Journal of Alloys and Compounds 626 (2015) 76–86

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Journal of Alloys and Compounds

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Precipitate phase stability and compositional dependence on alloyingadditions in c–c0–d–g Ni-base superalloys

http://dx.doi.org/10.1016/j.jallcom.2014.11.1550925-8388/� 2014 Elsevier B.V. All rights reserved.

⇑ Corresponding author.E-mail address: santonov@iit.edu (S. Antonov).

Stoichko Antonov a,⇑, Martin Detrois a, Randolph C. Helmink b, Sammy Tin a

a Illinois Institute of Technology, 10 W. 32nd Street, Chicago, IL 60616, USAb Rolls Royce Corporation, PO Box 420, Indianapolis, IN 46241, USA

a r t i c l e i n f o a b s t r a c t

Article history:Received 11 September 2014Received in revised form 22 November 2014Accepted 25 November 2014Available online 3 December 2014

Keywords:SuperalloyPrecipitateStabilityThermocalc

Ternary eutectic c–c0–d Ni-base superalloys have been the subject of a number of recent investigationsdue to the relatively high volume fraction of d phase providing composite strengthening. However, anadditional intermetallic precipitate phase, g, has been observed in Ni-base superalloys containing addi-tions of both Nb and Ti. As both precipitate phases d and g appear similar and are often mistaken for oneanother, a better fundamental understanding of their formation is required for developing innovativenew classes of Ni-base superalloys. The morphology, formation, and composition of precipitate phasesin a number of experimental alloys spanning a broad range of compositions were explored in this inves-tigation. Additionally, compositional relationships were developed to facilitate the design of c–c0–(d/g)Ni-base superalloys. When the alloy chemistry was observed to exhibit a compositional ratio of Al/(Nb + Ta + Ti) less than 0.85, d and/or g phase precipitates formed while a ratio greater than 0.85 resultedin conventional c–c0 microstructures. For alloys in which d and/or g phase precipitates were formed, theprevalent phase could be determined by evaluating the compositional ratio for (Nb + Ta)/(Al + Ti). Alloysthat had ratios greater than 1 were largely comprised of d phase precipitates, whereas a ratio less than 1resulted in the predominance of the g phase precipitates.

� 2014 Elsevier B.V. All rights reserved.

1. Introduction

Nickel-base superalloys are the materials of choice for the ‘‘hotsection’’ of modern gas turbine engines and power generationapplications and are likely going to remain so due to the uniquecombination of high strength and fatigue resistance of the materialwhen exposed to corrosive environments, high stress levels andelevated temperatures [1–9]. The superior high temperaturemechanical strength of Ni-based superalloys can largely be attrib-uted to the characteristic two phase microstructure, consisting ofordered L12 intermetallic precipitates c0 (Ni3Al) distributed coher-ently within a disordered FCC matrix c (Ni), that provides substan-tial order and coherency type strengthening by not permittingdislocation motion during plastic deformation. In modern poly-crystalline Ni-base superalloys, refractory alloying elements suchas W, Mo, Cr and Co provide strength to the c matrix phase throughsolid solution strengthening. Additionally concentrations of Al, Ti,Nb and Ta are utilized to produce an optimized volume fractionof the c0 strengthening phase. However, modest improvements inthe properties of the alloy can be attributed to the increased

volume fraction of the c0 precipitates and come at the expense ofprocessing capabilities, environmental resistance and overall alloycost. As a result, new and innovative alloying concepts and newalloy classes need to be considered in order to meet future designrequirements as the property limits of conventional c–c0 phasemicrostructure alloys are approached. One novel class of alloysunder consideration is based on the Ni–Al–Nb pseudo-ternaryeutectic c–c0–d alloy, which utilizes a large volume fraction of pri-mary intermetallic d (Ni3Nb) phase of D0a crystal structure dis-persed intragranularly to provide composite strengthening andthe c0 phase for precipitation strengthening.

First developed as a candidate material for gas turbine airfoilsduring the 1970s and 1980s, directionally solidified c–c0/d alloyswere shown to possess excellent high temperature strength andcreep resistance compared to the best performing c–c0 Ni-basesuperalloys available at the time. Despite possessing excellentmechanical properties at elevated temperatures, the applicationof these directionally solidified c/c0–d eutectic alloys for airfoilcomponents has been limited due to some of their intrinsic short-comings arising from the anisotropic nature of their crystal struc-tures, such as the relatively low d interfacial strength withsurrounding c–c0 matrix, limited transverse ductility, severelyreduced fatigue resistance and notch sensitivity, and the mismatch

S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86 77

of thermal expansion coefficient between the c matrix and the dprecipitates [10–12]. These limitations associated with direction-ally solidified structures were avoided by processing these alloysinto polycrystalline structures containing a uniform dispersion ofblocky d precipitates which made the alloy more amenable foruse at temperatures below 800 �C [13]. The resulting composite-like microstructure was found to avoid the directional characterof the c–c0–d Ni-base superalloys even though high fractions of dprecipitates often resulted in the accumulation of damage at thed/c–c0 interfaces [14]. In recent years, there has been a significantactivity on developing polycrystalline Ni-base superalloys contain-ing elevated levels of Nb [15–19]. The elevated levels of Nb addi-tions in these alloys have been shown to induce the formation ofthe c–c0–d microstructure that results in improved strength andtemperature capability, but it was also shown that significantreorientation of the d phase can occur during hot deformation ofthe material that leads to flow softening of the material [20].

At relatively low alloying levels in Ni-base superalloys, Nb, Taand Ti will preferentially partition to the c0 phase and serve tolower the anti-phase boundary energy and strengthen the precip-itates resistance to deformation. As the solubility limits areapproached, elevated levels of Ti additions will induce the forma-tion of a g phase precipitate [21,22]. This phase consists of a hex-agonal close packed (TCP) D024 crystal structure withstoichiometry of Ni3Ti, and although at first it was thought thatthe phase has no solubility for other elements, more recent inves-tigations suggest that g has solubility for Co, Al, and Cr and conse-quently it has also been reported to be Ni3Al0.5Nb0.5 [21,23,24].This phase exhibits a lamellar and/or rod like morphology whichis non-coherent with the matrix [25], and excessive precipitationwithin the microstructure has been considered to pose a negativeeffect on the mechanical properties of superalloys. However, itwas observed that the formation of g-Ni6AlNb occurred througha discontinuous precipitation reaction, which resulted in the for-mation of serrated grain boundaries, which can be used to enhancethe mechanical properties of the alloys [26,27]. Uneven tempera-ture and strain conditions through the cross-section of the discduring hot working has been shown to promote different aspectratios of g phase, which has been reported to improve dwell crackpropagation in some directions [28]. Additionally, g has beenreported to form at high temperatures but may undergo a solidstate transformation to c0 at lower temperatures. Traditionally,the ratio of Al/Ti in Ni-base superalloys was carefully balanced tomaintain phase stability and avoid precipitation of g [29]. Exces-sive alloying levels of Ta and Nb in Ni-base superalloys, on theother hand, both tend to induce the formation of d phase precipi-tates at elevated temperatures. At intermediate temperatures, ametastable c00 phase may form in alloys containing elevated levelsof Nb.

Currently, thermodynamic database models such as Thermocalcor Pandat by Computherm are extremely effective in predicting themulticomponent thermodynamic stability of complex phases incommercial and experimental alloys with compositions that fallwithin the range of alloys over which substantial experimentaldata exists. The predictive capability of these models becomeschallenged when experimental alloy compositions are extendedbeyond this range. Since limited experimental thermodynamicinformation on the c–c0–d–g alloy system exists and the precisionof the thermodynamic database models is lacking, an improvedfundamental understanding of how the different elemental alloyadditions affect the microstructure and the properties of this novelclass of Ni-base superalloys is needed in order to design and suc-cessfully implement Ni-base superalloys with improved proper-ties. The present report aims to elucidate compositional effectson the formation of d and g phase precipitates, to construct away to differentiate between d and g precipitates, and develop

relationships for the prediction of the long term stabilities of thetwo precipitate phases.

2. Experimental details

For Ni-base superalloys containing moderate to high levels of Nb (>6 at.%), thephase stability of both d and g precipitates has been found to vary considerablywith small changes in the alloy composition. Therefore, a set of experimental alloysspanning a range of Ti and Ta levels was explored. A total of 9 experimental alloys,with compositions listed in Table 1, were designed in order to span Ti and Ta levelsranging from 0 to 3 at.%. While the atomic percent of Cr, Mo and W were kept con-stant and the ratio of Nb over Al equal to 1, the atomic percent of Ti and Ta followedthe relation Ti[at.%] + Ta[at.%] = 1, 2, or 3 at.%. All of the alloys were prepared fromelemental Ni (99.95 pct.), Cr (99.5 pct.), Al (99.9 pct.), Ti (99.9 pct.), Nb (99.98 pct.),Ta (99.9 pct.), Mo (99.8 pct.) and W (99.9 pct.) melted in a MRF ABJ-900 (MRF Inc.,Suncook, NH, USA) arc melter under an argon atmosphere. Prior to the melting ofeach of the alloys, the arc melter chamber was purged under vacuum to a pressureof 10�3 mbar and backfilled with high purity argon (99.998 pct.). The process wasrepeated three times prior to melting to ensure cleanliness of the arc melt ingots.Additionally, each ingot was melted and remelted up to six times to convectivelymix and uniformly distribute the alloying additions throughout the ingot. Theingots were then homogenized at 1200 �C for 48 h and aged at 1100 �C for 38 h.

Metallographic samples were prepared for observation under a JEOL JSM-5900LV (JEOL Ltd., Tokyo, Japan) scanning electron microscope (SEM) in the as-cast,as-homogenized, and as-aged conditions. As a primary means of identification ofthe d and g phases, micrographs were taken using the back-scattered mode onthe SEM with an accelerating voltage of 20 kV and a spot size of 47. Confirmationof the identification of the phases as well as assessment of the composition of theprecipitate phases was performed using an Oxford Instrument X-Max energy dis-persive spectrometer (EDS) (Oxford Instruments, Oxfordshire, UK) connected tothe SEM. During characterization on all of the samples, the accelerating voltagewas kept constant to 20 kV and the working distance, spot size and acquisition timewere set to 10 mm, 46 and 60 s respectively for consistency. A minimum of threeprecipitates were considered for each alloy for precision of the reported valuesand the deviation was listed in the results. In order to better understand the mor-phology of each phase, the samples were electrochemically etched using a solutionof 90% methanol and 10% hydrochloric acid, for 45 min at 1.5 A and 4 V. Imageswere subsequently taken using the SEM under SEI mode with a spot size of 42, aworking distance of 10 mm and a voltage of 20 kV.

Finally, X-ray diffraction analysis was performed using Cu Ka radiation in aThermo Scientific ARL X’TRA powder diffractometer (Thermo Scientific, Pittsburgh,PA, USA). The step size, preset time, voltage and current were set to 0.02�, 1 s, 45 kVand 40 mA, respectively for the higher accuracy of peak position and peak-to noiseratio. The software package, Materials Analysis using Diffraction (MAUD) was usedto processed the experimentally measured diffraction patterns and calculate andrefine lattice parameters and error. CrystalMaker and CrystalDiffract were usedfor the visualization of the crystal structure and simulation of the diffraction pat-tern, respectively.

3. Results

The micrographs in Fig. 1 show the homogenized and agedmicrostructures of the 1Ti, 2Ti and 3Ti alloys. Following homogeni-zation, the microstructure of the 1Ti alloy consisted of c–c0 phasesonly while a modest amount of g phase precipitates were observedin the 2Ti and 3Ti alloys, with an increasing fraction of precipitatesas the Ti content increased. In addition to the g phase, a limitedfraction of blocky d phase precipitates were identified along thegrain boundaries of the 2Ti sample only. Ageing of the experimen-tal alloys at 1100 �C resulted in a significant increase in volumefraction of the g phase precipitates for all of the Ti series of alloyswith an increasing fraction of the g phase precipitates correspond-ing to increases in the overall Ti content. Following ageing, only afew small d precipitates were observed along the grain boundariesin the 1Ti and 2Ti alloys, Fig. 1. Preliminary identification of theprecipitate phases was based on the EDS results presented inTable 2 and the contrast difference in back-scattered SEM imagesthat emerged from the atomic number difference between thetwo precipitate phases making the d phase brighter than the gphase. Semi-quantitative EDS measurement of the precipitatechemistries was performed in order to observe the preferentialpartitioning of the alloying additions and differences in composi-tion between the d and g phase. A minimum of 4 precipitates were

Table 1Composition of the experimental alloys (at.%).

Alloy Nominal Actual

Ni Cr Al Ti Nb Ta Mo W Ni Cr Al Ti Nb Ta Mo W

1Ti Bal. 10.0 7.0 1.0 7.0 0.0 1.0 0.5 Bal. 10.2 7.3 1.0 7.7 0.0 1.2 0.42Ti Bal. 10.0 6.5 2.0 6.5 0.0 1.0 0.5 Bal. 10.1 6.6 2.1 7.1 0.0 1.1 0.53Ti Bal. 10.0 6.0 3.0 6.0 0.0 1.0 0.5 Bal. 10.1 6.3 3.1 6.7 0.0 1.2 0.51Ta Bal. 10.0 7.0 0.0 7.0 1.0 1.0 0.5 Bal. 10.1 6.7 0.0 7.4 1.2 1.1 0.62Ta Bal. 10.0 6.5 0.0 6.5 2.0 1.0 0.5 Bal. 10.2 6.5 0.0 7.2 2.3 1.2 0.83Ta Bal. 10.0 6.0 0.0 6.0 3.0 1.0 0.5 Bal. 10.2 6.1 0.0 6.4 3.4 1.2 0.91Ta–1Ti Bal. 10.0 6.5 1.0 6.5 1.0 1.0 0.5 Bal. 10.1 6.4 1.0 7.1 1.2 1.2 0.62Ta–1Ti Bal. 10.0 6.0 1.0 6.0 2.0 1.0 0.5 Bal. 10.0 5.9 1.0 6.6 2.3 1.1 0.72Ti–1Ta Bal. 10.0 6.0 2.0 6.0 1.0 1.0 0.5 Bal. 10.1 6.1 2.0 6.6 1.2 1.2 0.6

Fig. 1. SEM micrographs of the Ti only alloys after both homogenization and ageing.

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sampled in each alloy and the average compositions are reportedin Table 2. From the EDS results, the characteristic compositionof the g phase was measured to be �8 at.% Al, �7 at.% Cr, 8 at.%Nb and a Ti concentration that is proportional to the overall Ti con-centration in the alloy. Although trace amounts of blocky d precip-itates were observed along the grain boundaries in alloy 1Ti, theirsizes were extremely small and limited the ability to accuratelydetermine their composition. In the 2Ti alloy, the grain boundaryd phase was measured to have a composition that had a signifi-cantly higher concentration of Nb (�16.4 at.%) and lower concen-tration of Al (�1.3 at.%) when compared to the g phaseprecipitates. Interestingly, the Ni content of both phases was foundto remain nominally constant at approximately 75 at.%. In terms of

Table 2Composition of the d and g precipitate phases for the set of experimental alloys (at.%).

Alloy Precipitate phase Al Ti

1Ti d – –g 7.2 ± 0.2 1.0 ± 0.0

2Ti d 1.3 ± 0.1 2.2 ± 0.1g 8.7 ± 0.1 3.6 ± 0.3

3Ti g 8.2 ± 0.2 4.8 ± 0.21Ta d 1.5 ± 0.1 –2Ta d 1.3 ± 0.1 –3Ta d 1.2 ± 0.1 –1Ta–1Ti d 1.2 ± 0.2 1.3 ± 0.12Ta–1Ti d 1.3 ± 0.2 1.3 ± 0.12Ti–1Ta d 1.2 ± 0.1 2.5 ± 0.0

g 8.3 ± 0.1 3.4 ± 0.1

morphological differences, the g phase precipitates appeared aslong needles, resulting from the plate-like morphology, distributedwithin the c–c0 matrix.

The Ta series of alloys 1Ta, 2Ta and 3Ta were shown to containmicrostructures comprised of blocky as well as needle-like d phaseprecipitates following homogenization, Fig. 2. Despite exhibiting aneedle-like morphology that was similar to g phase precipitates,both the blocky and needle-like precipitates in the Ta series ofalloys were confirmed to be d phase and no g phase precipitateswere observed in these alloys, Table 2. Similar to the compositionof the d phase precipitates in the 2Ti alloy, the average compositionof the d phase precipitates in the Ta series of alloys was found toconsist of <1.5 at.% Al, �16 at.% Nb, 4 at.% Cr and a Ta content that

Cr Ni Nb Ta

– – – –10.7 ± 0.3 74.6 ± 0.2 6.5 ± 0.3 –5.1 ± 0.4 74.9 ± 0.2 16.4 ± 0.3 –3.9 ± 0.4 74.3 ± 0.1 9.4 ± 0.1 –4.8 ± 0.5 73.8 ± 0.1 8.5 ± 0.2 –3.9 ± 0.3 74.3 ± 0.3 17.7 ± 0.3 2.6 ± 0.03.3 ± 0.2 74.0 ± 0.5 16.2 ± 0.2 5.2 ± 0.13.1 ± 0.2 74.0 ± 0.4 14.3 ± 0.2 7.4 ± 0.13.3 ± 0.2 73.8 ± 0.3 17.5 ± 0.1 2.9 ± 0.13.1 ± 0.1 74.1 ± 0.3 14.9 ± 0.2 5.3 ± 0.13.5 ± 0.4 74.2 ± 0.2 15.9 ± 0.3 2.8 ± 0.04.7 ± 0.5 73.7 ± 0.1 7.9 ± 0.2 1.9 ± 0.0

S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86 79

was proportional to the overall Ta content in the alloys. As the Tacontent in the d phase increases, the Nb content decreases. Basedon these measurements, it appears that Al and Ti exhibit significantsolubility within the g phase, while the d phase has limited solu-bility for Al but a high solubility for Ta. In the as-homogenized con-dition, blocky d precipitates were found to be the predominantmorphology with precipitates located primarily along the grainboundaries while the needle-like d precipitates were largely con-tained within the grains. Increasing the Ta content in the alloysresulted in an increase in the fraction of the blocky d precipitates.Following ageing of the alloys, the volume fraction of the thin nee-dle or plate-like d precipitates increased considerably while thefraction of blocky d phase precipitates appeared to remainconstant.

The alloys containing both Ti and Ta additions, 1Ta–1Ti, 2Ta–1Tiand 2Ti–1Ta, displayed similar microstructures when compared tothe Ti and Ta series of alloys as shown in Fig. 3. Only d phase pre-cipitates were observed in both the homogenized and aged sam-ples of 1Ta–1Ti and 2Ta–1Ti. Again blocky d phase precipitateswere located primarily along grain boundaries while the thin, nee-dle or plate-like d phase precipitates were observed to form withinthe grains. A significantly higher density of transgranular d phaseprecipitates were observed in the samples after ageing at1100 �C. Interestingly, the microstructure of the 2Ti–1Ta samplecontained both d and g phase precipitates upon both homogeniza-tion and ageing. In this sample, the d phase precipitates were lar-gely confined along the grain boundaries while the thin, needleor plate-like precipitates were almost exclusively g phase. Compo-sitions of the d and g phase precipitates measured in these alloyswere consistent to those measured in the Ti and Ta series of alloys,Table 2.

In order to provide an accurate description of the various pre-cipitate morphologies present in the experimental alloys, deepetching techniques to remove the surrounding c–c0 matrix andreveal the 3-D morphologies of the d and the g phase precipitateswas performed. Fig. 4 shows the micrographs of the samples afterageing at 1100 �C in which the matrix surrounding the precipitateswas preferentially etched away. Characteristic differences betweenthe morphologies of each precipitate phase were then noted.Although the 3Ta alloy contained only d phase precipitates two dis-tinct morphologies were observed, Fig. 4a and Table 2. The major-ity of the d precipitates appeared as an interconnected network of

Fig. 2. SEM micrographs of the Ta only alloy

thin plates with large aspect ratios while the remainder was in theform of blocky precipitates with much smaller aspect ratios. Inter-estingly, some of the plates were observed to nucleate and extenddirectly off the ends of the blocky precipitates. The g phase precip-itates observed in alloy 3Ti, Fig. 4b, were found in the form of verythin plates only, sometimes stacked, with needle-like rods on theextremities. It should be noted that alloy 3Ti did not contain anyd precipitates, Table 2, and all of the elongated structures wereidentified to be g phase precipitates. Finally, Fig. 4c shows the typ-ical structure in alloy 2Ti–1Ta where both the d and g phase pre-cipitates are present, Table 2. At first glance the morphologies ofthe two phases are almost indistinguishable due to the similaritiesbetween the d and g platelet morphologies. However a detailedexamination reveals that the g platelets are thinner and moreirregular than the d plates. In addition, only plates or sheets of dwere observed to nucleate from the blocky d phases, while the gplates avoid them and usually appear as colonies comprised ofmultiple sheets and plates.

Finally, XRD analysis was performed on each of the experimen-tal alloys to compliment the EDS results and confirm the character-istic presence of the precipitate phases d and g. Simulated XRDpatterns, shown in Fig. 5 and obtained using CrystalDiffract, wereused to identify the peaks of the patterns obtained from the exper-iments. The diffraction patterns of the Ti series of alloys are plottedin Fig. 6. All of the alloys displayed peaks around 43�, 50�, 90� and117� corresponding to the [111], [200], [220], and [400] planes ofthe cubic system of the c–c0 matrix, respectively. For the Ti seriesof alloys, alloy 1Ti produced two g peaks at 42.6� and 43.7�, corre-sponding to the [021] and [004] planes. As the Ti content in thealloy increases, additional g peaks appeared and were located at46.4� and 90.5� corresponding to the [022] and [224] planesrespectively. These peaks were also observed on the diffractionpattern of the 2Ti and 3Ti alloys. Finally, the diffraction patternfor the 3Ti sample included the 52.9� and 81.8� peaks correspond-ing to the [023] and [026] planes, respectively. Interestingly, the43.4� and 74.3� peaks of the g phase was found to overlap withthe [111] and [220] planes of the cubic c–c0 matrix, respectively,which can be explained by the nature of the atomic stacking,where the c0 closed packed planes follow a ABCA stackingsequence, while the g phase follows a ABAC stacking sequence[21]. The diffraction patterns of the Ta-only alloys, 1Ta, 2Ta, and3Ta, are shown in Fig. 7 and contain peaks at 40.6�, 43.6�, 45.6�,

s after both homogenization and ageing.

Fig. 3. SEM micrographs of the Ta and Ti alloys after both homogenization and ageing.

80 S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86

46.2� and 90.6�, which correspond to the [201], [020], [012],[211] and [420] planes of the orthorhombic crystal structure,respectively. The intensities of these peaks were found to increasewith the volume fraction of d phase precipitates observed in Fig. 7.Consistent with the EDS results, the diffraction pattern for alloys1Ta–1Ti and 2Ta–1Ti exhibited the characteristic peaks that con-firmed that these alloys consist of predominantly d phase whilealloy 2Ti–1Ta, which was comprised of mixture of the two phasesin approximately equal proportions, had a corresponding XRD pat-tern consisting of peaks for both d and g precipitate phases, Fig. 8.For these analyses, the most important peaks that aided in theidentification of the two phases d and g were the d double peaklocated at 45� and 46� and the 42� and 52� g peaks.

The lattice parameters of the d and g phases for the experimen-tal alloys were calculated from the XRD patterns and the resultswere compiled in Table 3. The c–c0 phase lattice parameters werefound to be nominally constant at �0.359 nm for all of the alloys.The d phase lattice parameters ranged from 0.509 to 0.512 nm,0.413 to 0.415 nm and 0.452 to 0.0454 nm for a, b, and c respec-tively. Finally, the g lattice parameters measured in the alloys con-taining in excess of 2 at.% Ti were also nominally constant witha = �0.511 nm and c = �0.830 nm. On average, the error in thereported lattice parameters in Table 3 was determined to be�±0.0005 nm. These values are in agreement with reported valuesof �0.5093 nm and �0.8276 nm for a and c respectively [30].

4. Discussion

Polycrystalline c–c0–d Ni-base superalloys, with or without theadditional g phase, have been the focus of numerous investiga-tions due to the potential increase in the alloy properties from highlevels of precipitate strengthening. However, the two precipitatephases d and g are not well defined and are often mistakenly iden-tified and confused for one another due to their macroscopic sim-ilarities. Owing to the limited knowledge and information on thesetwo phases, thermodynamic database models, such as Thermocalc,often lack the ability to accurately predict the formation of the gphase precipitate. This investigation was aimed at elucidating themicrostructural and compositional differences between the twophases and establishing relationships for a better prediction ofthe stability of the precipitate phases.

From the 9 experimental alloys investigated with microstruc-tures represented Figs. 1–3, both d and g phase precipitates were

observed. Combining this microstructural information along withthat obtained from the deep-etched micrographs, Fig. 4, revealedthe characteristic morphologies of the d and g phase precipitates.The d phase precipitates may exhibit either a blocky or plate-likemorphology, while g phase precipitates appear as colonies ofthin-sheets or rods. Since the blocky d phase precipitates wereobserved within the microstructure in the as-cast condition, theseprecipitates are likely primary phases that form within the meltduring the latter stages of solidification. The high temperaturehomogenization heat treatment at 1200 �C was unable to solutionthese blocky d phase precipitates and they were present as discreteparticles that tended to reside along grain boundaries. The sheet orplate-like d and g phase precipitates form as part of a solid statetransformation at either the homogenization or ageing tempera-ture. Both of these precipitates possess very large aspect ratiosand form lamellar structures that extend across the grains. Owingto their identical crystal structures and chemistries, lamellar dphase precipitates were observed to nucleate from the blocky dphase precipitates. Compared to the lamellar d phase precipitates,the g phase precipitates were also found to exhibit a plate-likemorphology, but tended to be thinner and consist of multipleplates in close proximity to one another forming a colony-likestructure. Interestingly, the extremities of most g plates appearedto be lightly feathered or serrated while the edges of the d precip-itates were comparatively more linear. This can likely be attributedto the manner in which the g phase precipitates form as they ini-tially precipitate out in the form of needles which then coalesce toform sheets. From the results in Table 2, the apparent differencebetween the two phases is that Ti strongly partitions to the gphase, while Ta preferentially partitions to the d phase. Based onthese observations, it can be concluded that Ta additions preferen-tially stabilize the d phase and increase the concentration of Ta inthe d phase at the expense of Nb. Meanwhile, increases in the over-all level of Ti alloying additions prompted the formation of gphase, with no significant presence of d phase precipitates innon-Ta containing alloys. Within the range of compositions inves-tigated, it should be noted that the potency or effectiveness of Taadditions in forming d phase was stronger than the ability of Tiaddition to form g phase precipitates.

Using the phase compositions obtained with EDS, Table 2, thechemical formulas for the d and g phase precipitates can beexpressed as Ni3(Nb,Ta) and Ni3(Al,Ti)0.5Nb0.5, respectively. As theTa content increases in the d phase, the concentration of Nb must

Fig. 4. SEM images showing the morphologies of (a) the d phase in alloy 3Ta, (b) theg phase in alloy 3Ti, and (c) both the d and g phase precipitates in alloy 2Ti–1Ta.

Fig. 5. Generated XRD pattern for the Ni3Ti g phase and the Ni3Nb d phase usingCrystalDiffract.

Fig. 6. XRD pattern for the Ti series of experimental alloys.

Fig. 7. XRD pattern for the Ta series of alloys.

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decrease, in order for the sum of Nb + Ta to remain at �25 at.%, asthis trend is evident from the phase compositions of alloys 1Ta, 2Taand 3Ta in Table 2. On the other hand, as the Ti content in the alloyincreases along with the Ti content in the g phase, the Nb contentin the g phase must also increase in order for it to be proportionalto the Al + Ti content. This trend was observed in Table 2 for alloys1Ti, 2Ti and 3Ti. The crystal structures, already established in theliterature, were defined as being D0a for the d phase and D024 forthe g phase precipitates [31]. Using this information and Crystal-Maker Software, the crystal structures were modeled and shownin Fig. 9. The two variations of the d phase, Ni3Ta and Ni3Nb, areclassified as the D0a, crystal structure in Fig. 9a and the two d form-ing elements, Nb and Ta, occupy similar sites with proportions sim-ilar to the concentration of the element in the alloy [32,33]. Fig. 9b,

shows the g phase crystal structure, as defined by the Ni3Ti D024,however in this variation Al and Nb are also present in the crystallattice with the concentration of Nb being proportional to that ofAl + Ti. In addition, the niobium atoms have been reported tooccupy the 2a sites with Al and Ti atoms randomly filling the 2csites in Wyckoff notation [26].

Fig. 8. XRD pattern for the Ti + Ta series of alloys.

Table 3Lattice parameters for the set of experimental alloys.

Alloy c phase d phase g phase

a (nm) a (nm) b (nm) c (nm) a (nm) c (nm)

1Ti 0.359 – – – 0.511 0.8312Ti 0.359 – – – 0.511 0.8303Ti 0.359 – – – 0.511 0.8301Ta 0.359 0.510 0.413 0.454 – –2Ta 0.360 0.510 0.414 0.453 – –3Ta 0.360 0.511 0.414 0.453 – –1Ta–1Ti 0.359 0.511 0.413 0.452 – –2Ta–1Ti 0.359 0.512 0.415 0.452 – –2Ti–1Ta 0.359 0.509 0.415 0.452 0.510 0.830

Fig. 9. Crystal structures of the (a) d and (b) g phase precipitates.

82 S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86

Thermodynamic database models such as Thermocalc, allow forthe prediction of equilibrium phases that will form in the alloy sys-tem and the respective composition of each of the constituentphases. Fig. 10 shows the predicted phase fraction plots as a func-tion of temperature for the experimental alloys obtained usingThermocalc with the TCNi5 database. What can be immediatelynoticed is that the g phase, or Ni3Ti_D24 in the TCNI5 databaseof Thermocalc, was not predicted in any of the alloys while the dphase, Ni3Ta_D0a, was present in all of the fraction plots. In addi-tion, the c0 phase, FCC_L12#2, was predicted to form in all of thealloys, with especially high volume fractions in alloys 1Ti, 2Ti,and 3Ti. These predictions, however, contradict the experimentalobservations from this study where c0 tends to become suppressedby the formation of d and/or g and has a relatively low volumefraction. The predictions for 1Ti, 2Ti and 3Ti also indicate that dphase would form in these alloys albeit at temperatures signifi-cantly lower than the 1100 �C heat temperature that was actuallyused for these experimental alloys. From the microstructures inFigs. 1–3, the d phase was present in alloys 1Ti and 2Ti in verylow quantities while no evidence of d phase was observed in alloy3Ti. The phase fraction plot for alloy 2Ti–1Ta only predicted theformation of the d phase although the actual microstructure con-tained both d and g phase, Fig. 3.

The compositions of the d phase as predicted by Thermocalc at1100 �C for the experimental alloys are listed in Table 4. Interest-ingly, although some of the phase fraction plots showed the dphase forming below 1100 �C, Thermocalc was able to predict atheoretical composition at that temperature, which was fairly con-sistent throughout the given temperature range for the experimen-tal alloys. Comparison of the predicted compositions in Table 4with the measured compositions in Table 2, reveals that the

predictions are in fairly good agreement with the experimentaldata. Al, Ni and Ti showed little to no significant difference, Cr,however, was predicted to have limited solubility in d phase, whileit was observed to average about 3–4 at.% in the experimentalalloys. The predicted Ta concentration in the d phase was foundto be moderately lower than the experimental measured composi-tions and the difference increased as the overall Ta content in thealloy increased. Overall, the degree of Ta partitioning to the d phasewas higher than predicted, while the degree of Nb partitioning wasslightly lower than predicted.

In order to be able to better predict which phases will form andbe stable within the range of compositions investigated in this par-ticular range of alloy compositions, compositional relationshipswere drawn based on the characterization of the microstructurespresented in this article. Table 5 compiles some well-establishedand commonly used commercial c–c0 and c–c0–(d/g) alloys, inaddition to the experimental alloys, to support the relationships.At relatively low alloying levels, Nb, Ti and Ta, are all consideredto be c0 forming elements in Ni-base superalloys, since these alloy-ing additions are soluble in the c0 phase by substituting for Al inthe Ni3Al lattice. However, excessive additions of Nb, Ti and Ta inNi-based superalloys can promote the formation of other phasesbesides c and c0. To assess whether an additional phase, d and/org, will form in the alloy system, the ratio of Al/(Nb + Ta + Ti) wherethe concentrations are in at.%, should be considered. When thisratio is less than �0.85, the microstructure will tend to develop atendency to form d and/or g phase precipitates. This was observedfor all of the alloys investigated and also applies for commercialalloys such as Inconel 718 and Allvac 718Plus. Moreover, this ratiowas validated against and shown to be consistent with the micro-structural stability of another set of experimental alloys (RCH53,RCH49. RCH48 and LN8) in which the elevated levels of Nb + Ta + Tiresulted in the formation of both d and g phases. Contrary, inalloys that exhibit a ratio of Al/(Nb + Ta + Ti) that is greater than0.85, no d or g precipitates appear as the c–c0 microstructure ismore stable. Comparing the compositions of many of the commer-cial c–c0 alloys in Table 5, such as RR1000 (containing Ta but noNb), René95 (containing Nb but no Ta), René104 (containing bothNb and Ta) and Inconel 100 (that does not containing either Nbor Ta) and others, supports this compositional relationship.

In order to predict whether d, g or both phase precipitates willbe stable for alloys in which the Al/(Nb + Ta + Ti) is less than 0.85, asecond compositional relationship needs to be utilized. Based onthe overall alloy composition, the ratio of the d formers, in atomicpercent, (Nb and Ta) to the concentration of the g formers, again inatomic percent, (Al and Ti) was considered. When the ratio of

Fig. 10. Mass fraction plots of constituent phases for the experimental alloys as predicted by Thermocalc.

Table 4Composition of the d precipitate phase at 1100 �C as predicted by Thermocalc (at.%).

Alloy Al Ti Cr Ni Nb Ta

1Ti 1.18 1.62 0.98 74.25 20.76 –2Ti 0.92 2.87 0.98 74.23 19.71 –3Ti 0.75 4.00 1.01 74.19 18.70 –1Ta 1.58 – 0.90 74.33 19.79 2.332Ta 1.55 – 0.89 74.34 18.05 4.163Ta 1.33 – 0.87 74.36 16.69 5.721Ta–1Ti 1.09 1.47 0.93 74.28 18.70 2.442Ta–1Ti 1.07 1.51 0.90 74.29 16.99 4.192Ti–1Ta 0.82 2.52 0.95 74.23 17.74 2.61

S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86 83

(Nb + Ta)/(Al + Ti) is greater than 1, precipitation of the d phaseprecipitates was preferred over the g phase whereas a ratio lessthan 1 results in a microstructure predominantly comprised of gphase precipitates, as noted in Table 5. When the ratio is approxi-mately �1 the microstructure of the corresponding alloys exhib-ited an approximately equal volume fraction of g and d phases,evident from alloy 2Ti–1Ta in Fig. 3. Applying this relationship tosome commercial alloys with an established presence of d or g,the composition of Inconel 718 corresponds to a ratio of 1.55 whichis consistent with its microstructure containing d phase precipi-tates and no g phase [34]. Allvac 718Plus, on the other hand, hasa ratio of 0.86 and its microstructure contains g phase precipitatesand no d phase. In addition, for the set of experimental RCH53,

RCH49, RCH48 and LN8 alloys that contain both d and g phase pre-cipitates, the magnitude of the (Nb + Ta)/(Al + Ti) ratio was com-pared to respective volume fractions of the d and g phases [25].Remarkably, the fractions of precipitate phases were found to cor-relate closely to the magnitude of the ratio. For example, a ratio of1.50 was calculated for RCH53 which primarily contained d phaseprecipitates (37.5% d compared to 15.7% g). As this ratio decreasedgradually in the other RCH series alloys, the fraction of d alsodecreased while the fraction of g phase increased. For RCH48which had a (Nb + Ta)/(Al + Ti) ratio of �0.95, the fraction of dand g phase precipitates were nominally similar at 20.4% and21.6%, respectively. Finally, for LN8 that has a calculated ratio of0.82, only the g phase was observed to form.

Fig. 11 shows how the microstructure of various commercialand experimental Ni-base alloys varies as a function of the Al/(Nb + Ta + Ti) ratio. Based on the results in this figure, the alloyscan be categorized into two groups separated by a line passingthough the origin with a slope of 0.85. Alloys that reside abovethe Al/(Nb + Ta + Ti) = 0.85 ratio line exhibit a c–c0 microstructure,while alloys grouped below that line have a tendency to form d org phase precipitates. The characteristic grouping behavior canpotentially be explained by the statistical availability of Al atomsto promote the formation of the c0 phase. One potential explana-tion for this behavior is that when the overall concentration of Alatoms is significantly higher than the overall concentration ofNb + Ta + Ti, there is a high likelihood of Ni3Al c0 precipitate

Table 5Relative Al, Ti, Ta and Nb ratios.

Alloy AlNbþ Taþ Ti

Nbþ TaAlþ Ti

Predominant precipitatephase

Volume fraction of precipitate phase (%) Reference for composition and/orphase presence

d g

Experimental alloys1Ti 0.84 0.93 g – – n/a2Ti 0.72 0.82 g – – n/a3Ti 0.64 0.71 g – – n/a1Ta 0.77 1.29 d – – n/a2Ta 0.69 1.45 d – – n/a3Ta 0.62 1.62 d – – n/a1Ta–1Ti 0.69 1.11 d – – n/a2Ta–1Ti 0.59 1.29 d – – n/a2Ti–1Ta 0.61 0.97 g – – n/a

Commercial c–c0 alloysAlloy 10 1.36 n/a n/a n/a n/a [8,35]Astroloy 2.03 n/a n/a n/a n/a [2,36]Incoloy 800 1.77 n/a n/a n/a n/a [37]Incoloy 802 1.37 n/a n/a n/a n/a [37]Inconel 100 2.08 n/a n/a n/a n/a [36]Inconel 617 5.91 n/a n/a n/a n/a [8]Inconel 738 1.40 n/a n/a n/a n/a [8,38,39]Nimonic 90 1.06 n/a n/a n/a n/a [37]Nimonic 105 6.95 n/a n/a n/a n/a [37]Nimonic 115 2.35 n/a n/a n/a n/a [8]Nimonic PE16 1.77 n/a n/a n/a n/a [8,40]Nimonic PK33 1.77 n/a n/a n/a n/a [9,40]René 88DT 0.92 n/a n/a n/a n/a [2]René95 1.44 n/a n/a n/a n/a [8,36]René104 1.26 n/a n/a n/a n/a [8,41]RR1000 1.29 n/a n/a n/a n/a [8,42]Udimet 500 1.77 n/a n/a n/a n/a [2]Udimet 520 1.18 n/a n/a n/a n/a [2,40]Udimet 700 2.03 n/a n/a n/a n/a [2,40]Udimet 710 0.89 n/a n/a n/a n/a [2,40]Udimet 720 0.89 n/a n/a n/a n/a [2,43]Udimet 720LI 0.89 n/a n/a n/a n/a [2,43]

Commercial c–c0–(d and/or g) alloysIN718 0.21 1.55 d – n/a [24,26,34,38,44]Allvac718Plus 0.73 0.86 g n/a – [24,26,45]Inconel 625 0.17 3.34 d – n/a [9,38,46]Inconel 706 0.11 0.69 g n/a – [8,39,44,47]Inconel 740 0.56 0.30 g n/a – [8,39,48]Inconel X750 0.41 0.14 g n/a – [37,39]Nimonic 263 0.42 0.00 g n/a – [8,49]Nimonic 901 0.21 0.00 g n/a – [8,36]Waspaloy 0.77 0.00 g n/a – [2,44]

Other c–c0–d–g alloysV204A 0.41 2.46 d – n/a [13]V204B 0.42 2.40 d – n/a [13]V204G 0.45 1.89 d – n/a [13]V204H 0.50 1.88 d – n/a [14]RCH53 0.39 1.50 d 37.5 ± 3.9 15.7 ± 0.5 [25]RCH49 0.42 1.12 d 29.9 ± 1.8 20.5 ± 0.6 [25]RCH48 0.49 0.95 d–g 20.4 ± 1.9 21.6 ± 0.6 [25]LN8 0.70 0.82 g Negligible 34.2 ± 1.2 [25]

84 S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86

formation during which the Nb, Ti and Ta atoms are accommo-dated within the c0 phase. On the other hand, as the number ofNb + Ta + Ti atoms increases to approach or exceed the numberof Al atoms in the alloy, the likelihood of Ni3(Nb,Ta) d or Ni3Ti gphase precipitation increases relative to the overall concentrationof Nb + Ta + Ti. As the g phase exhibits extensive solubility for Al,the stoichiometry of the precipitates that were observed to formwithin this range of alloy compositions tended to be closer toNi3(Al,Ti)0.5Nb0.5 than the typical Ni3Ti g phase. It should be notedthat the overall alloy phase stability is also dependent on therelative concentrations of Nb, Ta and Ti in the alloy system. Somealloys like Rene 88DT, Nimonic 90, Udimet 720 and Waspaloy,which fall close to the 0.85 line, contain no Nb or Ta additions,while their Ti content is relatively high as compared to the Alconcentration. In such cases despite possessing a low ratio, thecombined concentration of the c0 forming elements is below

10 at.% and the degree of Nb, Ta and Ti supersaturation is toolow to induce the formation of stable g and/or d phases.

For Ni-base superalloys in which the Al/(Nb + Ta + Ti) composi-tional relationship predicts the corresponding microstructure toform phases other than c and c0, a separate relationship that con-siders the Al + Ti and Nb + Ta content in atomic percent can be usedto predict the likelihood of d and/or g phase precipitation. Thealloys shown in Fig. 11 that have characteristic compositionswhere Al/(Nb + Ta + Ti) < 0.85 are replotted on the compositionmap in Fig. 12 to highlight how varying the (Nb + Ta)/(Al + Ti) ratioinfluences the formation of d, g or both precipitate phases. Foralloy 2Ta–1Ti, increasing the Ta content increased the magnitudeof the ratio and therefore the alloy exhibited the formation of dphase only. Commercial alloys, such as Inconel 625 and Inconel718, also exhibited only d formation with ratios of 3.34 and 1.47,respectively. On the other hand, when the magnitude of the ratio

Fig. 11. Plot showing the microstructural phase stability of various commercial andexperimental alloys relative to Al/(Nb + Ta + Ti).

Fig. 12. Plot showing the occurrence of d and/or g phases relative to (Nb + Ta)/(Al + Ti).

S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86 85

is decreased, as in the 3Ti alloy, only g phase was observed. Thisappears to be consistent with observations on the stability of Inco-nel 706 and Allvac 718Plus which have been reported to exhibit gphase within their microstructure and have ratios of 0.69 and 0.86,respectively. For alloys where (Nb + Ta) and (Al + Ti) concentra-tions are relatively equal, a mixed microstructure containing bothd and g phase precipitates was exhibited as is the case with alloy2Ti–1Ta, RCH48 and RCH49. Once again it should be noted that therelative concentrations of the Al, Nb, Ta and Ti additions in thealloys should be considered as well, as in the case of RCH53 alloy,which formed both phases even though it falls closer to the d phaseonly region of the plot. This can most likely be attributed to thecomposition of RCH53 containing significant levels of both Nb(�10.2 at.%) and Ti (�2.2 at.%). Although high levels of Nb primar-ily serve to stabilize the d phase, Nb is also a constituent of theNi3(Al,Ti)0.5Nb0.5 g phase that is stabilized by the presence of Ti.

Within the set of alloys, RCH53, RCH49, RCH48 and LN8, Nb levelsvaries from 10.2 to 5.6 at.% while the Ti levels remained between2.1 and 3.2 at.%. The systematic increases in the Nb contentresulted in a gradual increase in the (Nb + Ta)/(Al + Ti) ratio andled to proportional increases in the fraction of the d phase as afunction of the Nb content. The boundary contours in Fig. 11 weredrawn such that increasing the Nb + Ta content while keeping theAl + Ti content constant resulted in a relative increase in theprobability of formation of the d phase. On the other hand,increasing the Al + Ti content with a constant Nb + Ta contentpromoted the formation of the g phase. From the alloys investi-gated, a Al + Ti content over �7.3 at.% resulted in the precipitationof the g phase. It should be noted that the composition map inFig. 11 was constructed primarily based on experimental resultsfrom high Nb (>5 at.%) and moderate Ta (<1.5 at.%) contentalloys. Significant changes in the relative proportions of Al, Ti, Taand Nb would be expected to alter the behavior and additionalalloying studies should be completed in order to complete thecomposition map.

5. Conclusions

Both the d and g phases were observed in multi-component Ni-base superalloy systems and studied based on their morphologies,compositions and fractions. Based on the results and observationspresented in this study, the following conclusions can be reported:

1. In addition to forming c and c0, the microstructures of theseexperimental Ni-base superalloys were composed of blockyand plate-like d phase precipitates and/or g phase precipitateswith a morphology consisting thin plates. The volume fractionof each precipitate phase was found to depend on their respec-tive compositions.

2. Ti was found to stabilize the g phase, meanwhile Ta stabilizedthe d phase. On an equiatomic basis, Ta was observed to be astronger d former than Ti was g former.

3. The composition of the g phase can be expressed asNi3(Al,Ti)0.5Nb0.5 which has a hexagonal D024 crystal structure,and the Nb content is proportional to the Al + Ti content. Thed phase possesses a chemical composition defined by Ni3(Nb,Ta)that has an orthorhombic D0a crystal structure with the propor-tions of each element as defined by the alloy compositionalconstituents.

4. XRD data calculations showed that the lattice parameters of d,g and the c–c0 matrix remain nominally constant throughoutthe range of alloy compositions investigated with average val-ues of a � 0.512 nm, b � 0.410 nm and c � 0.457 nm for the dphase, a � 0.510 nm and c � 0.832 nm for the g phase anda � 0.359 nm for the c matrix.

5. Although ThermoCalc predictions using the TCNi5 databasesuggested that d phase would form in some of the experimentalalloys, it was unable to correctly predict the formation and sta-bility of the g phase. The predicted compositions of d phasewere close to the measured compositions with the most signif-icant difference being that the predicted compositions con-tained higher Nb content at the expense of Ta content.

6. A compositional ratio of Al/(Nb + Ta + Ti) in atomic percent canbe used to predict whether the d and g phase precipitates willbe stable in the alloy system. A ratio less than 0.85 resulted inthe preferential stability of the d and g phases, while a ratiogreater than 0.85 produced conventional c–c0 microstructures.

7. For d and/or g containing alloys, a compositional ratio of(Nb + Ta)/(Al + Ti) in atomic percent can be used to predictwhich phase is likely to form. A ratio less than 1 resulted inthe predominance of the g phase, while a ratio greater than 1

86 S. Antonov et al. / Journal of Alloys and Compounds 626 (2015) 76–86

resulted in the preferential stability of the d phase. A ratio �1has the tendency to form a mixed microstructure of g and dof equal volume fractions.

Acknowledgements

Financial support for this work was provided by Rolls-RoyceNorth American Technologies, Rolls-Royce Corporation, Rolls-Royce plc., and NSF-DMR-1006953.

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