versión On-line ISSN 2411-9717
J. S. Afr. Inst. Min. Metall. vol.112 supl.1 Johannesburg ene. 2012
As cast and heat-treated alloys of the Pt-Al-V system at the Pt-rich corner
B.O. OderaI, II, III; L.A. CornishI, II, III; M.B. ShongweI, III; G.O. RadingII, IV; M.J. PapoIII, V
ISchool of Chemical and Metallurgical Engineering, University of the Witwatersrand
IIAfrican Materials Science and Engineering Network (AMSEN)
IIIDST/NRF Centre of Excellence in Strong Materials, hosted by University of the Witwatersrand
IVDepartment of Mechanical and Manufacturing Engineering, University of Nairobi
VAdvanced Materials Division, Mintek
The Pt-based alloys have a high potential for replacing some Ni-based superalloys (NBSAs) used in the highest temperature and most aggressive environments, and vanadium would be a beneficial alloying element.
Six alloys of average compositions Pt-26.6Al-9.1V, Pt-23.1Al-17.8V, Pt-9.6Al-21.1V, Pt-17.4Al-6.4V, Pt-22.3Al-7.9V, and Pt-6.0Al-10.1V (all in at.%) were manufactured by arc-melting and examined in the as-cast condition using scanning electron microscopy and X-ray diffraction (XRD). Five samples were heat treated at 1 000°C for 1 500 hours and water quenched, then examined. The phases were identified using energy-dispersive X-ray spectroscopy and the identities were confirmed by XRD. In most of the samples, the phases existing in the as-cast condition were different from those at 1 000°C after ~1500 hours heat treatment. A solidification projection, an isothermal section and a liquidus surface projection were determined for the Pt-rich corner.
Keywords: Phase diagram, Pt-Al-V system, scanning electron microscopy, X-ray diffraction.
Nickel-based superalloys (NBSAs) have been successfully used in turbine components since the late 1940s. Approximately 70 percent of the weight of a modern jet turbine comprises NBSAs, and their success is due to their high yield stress and excellent resistance to environmental attack at high temperatures1. Although the NBSAs have excellent properties, they are hindered by their maximum application temperature, which is limited by the melting point of the nickel solid-solution matrix. Currently, the maximum temperature at which NBSAs operate is ~1 100°C, which is approximately 85 percent of their melting temperature1. If the operating temperatures could be increased, there would be a number of advantages. Higher temperatures improve the efficiency of the turbine engines, and this enables greater thrust, improved fuel efficiency, and reduced pollution. Although thermal-barrier coatings can be used to increase the application temperature, the component is still restricted by the melting point of the substrate, so the maximum attainable temperatures are still limited if catastropic failure is to be avoided2.
There is increasing interest in using a different alloy system with a much higher melting point. Ceramics and selected intermetallic phases have been discussed because of their high melting points and high strength; however, both suffer from reduced toughness compared with NBSAs, and their processing costs are high2 One solution would be to base the new materials on alloys with high melting points and use the naturally occurring precipitates of that system. This should also help keep the processing cost to a minimum. Thus, the systems developed would be similar in structure to the NBSAs, with a matrix and a fine dispersion of small, preferably coherent and stable precipitates. Face-centered cubic (fcc) structures (like nickel) are advantageous because, being close-packed, they are more creep resistant. Refractory metals (niobium, molybdenum, and tungsten) have been considered because of their high melting points (2 469°C, 2 623°C, and 3 020°C respectively), but their more open body-centred cubic (bcc) structures are more susceptible to creep. Additionally, they are prone to rapid oxidation, even at relatively low temperatures3. Platinum-group metals (platinum, iridium, rhodium, and palladium) were targeted for investigation because, with the exception of palladium, they have high melting points, good environmental resistance, and a mostly fcc structure4.
Pt is similar to Ni in crystal structure (fcc) and chemistry. The important differences are the higher melting point (1769°C for Pt compared to 1 455°C for Ni) and improved corrosion resistances. These facts inspired Wolff's and Hill's research into the development of Pt-based analogues to Ni-based superalloys that could potentially serve in the most critical and demanding of the high-temperature applications*. Thus, phases similar to L12 Ni3Al could be used to obtain properties similar to those found in the NBSAs. Although platinum-based alloys are unlikely to replace all NBSAs on account of both higher price and density, it is likely that they can be used for the highest application temperature components.
Alloys based on Pt-Al have the highest potential, not only because of possible precipitation strengthening through Pt3Al, but also due to high oxidation and corrosion resistance. The resistance to oxidation and corrosion is attributed to the formation of a protective Al oxide scale6,7. At least 11 at.% Al is needed to form a thin protective oxide coating on the ternary and quartenary alloys at high temperatures». It is generally accepted that the L12 structure of the ã' phase is a prerequisite for the success of the Ni-based superalloys. Due to the close crystallographic relationship, precipitates with ordered fcc, L12 have coherent interfaces with the surrounding fcc matrix. Coherent ϒ/ϒ interfaces with small lattice misfit lead to low interfacial energies, thus prohibiting fast coarsening of ϒ8.
Under the auspices of the Platinum Development Initiative, platinum-based superalloys were investigated for high-temperature and special applications requiring good corrosion and oxidation resistance?. Ternary alloys based on Pt-Al, where the ternary additions comprised chromium, iridium, molybdenum, nickel, rhenium, ruthenium, tantalum, titanium, and tungsten were tested. As well as phase characterization work, mechanical and oxidation tests were undertaken. The best systems were found to be Pt-Al-Cr and Pt-Al-Ru. The microstructures were similar to NBSAs, and comprised -Pt3Al precipitates in a Pt-based matrix. However, the volume fraction of -Pt3A was only -40% instead of the -70% found in the NBSAs1?.
Chromium was found to stabilize the cubic form of the -Pt3Al phase, whereas ruthenium acted as a solid-solution strengthenerio. In order to exploit the additional benefits from the two ternary additions (Cr and Ru), quaternary alloys were made containing both. A number of these quaternary alloys have been investigated extensively. They exhibited a two-phase ã/ã' structures analogous to the NBSAs ii.
Much work has been done on the quarternary alloys of the Pt-Al-Cr-Ru system in an attempt to optimize the composition with respect to the precipitate volume fraction and hardness. The Vickers hardness of the Pt-Al-Cr-Ru alloys within the composition ranges selected was relatively independent of the chemical composition and the number of annealing stages and fell within the range HVi0 -400 to -430. The volume fraction of the precipitate, using image analysis, was estimated at approximately 25 to 30%i2.
There is a need to increase the hardness and consequently the strength, melting temperature, and volume fraction of the precipitates. It is also desirable to reduce density and cost. Although vanadium may not meet all these requirements, it is expected to reduce cost and density by replacing some Pt in a chosen Pt-Al-Cr-Ru alloy. It is also expected to act as a solid solution strengthener by going into solution in (Pt), as well as increase the melting temperature. The binary phase diagram of Pt-V shows that -20.5% V will go into solution in (Pt) at 700°C, increasing to -57% at 1 720°C, and the melting temperature of (Pt) has a maximum of -1 805°Ci3. Addition of Nb, Ta, and Ti to Pt-based alloys resulted in elevated strengths at high temperatures through precipitation-hardening. The polycrystalline Pt-based alloys containing 5 at.% Nb or Ta are stronger than the single-crystal NBSA CMSX-4 in (001) orientation above 1 200°Ci4. V is near Nb, Ta, and Ti in the periodic table and is smaller in size compared to Nb and Ta. Therefore, V may act as a precipitation-hardener in addition to solid-solution strengthening.
Six alloy buttons (Samples 1 to 6) weighing ~2 g each were prepared from Pt and Al of 99.9 percent purity and V of 99.6 percent purity. Each of the six alloys is named with respect to its average overall composition in at.% in the as-cast condition. The samples were manufactured by arc-melting under argon in a water-cooled copper hearth. Each button was turned over and re-melted three times to ensure thorough mixing of the elements.
Each of the samples was halved using a ceramic blade mounted in a Struers Secotom-10 cutting machine. The spindle speed was 3 500 r/min and a feed of 0.5 m/s was used. One half were prepared for metallographic examination, in the as-cast condition, by grinding on silicon carbide down to 1 200 grit, diamond polishing down to 1 ìçé and then final polishing by an oxide polishing system (OP-S).
Five of the other halves (samples 1H, 2H, 3H, 5H and 6H, where the number refers to the original as-cast sample and 'H' refers to the 1 000°C heat treatment) were placed in silica glass tubes 16 mm in diameter. The glass tubes were then made into vacuum ampoules by evacuating air and sealing while the vacuum pump was running. The ampoules with the samples were placed in a tube furnace and annealed at 1 000°C for 1 500 hours. Quenching was done by removing the glass ampoules from the furnace, dropping them in water and breaking the glass, as quickly as possible in order to retain the phases existing at 1 000°C. The samples were then prepared for metallographic examination in the same way as the as-cast samples.
The microstructures of all the alloys were analysed using a scanning electron microscope (SEM) model HR-NovaNano SEM200. Imaging in all the analyses was done using the backscattered electron (BSE) mode, while overall and phase compositions were determined by energy dispersive X-ray spectroscopy (EDX). Often it was difficult to discern between some of the phases, even in BSE mode because of the similarity in average atomic number of these phases. There was also the concern that orientation and different electron channelling would also change the contrast^, therefore the morphology of the phases was also used to differentiate them. Since the interaction volume generating the X-rays can be as much as 3 ìçé across and deep (especially for higher atomic number elements), the analyses for smaller regions than this are an indication only. Usually, an accuracy of ±1 at.% would be expected for this technique. Spot analyses were used on single phases and areal analyses on two-phase regions. The compositions of each phase and overall compositions were determined by taking EDX readings from at least five different areas and reporting the average and the standard deviation. X-ray diffraction (XRD) was used to confirm the phase identities.
As-cast Alloy 1, average overall composition Pt-26.6Al-9.1V (at.%)
Figure 1 shows the microstructure of the as-cast Alloy 1. It consists of medium grey dendrites surrounded by a dark thin layer, light needles, and a eutectic comprising the light phase and a barely resolveable two-phase dark region. The compositions of these phases are shown in Table I. The medium grey dendrites were identified as ~Pt3A and the light needles as ~Pt5Al3. The thin dark layer on the dendrites was identified as ~Pt3V. XRD confirmed the identities of these phases. Solidification started with the ~Pt3Al dendrites, followed by peritectic formation of ~Pt5Al3 which then solidified independently as needles. The ~Pt3Al dendrites and the ~Pt5Al3 needles reacted peritectoidally to form ~Pt2Al. The next reaction was eutectic, forming â + ~Pt5Al3 and the â decomposed eutectoidally to form ~PtAl + ~Pt5Al. The final reaction was peritectoid formation of a thin layer of ~Pt3V between the ~Pt3Al dendrites and the ~Pt5Al3 needles.
As-cast Alloy 2, average overall composition Pt-23.1Al-17.8V (at.%)
Figure 2 shows the microstructure of the as-cast Alloy 2 at different magnifications. It consisted of medium grey dendrites surrounded by a light phase, also appearing dendritic, and a complex eutectic. When the compositions of these phases were plotted (Table II), the medium grey dendrites were identified as ~Pt2V, the light phase as ~Pt5Al3, while the dark component of the eutectic was identified as PtAl, although the areas are likely to be affected by the nearby ~Pt2V and ~Pt5Al3. XRD confirmed the identities of ~Pt2V, ~Pt5Al3, and PtAl. Solidification started with (Pt). This was followed by a peritectic reaction between the (Pt) and the liquid, resulting in ~Pt5Al. There then followed a eutectic reaction with the liquid forming ~Pt5Al3 and â, with â decomposing at high temperature to form PtAl + ~Pt5Al3i3. The (Pt) dendrites ordered to ~Pt2V (identified by unique peaks of those phases amongst the common peaks). There was also a solid-state cellular precipitation of ~Pt2V / (Pt) in the ~Pt5Al3 phase, and these phases were too fine to analyse using SEM-EDX.
As-cast Alloy 3, average overall composition Pt-9.6Al-21.1V (at.%)
Figure 3 shows the microstructure of as-cast Alloy 3, average overall composition Pt-9.6Al-21.1V (at.%), at two different magnifications. It consists of dendrites, a light phase, and a dark phase between the dendrites. When the compositions of the phases in Table III were plotted, the dendrites were identified as ~Pt3V, and the light phase as ~Pt3Al. The dark phase was identified as ~Pt2Al, which formed peritectoidally at high temperature from ~Pt5Al3 and ~Pt3Al. Here, any ~Pt5Al3 that formed was consumed. Solidification started with (Pt), followed by a weak peritectic reaction between the liquid and (Pt) resulting in a sparse eutectic with the dark component ~Pt2Al in the light ~Pt3Al. The final reaction was (Pt) ordering to ~Pt3V. XRD confirmed the phase identifications.
As-cast Alloy 4, average overall composition Pt-17.4Al-6.4V (at.%)
Figure 4 shows that the as-cast Alloy 4, average overall composition Pt-17.4Al-6.4V (at.%), consisted of a single phase having grains at different orientations. The gain was increased to enhance the contrast. The compositions in Table IV are for different grains i.e. orientations, and hence different contrast. The plot of these compositions confirmed that the alloy was single phase ~Pt3Al, and XRD also confirmed the phase identity.
As-cast Alloy 5, average overall composition Pt-22.3Al-7.9V (at.%)
Figure 5 is a SEM-BSE image of the microstructure of as-cast Alloy 5, showing light- and medium-contrast dendrites with a thin dark contrast phase. Since the average atomic numbers of the phases were similar, the contrast between them was low, giving a poor image. Both the light and medium grey phases were identified as ~Pt3Al, from the close compositions and the same dendritic morphology, while the thin dark grey phase was identified as ~Pt2Al from the composition. Solidification started with ~Pt3Al, followed by a peritectic formation of a thin layer of ~Pt5Al3 which subsequently reacted peritectoidally with ~Pt3Al, to form ~Pt2Al.
As-cast Alloy 6, average overall composition Pt-6.0Al-10.1V (at.%)
Figure 6 shows the microstructure of as-cast Alloy 6, average overall composition Pt-6.0Al-10.1V (at.%), with a medium-contrast phase which was identified as (Pt) and a fine eutectic. The existence of the eutectic was confirmed after it coarsened during heat treatment for 1 500 hours at 1 000°C (described later). The eutectic and its components were too fine to be analysed accurately by EDX. Solidification started with (Pt), followed by the eutectic reaction producing (Pt) and ~Pt3Al. XRD confirmed the presence of (Pt) and ~Pt3Al (Figure 7).
Annealed Alloy 1H, average overall composition Pt-26.2Al-9.1V (at.%)
Figure 8 depicts the microstructure of annealed Alloy 1H, average overall composition Pt-26.2Al-9.1V (at.%), showing a medium-contrast phase, a light phase, a dark phase and a eutectic. In some cases, the medium contrast phase had a thin dark layer around it, indicating that these were the dendrites in the as-cast alloy although their composition had changed.
The original dendrites showed precipitates of a darker phase, presumed to be PtV considering the contrast and the morphology, which varied between coarse (easily visible) and fine (very difficult to resolve). The plot of the compositions identified the light phase as ~Pt5Al3, the dark phase as either ~PtV or ~Pt2V, while the medium grey phase was identified as -Pt3Al. The individual components of the eutectic were too small to be analysed accurately by EDX. XRD confirmed the identity of the light phase as ~Pt5Al3 and the dark phase as ~PtV (Figure 9).
Annealed Alloy 2H, average overall composition Pt-21.1Al-17.1V (at.%)
Figure 10 shows images of annealed alloy 2H, average overall composition Pt-21.1Al-17.1V (at.%), at different magnifications. Figure 10(b) shows a dark phase with light solid-state precipitates and a light phase with dark solid-state precipitates. Plotting identified the dark phase as ~PtV and the light phase as ~Pt5Al3. The precipitates were too small to be analysed accurately by EDX. XRD confirmed the presence of -PtV and ~Pt5Al3. Compared to the as-cast sample, the ~PtV dendrites had reduced in volume with precipitation of ~Pt5Al3, and the light matrix had aligned precipitates of ~PtV.
Annealed Alloy 3H, average overall composition Pt-10.8Al-25.5V (at.%)
Figure 11 shows Alloy 3H, average overall composition Pt-10.8Al-25.5V (at.%), at different magnifications. Figure11(a) shows a medium grey phase, a light phase, and a dark phase. Comparison with the as-cast structure (Figure 3) shows that the amount of ~Pt2V dendrites had been reduced, and a lighter phase, ~Pt3Al, had precipitated within the dendrites (Figure 11 (b)). The phase was too small to analyse accurately, although larger areas of the medium phase were analysed. Plotting identified the medium phase as ~Pt2V, the dark phase as ~PtV, while the light phase was identified as ~Pt3Al. XRD confirmed these identities.
Annealed Alloy 5H, average overall composition Pt-23.2Al-7.1V (at.%)
Alloy 5H, average overall composition Pt-23.2Al-7.1V (at.%), contained a dark and a medium contrast phase of the same morphology and a small amount of a light phase. EDX compositions of the phases showed that the dark and medium contrast phases were the same phase, ~Pt3Al, and the appearance was likely due to different orientations. The light phase was identified as ~Pt2Al. XRD confirmed the presence of both phases (Figure 11).
Annealed Alloy 6H, average overall composition Pt-5.8Al-10.3V (at.%)
Figure 13 shows annealed Alloy 6H, average overall composition, Pt-5.8Al-10.3V (at.%), at different magnifications. Both have poor contrast owing to the close average atomic numbers of the phases. Figure 13(a) shows a light contrast phase and a dark phase. From the plotting, the light phase was identified as (Pt), and the dark phase as ~ Pt2V, which was not observed in the as-cast sample. The coarsened eutectic could just be discerned in Figure 13(a) and was more clearly resolved in Figure 13 (b). The components of the eutectic were too small to be analysed accurately by EDX. Extrapolation indicated that the eutectic is between (Pt) and ~Pt3Al. XRD confirmed the presence of ~Pt2V, ~Pt3Al, and (Pt).
Although some of the phases were very difficult to resolve, mainly because they had very similar average atomic numbers, the morphologies, analyses, and XRD results were consistent, even if the alloys were not fully homogeneous.
The solidification projection, obtained from the as-cast samples, at the Pt-rich corner is given in Figure 14, although changes to the phase field boundaries are expected as more alloys are studied. All the phases identified in the as-cast alloys were extensions of the binary phases. The phase ~Pt3Al extended the most, particularly in the annealed condition (Figure 15).
The (Pt) phase had a greater solubility for Al above ~20 at.% V than between 10-20 at.% V. The annealed results were used to obtain the isothermal section in Figure 15. There were similar shapes for the phase field boundaries in the solidification projection and the isothermal section, particularly for (Pt) between 10-20 at.% V. It is not yet possible to discern the shape of the ~Pt3Al phase field boundary on the Pt-rich side, but new samples are in preparation.
Alloy 1 had four visible phases, ~Pt3Al, ~Pt2Al, ~Pt5Al3, and ~Pt3V, in the as-cast condition, while a fifth phase, PtAl,
The Journal of The Southern African institute of Mining and Metallurgy
could not be resolved. In the annealed condition Alloy 1 contained ~PtV, ~Pt5Al3 and ~Pt3Al. The light phase was deduced to be to be ~Pt5Al3 on the basis of not showing the laths as found by Tshawe et al. 16 in the Pt-Al alloys as well as the XRD pattern. The major change is that in the as-cast condition the alloy contained ~Pt3V, while in the annealed condition it contained ~PtV. The change can be attributed to diffusion as the phase compositions change, or a reaction in which ~Pt3V changed to ~PtV (although the involvement of the other phases cannot be identified yet). The microstructure of the annealed alloy was also coarser than that of the as-cast, which is expected.
Alloy 2 contained three phases in the as-cast condition, namely ~Pt2V, ~Pt5Al3, and PtAl. There was a lamellar structure which consisted of a solid-state cellular precipitation of ~Pt2V / (Pt) in ~Pt2Al, which shows that the ~Pt5Al5 solvus retreats at lower temperatures. The complex eutectic was originally ~Pt5Al3 + â, and the â decomposed eutectoidally13 to give almost unresolveable ~Pt5Al3 + ~PtAl. The annealed Alloy 2H contained two phases, ~PtV, ~Pt5Al3, and the remnants of a eutectic. The precipitation of ~PtV within the ~Pt2Al matrix (Figure 9 (b)) shows that the ~Pt5Al3 solvus range decreases markedly with lower temperatures, as seen by the cellular precipitation in the as-cast alloy (Figure 2 (b)). Additionally, there was precipitation of ~Pt5Al3 within the ~PtV dendrites, which shows that the ~PtV solvus retreats substantially with decreasing temperature.
Alloy 3 contained three phases, ~Pt3V, ~Pt3Al, and ~Pt2Al (in a small proportion), in the as-cast condition. In the annealed condition, there was ~Pt2V, ~PtV, and ~Pt3Al, and precipitation of ~Pt3Al in the ~Pt2V phase.
Alloy 4, which had an average overall composition Pt-17.4Al-6.4V, was single-phase ~Pt3Al in the as-cast condition. This alloy has not yet been heat treated. Alloy 5 had an average overall composition Pt-22.3Al-7.9V, with two phases, ~Pt3Al and ~Pt2Al in the as-cast condition and the same two phases after annealing.
Alloy 6 contained (Pt) dendrites and a (Pt) + ~Pt3Al eutectic in the as-cast condition. After annealing, the alloy had the same (Pt) and the (Pt) + ~Pt3Al eutectic as well as a third phase, ~Pt2V which must have formed as a result of a reaction during annealing.
The solidification projection, an isothermal section, and a liquidus surface projection for the Pt-rich corner of the Pt-AlV were found to be consistent with each other. Using the solidification sequences from the phases and their morphologies, a tentative liquidus surface was derived. The ~Pt2Al phase forms only in the solid state, and so does not appear on the liquidus surface.
Figure 16 (b) shows the solidification reactions for each of the as-cast alloys. These are schematic, since the analysed primary phase is assumed not to change on further solidification. Solidification starts with the composition moving directly away from the primary solid (on that primary liquidus surface) until a boundary with another liquidus surface is met, and the liquid composition then follows the boundary between the two liquidus surfaces for either a eutectic or a peritectic reaction (ignoring the fact that on non-equilibrium cooling, the liquid composition can sometimes run across the liquidus surface of the peritectic component. For Alloy 1, the second phase was also used to derive the direction of the liquid composition as the ~Pt5Al3 formed peritectically. Unfortunately the extrapolation of this falls on the direction of the liquidus for Alloy 5, but this is of no significance and is coincidental. Of the reactions observed, all had been reported in the Pt-Al binary, except for the peritectoid reaction ~Pt3Al + ~Pt5Al3 Â ~Pt3V, which formed as a thin layer.
Several phases showed solid state-precipitation, indicating that the solvus of the original phase retreats noticeably with decreasing temperature. These included ~Pt3Al, ~Pt5Al3, ~Pt2V, and ~PtV. The solidification projection and the isothermal section were consistent with each other.
The authors would like to thank AMSEN, the DST/NRF Centre of Excellence in Strong Materials, National Research Foundation and Department of Science and Technology, Mintek, Dr. Elma van der Lingen, Mr. Mokae Bambo, Mr. Richard Couperthwaite, and Mr. Edson Muhuma for their support and assistance at various stages in the project.
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©The Southern African Institute of Mining and Metallurgy, 2012. SA ISSN2225-6253. This paper was first presented at the ZrTa2011 New Metals Development Network Conference, 12-14 October 2011, Mount Grace Country House & Spa, Magaliesburg.