to equiaxed by annealing wrought Ti64-0.1B alloy in the α+β phase region resulted in improvement in strength by ~4.5% with no effect on ductility [48]. The observed improvement in strength due to trace addition of B in wrought alloys was attributed to the presence of TiB whiskers, as strengthening from Hall-Petch mechanism would be expected to be similar in base and B-modified alloys. Chandravanshi et al. [47] and Prasad et al. [30] quantified strengthening in wrought Ti1100-0.1B and Ti834-0.2B alloys by considering them as MMCs with aligned but fragmented TiB whiskers. Load-transfer based model originally proposed by Curtin and Zhou [113] and modified by Boehlert et al. [114] was used to rationalize the strength increase due to TiB. In this model, the whiskers are loaded in tension by means of shear stresses along the whisker/matrix interfaces [114]. As-fabricated flaws present in the TiB whiskers like kinks, steps results in tensile damage in the form of cracking along the lateral side of the whiskers [114]. This is followed by matrix yielding at these cracked locations [114]. TiB whiskers fragment further into smaller pieces until critical whisker length, δc, is reached, at which they no longer can sustain further load [114]. Failure of the composite then occurs through matrix fracture. In view of the limited work hardening, and hence near elastic-perfectly plastic nature of the matrix material, Boehlert et al. [114] suggested simplified expression for the ultimate tensile strength of the composite: 𝜎𝑢𝑡𝑐 = 𝑉𝑇𝑖𝐵𝜎𝑐𝜉(𝜌𝑜,𝑚𝑊)+(1 ― 𝑉𝑇𝑖𝐵)𝜎𝑚𝜀𝑢 …………………..

(5)

where 𝜉(𝜌𝑜,𝑚𝑊) is a numerical parameter that depends only on the dimensionless 𝛿𝑐

initial whisker length 𝜌𝑜 = 𝐿𝑊 =

𝑟𝑊𝜎𝑐 𝐿𝑊𝜏

and the Weibull modulus, mW. The ultimate tensile

strength of the composite is thus directly proportional to characteristic TiB strength σc and

𝜉(𝜌𝑜,𝑚𝑊) [114].𝜉(𝜌𝑜,𝑚𝑊)was

further

25

simplified

as

𝜉(𝜌𝑜,𝑚𝑊) =

( ) ―2/3, where 𝛾(𝑚𝑊) is a fitting parameter. The values of σutc, LW, rW, τ (𝜌3/2 𝑜 + 𝛾 𝑚𝑊 ) and σm were determined experimentally from the tensile test results and inserted into Eq. (5) to solve it analytically for various possible values of mW and 𝜌𝑜. With low initial damage of the TiB whiskers, predicted ductility from the stress-strain curve matches well with the experimentally determined value for mW = 2 and 𝛾(𝑚𝑊) = 1.975. This gave σc = 8 GPa, 𝜌𝑜~0.5 and δc = 11 μm [114]. For Ti-B alloys with significant initial whisker damage during processing like during α+β processing where average flow stress is much higher than β processing, in this case 𝜌𝑜 =

𝛿𝑐

𝐿𝑊 >>1, making

contribution of strength from TiB whiskers negligible as 𝜉(𝜌𝑜,𝑚𝑊) scales inversely with 𝜌𝑜. Only in the case where initial whisker damage is much lesser, strengthening from B addition was observed [30,47]. It was also found that elongation-to-failure of the composite is sensitive to the whisker strength statistics and 𝜌𝑜[114]; ductility is higher when 𝜌𝑜exceeds a critical value of ~0.5 for TiB whiskers with σc = 8 GPa, 𝜌𝑜~ 0.5 and mW = 2, [114]. Chandravanshi et al. [47] and Prasad et al. [30], who assumed the whisker strength characteristics to be the same as in [114], estimated 𝜌𝑜to be more than 1. Thus, the obtained elongation-to-failure values in their works were consistent with the predictions made from the above model. From the above discussion, it is clear that strength and elongation-to-failure values of B-modified Ti alloys predicted through the model proposed by Boehlert et al. [114] are sensitive to initial TiB whisker length through 𝜌𝑜. For example, a change in the processing temperature or route like from cast to wrought will give different values of 𝜌𝑜. 5.2.4 High temperature tensile properties High temperature tensile properties of near-α and α+β B-modified Ti alloys were investigated by several researchers [30–32,95–97,115] with the primary purpose of examining whether the trace addition of B is beneficial in improving the maximum service temperature of the alloy. Values of E, σy, σu, and εf at different temperatures reported in these works are summarized in Table 2. As seen from it, addition of B enhances the high temperature strength as well. These variations are consistent with the trends in strength and ductility with the B content at RT [30,31,44,47,100,111]. Singh et al. [97] investigated tensile properties of the as-cast Ti64-B alloys in 475-550 °C

26

temperature range. Representative tensile stress-strain responses of as-cast Ti64-xB alloys at 500 °C are shown in Fig. 21. All the alloys exhibit elastically-perfectly plastic behavior with the stress levels of B-modified alloys significantly higher in comparison to the base alloy [97]. At a fixed temperature, σy and σu of Ti64 alloy increases with the B content. With respect to ductility, εf at high temperatures were substantially higher than the corresponding values at RT (see Tables 1 and 2) and display negligible dependence on B content at temperatures above 500 °C. Boehlert and Chen [95] examined tensile properties of as-cast and extruded Ti64 alloys with B content up to 1 wt.% at 455 °C. In comparison to the as-cast alloys, the extruded alloys showed significantly higher strength, but not sensitive to the B content. This was believed to be result of development of strong (0002) α-phase texture in them during β extrusion. The XRD data further indicates (0002) α-phase peak weakens with increasing B content, and this factor was primarily responsible for tensile strength of Ti64 alloy comparable with B-modified alloys. Thus, strengthening from load-sharing mechanism by TiB whiskers in B-modified alloys plays a minor role compared to texture strengthening in extruded alloys. B additions in near-α Ti alloys like Ti6242S [96], Ti-1100 [32], Ti685 [31] and Ti834 [30] increases their elevated temperature strength. At 455 °C, Ti6242S-0.1B alloy showed a maximum εf of ~10%, further increase in B content to 0.4 and 1 wt.% resulted in brittle fracture with εf ~ 1% [96]. Tensile properties of wrought Ti-1100 and Ti834 alloys at 600 °C show that the addition of 0.2 wt.% B improves the strength with a marginal reduction in ductility as compared to the B-free counterpart in similar heattreatment condition [30,32]. Unlike Ti64, Ti-1100, Ti834, Ti6242S alloys, the extent of high temperature strengthening with B addition (up to 0.5 wt.%) was not significant in as-cast Ti685 alloy [31]. It was suggested that strengthening from load-sharing mechanism from TiB whiskers is offset by the increase in the α-lath thickness with B content in transformed β microstructure of Ti685. 5.3 Hardness and compressive properties The compressive yield strength of B-modified alloys both at room and elevated temperatures improves with B addition [33,116,117]. Commensurate with it, Hv, which is directly related to its yield strength or flow stress in compression through the Tabor equation Hv ~ Cvσy [118], where Cv is a constraint factor approximately equal to 3, also

27

gets enhanced with the B addition [33,100,116]. Good correlation between Hv and both dβ-1/2 and dα-1/2 was obtained, see Fig. 22 [100]. This indicates increase in Hv with B addition is due to refinement in dβ and dα according to Hall-Petch mechanism. However, Sen et al. [100] pointed out that the relative enhancement in Hv with B addition vis-àvis σy was much larger, due possibly to the TiB needles (Hv ~ 700 VHN [119]) network present at the grain boundaries [100]. 5.4 High-cycle fatigue Sen et al. [44] and Singh et al. [112] investigated HCF behavior of Ti64 alloy in the ascast and wrought conditions with B contents of up to 0.55 and 0.09 wt.%, respectively. Representative Smax vs. Nf plots of Ti64 alloy with 0.09 wt.% B in as-cast and wrought conditions are displayed in Fig. 23(a). Large scatter in Nf values were observed at intermediate stress levels, which could be due to the presence of porosity in the as-cast samples, while it might be due to competing mechanism of crack initiation at the surface or subsurface in the wrought samples [98,120]. The variation of fatigue strength, σFS, (at 106 cycles) with B content is plotted in Fig. 23(b). Addition of B in Ti64 enhances its fatigue performance both in the as-cast and the wrought conditions. For instance, addition of 0.55 wt.% B in as-cast Ti64 alloy enhances its σFS from 350 to 550 MPa. Marginal improvement in σFS of wrought alloys over as-cast alloys is also observed. This was attributed to improvement in strength and healing of porosity during processing [112]. Surprisingly, a slight reduction in σFS of 0.04B alloy in both as-cast and wrought conditions is noted. Detailed SEM fractographic analyses suggest that it may be due to premature crack nucleation at the soft GB-α phase, see inset of Fig. 23(b). The HCF strength of B-modified Ti64 alloy can be further improved by changing the morphology of the α phase via deformation and heat-treatment [48]. The Smax vs. Nf curves of wrought Ti64 with 0 and 0.1 wt.% B in lamellar and equiaxed forms are compared in Fig. 24. It clearly reveals B-modified alloys show considerably higher HCF strength as compared to the B-free alloy. The fatigue strength of equiaxed form was superior than lamellar form due to much finer equiaxed-α grain size in comparison to α-colony size in the lamellar microstructure [48]. Fractographic analysis of the fatigue fractured samples show no evidence of either TiB needle or TiB-matrix interface cracking as crack nucleation sites (Fig. 25) [44,112].

28

Sen et al. [44] and Hagiwara et al. [48] offered the following rationale for the observed improvements in the σFS of as-cast and wrought Ti64 alloys. For alloys with lean B content ≤ 0.1 wt.%, VTiB is too low and its effect on fatigue properties can be neglected. The marginal improvement in σFS with 0.1 wt.% B addition is due to microstructural refinements. Nf is statistically dependent upon the location of the inhomogeneity (like GB-α, inclusion, porosity, α colony etc.), its size and orientation with the loading axis. On the other hand in higher B containing alloys, Vα, dβ, c, and dα remain unaltered. However, VTiB increases significantly, for e.g. from 0.5% in 0.09B alloy to 2.7% in 0.55B alloy. Formation of a strong interface between TiB needle/matrix and removal of GB-α phase network with much stiffer necklace like arrangement of TiB needles in higher B containing alloys improve σFS. It is important to mention these would be less effective in P/M processed alloys due to low STiB and weaker TiB-matrix interface. The HCF behavior of B-modified Ti64 alloy in the as-cast, extruded, P/M rolled and P/M extruded and Ti-6242S in as-cast forms were investigated at 455 °C by Boehlert's group [121–124]. The Smax vs. Nf curves of Ti64 and Ti6242S alloys with varying B contents and forms are shown in Figs. 26. In the as-cast condition (Figs. 26 a and b), for stress levels above Smax > 350 MPa, Ti64 and Ti6242S with 0.1 wt.% B content exhibited higher average Nf than Ti64 with 1 wt.% B and Ti6242S with 0.4 and 1 wt.% B contents. Here most of the fatigue lives were considered to be in the LCF regime (< 105 cycles), where greater εf and crack propagation resistance corresponds to greater fatigue lives [123,124]. Significantly higher elongation-to-failure (εf ~11.8% for Ti64-0.1B and εf ~10.3% for Ti6242S-0.1B) and more tortuous crack front morphology were attributed for the improved fatigue lives of Ti64 [123] and Ti6242S with 0.1 wt.% B content [124]. Thus, trace addition of 0.1 wt.% B was found to be beneficial in improving the high temperature fatigue lives as well. Reduction in effective slip length due to microstructural refinement with the B addition, load sharing by TiB whiskers and obstacle to crack propagation by the stiff TiB needles, mechanisms similar to observed by Sen et al at RT, were considered to be responsible for improved fatigue strength in B-modified alloys [121–124]. For the extruded alloys, B-free Ti64 alloy exhibited higher average fatigue lives than the B-modified alloys (Fig. 26c) [123]. Fig. 26(d) shows Smax vs. Nf curves of P/M processed Ti64-1B alloy in rolled and extruded forms. A significantly larger average

29

fatigue lives of the latter are due to much finer equiaxed α grain size and pronounced (0002) α-phase texture. 5.5 Low-cycle fatigue The stabilized CSSR responses of B-modified Ti64 alloys lie below their respective monotonic stress-strain curves, indicating cyclic softening prevailing in B-modified alloys [44]; see Fig. 27(a) for a representative Ti64-0.09B alloy. Variation of Σ𝑆 and ∆𝜎 is plotted against B content in Fig. 27(b). The extent of cyclic softening, Σ𝑆 (= 𝜎𝑦 ― 𝜎′𝑦), increases with the B content and is directly related to strain-hardening behavior, gauged by ∆𝜎 (= 𝜎𝑢 ― 𝜎𝑦), under monotonic loading [44]. TEM observations indicate that during cyclic deformation, pile up of dislocations, which are generated due to the strain incompatibility between the TiB needles and the alloy matrix during cyclic loading [44], against α-β and α-TiB interfaces leads to internal stresses, which, in turn lead cyclic softening [44,125]. It is expected that internal stresses increases with increase in the density of α-β and α-TiB interfaces with B addition, thereby increasing cyclic softening with B addition. Strain-controlled LCF behavior of B-modified Ti64 alloys were investigated by Singh et al. [126]. The ΔεT/2 vs. Nf plot displayed in Fig. 28 shows for ΔεT/2 ≤ 0.75%, B addition in Ti64 improves Nf, with higher B content leading to better Nf. However at ΔεT/2=1%, 0 and 0.06B containing alloys showed better LCF performance in comparison to 0.11B alloy. It was determined that for ΔεT/2 ≤ 0.75%, fatigue life is within the elastic regime and Nf is controlled by the strength of the alloy. Improvement in the fatigue lives of B-modified alloys in this regime (ΔεT/2 ≤ 0.75%) was attributed to enhanced σy and σu due to microstructural refinements. Only in the case of ΔεT/2 = 1% where plastic strain is dominant, Nf is controlled by the ductility of the alloy. Low Nf of 0.11B alloy at ΔεT/2 = 1% was mainly attributed to its poor ductility. Multiple cracking and decohesion of the TiB needles were observed on the fracture surface of 0.11B alloy tested at ΔεT/2 = 1% (Fig. 29). The cyclic stress responses at ΔεT/2 = 1% of wrought Ti834 alloy with 0 and 0.2 wt.% B at RT and 600 °C are displayed in Fig. 30 [46]. The LCF lives of B-modified and B-free Ti834 alloys were similar and showed negligible dependence on temperature and B content [46]. The cyclic flow stress of B-modified alloy was significantly higher than the B-free alloy at RT while at 600 °C it was comparable. Both the alloys exhibited

30

mild cyclic softening until failure at both the temperatures investigated. Cyclic softening in Ti834 alloy at room and elevated temperatures are already reported in the literature [127,128]. Shearing of the ordered Ti3Al precipitates, which results in planar slip, is one of the causes. Subsequent TEM analysis performed on the failed LCF specimens does not indicate to B addition or TiB needles effecting slip planarity in any manner. Nf was found to be controlled by the ductility of the alloys. Due to nearly identical microstructural parameters and matrix microstructure in the wrought condition unlike as-cast alloys, εf values of wrought B-free and B-modified alloys are similar at RT and 600 °C (see Tables 1 and 2) with no significant influence of TiB needles. This explains similar LCF lives for both the alloys at RT and 600 °C. To summarize, improvement in the fatigue lives of Ti alloys with B addition was found as long as the strains are within the elastic regime. This is because fatigue is strength controlled, which improves with B addition, in this regime. Although TiB needle cracking and decohesion were observed, post-mortem examination of the failed LCF specimens do not point to TiB needles at the crack nucleation sites both at room and elevated temperatures. In fact, the weakest link mechanism for the fatigue crack nucleation were porosity, shear fracture of α colonies, α-β interfaces, GB-α phase and α-case at the surface. On the other hand, when plastic strain is dominant during fatigue, the fatigue response of the B-modified alloys is dependent upon its ductility. Higher B addition lead to adverse effects like lower Nf and poor FCG resistance. 5.6 Creep resistance Creep properties of B-modified Ti alloys, in as-cast, extruded and P/M processed forms and in 400-550 °C temperature range were extensively examined [30,59,95–97,115]. Representative creep curves in terms of variation of ε vs. t for the alloys are shown in Figs. 31 (a) to (c). Measured values of creep parameters like creep strain after 100 h, 𝜀100ℎ, minimum creep rate, 𝜀𝑚𝑖𝑛, and time-to-rupture, tr (where reported) are summarized in Table 4. At all the T - σ investigated in as-cast and P/M processed Ti64 alloys (Fig. 31 a and c), tr increases while 𝜀𝑚𝑖𝑛decreases with increasing B content. Therefore, it can be inferred that addition of B enhances creep resistance. Improvement in the creep resistance of as-cast Ti64 with 0.1 wt.% B addition was noteworthy at 475 °C; tr improves by 30 times and 𝜀𝑚𝑖𝑛 decreases by an order of magnitude over B-free Ti64 alloy, see Table 4 [97]. The 𝜀𝑚𝑖𝑛of Ti64-1B alloy in P/M processed and as-cast

31

forms were similar (2.1 x 10-4 h-1). For the cast-then-extruded alloys, Ti64 alloy exhibit the lowest 𝜀𝑚𝑖𝑛 followed by Ti64-1B and Ti64-0.1B alloys. On the basis of the creep exponent and apparent activation energy for creep deformation in the T and σ range investigated in Refs. [59,96,97], it was concluded that the dislocation creep controlled by the climb process was the operating creep mechanism. Following are the possible reasons offered for the observed enhancement in the creep resistance of Ti64 alloy with B addition [59,95,97]: 1. In the power-law creep regime, which is the operative mechanism in the T- σ range investigated, dependence of 𝜀𝑚𝑖𝑛 on σ and E is described by the following relation [129]: 𝜀𝑚𝑖𝑛 𝐷𝐿

∝𝐴

𝜎 𝑛𝑐 …………………………… 𝐸

()

(6)

As discussed in section 5.1, addition of B in Ti64 enhances its high temperature E, therefore improves its creep strength. 2. The number density of interfaces (α-α, α-β, and α-TiB) in the microstructure increase with B content. For the same nominal B content, P/M processed and extruded alloys have much finer grain sizes and increased density of interfaces in comparison to ascast alloy. These interfaces, as shown in Fig. 32, provide obstacle to dislocation movement during creep. However, it was observed that extent of creep strengthening with B addition becomes marginal at T > 500 °C. This behavior was rationalized in terms of increased contribution to creep strain through interface sliding which offsets dislocation-barrier effect of the α-β interfaces. 3. Limited strengthening through load-sharing mechanism by the strong and stiff TiB whiskers. This was further corroborated through high-temperature in-situ creep experiments by Boehlert [59], where crept specimens showed cracking of the TiB whiskers, indicating load-sharing mechanism is operating under creep condition also. It was also found that TiB whisker cracking initiated at stress levels well below global yield stress with the extent of cracking increasing with increasing creep displacement. 4. In the extruded alloys, additional creep strengthening arises from strong (0002) αphase texture, which weakens with B addition.

32

B addition in near-α Ti alloys such as as-cast Ti6242S [96], wrought Ti-1100 [115] and Ti834 [30] improves their creep resistance. These alloys are commonly used for compressor disks in the aero-engines with a maximum exposure temperature of up to 550 °C. Chen and Boehlert [96] investigated the creep properties of as-cast Bmodified Ti6242S alloys and found that the extent of the enhancement in its creep resistance with B addition is not as significant as in Ti64, see Table 4. EPMA data indicates redistribution of Zr and Mo from the Ti6242S matrix to TiB whiskers [96]. This factor is likely to reduce the solid solution strengthening effect of the matrix and thus creep resistance. Prasad et al. [30] examined creep properties of wrought Ti834 alloy with 0 and 0.2 wt.% B in different solution heat-treated conditions. Creep properties like 𝜀𝑚𝑖𝑛 and 𝜀100ℎdecrease with an increase in the solutionizing temperature. This was attributed to an increase in the volume fraction of much stronger transformed β phase as compared with primary α with temperature [130–132]. At all the solutionizing temperatures examined, the B-modified alloy has lower 𝜀𝑚𝑖𝑛 and ε100h as compared to the B-free alloy, and was attributed to improved strength and limited strengthening from load-sharing mechanism by the TiB whiskers. Chandravanshi et al. [115] studied the effect of α+β and β processing, followed by solution heat-treatment at Tβ - 100, Tβ - 50, Tβ - 30 and Tβ + 10 °C, on the creep properties of wrought Ti-1100 alloy with 0.2 wt.% B. The objective of their work was to study the influence of different morphologies of α, generated as a result of processing and heat-treatments, on the creep properties. β processed alloys exhibited better creep resistance in comparison to α+β processed alloys for solution heat-treatment temperatures below Tβ. The morphology of α obtained in β processed condition was fine and lenticular whereas it is coarse and equiaxed in α+β processed condition. The former exhibits better creep resistance in comparison to the latter due to much finer slip length, limited GBS, and alloying elements partitioning effect [133]. However, 𝜀𝑚𝑖𝑛 and 𝜀100ℎ are approximately similar in alloys, that were solution heat-treated slightly above Tβ (at Tβ +10 °C) due to the fact that transformed β microstructure is obtained after solutionizing irrespective of α+β or β processing. Fractographic and microstructural analyses of the longitudinal section of the crept specimens indicate ductile fracture surface appearance with presence of dimples (Fig. 33a) [97]. Transverse cracking of TiB needles along with interfacial decohesion

33

between TiB needles and matrix was also observed. In some instances, α-case cracking, cavities formation along prior β grain boundaries, TiB needle and α-β interfaces are also reported (Fig. 33b) [95,97]. 5.7 Damage tolerance 5.7.1 Fatigue crack growth Sen et al. [100] studied the FCG behavior and plane strain fracture toughness, KIC, and crack growth characteristics of Ti64-B alloys so to determine the key microstructural length scale 𝑙 ∗ that controls these properties. The da/dN, vs. ΔK plots for the as-cast Ti64-xB alloys are shown in Fig. 34(a). In the near-threshold regime (shown with arrows in Fig. 34a), there exist a threshold, ΔKo, at which crack growth occurs at an undetectable rate (da/dN < 10-9 m/cycle) [134]. In the intermediate crack growth regime, also known as Paris regime, linear relationship between da/dN vs. ΔK on log-log scale holds, 𝑑𝑎 𝑑𝑁

= 𝐶(∆𝐾)𝑚

……………………….

(7)

It was found that the FCG rates of B-modified Ti64 alloys are nearly an order of magnitude faster in comparison to B-free Ti64 alloy in the Paris regime [100]. The m values obtained for the B-modified and B-free Ti64 alloys are nearly identical (m ~ 4 for ΔK in the range 10-27 MPa 𝑚), this is due to negligible microstructural effect on the FCG rate in Paris regime [134]. The near-threshold regime is highly sensitive to microstructure with ΔKo, as shown in the inset of Fig. 34(a), decreasing with increasing B content. The dependence of 𝑙 ∗ on ΔKo was rationalized by Sen et al. using micromechanical model based on the transition from microstructure-sensitive nearthreshold to microstructure-insensitive Paris regimes [135]: ΔKo ~ 𝜎′𝑦 𝑙 ∗

……………………………

(8)

where 𝜎′𝑦 is the cyclic yield strength which is assumed to scale with σy. An excellent correlation between ΔKo and σy 𝑑𝛼, indicates that it is the α lath size which controls the near-threshold FCG behavior in Ti64-xB alloys. Higher ΔKo of B-free Ti64 alloy is mainly due to large plastic zone size ahead of the crack tip and coarser α lath size which results in crack closure and excessive deflection around α lath boundaries [100]. For the B-modified Ti64 alloys, increase in the FCG rate and reduction in ΔKo were attributed solely to the microstructural refinement, i.e., reduction in α lath size, that occurs with the B addition. Notably, an increased interaction of cracks with TiB needles, 34

as observed in Ti64-0.4B alloy in Fig. 34(b) [100], leads to TiB needle cracking and in few cases debonding also. However, the VTiB is too low (~0.021 in Ti64-0.4B alloy) to retard FCG rate in a significant manner. 5.7.2 Fracture toughness The values of KIC for as-cast Ti64 [100] & Ti-15-3 [29] and wrought VT8 [111], Ti834 [46], Ti-1100 [47] alloys with varying B contents are listed in Table 5. As seen from

Table 5, the extent of decrease in KIC with B addition is much larger in cast alloys as compared to wrought ones. The observed reduction in KIC of as-cast Ti64 and Ti-15-3 alloys with B addition was primarily due to decrease in the fracture surface roughness with increasing B content [29,100]. The higher KIC values in the B-free as-cast Ti64 and Ti-15-3 alloys was mainly attributed to coarse starting structure, which results in more tortuous crack front morphology with excessive crack deflection during its propagation [29,100].

The observed reduction in KIC with the increasing B content in as-cast Ti64 alloys was rationalized by Sen et al. by using the stress-controlled brittle cleavage model proposed by Ritchie, Knott and Rice (RKR), which incorporates the microstructural length scale, 𝑙 ∗ , in determining the toughness. According to the RKR criterion, slip-initiated onset of brittle transgranular cleavage fracture occurs when local opening stress at the crack tip exceeds critical fracture stress over a microstructurally significant length scale, 𝑙 ∗ [136]. Assuming that Ti64-xB alloys are elastic-perfectly plastic with negligible strain-hardening and fracture stress remains invariant with the B content, the following simplified relationship between KIC and 𝑙 ∗ was obtained [100]: KIC ~ 𝑙 ∗ 𝜎𝑦 A good correlation between KIC and

…………………………

(9)

𝑑𝛼𝜎𝑦 led Sen et al. to conclude that dα is the

fracture-controlling length scale in Ti64-xB alloys. Unlike the as-cast alloys, matrix microstructure and microstructural parameters of B-modified alloys in the wrought condition are similar to B-free alloys except for the presence of TiB but with low VTiB. Therefore, decrease in KIC due to reduction in fracture surface roughness with B addition is ruled out in wrought alloys. Thus, the observed marginal decrease in the KIC of wrought Ti834 and VT8 alloys with B addition is mainly due to TiB whiskers. It is widely accepted that toughness of the material depends upon its intrinsic ductility through 𝛿𝐶𝑂𝐷 [137]. Presence of TiB needles ahead

35

of the crack-tip could trigger localization, hence reduces 𝛿𝐶𝑂𝐷and, in turn, KIC [100]. Cracking of TiB needles observed on the fracture surface of wrought Ti834-0.2B alloy supports this hypothesis [46], although the maximum VTiB in wrought Ti834 (VTiB ~ 0.015), VT8 (VTiB ~ 0.01) and as-cast Ti64 (VTiB ~ 0.021 in 0.4B alloy) alloys are too low to affect KIC in a significant manner. Thus, KIC in B-modified as-cast and wrought alloys is more controlled by the fracture resistance of the matrix. The effect of solution heat-treatment at different temperatures in the α+β and β phase fields on the KIC of wrought Ti-1100 alloy with 0 and 0.1 wt.% B is also examined in Ref. [47]. It was observed that KIC of wrought Ti-1100-0.1B alloy showed insignificant effect to solution heat-treatment temperature and was marginally lower than that of B-free alloy. The decrease in KIC was attributed to TiB whiskers [47] and is in line with the observed trend in other wrought alloys like Ti834 and VT8. Before closing this section, it is important to mention that damage tolerant properties like KIC and ΔKo was found to be higher in alloys with coarser microstructures, see dependence of KIC and ΔKo on 𝑙 ∗ in Eqs. (8) and (9). However, in many practical applications subjected to cyclic loading, resistance against HCF loading is of utmost importance, and it improves with refinement in grain size. Therefore, there exist a trade off in B content to achieve optimum damage tolerant and unnotched HCF properties. Nevertheless, from the discussions in the preceding sections it is clear that addition of 0.1 wt.% B in Ti alloys exhibit good HCF fatigue properties with adequate fracture toughness. 5.8. Dynamic mechanical behavior Yu et al. [138] investigated effect of B on the compressive dynamic properties of wrought Ti64 alloy in different microstructural conditions using Split-Hopkinson pressure bar. The dynamic true stress-strain curves of wrought Ti64 with 0 and 0.1 wt.% B in equiaxed, bi-modal and Widmansttäten microstructures are shown in Fig. 35(a). The average dynamic flow stress and failure strain in the as-cast condition for the B-modified alloy was superior in comparison to B-free Ti64 alloy. In the wrought conditions, the average dynamic flow stresses were comparable and showed negligible dependence on B content and microstructure. This behavior was similar to observed under quasi-static loading rates [48]. Failure strains for B-free Ti64 alloy in all the three microstructural conditions were higher in comparison to Ti64-0.1B alloy. Overall, bi-

36

modal and equiaxed microstructures gave best combination of dynamic properties for B-modified Ti64 alloy. It is well known that Ti alloys under dynamic loading fail due mainly to the formation of ASBs during deformation [139,140]. Adiabatic shear banding occurs as a result of competitive processes of strain/strain rate hardening, which does not allow for flow localization, and adiabatic material softening that favors it [141]. Fig. 35(b) shows TiB needles cracking and cavity formation within ASB in wrought Ti64-0.1B alloy in Bi-modal condition tested at an 𝜀 = 3100 s-1. Since strain/strain rate hardening behavior of B-modified alloys in wrought condition are expected to be similar to B-free alloys, however how TiB needles effect the adiabatic softening and flow localization under impact loading need further investigation. 5.9. Mechanical properties of AM Ti-B products Figs. 36 (a) and (b) show the variations of different tensile properties and their anisotropy in DED-L Ti64-B alloy [56]. Both σy and σu increase while εf gradually decreases with increasing B content in both the directions ( ∥ BD and ⊥ BD). Overall, addition of 0.08 wt.% B in DED-L Ti64 gave the best combination of tensile properties. The measured anisotropy in strength and elongation is plotted against B content in Fig. 36(c); here anisotropy is measured as (𝑟 ⊥ BD - 𝑟 ∥ BD)/(𝑟 ⊥ BD), where 𝑟 is the tensile property value like σy, σu and εf. Strength anisotropy decreases with increasing B content while anisotropy in elongation initially increases up to 0.17 wt.% B and then decreases significantly with 0.25 wt.% B addition. The observed lowest anisotropy in strength and elongation in DED-L Ti64-0.25B alloy was related to maximum refinement in β grain size in both ∥ BD and ⊥ BD, more equiaxed form and uniform distribution of TiB needles [56]. Rashid et al. [71] examined clad hardness of DED-L Ti64 with 0.04 wt.% B. The clad hardness improved by 30% over B-free clad as a result of much finer α-laths and TiB needles. The compressive mechanical properties of WAAM Ti64 alloy with B content up to 0.13 wt.% in as-deposited and heat-treated conditions (as-deposited + heat treatment @1050°C/FC) is examined by Bermingham et al [39]. The aim of the heat-treatment was to improve ductility through coarsening of α laths and removal of martensitic 𝛼′ phase which is usually present in as-deposited state. The maximum compressive strength in as-deposited and heat-treated conditions improved very little with B addition. On the other hand, compressive strain at maximum

37

stress increases with B addition and was much higher in heat-treated condition. To summarize, addition of trace amount of B in AM Ti alloys offers unique advantages in terms of morphology change of β grains from columnar to less columnar or equiaxed, refinements in β grain size and α-lath size length, minimizing anisotropy in tensile properties. These beneficial effects could not be achieved by optimizing process parameters during deposition, which has a very narrow window. 7. Effect of B addition on Tβ The Tβ at which hcp-α phase transforms completely to bcc-β phase, is an important parameter during thermo-mechanical processing of Ti alloys [43,142,143]. Microstructure evolved during processing along with optimum α/β phase content depend sensitively upon the processing temperature relative to Tβ [144], which depends on the α and β stabilizer contents present in the alloy [94]. α stabilizers like Al, O, and N raise Tβ with respect to that of pure Ti, whereas β stabilizers like V, Mo, Cr, and Fe reduce Tβ [1]. The following empirical relation is often utilized to estimate Tβ on the basis of alloy's composition [145]: Tβ (°C) = 882 + 21.1[Al] - 9.5[Mo] + 4.2[Sn] - 6.9[Zr] - 11.8[V] 12.1[Cr] -15.4[Fe] + 23.3[Si] + 123.0[O].

…….... (10)

It is well established that Tβ of Ti alloys is strongly influenced by the interstitial O content, this is also evident from the high [O] pre-factor in Eq. (10). For example, increasing O content from 1100 (ELI grade) to 1800 ppm (commercial grade) can increase Tβ by ~50 °C [146]. Several researchers examined effect of B on the Tβ of Ti64 in as-cast [23] and P/M processed [143] conditions, as well as in wrought Ti834 [46], Ti-1100 [32] alloys. Techniques like metallography of the rapidly quenched samples, differential thermal analysis, dilatometry and electrical resistivity were explored to measure Tβ. Reported values of Tβ are plotted in Fig. 37 as function of the B content in three different Ti alloys. Tβ of the B-modified as-cast Ti64 alloys are in the range of 1000 to 1025 °C indicating an insignificant increase in Tβ over B-free Ti64 alloy [23]. For the wrought alloys of Ti834 and Ti-1100, the addition of 0.2 wt.% B increases Tβ by just 30 and 15 °C respectively. B is an α stabilizer [25,92] and should be expected to increase Tβ. However,

38

the experimental trends of Tβ with B content show it does not affect Tβ significantly at an identical [O] level. This was expected as the solid solubility of B in α and β phases is less than < 0.02 at. % under equilibrium conditions [18]. For the P/M processed Ti64 alloy, addition of 1.7 wt.% B shifts Tβ by ~60 °C [143]. This was attributed to equilibrium as well as supersaturated B in solid solution as a result of the rapid solidification [143]. When the P/M processed alloy was subsequently heat-treated, which results in achieving equilibrium microstructure and TiB formation, the measured Tβ decreases and becomes comparable to B-free P/M processed alloy [143]. Ivasishin et al. [23] attributed marginal increase in the Tβ with B addition in ascast Ti64 to redistribution of Al and V concentrations. Quantitative chemical analyses of α, β and TiB phases performed using EPMA indicate that the addition of B results in consumption of Ti from the matrix with corresponding increase in the concentration of Al in the α phase and V in the β phase [23,41]. This relative increase can be offset if Al and V are soluble in the TiB phase. However, Ivasishin et al. [23] observed only an average V content of 3.3 wt.% in TiB phase while Al could not be detected. Therefore, marginal shift in the Tβ with B addition can be related to the relative enrichment of α phase with Al. This was further confirmed on the basis of the empirical relationship that 1 wt.% Al increases Tβ by 16.7 °C while 1 wt.% V decreases Tβ by 9.1 °C [147]. Change in the concentrations of the alloying elements with B addition was used to predict the shift in Tβ based on Eq. (10) and is shown with solid line in Fig. 37. Good agreement was observed between the predicted and the experimentally determined Tβ. For the wrought Ti834 and Ti-1100 alloys with their chemical compositions different than that of Ti64, which alloying element(s) plays a significant role in shifting Tβ still need to be determined. 8. Thermo-mechanical processing conditions Trace addition of B in Ti alloys produces fine-grained billets, which has the potential to completely eliminate or reduce ingot breakdown steps, Fig. 2(b). Several researchers explored this possibility where direct rolling of Ti64-0.1B [35,148–150] and forging of CP-Ti-0.2B [151] alloys were performed in the as-cast ingot stage itself. For comparison purposes, B-free as-cast and mill products were also rolled/forged under identical conditions. Figs. 38(a)-(c) show the images of the B-free and B-modified Ti64 plates rolled directly from the ingot blanks. At all the rolling temperatures examined, non-B containing plates showed large evidence of surface and edge cracking while the 39

Ti64-0.1B alloy plates where crack-free with good surface finish. PAM and ISM Ti640.1B plates were further successfully rolled into sheets with tensile properties equivalent or superior to AMS 4911 grade Ti64 sheet [35]. Poor hot workability of ascast Ti64 ingots was attributed to coarse starting prior β grain size. The rolling response of Ti64 in mill form was found to be similar to Ti64-0.1B alloy due to fine starting equiaxed α grain size of ~15 μm [35]. It was also observed that spherodization of α phase and cracking of TiB needles as a result of imposed deformation occurred during rolling. At higher rolling temperatures above 950 °C, the space between broken TiB needles is filled by the ductile matrix (Fig. 39a). On the other hand, rolling at lower deformation temperatures of 750 and 850 °C, the matrix is not sufficiently ductile, which resulted in void formation around TiB needles (Fig. 39b). Similarly, forging behavior of as-cast CP-Ti ingot with 0.2 wt.% B was investigated in the temperature range 600-900 °C [151]. It was found that its deformation behavior was homogeneous and similar to mill form with the formation of very fine equiaxed α grains. In cast CPTi ingot without B, deformation was heterogeneous with alignment of β grains along the direction of flow. Forming limit diagrams for B-modified Ti alloys sheets have not been constructed yet. Other simpler sheet formability tests like double bend test (ASTM E 290), which gives qualitative estimation of cold workability, were performed and reported in [35]. PAM Ti64 & Ti64-0.1B and ISM Ti64-0.1B sheets were bent into “Z” shape, with bending axis parallel to the RD. The bend factor (diameter of bend divided by sheet thickness) during the test was ~8. The final images of the sheets after double bent tests are shown in Figs. 40(a) to (c). The PAM sheet of Ti64-0.1B was successfully bent into “Z” shape without any evidence of cracking whereas ISM Ti64-0.1B sheet broke during the second bend. The PAM sheet of Ti64 failed on both bends (Fig. 40c). Despite microstructural refinement due to B addition, observed reduced formability of Ti64-0.1B sheet in ISM processing route in comparison to PAM was related to higher interstitial impurity contents in former [35]. The optimum hot working zone (T and 𝜀) of as-cast Ti64-B alloys were identified by Sen et al. [152] with the aid of the processing map approach developed by Prasad and co-workers [153]. Power dissipation (η) and instability (ζ) maps are generated at a ε = 0.5 and one representative map for Ti64-0.09B alloy is displayed in Fig. 41. The numbers in the contour map of power dissipation represent numerical values of η (%) and in the instability map ζ values are in fraction. Here, η represents 40

how the work piece material dissipates power through various metallurgical processes like DRX, DRV, superplasticity etc [154]. A high value of η is generally associated with DRX mechanism and is the best choice for hot-working as it provides relatively stable flow stress with low work hardening rate. On the other hand, negative values of ζ indicate flow instability like cracking, ASB formation, kinking, etc [154]. A domain in these maps is defined as the window in the T and 𝜀 space, over which a subtle changes in the values of η and ζ are observed and can be correlated with specific deformation mechanism [153]. Sen et al. [152] found that processing maps of B-modified Ti64 alloys can be broadly classified into four domains with Domain-II (T = 900-1000 °C and 𝜀 = 10-310-2 s-1) being the optimum for hot working with η = 40-56% and a positive ζ value ~0.2, see Fig. 41. Microstructural characterization of the alloys deformed in this domain indicates DRX as the operation deformation mechanism with complete spherodization of primary α laths. One such microstructure of deformed Ti64-0.3B alloy is shown in Fig. 42(a), where spherodized α with dα ~8 μm and Sα ~1.25 in a transformed β matrix is observed. It was also found that microstructural features of B-free and B containing Ti64 alloys are essentially similar in this domain. Therefore, optimum hot working zone for the occurrence of spherodization of α laths remains unaltered with B addition in Ti64. This was attributed to two factors that are briefly discussed below [152]. Firstly, the main driving force for spherodization is the formation of energy intense deformation bands within α laths [155,156]. Flow instability that arises at the α-α interface - β lamellae intersections leads to the penetration of the α plates by the β phase (or vice-versa) driven by the surface tension. Further, spherodization occurs through diffusion-controlled processes to minimize interface energy. As the α lath size does not changes significantly with B addition, diffusion distance required for spherodization remains similar in B-free and B-modified alloys. Secondly, DRX process is controlled by power-law creep (dislocation glide/climb) mechanism, which is independent of prior β grain size. Microstructural characterization of the alloys deformed in the unstable domainIII (T = 750-850 °C and 𝜀 = 10+0-10+1 s-1) with ζ < 0 shows flow instabilities in the form of α lath bending, void formation and cracking of TiB needles, see Fig. 42(b). SFT is widely used for Ti alloys like Ti64 to produce near net shape sheet metal components, which are otherwise, difficult to form by conventional sheet metal forming operations [157]. For a material to exhibit superplasticity, it should maintain fine grain 41

size during deformation, display elongation and SRS values greater than 200% and 0.3 respectively [158–160]. The superplastic deformation behavior of cross-rolled Ti640.1B alloy sheet with a starting equiaxed grain size of dα~ 5 μm is examined in the temperature range 700-950 °C and 𝜀 = 10-5-10-2 s-1 by Sinha et al. [161]. The values of SRS obtained from the strain rate jump tests for Ti64-0.1B sheet are greater than 0.3 in the T range 775-900 °C and 𝜀 = 10-5-10-3 s-1, indicating superplastic behavior. The optimum temperature and strain rate for superplastic forming of conventional Ti64 sheet with staring grain size of 6 μm were found to be 900 °C and 𝜀 = 3 x 10-4 s-1, respectively [157]. At 900 °C, the SRS value for Ti64-0.1B sheet was 0.37 at a strain rate of 5x10-3 s-1, approximately an order of magnitude higher than that reported for a Ti64 sheet. Thus, Ti64-0.1B sheet can be superplastically formed at a higher forming speed. This is beneficial in terms of reducing production time along with lower susceptibility to α-case formation. In addition to SRS measurements, Sinha et al. performed tensile tests at a constant 𝜀 = 3 x 10-4 s-1 in the T range 700-950 °C. Tensile elongation of Ti64-0.1B sheet is plotted as a function of temperature in Fig. 43(a). For the T range of 725-950 °C, elongation is always greater than ≥ 200% with peak elongation value of 646% at 900 °C. Fig. 43(b) compares the tensile true stress-true strain responses of Ti64-0.1B and Ti64 sheet materials [157] at T = 900°C and 𝜀 = 3 x 10-4 s-1. The flow curve of Ti64-0.1B sheet has a relatively stable flow region with lower stress values. Since the flow stress required for superplastic forming is directly related to the Ar gas pressure, Ti64-0.1B sheet will require less pressure to form sheet components. Microstructural analysis of superplastically deformed Ti64-0.1B indicates coarsening of α grains along with the formation of cavities along α-α, α-β and α or βTiB interfaces (Fig. 43c). No evidence was found to suggest that TiB needles were the preferential sites for cavity nucleation during superplastic forming [79,161]. In fact, the inter-particle spacing between TiB needles is much larger for cavities to coalesce at these locations to form cracks. The observed elongation during superplastic deformation is limited by the grain coarsening and cavity nucleation and its coalescence around α-α and α-β interfaces [161]. Nevertheless, the presence of TiB needles in the Ti64 matrix can affect the material's response in two opposing ways. First, they can restrict grain growth (see section 4.1) during superplastic deformation at high temperatures. This will enhance

42

superplasticity and improves mechanical properties of post superplastically formed sheets. Measured activation energy for deformation indicates GBS as the rate controlling mechanism for superplasticity. Second, and in contrast, TiB needles present at the grain boundaries will resist GBS and hence reduce the overall strain that can be achieved via superplastic deformation. As the VTiB in the work of Sinha et al. [161] is low (VTiB ~ 0.005), these two factors will offset each other and make superplastic forming characteristics of Ti64-0.1B sheet similar to conventional Ti64. 9. Joining Ti alloys are joined conventionally using fusion welding, solid-state welding and Brazing [10]. Only the effect of B on the fusion welding of Ti alloys is explored in this section as the grain size refinement effect due to B addition would be ineffective in solid state welding and Brazing, as these processes involve no direct melting of the BMs to be joined. Trace additions of B have been shown to refine the β grain size in the FZ in a variety of fusion welding processes like laser welding [162], WAAM [39] and GTAW [73]. The main aim of these works was to take advantage of the grain refinement effect of B and Zener pinning effect of TiB needles at the β grain boundaries to prevent excessive grain coarsening in the FZ and HAZ, which otherwise would have degrade strength and hardness of the weld. Tolvanen et al. [162] examined laser weldability of as-cast Ti64 alloy plates with 0, 0.06 and 0.11 wt.% B content. Fig. 44 shows a cross-sectional images of the laser welded as-cast Ti64-B plates [162]. EBSD generated IPF map of α phase and β reconstruction maps (based on BOR between parent β and daughter α phase [163]) of the BM, HAZ and FZ are shown in Fig. 45. Based on Figs. 44 and 45, following observations were made [162]. In the B-free plate, coarse columnar β grains, with size in the order of few ~ mm, was observed. In contrast, B-containing laser welded plates showed much smaller FZ size with refined columnar β grains. Zener pinning effect of TiB needles restricts β grain growth in the HAZ of B-containing plates during welding. As explained by Bermingham et al. [39] and Tolvanen et al. [162], growth restriction effect of the B solutes due to CS prevents lateral growth of the columnar β grain. This allows more nucleation and growth of neighboring columnar grains resulting in much finer columnar β grains in the FZ of B-alloyed laser welded plates. The hardness of the FZ, HAZ and BM was found to be higher in the B-alloyed welds [162]. Although the addition of B seems to be beneficial during laser welding of as-cast Ti64 alloy, 43

Tolvanen et al. highlight the potential for increased susceptibility to liquation cracking in the HAZ. It was observed that at the vicinity of the fusion line, TiB needles were partially melted/dissolved while the β grains remains unmelted. Ti-B eutectic has a lower melting point (1540 °C) than pure Ti (1670 °C, see Fig. 3) and forms at the prior β grain boundaries during solidification. If the local temperature in the HAZ exceeds Ti-B eutectic temperature, such partially melted eutectic phase may lead to the formation of a liquid film along the prior β grain boundaries in the HAZ during welding. Although, no cracks were observed in the study of Tolvanen et al., welding parameters like welding speed and heat input should be carefully optimized to minimize the scenario of liquation cracking. Anis et al. [73] explored the use of CP-Ti fillers with varying B contents (0, 0.5 and 1 wt.%) to improve the weldability and mechanical properties during GTAW of metastable β Ti alloy, Ti-15-3, sheet. For comparison purposes, autogenous (no filler) Ti-15-3 welds made using identical welding parameters were also examined. Fig. 46 shows the microstructures in the FZ of Ti-15-3 autogenous weld and welds prepared using CP-Ti fillers with varying B contents. It was found that microstructures of autogenous Ti-15-3 weld and weld prepared using CP-Ti filler are similar and showed coarse columnar dendritic β grains in the FZ (Fig. 46 a and b). On the other hand, significant grain refinement in the FZ was observed in the welds prepared using the latter (Fig. 46 c and d). An additional important observation was the change in the morphology of columnar β grains from flat in autogenous and weld prepared using CPTi filler to equiaxed dendrite in welds prepared using B-modified CP-Ti fillers. Fig. 47(a) show the variation of hardness across the weldline [73]. Welds prepared using CP-Ti fillers with 0.5 and 1 wt.% B showed maximum hardness in the FZ while its values remains similar in the BM and HAZ. Ti-15-3 autogenous weld and weld prepared using CP-Ti filler showed similar hardness values in the BM, HAZ and FZ. Tensile tests performed on the weld samples and BM indicates CP-Ti filler with 0.5 wt.% B gave the best combination of strength and ductility, see Fig. 47(b) [73]. Marginal loos in the ductility of weld prepared using CP-Ti filler with 1 wt.% B was observed due to increased VTiB in the FZ. Therefore, a good potential exists to refine the grain size of Ti alloys in the FZ with subsequent improvement in the mechanical properties using optimum amount of B. Post weld heat-treatments can further improve the strength and ductility of the welds

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either through precipitation or avoidance of grain coarsening due to Zener pinning effect of TiB needles present in the FZ. 10. Wear Dixit et al. [164] examined pin-on-disc dry sliding wear behavior of as-cast Ti64 alloys with varying B contents against hardened steel at temperatures between 20-300 °C. Fig. 48(a) shows wear rate as a function of B content at different temperatures. They report that B additions up to 0.3 wt.% decrease the wear rate at all the examined temperatures. This was expected as the wear resistance is inversely related to material’s hardness [165]. Although the hardness of Ti64 alloy increases with B content (see section 5.3) but adding B beyond 0.3 wt.% increases wear rate. To determine the underlying wear mechanism, the worn surfaces morphologies of as-cast Ti64-B alloys tested at room and elevated temperatures are examined inside SEM [165]. At RT, samples showed severe plastic deformation and wear mechanism was predominately adhesion followed by abrasive (Fig. 48b). At high temperatures, cutting and plowing of the surfaces with deeper grooves along with sub-surface delamination was more prominent (Fig. 48c). The initial decrease in the wear rate with B addition up to 0.3 wt.% at all the temperatures investigated was attributed to refinement in dβ and dα [164]. The increased fraction of brittle TiB needles in 0.55B alloy (increases from 1.8% in 0.3B to 2.7% in 0.55B alloy) underwent debonding from the matrix as a result of contact loading dominated by shear stress. These hard debonded TiB needles penetrate deeper into the soft matrix, accelerate the wear process and eventually change the wear mechanism from two-body to three-body [164]. 11. Oxidation resistance The high temperature capability of Ti64 alloy is often limited to 300 °C mainly due to its poor oxidation resistance [166]. During prolonged exposure of Ti alloys at elevated temperatures, TiO2 oxide scale present at the surface loses its protective nature with the formation of O enriched layer beneath the scale, commonly known as α-case [167,168]. Subsequent growth of the oxide scale results in cracking, decohesion at the oxide layersubstrate and ultimately loss in load bearing capacity with material failure. To examine whether addition of B is beneficial in improving the oxidation resistance of Ti64 alloy, Brice et al. [169] studied the effect of 1 wt.% B on the oxidation behavior of Ti64 sheet in atmospheric air at temperatures between 650-950 °C for exposure times of 25 and

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50 h. The mass gain per unit surface area as a result of ingestion of O was measured and compared with B-free Ti64 alloy in Fig. 49 [169]. At 650 °C, weight gain as a result of oxidation was found to be negligible. There was a marked increase in weight gain at oxidation temperatures above 750 °C. Clearly at a given temperature and exposure time, addition of 1 wt.% B in Ti64 retards weight gain and improves its oxidation resistance. This is consistent with the finding of Luan et al [170], where improvement in the oxidation resistance of Ti64 alloy with just as 0.02 wt.% B addition was reported at 400 °C for exposure times up to 650 h. SEM micrographs of the oxide scale-substrate of Ti64 with 0 and 1 wt.% B after exposure at 950 °C/50 h are shown in Figs. 50 (a) and (b) [169]. As seen from Fig. 50(a), the extent of cracking in the oxide layer and spallation is much higher in B-free Ti64 alloy. On the other hand, the thickness of the oxide layer is smaller and well adhered to the substrate in Ti64-1B (Fig. 50b). Brice et al. [169] also observed formation of decomposed structure beneath the oxide/substrate interface with depleted β phase content (Fig. 50c). This decomposed structure forms as a result of chemical changes due to ingestion of O [169]. Detailed characterization of the decomposed structure shows bright grey phase in Fig. 50(c) is O-rich α and the dark grey phase is Al rich Ti3Al. The depth of this decomposed structure was found to be smaller in Ti641B alloy. From the work of Brice et al., it is clear that addition of B is beneficial in improving the oxidation resistance of Ti64 alloy in terms of lower weight gain, lower oxide scale growth and smaller depth of decomposed structure. However, the exact mechanism of how B or TiB improves oxidation resistance of Ti alloys is not very well understood in the literature [171–174]. The oxidation of TiB produces TiO2 and B2O3 [171,172]. B2O3 evaporates at temperatures above 750 °C and creates pores at the TiBTi interfaces [172]. This should decrease the oxidation resistance as cracks and pores will provide easy diffusion path for O. Surprisingly, no pores or crack formation around TiB-Ti interfaces was detected in [169]. Brice et al postulated that B somehow lowers volume and grain boundary diffusion of O either in the oxide layer or BM necessary for oxide scale growth. Much thinner oxide scale with fewer cracks in Ti64-1B alloy will further limit O diffusion, thereby reducing O ingestion and formation of decomposed structure.

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12. Machining Ti and its alloys are generally considered difficult materials to machine in comparison to steel or Al alloys owing to several inherent material properties like low thermal conductivity, high chemical affinity with all cutting tool materials and spring back effect [175]. At the time of writing this paper there is no reported work on the effect of B on the machinability of Ti alloys using conventional methods like chip forming machining, only non-conventional machining method like EDM is examined [176]. There are several unique advantages of EDM of hard materials like Ti such as no physical contact between the tool and the workpiece, can machine complex geometries, burr-free surfaces, less sub-surface damage, distortion etc [177,178]. Main output parameters of EDM like MRR and TWR of as-cast Ti64 alloy with 0, 0.04 and 0.09 wt.% B are determined by Sen and co-workers at various discharge energies using Cu as an electrode or tool material [176]. It was found that MRR remains invariant while TWR improves significantly with B addition in Ti64. For example, a 0.09 wt.% B addition in Ti64 decreases TWR by ~80% at 500 μJ energy level. MRR during EDM operation depends upon the hardness and the melting point of the material [176]. Conventionally, MRR decreases with increase in hardness and melting point of the material. In the work of Sen et al. [176], addition of 0.09 wt.% B in Ti64 increases its hardness just by ~1.6% while melting point decreases merely by ~10 °C. These two factors, although changes negligibly with B addition, will tend to offset each other and make MRR invariant with B addition [176]. On the other hand, TWR improves with enhancement in strength of the workpiece. Strength improvement with microstructural refinement with B addition and shielding effect of TiB needles with higher melting point (~2200 °C) and hardness (Hv ~ 700 VHN [119]) was attributed to lesser tool wear in B-modified alloys [176]. However, no attempts were made to investigate how B influences surface integrity of the alloys after EDM. Nevertheless, reduction in TWR during EDM of B-modified Ti64 alloys can offer unique advantages in terms of tool service life and cost. 13. Hydrogen embrittlement In certain applications, Ti alloys components are exposed to hydrogen containing environments. Examples are hydrogen fuel containers in rocket engines, and during descaling and pickling operations to remove oxide scales [179]. In such cases, the

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presence of H either as interstitial solid solution or in the form of δ-hydride (TiH2), which forms when the H concentration exceeds critical level like ~0.3 wt.% in Ti64 [180], can have a severe adverse effect on both strength and ductility of Ti alloys [179– 182]. This will limit the alloy’s applicability in such scenarios. The hydrogen embrittlement behavior of B-modified Ti64 alloys is addressed in [183] with the main objective to determine whether B addition improves relative mechanical performance of Ti64 in hydrogen environments. For this purpose, as-cast Ti64-xB (x = 0, 0.04, 0.09, 0.3 and 0.55 wt.%) alloys were hydrogen charged at 500 and 700 °C for times, t, up to 240 and 120 minutes respectively. High incubation period for hydrogen absorption at 500 °C was noted while hydrogen uptake increases substantially with increasing temperature to 700 °C. Combined TEM and XRD analyses indicate formation of δ-hydride phase only in the alloys charged at 700 °C for 2 h. The δ-hydride phase has a needle like morphology (see Fig. 51) and nucleates at the α-β interface and grows within α laths through diffusion-controlled processes [180]. Low volume misfit ΔV/V associated with α→ δ-hydride transformation [184] and comparatively lower solubility of hydrogen in α phase [185] are possible factors promoting the δ-hydride formation in α phase of Ti64-xB alloys. Tensile tests performed on the hydrogen charged samples at RT indicate marginal enhancement in strength for t up to 4 h at 500 °C and up to t = 30 min at 700 °C. This enhancement in strength was attributed to solid solution strengthening by hydrogen. Ductility drops in all the cases due to hydrogen induced localized plastic deformation where it pins mobile dislocations. No specific trends in σy, σu and εf with B content in hydrogen charged alloys was observed [183]. In contrast, the samples charged at 700 °C for t = 2 h failed in a brittle manner, aided by the δ-hydride phase. The relative reduction in the strength of 0.3B and 0.55B alloys hydrogen charged at 700 °C for t = 2 h, as compared to their un-hydrogenated counterparts, is smaller in comparison to lean B containing alloys ≤ [𝜎𝑢]𝑡 = 2ℎ ― [𝜎𝑢]𝑡 = 0ℎ

0.1 wt.%. This is illustrated in Fig. 52 where ∆𝜎𝑢(=

[𝜎𝑢]𝑡 = 0ℎ

) is plotted against

VTiB, here [𝜎𝑢]𝑡 = 2ℎ and [𝜎𝑢]𝑡 = 0ℎ are the tensile strength of Ti64-xB alloys hydrogen charged at 700 °C for t = 2 h and in un-hydrogenated conditions respectively. This observation implies that a higher VTiB in the alloy makes it relatively less susceptible to hydrogen embrittlement.

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14. Tensile properties at cryogenic temperatures The applicability of α+β Ti alloys like Ti64 is limited to 77 K mainly due to their increased sensitivity to notches and low KIC at temperatures below 77 K [186–189]. Whether addition of B improves the tensile properties of Ti64 alloy at cryogenic temperatures was explored by Singh et al. [41] by testing unnotched as well as notched Ti64-B alloy bars at 77 and 20 K. Variations of σy, σu, and εf with B content are shown in Figs. 53 (a) and (b) [41]. RT tensile properties from [44] are also plotted in the same figure. Ti64-0.55B alloy failed in a brittle manner without any yielding at both the examined temperatures (Fig. 53a). The other four alloys showed some amount of plasticity as they exhibited yield before failure. Both σy and σu increases with decreasing T and/or increasing B content. This is consistent with the trend observed at RT by Sen et al. [44]. As expected, εf falls sharply with decreasing temperature (Fig. 53b). Interestingly, εf values of B-modified Ti64 alloys at 20 K were found to be insensitive to B content and were significantly lower than that of B-free Ti64 alloy, which exhibited exceptionally high ductility (~5.85 %). Post-deformation microstructural characterisation of the alloys with the aid of EBSD and TEM revealed activation of deformation twinning at 20 K. No evidence of twinning was detected in samples tested at 77 K and its deformation was more dominated by slip. Extensive twinning was observed in B-free Ti64 alloy due to its coarser microstructure in comparison to B-containing alloys where twinning activity was very limited. Operating twin modes were further identified based on the unique axis/angle misorientation pair between the twin and the matrix. {1012}, {1121} and {5 613} twinning modes were indexed in B-free Ti64 alloy while only {1012} mode in the B-containing alloys (Fig. 54). Singh et al. [41] attributed high εf and low σy at 20 K to the activation of these additional twin modes in B-free Ti64 alloy. Due to low fraction of only detected {1012} twins in B-containing alloys, its contribution to overall plastic strain was found to be insignificant. Microstructural refinements and decrease in the Sα with B addition were possible reasons proposed for suppression of these additional {11 21} and {5613} twinning modes in B-containing alloys. The variation of NSR (= σn/σu) with B content at 77 and 20 K is displayed in Fig. 55 [41]. The NSR gives a qualitative indication of notch-sensitivity of the material. The transition from notch-sensitive to notch-insensitive condition occurs when NSR > 1. Fig. 55 shows that NSR decreases with increasing B content and temperature. NSR’s

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of Ti64 alloys with B content ≤ 0.09 wt.% were greater than 1 at 77 K but slightly lower than 1 at 20 K with NSR = 0.83 being the worst case for 0.09B alloy. The values of NSR for 0.3B and 0.55 B alloys were quite low and varied between 0.55-0.65. The increased fraction of brittle TiB needles in 0.3B and 0.55B alloys makes these alloys brittle and unsuitable for cryogenic applications. To summarize, results of the smooth and notch tensile tests clearly indicates addition of B up to 0.1 wt.% in as-cast Ti64 alloy enhances its cryogenic strength with no detrimental effect on the ductility. 15. Biomedical properties Excellent corrosion resistance in human body fluid, biocompatibility, osseointegration and biofunctionality of Ti makes it ideal choice of material for medical and dental applications [3,190–192]. Elimination or reduction of primary hot-working steps with B addition in Ti alloys can be exploited to reduce the high manufacturing cost associated with fabrication of implants. The potential of Ti64-B alloys for possible biomedical applications was explored by Bahl et al. [193]. The corrosion behavior of Ti64-0.1B alloy in forged and extruded forms was examined in SBF and compared with B-free Ti64 alloy in identical conditions. The experiments carried out by Bahl et al. revealed that corrosion rate of forged and extruded Ti64-0.1B alloys were higher compared to B-free counterparts. They proposed that both α and β phases becomes anodic with respect to TiB needles, leading to their preferential dissolution with increased corrosion rate. Biocompatibility of the alloys measured using osteoblast response was found to be unaffected by the B content or processing route [193]. β-Ti alloys with fully biocompatible elements like Nb, Ta, Zr etc. [194] and Young’s modulus close to that of bone (10-30 GPa) [192] are being intensively developed for biomedical applications. Sliding wear and mechanical properties of Ti13Nb-13Zr alloy have been reported to improve with addition of 0.5 wt.% B [195]. In an another work, Málek et al. [196] performed cytocompatibility tests on Ti-35Nb-6TaxB wires. It was found that B content up to 0.05 wt.% exhibits no cytotoxicity, i.e., the number of living cells is independent of dilution rate. Increasing B content above a certain level (> 0.1 wt.%), the alloy possesses slight cytotoxicity (approximately 81% of the living cells). Nevertheless, further studies are clearly needed to evaluate complete biological impact of B in Ti alloys.

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16. Summary and outlook The addition of trace B addition to commercial Ti alloys like Ti64 refines both the prior β grain size and α-colony size by roughly an order of magnitude in the cast state as a result of constitutional supercooling. Results of the microstructural and mechanical property studies, reviewed in this paper, show that the optimum quantity for best property combinations is ~0.1 wt.% of B. While the order of magnitude microstructural refinements enhance the as-cast mechanical properties vis-á-vis B free alloys, the improvements in wrought products are not significant. The main advantage here would be the marked reductions in the number of themo-mechanical processing steps required for breaking the (otherwise) coarse primary ingot structure, which could translate directly into substantial material and cost savings. For high temperature applications one may consider B additions higher than 0.1 wt.%, as the network of TiB needles present at the grain boundaries with necklace like arrangement can provide extra stiffening to the matrix and restrict grain growth at high temperatures. Trace addition of B during fusion welding and additive manufacturing of Ti alloys was found to be beneficial not only in refining the grain size but also for changing the morphology of β from columnar to equiaxed. In all, it is quite clear that a small addition of B to Ti alloys can only be beneficial or, at worst, neutral in terms of microstructures and properties. Since the trace additions do not require new manufacturing infrastructures or component design methodologies (as compared to, say, composites with anisotropic properties or machining issues), it is surprising to see that large scale adaptation of B-modified Ti alloys has not been implemented yet. The arguments of industrial practitioners that segregation of such modified alloys' scrap (such as machining chips), especially the high grade ones, becomes paramount as inadvertent mixing of B-modified and B-free alloys' waste could cause problems, while valid, does not appear to be insurmountable. Thus, it is hoped that B-modified Ti alloys will start seeing large scale applications in various industries, given the high cost of Ti alloys' finished products.

Acknowledgements The authors would like to acknowledge financial support provided by the Boeing Company, USA through research grant SID/PC36013. GS would like to thank EPSRC

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for financial support during his post-doctoral period through Light Alloys for Sustainable Transport 2nd Generation (LATEST2) [EP/H020047/1] grant.

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List of Tables Table 1. Room temperature tensile properties and microstructure parameters of Bmodified Ti alloys Alloy

B content (wt.%)

Processing technique

Microstructural parameters

E (GPa)

Tensile properties σy σu (MPa) (MPa)

εf (%)

Ref.

Near α-Ti alloys 0 0.2

Ti-1100

0

0.2

As-cast+ ST/1095 °C/2h/AC +aged/595°C/8h/AC Wrought (VAR + α+β rolled@875°C+ ST/1095 °C/2h/AC +aged/595°C/8h/AC) Wrought (VAR + α+β rolled@980°C +ST/1095 °C/2h/AC +aged/595°C/8h/AC)

…. STiB = 18, VTiB = 1.7% ….

STiB = 6, VTiB = 1.7%

112

937

1050

3.9

125

938

1037

1.3

114

931

1027

3.0

128

963

1112

7.5

977

1135

9.2

1044

1196

8.2

1033

1180

8.1

1085

1232

7.3

1092

1226

7.6

1153

1283

6.4

….

1125

1248

5.8

….

1183

1301

3.3

….

826

920

5

….

828

917

2.5

….

850

864

0.25

[32]

Wrought (α+β rolled +ST/1100°C/2h/FC + ST/ x °C/2h/OQ + aged/700 °C/2h/AC)

Ti834

0

x = 945

0.2

x = 975

0

x = 995

0.2

x = 1025

0

x = 1025

0.2

x = 1055

0

x = 1055

0.2

x = 1085

0 Ti685

0.2 0.5

As-cast + heat treated (VAR+ST/1050°C/30 min/OQ + aged/550°C/24 h/AC)

dβ = 28 μm, dα = 2.122.15 μm dβ = 20 μm, dα = 2.132.27 μm dβ = 43 μm, dα = 2.122.31 μm dβ = 31 μm, dα = 2.2-2.9 μm dβ = 55 μm, dα = 2.12.25 μm dβ = 42 μm, dα = 2.22.30 μm dβ = 108 μm, dα = 2.22.40 μm dβ = 92 μm, dα = 2.32.36 μm dβ = 1520 μm, dα = 1.49 μm dβ = 270 μm, dα = 5.20 μm, VTiB = 1.44%, STiB = 11.85 dβ = 225 μm, dα = 6.65 μm, VTiB = 4.11%, STiB = 15.86

66

…. …. …. …. …. ….

[30] [46]

[31]

Alloy

B content (wt.%)

0 0.1 Ti-1100

0 0.1

Processing technique

Microstructural parameters

Wrought (VAR + β forged + α+β rolled + ST/x °C/2h/AC +aged/595°C/8h/AC) x = 975 x = 1000 x = 1050 x = 975 x = 1000 x = 1050

….

Wrought (VAR + β forged + β rolled + ST/x °C/2h/AC +aged/595°C/8h/AC) x = 975 x = 1000 x = 1050 x = 975 x = 1000 x = 1050

E (GPa)

Tensile properties σy σu (MPa) (MPa)

εf (%)

…. …. …. …. …. ….

896 956 917 938 899 902

988 1032 1016 1011 1013 972

11.8 15.5 6.7 14.4 15.8 15.7

…. …. …. …. …. ….

806 845 831 932 924 952

902 933 933 1064 1055 1084

11.9 14.0 13.3 10.5 12.6 12.7

113

784.8

830.6

13.2

121

808.2

875.1

15.2

114

828.4

912.3

17.6

120

858.2

948.9

15.1

126

876.3

964

12.4

….

….

1007

1098

7.8

dβ = 100 μm, c = 50 μm, STiB = 10

….

939

1034

19.3

dα = 8 μm

….

979

1053

18.8

dα = 8 μm, STiB = 10

….

998

1080

19.4

dβ = 2386 μm, c = 250.8 μm, dα = 2.3 μm dβ = 777.4 μm, c = 62.7 μm, dα = 3.8 μm, VTiB = 0.2%

….

890.7

991

15.6

811.6

924

13

VTiB = 0.7%, STiB = 4.3

…. VTiB = 0.7%, STiB = 7.4

Ref.

[47]

α+β-Ti alloys 0 0.04 Ti64 0.09

As-cast (ISM+HIPed@900°C/2 h)

0.30 0.55 0 0.1 Ti64 0 0.1

0 Ti64 0.04

Wrought (Levitation melted+β/α+β processed + ST/1050°C/1h/0.03Ks1) Wrought (Levitation melted+β/α+β processed + ST/930°C/0.03Ks-1) Wrought (ISM+HIPed@900°C+ α+β forged)

dβ = 2386 μm, c = 244.4 μm, dα = 2.48 μm dβ = 733 μm, c = 22.2 μm, dα = 6.05 μm, STiB = 9-10 VTiB = too low to measure dβ = 224 μm, c = 37.8 μm, dα = 4.52 μm, STiB = 9-10, VTiB = 0.5% dβ = 121 μm, c = 23.7 μm, dα = 4.69 μm, STiB = 9-10, VTiB = 1.8% dβ = 100 μm, c = 21 μm, dα = 5.29 μm, STiB = 910, VTiB = 2.9%

67

[44,92]

….

[48]

[112]

Alloy

B content (wt.%)

Processing technique

0.09

Ti64

1.0

0

VT8

0.2 0 0.2

P/M Extruded P/M Extruded +MA (735 °C/2h/AC) P/M Extruded + DA (940 °C/10 min/AC + 675 °C/4 h/AC) P/M Extruded + STA (940 °C/10 min/WQ + 525 °C/4h/AC) P/M Extruded + BA (1035 °C/30 min/AC + 730 °C/2h/AC) P/M Extruded +HBA (1200 °C/30 min/AC + 730 °C/2h/AC) 3D forging (650700°C)+ β-heat treatment (1030°C/30 min/FC) 3D forging (650700°C)+ α+β-heat treatment (950°C/1h/FC+590°C/ 2h/FC)

Microstructural parameters

E (GPa)

Tensile properties σy σu (MPa) (MPa)

εf (%)

dβ = 216 μm, c = 24.1 μm, dα = 2.6 μm, VTiB = 0.5%

….

831

927

15.6

VTiB = 6%

142

1070

1219

13.3

VTiB = 6%

138

1089

1168

11.7

VTiB = 6%

141

1078

1186

13.4

VTiB = 6%

138

1259

1366

9.6

132

986

1191

11.6

128

985

1186

10.3

….

….

902.3

9

….

….

970.5

19.22

….

….

938.3

19.1

dα = 5 μm, VTiB = 1.0%

….

….

963.7

17.1

VTiB = 6% VTiB = 6%, c = 10 μm, dα = 2 μm dβ = 600 μm, c = 230 μm, dα = 3.4 μm dβ = 60 μm, c = 20 μm, dα = 3.5 μm , VTiB = 1.0% dα = 5 μm

Ref.

[58]

[111]

β-Ti alloys Beta 21S Ti5553

0 0.1

As-cast (ISM + HIPed@900°C/2h)

dβ = 150 μm dβ = 50 μm, STiB = 3.8-9

89 92

944 951

965 963

2.3 5.5

0 0.1

As-cast (ISM + HIPed@900°C/2h)

dβ = 350 μm dβ = 50 μm, STiB = 2-3

105 112

966 1033

1071 1112

8.4 6.0

0

Wrought (VAR + Rolled@800 °C + ST/900 °C/30 min/OQ)

dβ = 211 μm

…. ….

567

734

34

706

801

21.7

0.2 Ti-15-3 0 0.2

Wrought (VAR + Rolled@800 °C + ST/900 °C/30 min/OQ + aged/535 °C/8h/AC)

dβ = 135 μm, VTiB = 1%, STiB = 4.1 ….

….

811

903

14.8

VTiB = 1%, STiB = 4.1

….

905

1092

12.1

68

[36]

[42]

Table 2. High temperature tensile properties of B-modified Ti alloys B content Alloy

(wt.%)

Test Processing technique

E

σy

σu

εf

(GPa)

(MPa)

(MPa)

(%)

90

379

435

9.7

94

435

552

11.8

103

513

614

2.5

102

676

750

10.7

105

640

747

10.8

113

663

785

8.2

475 500 550 475 500 550 475 500 550

96.1 94.6 93 100.8 99.1 96.2 97.3 95.5 94

395.5 390.7 343 459.6 436 411.6 464.9 463 459.8

467.4 437.2 400.5 535.2 524.7 470.4 586 536.2 476.2

19.9 16.2 20.9 22.7 20.2 21.0 14.5 20.2 22.0

455 455 455 510 565 455 510 565

97 95 97 …. …. 112 …. ….

542 587 663 526 500 660 552 500

569 660 667 669 667 707 645 652

8.2 10.3 1.1 …. …. 1.3 …. ….

513 526 513 537

640 644 636 695

23.3 17.7 21.6 14.1

545 547 571 594

634 683 711 730

12.7 15.3 12.0 12.1

temperature (°C)

0 0.1

As-cast (ISM+HIPed@900°C/ 2 h)

1.0 Ti64

455

0 0.1

As-cast + extruded@1100 °C

1

0 Ti64

0.06

As-cast (ISM+HIPed@900°C/2 h)

0.11 0 0.1 Ti6242S

0.4

As-cast + HIPed@900 °C

1.0

Ref.

[95]

[97]

[96]

VAR + β forged@1100 °C + (α+β) rolled + ST/x °C/2 h + aged/595 °C/8h/AC x = 915 Ti-1100

0

0.2

x = 965 x = 985 x = 1025 600

x = 930 x = 980 x = 1000 x = 1040

69

….

[32]

B content Alloy

(wt.%)

Test Processing technique

σy

σu

εf

(GPa)

(MPa)

(MPa)

(%)

300 400 500

…. …. ….

…. …. ….

650.8 586.8 545.5

12.6 12 11.6

300 400 500

…. …. ….

…. …. ….

781.9 723.1 666.1

20.3 22.1 22.6

300 400 500

…. …. ….

…. …. ….

773.8 727 676.4

19.2 18.1 24.3

300 400 500

…. …. ….

…. …. ….

802.9 761.5 703.8

18.4 19.5 25.7

617

721

9.2

(°C)

0

0.2

E

temperature

3D forging (650700°C)+ β-heat treatment (1030°C/30 min/FC)

VT8

Ref.

[111] 0

0.2

0

Ti834 0.2

3D forging (650700°C)+ (α+β)-heat treatment (950°C/1h/FC+590°C/ 2h/FC)

VAR+β forged@1100°C+ (α+β) rolled+ ST/1025°C/2h/OQ +aged/700°C/2h/AC VAR+β forged@1100°C+ (α+β) rolled+ ST/1055°C/2h/OQ +aged/700°C/2h/AC

600

….

[30] 641

753

8.3

530

622

13

531

634

4

565

653

2

0 Ti685

0.2

VAR + ST/1050 °C/30 min/OQ + aged/550 °C/24h/AC

600

0.5

70

….

[31]

Table 3. Theoretical and experimental comparison of tensile yield strength in as-cast Ti64 alloy with varying B contents B content (wt.%)

∆𝜎𝐿 ― 𝑆

∆𝜎𝐻 ― 𝑃

0

….

….

….

Experimental Yield strength (MPa) 784.8

0.04



37.4

822.2

808.2

0.09

9.8

41.8

836.4

828.4

0.30

31.4

60

876.2

858.2

0.55

57

65.4

907.2

876.3

Theoretical Δσy (MPa)

Theoretical Yield strength (MPa)

 Negligible due to very low volume fraction of TiB needles  Values taken from Ref. [44].

71

Table 4. Measured creep parameters like creep strain at 100 h, ε100h, minimum creep rate, 𝜀𝑚𝑖𝑛, and time-to-rupture, tr in various B-modified Ti alloys. Alloy

B content (wt.%)

Processing technique

Test parameters (T/σ)

ε100h (%)

𝜀𝑚𝑖𝑛 (h-1)

tr (hrs)

Ref.

As-cast (ISM+HIPed@900°C/2h)

475°C /300 MPa 500°C /300 MPa 550°C /300 MPa 475°C /300 MPa 500°C /300 MPa 550°C /300 MPa 475°C /300 MPa 500°C /300 MPa 550°C /300 MPa

…. …. …. …. …. …. …. …. ….

1.67 x10-4 9.72 x10-4 1.09 x10-2 9.18 x10-5 4.9 x10-4 9.18 x10-3 1.59 x10-5 3.82 x10-4 8.1 x10-3

148.3 46.6 2.86 750 154.7 6.88 4722.3 185.3 7.66

[97]

0.72 0.39 0.35 0.13 0.24 0.22 0.18 0.08

4.3 x10-5 2.4 x10-5 1.8 x10-5 5.2 x10-6 6.8 x10-6 4.6 x10-6 2.4 x10-6 6.1 x10-7

…. …. …. …. …. …. …. ….

4.26 1.64 0.72 0.21

3.5 x10-4 1.48 x10-4 4.36 x10-5 5.30 x10-6

…. …. …. ….

0.0

Ti64

0.06

0.11

Ti834

0.0

0.2

Ti-1100 0.2

Ti6242S

Ti64 Ti64

0 0.1 0.4 1 0 0.1 1 1

VAR+β forged@1100°C+ (α+β) rolled+ ST/x °C/2h/OQ +aged/700°C/2h/AC x = 945 x = 995 x = 1025 x = 1055 x = 975 x = 1025 x = 1055 x = 1085 VAR + β forged@1100 °C + (α+β) rolled + ST/x °C/2 h + aged/595 °C/8h/AC x = 930 x = 980 x = 1000 x = 1040 VAR + β forged@1100 °C + β rolled + ST/x °C/2 h + aged/595 °C/8h/AC x = 930 x = 980 x = 1000 x = 1040

600°C/150 MPa

[30]

[115] 600°C/150 MPa

As-cast

455 °C/450 MPa

As-cast (ISM+HIPed@900°C/2h)

455 °C/450 MPa

P/M processed

455 °C/450 MPa

72

0.96 0.37 0.27 0.22 …. …. …. …. …. …. …. ….

2.58 x10-5 3.85 x10-6 9.50 x10-6 4.45 x10-6 1.19 x10-6 1.84 x10-6 1.49 x10-6 8.06 x10-7 4.46 x10-3 1.46 x10-3 2.1 x10-4 2.1 x10-4

…. …. …. …. …. …. …. …. …. …. …. ….

[96]

[95,96] [59]

Alloy

B content (wt.%)

Processing technique

Ti64

0 0.1 1

Cast-then-extruded (ISM+HIPed@900°C/2h+βextrusion@1100°C)

Test parameters (T/σ)

ε100h (%)

𝜀𝑚𝑖𝑛 (h-1)

tr (hrs)

Ref.

400 °C/450 MPa

…. …. ….

2.0 x10-6 5.8 x10-6 3.77 x10-6

…. …. ….

[95]

73

Table 5. Plain strain fracture toughness, KIC, values of B-modified Ti alloys. Alloy

Ti64

VT8

Ti834

Ti-1100

B content (wt.%) 0 0.05 0.1 0.4 0 0.2 0 0.2 0 0.2

0 0.1 0 0.1 0 0.1

Processing technique As-cast (ISM+HIPed@900°C/ 2 h) Forged (650-700°C)+ β heattreatment Forged (650-700°C)+ (α+β) heat-treatment Forged 1100°C + ST (Tβ20°C)/2 h/OQ +aged 700°C/2h/AC (α+β) rolled + ST (x°C/2h/AC)+aged 595°C/8h/AC x = 975 x =1000 x =1050

KIC (MPa 𝒎) 126* 73±4.4 63±6.4 38±2.6 69±2 64±2 71±2 67±2 37.4 32.3

50.9 39.6 53 41.4 58.7 44.1

Ref. [100]

[111]

[46]

[47]

β rolled + ST (x°C/2h/AC)+aged 595°C/8h/AC

0 59.9 x = 975 0.1 44.3 0 63.4 x =1000 0.1 45.3 0 70.6 x =1050 0.1 48.1 0 51±4 As-cast + ST 900°C/1h/WQ + Ti-15-3 [29] aged 535°C/8h/AC 0.2 36.6±1 * - Not valid plain strain fracture toughness test Compact-tension specimens were extracted from L-R direction for Ti834 and Ti-15-3 and L-T direction for Ti-1100.

74

List of Figures

Figure 1. Use of Ti alloys by weight in commercial aircrafts over the last several decades [6–8]. The size of the rectangles and circles refers to aircraft average passenger capacity.

75

Figure 2. (a) Cost breakdown of Ti at various stages of component fabrication [9]. (b) Typical steps involved during production of commercial Ti alloys [10]. Trace B addition in Ti alloys can potentially reduce or eliminate β working steps of as-cast ingot.

76

Figure 3. Binary Ti-B phase diagram with B content up to 50 at.% [18].

77

Figure 4. BF TEM micrograph of TiB needle showing clean interfaces between TiB and the matrix. The transverse cross-section image of TiB needle is shown in the upper inset. Images reproduced with permission from [19].

78

Figure 5. Macrostructures of the transverse section of Ti-1100 ingots with (a) 0 and (b) 0.2 wt.% B content. Regions I and II indicated with arrows in (a) shows columnar and equiaxed regions of the ingot. Images reproduced with permission from [32].

79

Figure 6. Variation of as-cast prior β grain size with B content in various commercial Ti alloys [17,30,31,36].

80

Figure 7. Illustration of grain refinement mechanism and solidification route of Ti alloys with B content in the hypoeutectic regime. Image reproduced with permission from [17].

81

Figure 8. Microstructures of as-cast Ti64 and Beta 21S alloys with 0 and 0.1 wt.% B. (a) Ti64, (b) Ti64-0.1B, (c) Beta 21S and (d) Beta 21S-0.1B. Images reproduced with permission from [36,41].

82

Figure 9. Variation of prior β grain size, dβ, and α colony size, c, with B content in as-cast Ti64 alloy [44].

83

Figure 10. Pseudo-colored OM maps of Ti5553-0.5B alloy showing (a) equiaxed and (b) lathlike morphology of α precipitates around TiB needles. Images reproduced with permission from [50].

84

Figure 11. Production routes of B-modified Ti alloys [17,39,55,56].

85

Figure 12. Longitudinal cross-sectional ( ∥ BD) microstructures of DED-L Ti64 alloy with (a) 0 and (b) 0.17 wt.% B content. Images reproduced with permission from [56]

86

Figure 13. Beta grain size, dβ vs. annealing time plot of as-cast Beta 21S alloy with 0 and 0.1 wt.% B content at different annealing temperatures. Image reproduced with permission from [80].

87

Figure 14. (a) Auger spectra obtained from the grain boundary (black) and transgranular regions (blue) of fractured Ti64-0.05B specimen. (b) SEM fractographs obtained from the tensile tested (b) Ti64 and (c) Ti64-0.05B specimens. Images reproduced with permission from [91].

88

Figure 15. Variation of experimentally determined elastic modulus, E, with B content in as-cast Ti64 alloy. Estimated values of E determined using isostrain ROM are shown with blue squares. Necklace like arrangement of TiB needles at the prior β grain boundaries in Ti64-0.55B alloy is shown in the inset of the figure. Image reproduced with permission from [92].

89

Figure 16. Variation of (a) ultimate tensile strength and (b) strain-to-failure in various as-cast Ti alloys with B content at RT.

90

Figure 17. SEM fractographs of Beta 21S alloy with (a) 0 and (b) 0.1 wt.% B content tensile tested at RT. Image reproduced with permission from [36].

91

Figure 18. Microstructures of P/M processed-then-extruded Ti64 and Ti64-1B alloys after various heat-treatments. See Table 1 for heat-treatment details. Images reproduced with permission from [58].

92

Figure 19. Variation of yield and ultimate tensile strength in P/M processed-thenextruded Ti64 and Ti64-1B alloys in different heat-treatment conditions [58].

93

Figure 20. Variation of yield strength, ultimate tensile strength and strain-to-failure with solution heat-treatment temperature in wrought Ti834 alloy with 0 and 0.2 wt.% B content [30]. B-free and B-modified Ti834 alloys are shown with solid and broken lines respectively.

94

Figure 21. Engineering stress vs. strain curves of as-cast Ti64 alloy with 0, 0.06 and 0.11 wt.% B content tensile tested at 500 °C. Image reproduced with permission from [97].

95

Figure 22. Variation of Vickers hardness in as-cast Ti64-B alloys as a function of inverse of square root of prior β grain size, dβ-1/2 and α-lath size, dα-1/2 [100]. Vickers hardness tests were performed at a load of 0.5 Kg and hold time of 10 s at the peak load.

96

Figure 23. (a) Maximum stress, Smax, vs. Number of cycles to failure, Nf, plot of Ti640.09B alloy in as-cast and wrought conditions [44,112]. Arrows indicate run-out specimens. (b) Variation of fatigue strength, σFS, at 106 cycles vs. B content in ascast and wrought Ti64. Image in the inset shows crack front profile in wrought Ti640.04B alloy fatigue tested at Smax = 350 MPa and survived Nf = 4.8 x 106 cycles. Crack propagation along GB-α phase is shown with white arrow. Images reproduced with permission from [112]. Rotating bending fatigue tests were performed at R = -1, Kt =1, f = 100 Hz and T = 20 °C.

97

Figure 24. Maximum stress, Smax, vs. Number of cycles to failure, Nf, plots of wrought Ti64 alloy with 0 and 0.1 wt.% B content in lamellar and equiaxed forms [48]. Fatigue tests were performed at R = 0.1, Kt =1, f = 10 Hz and T = 20 °C.

98

Figure 25. SEM fractograph showing subsurface fatigue crack initiation site (highlighted with circle) in as-cast Ti64-0.3B alloy specimen (Smax = 425 MPa and Nf = 1.04 x 107 cycles). Image reproduced with permission from [44].

99

Figure 26. Maximum stress, Smax, vs. Number of cycles to failure, Nf, plots of Ti64 and Ti6242S alloys in different conditions. (a) Ti64 with 0, 0.1 and 1 wt.% B in ascast condition [123]. (b) Ti6242S with 0, 0.1, 0.4 and 1 wt.% B in as-cast condition [124]. (c) Ti64 with 0, 0.1 and 1 wt.% B in cast-then-extruded condition [123]. (d) P/M processed Ti64-1B alloy in rolled and extruded forms [121]. Test parameters T = 455 °C, R = 0.1 and f = 5 Hz. Images reproduced with permission from [121,123,124].

100

Figure 27. (a) Cyclic and monotonic stress-strain curves of as-cast Ti64 alloy with 0.09 wt.% B content. Image in the inset shows degree of cyclic softening (Σ𝑆 = 𝜎𝑦 ― 𝜎′𝑦). (b) Variation of Σ𝑆 and ∆𝜎 (=𝜎𝑢 ― 𝜎𝑦) with B content. Images reproduced with permission from [44]. Cyclic stress-strain curve is generated at R = -1, T = 20 °C, ΔεT/2 = 1.0% and f = 0.1 Hz.

101

Figure 28. ΔεT/2 vs. Nf plots of as-cast Ti64 alloy with 0, 0.06 and 0.11 wt.% B contents. Fatigue test parameters R = 0, ΔεT/2 = 0.25-1%, Kt = 1, T = 20 °C, f = 0.25 Hz. Image reproduced with permission from [126].

102

Figure 29. SEM fractograph obtained from the Ti64-0.11B alloy specimen LCF tested at ΔεT/2 = 1%. Cracking and decohesion of the TiB needle is shown with arrow. Image reproduced with permission from [126].

103

Figure 30. Cyclic stress response curves of wrought Ti834 alloy with 0 (base) and 0.2 wt.% B (modified) content at RT and 600 °C. Fatigue test parameters R = -1, ΔεT/2 = 1%, Kt = 1. Image reproduced with permission from [46].

104

Figure 31. Typical creep curves of Ti64-B alloys in different forms. (a) as-cast, T = 475°C, σ = 300 MPa [97]. (b) cast-then-extruded, T = 455°C, σ = 400 MPa [95]. (c) as-cast and P/M processed, T = 455°C, σ = 450 MPa [59,95]. Note unit of time in (b) is in ks. Images reproduced with permissions from [95,97].

105

Figure 32. (a) BF TEM micrograph of as-cast Ti64-0.06B alloy creep tested at T = 500 °C and σ = 300 MPa. Arrow indicates pile-up of dislocations at the α-β interface. (b) Heavy dislocation activity around TiB whisker/matrix interface in as-cast Ti640.11B alloy creep tested at T = 475 °C and σ = 300 MPa. Stacking faults are visible within TiB needles. Images reproduced with permission from [97].

106

Figure 33. (a) SEM fractograph of as-cast Ti64-0.11B alloy creep tested at T = 475 °C and σ = 300 MPa. Arrows indicate TiB needle cracking and decohesion. (b) Longitudinal section image of as-cast Ti64-0.06B alloy creep tested at T = 500 °C and σ = 300 MPa. Cavity formations along TiB needle and α-β interfaces are shown with circles. Images reproduced with permissions from [97].

107

Figure 34. Fatigue crack propagation curves, da/dN vs. ΔK, for as-cast Ti64 alloy with 0, 0.05, 0.1 and 0.4 wt.% B contents. Arrows indicate fatigue threshold ΔKo (da/dN < 10-9 m/cycle) and its values are shown in the inset of the figure. (b) FCG in the near-threshold regime of as-cast Ti64-0.4B alloy showing interaction of cracks with TiB needles. The FCG experiments were carried out T = 20 °C, R = 0.1 and f = 10 Hz. Images reproduced with permission from [100].

108

Figure 35. Room temperature dynamic compressive true stress vs. true strain plots of Ti64 alloy in different microstructural conditions. Tests performed at average strain rates, , varying between 3100-4000 s-1 [138]. (b) Adiabatic shear band formation, TiB needles cracking and cavity formation (shown with circle) in wrought Ti64-0.1B alloy in Bi-modal microstructural condition tested at a strain rate of 𝜀 = 3100 s-1 [138].

109

Figure 36. Variation of σy, σu and εf in DED-L Ti64 alloy with B contents in (a) ⊥ BD and (b) ∥ BD. (c) Percentage anisotropy in σy, σu and εf with B content in DED-L Ti64 alloy. Images reproduced with permission from [56].

110

Figure 37. Beta transus, Tβ, temperature of Ti64, Ti834 and Ti-1100 alloys as a function of B content in different forms [23,32,46,143]. Solid line shows predicted Tβ in as-cast Ti64 alloy.

111

Figure 38. As-rolled plates produced directly from (a) PAM Ti64, (b) PAM Ti640.1B and (c) ISM Ti64-0.1B ingots. PAM and ISM ingots rolled at 982 and 954 °C respectively. Arrows in (a) shows surface and edge cracking. Images reproduced with permission from [35].

112

Figure 39. (a) Microstructure obtained after direct rolling of ISM Ti64-0.1B ingot to 75% reduction in thickness at 982 °C. Alignment of TiB needles along RD. (b) Void formation around TiB needles in Ti64-0.1B alloy rolled at a lower temperature of 750 °C (black arrows). White arrow in (a) and (b) indicate RD. Images reproduced with permission from [35,148].

113

Figure 40. Images of the bent (a) PAM Ti64-0.1B, (b) ISM Ti64-0.1B and (c) PAM Ti64 sheets after double bent tests. Images reproduced with permission from [35].

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Figure 41. (a) Power dissipation (η) and (b) instability (ζ) maps of as-cast Ti64-0.09B alloy. Maps generated at a ε = 0.5. Images reproduced with permission from [152].

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Figure 42. Representative microstructures of Ti64-0.3B alloy deformed in (a) stable Domain-II (T = 900-1000 °C and 𝜀 = 10-3-10-2 s-1) and (b) unstable Domain-III (T = 750-850 °C and 𝜀 = 10+0-10+1 s-1). Images reproduced with permission from [152].

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Figure 43. (a) Variation of tensile elongation with temperature in Ti64-0.1B sheet. Tensile tests performed at a constant 𝜀 = 3 x 10-4 s-1. Dashed line is the minimum elongation required for a material to exhibit superplastic behavior. (b) True stress vs. true strain curves of Ti64-0.1B [161] and conventional Ti64 sheets [157] in tension at test temperature of T = 900°C and 𝜀 = 3 x 10-4 s-1. (c) Back scattered SEM image of superplastically deformed Ti64-0.1B sheet at 900 °C and 𝜀 = 3 x 10-4 s-1. Cavity formation shown with circles. Images reproduced with permission from [161].

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Figure 44. Macrostructures of the laser welded as-cast Ti64 plates with (a) 0, (b) 0.06 and (c) 0.11 wt.% B content [162]. BM, HAZ and FZ in laser welded Ti640.11B alloy plate are shown in (c).

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Figure 45. (a) IPF map of α phase in BM, HAZ and FZ of laser welded Ti64-0.06B alloy plate, corresponding β reconstruction map of (a) is shown in (b). (c) β reconstruction map of laser welded B-free Ti64 alloy plate [162]. The orientation code for hcp-α and bcc-β phase is given in (d). For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.

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Figure 46. Microstructures in the FZ of Ti-15-3 welds. Columnar dendritic morphology is observed in (a) autogenous weld (no filler) and (b) weld prepared using CP-Ti filler whereas equiaxed dendritic morphology in welds prepared using CP-Ti fillers with (c) 0.5 and (d) 1 wt.% B. Images reproduced with permission from [73].

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Figure 47. (a) Hardness profiles across the weld cross-section and (b) Engineering stress-strain curves of Ti-15-3 autogenous weld and welds prepared using CP-Ti fillers with 0, 0.5 and 1 wt.% B contents. Summary of the tensile properties is shown in the inset of (b). Hardness tests performed at a load of 0.5 Kg. Images reproduced with permission from [73].

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Figure 48. (a) Wear rate as a function of B content in as-cast Ti64 alloy at different temperatures. SEM images showing worn morphologies of Ti64-0.3B alloy at (b) 20 and (c) 300 °C. Images reproduced with permission from [164].

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Figure 49. Mass gain per unit area as a function of oxidation temperature in Ti64 alloy sheet with 0 and 1 wt.% B [169].

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Figure 50. Oxide layer formation in (a) Ti64 and (b) Ti64-1B alloys after exposure at 950 °C for 50 h. (c) Formation of decomposed structure beneath oxide layersubstrate interface in Ti64-1B alloy after exposure at 950 °C for 50 h [169].

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Figure 51. BF TEM micrograph showing δ-TiH2 hydride needles in Ti64-0.09B alloy hydrogen charged at 700 °C for t = 2 h. Corresponding SAED pattern is shown in the inset of the figure. Image reproduced with permission from [183].

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Figure 52. Variation of Δσu with VTiB in Ti64-xB alloys hydrogen charged at 700 °C for t = 2 h. Image reproduced with permission from [183].

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Figure 53. Variation of (a) strength (σy & σu) and (b) strain-to-failure with B content in as-cast Ti64 alloy. Images reproduced with permission from [41].

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Figure 54. IPF maps of (a) Ti64 and (b) Ti64-0.09B alloys obtained after tensile testing at 20 K. Point-to-point misorientation analysis performed along solid lines A and broken line C indicates activation of {1012} and {5613} twins respectively. Line profile of the point-to-point misorientation along C is shown in the inset of (a). Images reproduced with permission from [41]. For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.

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Figure 55. Variation of notch strength ratio (σn/σu) with B content at 77 and 20 K. Image reproduced with permission from [41].

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