DOI: 10.1002/adem.201400336

Grain Growth and b to a Transformation Behavior of ab-Solidifying TiAl Alloy**

By Thomas Klein,* Salar Niknafs, Rian Dippenaar, Helmut Clemens and Svea Mayer

Advanced intermetallic g-TiAl-based alloys such as multi-phase TNM alloys exhibiting a compositional range of Ti–(42–44)Al–(3–5)Nb–(0.1–2)Mo–(0.1–1)B (all compositions are giv-en in at%, unless otherwise indicated) have attracted attentionin the automotive and aircraft industries due to their potentialas light-weightmaterials for high temperature-applications.[1–5] In the acronym TNM, T stands for TiAl, N and M stand forthe alloying elements Nb and Mo, respectively. A majordriving force for the development of advanced structuralmaterials in propulsion systems is the increasing globalecological awareness. In particular, a significant reduction ofthe pollutants CO2 and NOx is required, because of theirsevere impact on the quality of the atmosphere and theenvironment. Due to the low density of titanium aluminidesin comparison to Ni-based superalloys, a replacement of theseheavy alloys for certain applications in aero and automotiveengines is expected to reduce their mass, in particular that ofrotating components, leading to a higher efficiency andreduced fuel consumption.[1–6]

The present study aimed to analyze the microstructuraldevelopment of a TNM alloy showing a nominal compositionof Ti–43.5Al–4Nb–1Mo–0.1B. At room temperature (RT), thisalloy consists of mainly g phase (L10 structure), a2 phase (D019structure) and a small amount of bo phase (B2 structure). Dueto the intermetallic character of all the constituent phases andthe related inherent brittleness and low ductility, enhancedmicrostructural control is required to gain a fine-grained,isotropic, and homogeneous microstructure. The most crucialcontributing factors that influence the final microstructure are

different solid state phase transformations as well as graingrowth processes occurring during annealing and coolingcycles.[1,7,8] Starting from solidification, the evolution of thephases in the TNM alloy under consideration can besummarized as follows, neglecting for the moment, theformation and presence of borides: L!Lþb !b!b þa!

aþbþ g!aþbþbo þ g!a þ boþ g ! aþa2þboþ g!

a2þboþ g.[9] The TNM alloy solidifies via the disordered b

phase showing a body-centered cubic structure. As a commonfeature, b-solidifying alloys show a fine-grained microstruc-ture, whereas peritectic solidifying g-TiAl-based alloys (L!Lþb!Lþbþa!bþa…) are prone to coarse colony sizesexhibiting a strong solidification texture and segregation asdescribed in ref.[5,10] After solidification and upon furthercooling, TNM alloys pass through a relatively large single b

phase region of about 100 °C.[9] Therefore, in this microstruc-tural regime fast coarsening occurs due to the high state ofthermal activation of atomic movement, the presence ofdisordered b phase resulting in a high diffusivity and theabsence of a second phase. A better understanding of themicrostructural changes during the coarsening process willpermit the design of cooling sequences so as to enable uncriticalpassing of this single phase regionwithout rapid grain growth.On cooling to the b transus temperature (Tb), the disorderedhexagonal a phase nucleates heterogeneously at grain bound-aries, triple points, and titanium borides. This b to a phasetransformation is known to critically influence the finalmicrostructure as the a grain size determines the size as wellas the shape of a2/g-colonies.[11,12] On further cooling the g

phase is formed and ordering of the b and a phases occurs.[13]

The main aim of this study was to visualize the kinetics ofgrain growth in the singlebphasefield regionand to contributeto the fundamental understanding of the b to a phase trans-formation in the TNM alloy system. For this purpose, a high-temperature laser-scanning confocal microscope (HTLSCM)was employed, which allows the visualization of the micro-structural evolution at high temperature in situ and in real time.Since the microstructural development is recorded on videoin real time, it was possible to determine continuous coolingtransformation (CCT) diagrams from the captured sequencesand the corresponding cooling data. Following this approach,these diagramswerederived for theTNMalloy and the carbon-containing variant, the TN1.5Mo0.5C alloy. This alloy has anominal composition of Ti–43Al–4Nb–1.5Mo–0.1B–0.5C andshows a similar transformation pathway, yet shifted to highertemperatures as reported in ref.[14]

*[*] T. Klein, Prof. H. Clemens, Prof. S. MayerDepartment of Physical Metallurgy and Materials Testing,Montanuniversitaet Leoben, 8700 Leoben, AustriaE-mail: thomas.klein@unileoben.ac.atDr. S. Niknafs, Prof. R. DippenaarFaculty of Engineering, University of Wollongong, Wollon-gong, NSW 2522, Australia

[**] A part of the work in this paper was conducted within theframework of the following project: FFG project 832040“energy-drive”, Research Studios Austria, Austria. Experi-mental and financial support by the University of Wollon-gong, Australia, is gratefully acknowledged. T. K. wants tothank the Materials Process Engineering Group of theUniversity of Wollongong for their hospitality during thestay and the fruitful collaboration in the framework of thisjoint project.

DOI: 10.1002/adem.201400336 © 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim wileyonlinelibrary.com 1ADVANCED ENGINEERING MATERIALS 2014,

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1. Experimental

1.1. MaterialThe TNM alloy was produced by GfE Metalle und

Materialien GmbH (Nuremberg, Germany), using the “ad-vanced beta process”, which involves two interlacing series ofvacuum arc remelting (VAR).[15] Subsequently, the materialwas melted using a VAR skull melter and cast by a centrifugalcasting process.[1] The TN1.5Mo0.5C alloy was processed insimilar fashion, but only a single VAR step was applied. Inorder to close residual porosity, both alloys were subsequentlyhot isostatically pressed (HIPed) for 4 h at 1200 °C at apressure of 200MPa.

1.2. High-Temperature Laser-Scanning Confocal MicroscopyFor the HTLSCM investigations, thin bars of 8mm in

diameter were eroded from HIPed material. These barswere cut into disks with a thickness of 1.4mm, using anAccutom 3500 diamond cutter followed by grinding on 800,1200, 2400, and 4000 abrasive papers in sequence beforebeing polished by standard metallographic techniques,using 6 and 1mm diamond paste. These disks were thenplaced into yttria (Y2O3) crucibles, positioned in a platinumsample holder and placed in the infra-red furnace of theHTLSCM. The sample holder is equipped with a B-typethermocouple. The HTLSCM technique has been fullydescribed in ref.[16–18] In order to avoid reaction betweenthe crucible and the titanium aluminide alloys, high purityyttria was chosen as crucible material because of itsthermal and chemical stability. All experiments werecarried out in a type 1LM21H laser-scanning confocalmicroscope, supplied by Lasertec Corporation (Japan).Temperature calibration was carried out using the nucle-ation temperature of d ferrite in a low-carbon steel showinga well known transformation temperature.[19] Because thespecimen within the infra-red furnace is heated byradiation, which is concentrated in the center, there is atemperature gradient across the specimen. The temperaturedistribution is constant only within a radius of approxi-mately 1mm from the center. For this reason, the specimencenter was marked with microindents. All images weretaken within this region of constant temperature becausethe temperature calibration is only valid for this region ofthe specimen. High purity argon was used as inertatmosphere, which was additionally cleaned by a purifica-tion train containing a stainless steel tube filled withzirconium filings, which was placed within a furnace andheld at 900 °C. Before each experiment, the HTLSCMchamber was purged 12 times in succession, so that tracelevels of oxygen were sufficiently removed.[17]

Following the in situ observation of the microstructuraldevelopment at high temperature, the same specimens wereexamined at RT in a field-emission scanning electronmicroscope (FESEM) JSM-7001F from Jeol Ltd. (Japan). Allimages were taken in back scattered electron (BSE) mode at avoltage of 15 kV and a current of 3.8mA.

2. Results and Discussion

2.1. Grain Growth in the Single b-Phase Field RegionSpecimens were heated to a temperature of 1405 °C at a

heating rate of 1.67 °C s� 1 and annealed within the infra-redfurnace of the HTLSCM for selected times as shown inFigure 1 marked with (a). At this temperature exclusively b

phase is present, in agreement with findings of Schwaighoferet al.[9] The grain boundaries are clearly delineated as shownin the inset of Figure 2. The roughness observed in this figureis due to prior events, which occurred during heat-up (this issimply a characteristic of confocal microscopy; e.g. a grainboundary groove that was formed at an earlier stage willremain visible as a groove even though grain growth hadoccurred and the newgrain boundaries are located in differentareas).[20] A square of about 2mm side length, with ahomogeneous temperature distribution, was mapped in pre-selected time intervals in order to gain statistically relevantdata of grain size. Subsequently, grain size determination wascarried out by using the mean linear intercept method. Theevolution of the grain size with respect to time is shown inFigure 2. It is evident that the grain size increases with time,initially showing a slow rate. After annealing for 1200 s at1405 °C the grain size has approximately doubled from itsinitial value. At this stage the in situ determination of grainsize was terminated as there were too few grains visible in thefield of view to ensure a statistically valid analysis. Evidently,after a period of initiation, where the grain size stays almoststable, grain growth occurs much faster. However, due to theabove-mentioned experimental limitations the further courseof the grain growth curve cannot be evaluated. An image ofthe microstructure after annealing for 480 s at 1405 °C isshown in the inset in Figure 2. Grain boundaries and somesurface roughness are visible. Grain boundaries are observedbecause the free surface is thermally etched. However, grainboundary grooves that had been left behind following further

Fig. 1. Sketch of the temperature profiles applied in this study: Grain growth wasobserved in the single b phase field region at T1 at a pre-selected time t1 (marked witha). Following annealing for t1 the specimens were cooled at different cooling ratesbetween 0.03 °C s� 1 and 12 °C s� 1 in order to study the b to a phase transformation(marked with b). Tb is the b transus temperature.

T. Klein et al./Morphological Evolution of a b-solidifying TiAl Alloy

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grain growth are often observed and have to be distinguishedfrom actual grain boundaries. Grain boundaries werefrequently observed to move with a staggered motion ratherthan moving continuously in agreement with findings ofPhelan,[18] who explained this phenomenon by reference toearlier analysis by Mullins.[21] On many occasions grainboundaries were temporarily pinned by particles, whichpresumably are borides.[1]

2.2. b to a Phase TransformationThe microstructural development of disk shaped speci-

mens of the Ti–43.5Al–4Nb–1Mo–0.1B alloywas studied in thecourse of heat treatments applied within the infra-red furnaceof the HTLSCM (see (b) in Figure 1). As oxidation of titaniumalloys at high temperatures is always problematical, prelimi-nary experiments were conducted to assess the likelihood ofoxide formation. However, neither during these preliminaryexperiments nor during the actual experiments any signs ofoxide formation were detected. The solution time t1 requiredfor ensuring the formation of an equilibrium and stable b

structure, was determined in preliminary experiments. In thecase of the transformation study, the TNM alloy was annealedfor 300 s at 1405 °C. The shortest possible solution heat-treatment was selected because it is known that large b grainspromote the formation of Widmannst€atten-type structures,which we intended to avoid in the initial analysis.[22] Thespecimens were cooled at different rates ranging from0.03 °C s� 1 to 12 °C s� 1 allowing the b to a transformationto take place. In contrast to grain growth experiments, whichare sensitive to specimen thickness and influenced bythe surface, phase transformations, visualized within theHTLSCM on the surface, are generally not affected by the freesurface and are fully representative of bulk behavior.[18]

The transformation of b to a phase in a TNM alloy oncooling at a rate of 0.08 °C s� 1 is shown in the image sequencedisplayed in Figure 3. Figure 3a shows the structure at 1373 °C,

2 °C below the temperature at which the first a-nucleiformed. Figure 3b shows the microstructure at 1368 °C andFigure 3c the microstructure at 1363 °C. The b phase can beobserved in Figure 3a; grain boundaries are visible and itappears as if the grain boundaries were pinned by obstacles,which are visible in the center of Figure 3a. These particlesare presumably borides, which are thermodynamicallystable at this temperature.[1,2,9] Most of the evolving a grainsgrow in an equiaxed way and do not show distinct growthorientations. The roughness and vague thermal grooves aredue to the occurrence of microstructural changes anddifferent thermal expansion of the microstructural constit-uents prior to cooling. It appears that the a phase nucleatespreferentially on grain boundaries, triple points, and pre-existing particles, presumably borides. Because the newlyformed a grains nucleate heterogeneously at numerousdifferent sites, significant grain refinement is achieved. Forthe sake of comparison, an image of the evolving a phasewhile cooling at a rate of 12 °C s� 1 is displayed in Figure 4,displaying morphological features different to cooling at

Fig. 2. Evolution of grain size with time, determined by means of the mean linearintercept method at 1405 °C. Zero time corresponds to the time when a temperature of1405 °C was reached. The inset shows a region in the vicinity of the sample center.Although there is significant surface roughening, grain boundaries are clearly visible.

Fig. 3. HTLSCM image sequence of the formation of a phase as obtained at a coolingrate of 0.08 °C s� 1. Image (a) was taken at 1373 °C, image (b) at 1368 °C, and image(c) at 1363 °C. In (a) grain boundaries in the b microstructure are visible. Nucleationof the first a grains takes place in the lower right corner (arrows). In (b) and (c)progress of the transformation process is visible (see text).

T. Klein et al./Morphological Evolution of a b-solidifying TiAl Alloy

DOI: 10.1002/adem.201400336 © 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim http://www.aem-journal.com 3ADVANCED ENGINEERING MATERIALS 2014,

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lower rates. In this instance, a Widmannst€atten-typestructure is formed and the laths grow with a specificangular relationship along preferred crystallographic ori-entations. It is well known that in case of the transformationof the body-centered cubic b phase into the hexagonal-close-packed a phase, the crystallographic orientation relationshipis governed by the so-called Burgers orientation relation-ship.[23] The strong dependence of the preferred a nucleationsites as well as the occurrence of an orientation relationshipwith the b parental phase on the cooling rate wascomprehensively investigated by Oehring et al.[12] usingdifferent methods. In the present experiments, no indicationof a massive-type transformation was observed, not even atthe highest cooling rate imposed. However, an interestingphenomenon was observed. At a cooling rate of 12 °C s� 1, asignificant number of a grains nucleated beneath the surface.This observation might be attributed to an effect of specimenthickness: because of the position of the sample within theinfra-red furnace, a temperature gradient along the thicknessof the specimen is generated, where the bottom is colder thanthe top surface that is being observed. Hence, it is expectedthat nucleation might preferentially occur at positions belowthe surface of observation, especially at high cooling rates.

Following in situ observations within the confocal micro-scope, specimens were cooled at a rate of 1.67 °C s� 1 andmetallographically prepared for analyses in a FESEM. Thepresence and morphology of previous a grains, now trans-formed to a2/g-colonies, confirmed the morphologicalevolution of the a phase observed in the HTLSCM. TheFESEM observations also confirmed that the observations inthe high-temperature microscope are representative of bulk-behavior and ruled out any significant influence of the freesurface, which is in agreement with the findings of Phelan.[18]

2.3. Continuous Cooling Transformation DiagramsIn the practical use of alloys in component design,

the kinetics of the pertaining transformations are ofcritical importance to control the microstructural evolutionand, hence, the mechanical properties. The kinetics of thepertaining phase transformations as presented in CCTdiagrams are therefore of great value to engineers. For these

reasons, CCT diagrams were constructed for the initiation ofthe b to a phase transformation in the TNM alloy and theTN1.5Mo0.5C alloy by utilizing the information obtainedfrom the image sequences and the concomitant temperaturesduring the in situ observations in the high-temperaturemicroscope. The CCT diagram for the TNM alloy is shown inFigure 5a and the CCT diagram for the TN1.5Mo0.5C alloy isdepicted in Figure 5b. The cooling curves correspond to theexperimentally measured cooling curves and are labeled inFigure 5. Black crosses at each cooling rate refer to theinitiation temperature of a phase formation as visuallyobserved in the HTLSCM. It appears that the nucleationtemperature of the b to a phase transformation in theTN1.5Mo0.5C alloy is approximately 50 °C higher than that ofthe TNM alloy. The transformation temperatures determinedat the lowest cooling rate of 0.03 °C s� 1 are �1380 °C and�1430 °C for the TNM and the TN1.5Mo0.5C alloy, respec-tively. In the case of the TNM alloy, this temperature agreeswell with the one determined by Schwaighofer et al.[14] using adifferent method. This suggests that calibration and methodof transformation temperature determination are reliable and,hence, also the determined transformation temperature forthe TN1.5Mo0.5C alloy is trustworthy. The necessary

Fig. 4. HTLSCM image of a phase formation at a cooling rate of 12 °C s� 1. Thedepicted image was taken at 1325 °C. Lath-like structures are observed that show adistinct path of progression as well as a crystallographic relationship between thea and the b phase.

Fig. 5. CCTdiagramsof the initiationofaphase formation for (a) theTNMalloyand (b)theTN1.5Mo0.5C alloy.Note the different scaling of the temperature axes. Black crossesrepresent the onset ofa phase formation. The transformation temperature is about 50 °Chigher in the TN1.5Mo0.5C alloy than in the TNM alloy. In (a) Tb determined bySchwaighofer et al.[14] is indicated. Parameters of the solution heat treatment can befound in the right upper corner of each diagram; further the phase regions are indicated.Cooling rates are given at the lower end of each corresponding line.

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undercooling to induce nucleation of the a phase increaseswith increasing cooling rate. With respect to the lowestcooling rate applied, the highest undercooling needed for theTNMalloy and the TN1.5Mo0.5C alloy are�45 °C and�30 °C,respectively.

3. Conclusion

The present investigation was launched because of thecritical role played by the evolution of b grains and theirsubsequent transformation to a phase in the design ofprocessing techniques and, hence, microstructural controlof intermetallic titanium aluminides. By utilizing the uniquecharacteristics of high-temperature laser-scanning confocalmicroscopy, the evolution of b grains in a TNM alloy wasdetermined at 1405 °C, in situ and in real time. Following thesolution heat treatment of the b phase at this temperature, themorphology and rate of formation of theaphaseupon coolingwere determined. The evolution of b grainswas quantified bymapping an area of homogeneous temperature within thehigh-temperature microscope. Grain boundaries were oftenmoving by staggered motion, probably due to pinning bythermal grooves and/or particles.

On cooling specimens from the single b phase region totemperatures below the b transus temperature, the formationandgrowthof theaphasecouldbeclearlyobserved in real time.Nucleation preferentially occurred at grain boundaries, triplepoints and on pre-existing particles, presumably borides.Significant grain refinement was achieved because many a

grains nucleated on the boundaries of a singleb grain and grewinto thebparentalgrain.Theexactmorphologicaldevelopmentwas a strong function of the cooling rate. Low cooling ratesrendered equiaxed structures with no preferred growthdirection, whereas high cooling rates led to the formation oflath-like structures displaying a crystallographic relationshipbetween the two phases. The kinetics of the b to a phasetransformation was assessed by determining the initiation ofthis phase transformation as a functionof time and temperatureand plotting this data in the form of CCT diagrams. The b to a

phase transformation in theTNMalloy initiates at�1380 °Cat acooling rate of 0.03 °C s� 1 anddecreases to�1335 °Cat a coolingrate of 12 °C s1. TheTN1.5Mo0.5Calloy transformed in a similarfashion, except that the actual transformation temperature ishigher by about 50 °C

Received: July 15, 2014Final Version: September 5, 2014

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