PROYECTOS EN LOS QUE EL INVESTIGADOR PRINCIPAL PERTENECE A OTRO ORGANISMO

5.1. Introduction

During aluminising of high pressure turbines blades a slurry is used to mask part of the external surface to prevent the deposition in key areas. The procedure of ‘stop-off’ involves application of slurries to the part via paint brush before drying in a tunnel oven at ca 100°C.

These slurries consist of a nickel – chromium – aluminium alloy (M1™) and ceramic powder in a liquid binder, which contains a polymer dissolved in organic solvents that evaporate leaving behind the polymer to hold the paint to the surface. Typically two coats of M1™

‘stop-off’ are applied before a final two layers of M7™ ‘stop-off’, a different composition, are used to cap the surface. These maskants are manufactured by Akron Paints and Varnish (APV) and are bought in for use in the coating facility during aluminising. After processing in the chemical vapour deposition (CVD) rig the internal layers become a powder, which helps removal of the part from the maskant, this transformation is due to the breakdown of the polymer binder at temperatures during the process of 1030°C coupled with a nickel –

chromium – aluminium metal powder mixture (M7™) with sufficient ceramic such that it does not sinter. The outer layer forms a harder shell that forms a gas-tight seal and contains the powder during the CVD process.

The undesired result of this system is a sporadic appearance of small carbide precipitates beneath the surface correlating with those regions coated in slurry. These precipitates cannot be tolerated within single-crystal components as it is likely they pose a detriment to the physical properties due to the deviation from the intended microstructure.

Carbon added to an alloy composition prior to casting is distinctly different from the later addition of carbon via surface treatment. Removal of carbon from superalloys enabled the use of tailored heat treatments for effective homogenisation. This work concerns carbon addition during coating processing but much of the literature is not directly comparable, it does however contain information regarding carbide formation in nickel-based superalloys, but usually during casting. Before the introduction of single-crystal casting, superalloys were conventionally cast leading to polycrystalline structures with grain boundaries. Carbide formation upon grain boundaries enabled an increase in high temperature creep properties.

The alloy, Mar-M002, features ca 1.7 wt. % hafnium and 0.15% carbon for the purpose of hafnium-carbide formation.^108–110

103 5.1.1. Carbide types

There is an array of commonly found carbides within nickel-based superalloys, these include MC, M23C6, M7C3 and M6C, where M stands for a combination of metallic elements with the non-metallic element carbon (C).^108,111 The MC-type carbide usually displays a face centred cubic (NaCl) structure.¹⁰⁸ The types of carbides observed depend upon alloy

composition and high-temperature exposure. In general the MC-type carbide is considered stable from 850°C to 950°C, the M6C carbide is most stable above 1000°C and the M23C6

carbide appears in the 950°C – 1000°C temperature range. Above 1000°C script-like MC carbide decomposes into alternative carbides rapidly.^112,113

In work by Liu et al. ¹¹⁴ carbon was used in single-crystal alloys to help purify the melt and lend strength to sub-grain boundaries. Castings of large components such as nozzle guide vanes often feature carbon as an insurance, since to guarantee a single-crystal casting of such size is difficult. Liu et al. determines the benefits that MC-type carbides can bestow upon single-crystal castings include providing creep resistance. EDS work by Liu et al. determined that the carbides were rich in two elements, tantalum and titanium. Following a heat treatment, to relieve elemental segregation, Liu et al. found that some of the MC-type carbides had decomposed and formed M6C carbides, rich in molybdenum, tungsten and chromium, via the following reaction:

𝑀𝐶 + 𝛾 → 𝑀₆𝐶 + 𝛾^′

These samples were creep tested to determine the effect of the carbides. At 870°C the decomposition reaction was not activated and the MC-type carbides remained, these

carbides were found to contribute to crack initiation. The M6C carbides had two origins: firstly the decomposition reaction that saw MC-type carbides transform and second the

precipitation of M6C fresh from carbon latent within the crystal structure.¹¹⁴ The morphology of the M6C carbides that degrade from MC is irregular whilst the freshly precipitated M6C carbides feature an octahedral habit with sides of between 400 and 500nm, are coherent and feature an orientation relationship with the superalloy matrix. Liu et al. also determines the existence of M23C6 carbides which form during heat treatments at lower temperatures between 760 and 980°C. These M²³C6 carbides form preferentially in the γ channels and could be helpful in providing creep resistance yet their presence removes elements added to the alloy as solid solution strengtheners as well as the element chromium.

104 5.1.2. Carbide compositions

EDS analysis has revealed that script-like and angular carbides were generally rich in hafnium, tantalum, titanium and tungsten identifying them as MC-type. The occasional presence of other elements (nickel, chromium and cobalt) in the EDS analysis of the smaller particles was attributed to surrounding matrix contribution to the signal.¹⁰⁸ The individual compositions of the MC-type carbides were seen to vary from hafnium-lean (Ti-0.4-Ta-0.35-W-0.2-Hf0.05)C to hafnium-rich (Ti-0.07-Ta-0.12-W-0.06-Hf-0.75)C.¹⁰⁸ The compositions of needle shaped MC-type carbides were found to be tantalum – rich (ca 46 at. %), with niobium (ca 22 at. %) and titanium (ca 11 at. %). Hafnium, molybdenum and nickel concentrations lie around ca 5 at. %. 109,111,112

Finer carbide particles have been observed within the γ/γ’ eutectic and the dendritic regions. These were identified as chromium-rich M23C6.^108,115

EDS analysis of a cubic phase identified using TEM to be an M6C carbide was shown to be rich in molybdenum and tungsten.¹¹⁴ It is claimed that the M6C appears only when the molybdenum content of nickel- or cobalt- base superalloys exceeds ca 6 to 8 wt % (or its equivalent in tungsten atoms).^111,115

5.1.3. Orientation relationships

105 5.1.4. Spatial locations

Cubic shaped carbide precipitates, with a size between 50 and 500 nm, were found to have a strong tendency to nucleate within the γ matrix channels of superalloys with

prolonged high-temperature exposure. Discrete M23C6 particles formed within the γ channels may be beneficial to creep properties through blocking dislocation movement but they also withdraw solid solution strengthening elements from the matrix.¹¹⁶

Three types of carbides; MC, M23C6 and M6C were found upon grain boundaries. MC carbides rich in tantalum and titanium are mostly present in the blocky or needle form with finer M23C6 particles present at grain boundaries as well as in γ matrix channels. At the crossing point of two γ channels the M23C6 particles precipitated are larger than those only inside a single γ channel.^116,117 In the interior of grains, precipitation of the secondary M23C6

carbide occurred preferentially on crystal imperfections such as dislocations and stacking faults. Furthermore with increasing aging time the precipitates that nucleated upon the dislocations coarsened and coalesced into chains, while precipitates on the stacking faults grew into a lath shape.¹¹¹ Secondary carbides are chromium-rich M23C6 and tungsten-rich M6C.

There are a number of blocky, primary MC-type carbides precipitated along grain

boundaries forming coarse (2 to 10 μm in diameter) as well as elongated tantalum, hafnium, tungsten enriched precipitates within the grain. Fine (0.2 to 0.8 μm in diameter) globular M23C6 carbides were present in the superalloy microstructure after heat treatment. ^118,119

Tantalum and hafnium enriched MC-type carbides predominated within the matrix and chromium – rich M23C6 carbides precipitate preferentially at the periphery of the MC carbides and within the (γ + γ’) eutectic at grain boundaries.¹¹⁸ M23C6 particles are either formed aligning along grain boundaries or isolated to precipitate in globular morphology.¹¹² 5.1.5. Morphologies

Four tantalum–rich MC-type carbide morphologies were identified; discrete blocky dispersion, well-distributed script-like, needle and nodular.^11,112 Small, spherical carbides were also found to be tantalum – rich and were, therefore, considered to also be MC-type carbides. MC carbides tie up elements such as tantalum, titanium and tungsten, which can have an effect on γ/γ′ lattice misfit and therefore coarsening behaviour.¹²⁰

106

Fig. 5.1 – micrographs of tantalum-rich MC carbide morphologies observed in Rene N5 and N6 based alloys, blocky (a), script (b) and nodular (c) ¹¹

5.1.5.1. Blocky

The blocky MC precipitates have a larger misfit with the matrix than the needle MC carbides. The larger lattice misfit produces higher interface energy per unit area due to strain energy. There is no doubt that large misfit energy tends to reduce the surface area of a precipitate.¹⁰⁹ For the same volume, blocky MC-type carbides have a smaller surface area than needle-shaped carbides. Hence, the MC-type carbides tend to be blocky to reduce their total interface energy.¹⁰⁹

5.1.5.2. Octahedral

The octahedron-shape is a leading carbide morphology, and is considered by some to be the ’master’ carbide. At slow carbide growth rates, these carbides grow as layers on the faceted surface to maintain the equilibrium octahedron shape. The octahedron with {111}

faces is the equilibrium carbide shape due to a minimum interface energy.¹⁰⁹ 5.1.5.3. Acicular

In addition a needle-like phase is found which is inclined at an angle of 45° to the [001]

direction. EDS analysis shows that this acicular phase is rich in tungsten and molybdenum, so its morphology and composition indicate a TCP phase however, electron diffraction and chemical analyses prove that the needle-like phase is M6C carbide precipitate. Careful examination of the needle-like particles on various sections of the specimens shows that they are platelets, parallel to {110} planes of the matrix, which are aligned to 45° to [001]

orientation.¹¹⁴

5.1.5.4. Platelets

Studies of intragranular carbides showed that M23C6 carbides precipitated as platelets on {111} planes of the matrix.^112,121

5.1.5.5. Rods

M7C3 carbide is in the form of rods or irregular aggregates.¹¹¹

107 5.1.6. Internal structure

Some carbides were observed to have inner cores that etched differently yielding a central region with an optically different colour, additionally this region displayed darker (lower Z) contrast than the surrounding material in the SEM under backscatter mode.¹¹⁹ EDS spectra suggests these cores are carbo-nitrides, or more specifically (TiTa)CN. This phase is consistently observed at the cores of the MC-type carbides, it is concluded that it must form at higher temperatures during the solidification process than the MC.¹¹⁹ Further evidence of these titanium – rich MC-type carbides growing on Ti(CN) was identified in XRD spectra of particles chemically extracted from the superalloy material but again these were not detected as discrete particles with the TEM or SEM, supporting the conclusion that Ti(CN) (or the similar (Ti-0.5-W-0.5)C) exists as particles within the larger MC-type carbides and that in this work the formation of the ‘core’ without the subsequent outer layers is not observed.108,122,123

Fig. 5.2 – Colour light optical micrograph showing an ‘orange’ core in the MC-type carbide of the alloy IN738 (a) and SEM-BSE of one of the MC-type carbides containing a dark, low Z phase in the central core 108,122,123

The existence of the prior nitride or carbo-nitride has a very strong effect on the nucleation and growth behaviour of the subsequent carbide. When the size of nitride or carbo-nitride particles exceeds a critical carbide nucleation size, MC-type carbides begin to grow on their surface and envelop them. Therefore, the newly precipitated nitride or

carbonitride core within the carbide remains very small and cannot be detected by techniques including EPMA.¹⁰⁹ Final proof of this phenomena is that a carbide edge

composition has been seen to differ from the carbide centre composition. The carbide edge contains less titanium, but is enriched with hafnium, while tungsten and tantalum contents often remain constant through the entire carbide.¹⁰⁹

108

5.2. Experimental

5.2.1. Methodology

This investigation examined one design of high pressure turbine blade, referred to as blade Y, at various stages during manufacture and coating application, as detailed in Fig.

5.5. Blade Y, cast from CMSX-4, which was subject to various coating deposition techniques including platinum plating, vapour aluminising and TBC application through EB-PVD. Blade Y displayed a precipitate-rich band in the superalloy sub-surface region typically extending from the external surface to a depth of ca 100 – 200 μm deep. This feature, observed following vapour phase aluminising, detrimentally affected production yield since this new feature was unknown and potentially harmful.

Fig. 5.3 – macro images detailing research and development chemical vapour deposition furnace used in this work, outer bell retort (a), inner reaction vessel / retort (b), gas flow and furnace control (c), rig setup (d)

a b

c d

109

Fig. 5.4 – macro images providing further detail regarding CVD rig, outer retort immediately following a high temperature aluminising run (a), upturned lid of the inner retort / reaction vessel showing thermocouple inlet (b), inside base of the inner retort showing aluminising chips plus halide activator beneath inert gas distribution

manifold (c), inner retort torturous gas seal with 20 mesh alumina– note discolouration following reaction (d)

It was therefore necessary to determine which materials came into contact with blade Y components during this process and how the precipitates form. Aluminising in this instance was carried out using the above-the-pack method, this is where chromium-aluminium chips are placed within the inner retort and coated in a small layer of aluminising activator, in this instance AlF3. This set-up is seen in Fig. 5.4 (c) where the chips are sat beneath the inert gas distribution manifold which blasts the chips during the high-temperature process to aid circulation and agitation of the aluminising vapour cloud formed. This vapour cloud envelops the test pieces or components which are situated above the chips on racking or specifically designed tooling.

a

b

c d

110

Fig. 5.5 –coating processing for a single-crystal turbine blade

Aqueous cleanVacuum burnoutBlastMask for platingPlatinum plateDiffusion heat treat Machining of cooling holes

Stop off internal aluminising Aqueous cleanPenetrant processAqueous cleanDiffusion heat treatEB-PVD coatingBlastDiffusion heat treat

Vapour aluminising internals Super PolishDegreaseBlastHVOFAirflow and inspectPlastic bead blastHeat tintInspect