salt fluxing degradation final report- gary bywater

8/8/2019 Salt Fluxing Degradation Final Report- Gary Bywater

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University of Manchester

Student ID: 76041250

Materials Performance Module

Degradation Mechanism Final Report

Salt Fluxing Hot Corrosion Mechanism of Nickel-Base

Superalloy Gas Turbine Blades

Dr T J Marrow

March 2010





Hot Corrosion

Oxidation is the most important hot corrosion mechanism; metals/alloys will be oxidized when

heated to elevated temperatures or when operating in highly oxidizing environments such as a

combustion atmosphere. It is important to know that upon oxidation of a gas turbine blade an oxide

layer forms which can be protective or non protective depending on its chemical morphology and

adherence to the substrate. Initially an oxide layer which forms is sacrificial, regenerative, and

protective against other hot corrosion attack such as sulphidation, carburization, and as the concern

of this report molten salt fluxing attack. The alloy additions of chromium and aluminium allow the

superalloy to from a protective chromium/aluminium oxide layer.

Over time however due to the various hot corrosion mechanisms the oxide layer can either be

degraded exposing the substrate or, can be chemically changed to no longer be protective and will

result in a degrading porous oxide formation under the surface layer of the superalloy. When hot

corrosion of the turbine blade is discussed within this report it is assuming there is no specialised

surface coating present other than the protective oxide layer which is being degraded to expose the

superalloy for direct hot corrosion attack. The role of coatings and how they protect against hot

corrosion will be discussed later in this report.

The salt fluxing mechanism of hot corrosion involves the dissolution of the protective oxide layer at

the oxide/salt interface resulting in non-protective precipitates (such as chromium sulphides). Oxide

layer dissolution occurs by either the combination of oxides with O2-

to form anions (via basic

fluxing) or by decomposition of oxides into cations and O2-

(via acidic fluxing) depending on the salt

composition. Salts high in SO3 are acidic and basic when low in SO3. Sodium sulphate is the

dominant salt of hot corrosion due to its high thermodynamic stability. Molten salt hot corrosion

occurs predominantly in the range of 700-900˚C involving sulphur from jet fuel which reacts with

sodium chloride (NaCl) originating from ingested air during combustion in the combustor to form

molten sodium sulphate (Na2SO4) which deposits on the turbine blade. Firstly the protective oxide

layer is attacked and once degraded the superalloy substrate is attacked. Other salt contaminants

can be formed from the environment; oxygen and sulphur will combine with sodium, potassium,

vanadium, and chlorine within the gas turbine engine to create sodium chloride (NaCl), vanadium

oxide (V2O5), and potassium sulphate (K2SO4) (2) (3). Several chemical reactions can occur to form

sodium sulphate.

2NaCl + SO3 + 1/2O2 Na2SO4 + Cl2

2NaCl + SO3 + H2O Na2SO4 + 2HCl

2NaCl + SO2 + O2 Na2SO4 + Cl2

Hot corrosion can be separated into two forms, type I hot corrosion known has high temperature

hot corrosion (HTHC), and type II hot corrosion known as low temperature hot corrosion (LTHC).

HTHC occurs in the temperature range of 850-950˚C where pure sodium sulphate is above its

melting temperature. Sodium sulphate will form a mixture with the other salt contaminants such as

sodium chloride; this lowers the melting temperature as a eutectic is formed and broadens the

molten salt attack range. In this attack sulphur is released from the reduction of sodium sulphate

which diffuses into the superalloy to form chromium sulphides as in the reaction below (3).





Cr2O3 + O2-

2CrO2-

This reaction depletes the superalloy matrix of chromium allowing the oxidation of the base metal to

be accelerated and a porous non-protective oxide layer will form. It is found that

chromium/titanium sulphide particles deposit in the depleted matrix zones. As the reaction

proceeds released sulphur diffuses deeper into the substrate to form more sulphides inducing

further corrosion. The further the oxidation attack penetrates the larger the loss of structural

material will occur and failure will be inevitable. HTHC is a form of intergranular attack and is basic

fluxing. Figure 1 below shows the near surface morphology of HTHC attack.

Figure 1: High temperature hot corrosion (HTHC) (4)

Type II LTHC involves acidic fluxing and occurs approximately in the range of 600-800˚C (there is a

cross over range of HTHC and LTHC forming a transition attack). This form of corrosion again

involves sodium sulphate however eutectic mixtures are formed with nickel sulphate (NiSO4)

generating a eutectic mixture with a much lower melting temperature than that of HTHC. LTHC

sulphide eutectic mixtures are dependent upon the partial pressure of SO3 in a gaseous phase and

typically show no chromium depleted matrix and little intergranular attack. With a high partial

pressure of SO3 a high rate of attack is the result causing pitting of the surface. It is found that

chromium and titanium sulphides form a continuous layer. Figure 2 below shows surface

morphology of LTHC attack.

Figure 2: Low temperature hot corrosion (4)

It is common that a combination of HTHC and LTHC can occur resulting in a surface morphology

depicted in Figure 3.





Figure 3: Transition hot corrosion (4)

Influence of superalloy composition and microstructure

Considering the gas turbine blade independent of surface coating, the composition and

microstructure are directly responsible for the severity of hot corrosion attack (the former being

more influential due to significance of electrochemistry). Generally, the addition of alloying

elements such as aluminium, chromium and silicon are used to form protective oxide layers to resistattack from hot corrosion. Increased chromium content will improve HTHC resistance, with a deeper

chromium matrix reservoir to deplete (see Figure 4). However, excess amounts of chromium will

result in the precipitation of a topologically close-packed (TCP) phase which embrittles the alloy,

reducing ductility and the high temperature strength (1) (5).

Figure 4: Corrosion rate of various superalloys with chromium content (1)

An example of alloy development is shown in Figure 5. Pratt & Whitney improved a single crystal Ni-

base superalloy for gas turbine blade application by increasing the aluminium content from 5-

5.6wt% and reducing the titanium content to close to zero. The reason for this is it was found that

titanium ions introduced vacancies into alumina lattices increasing ionic mobility’s making the

previous superalloy more susceptible to fluxing.





Figure 5: Pratt & Whitney data of improved alloy, PWA1484

Critical alloy additions to be aware of are molybdenum (Mo) and tungsten (W). They are capable of causing catastrophic self-sustaining degradation as they form oxides which react with sodium

sulphate to form acidic salts high in SO2. This is a classic example of an acidic fluxing mechanism

forming compounds such as Na2MoO4 which posses a high solubility of the protective Al2O3 and

Cr2O3 layers. Molybdenum and tungsten are useful for enhancing mechanical properties but there is

a clear trade off in properties for hot corrosion resistance and vice versa (1) (2) (5).

Table 3: Role of alloying elements in superalloys (6)

The addition of rare earth elements have been found useful. Hafnium (Hf), lanthanum (La), and

yttrium (Y) bind strongly with sulphur which is responsible for sulphidation attack leading to

oxidation attack, as mentioned these mechanism are closely related to salt fluxing. Figure 6 shows

example test data of rare-earth metal doping.





Figure 6: Weight change data due to cyclic oxidation of a CMSX-4 single-crystal superalloy with a peak temperature of

1093C, and with variants of the superalloy doped with rare-earth elements La, Ce, and Y.

Ni-base superalloys are highly alloyed components. The hot corrosion behaviour is very complex

and there is the constant trade off of mechanical properties for hot corrosion/oxidation resistance

when varying composition and the behaviour is made even more complex when considering the

numerous mechanisms of hot corrosion attack and yet again, the behaviour of different superalloy

composition to each mechanism. From Figure 7 it can be noticed that later day generation single-

crystal superalloys have a lower hot corrosion resistance than their elders. This is due to the

development of surface coating technology to protect the superalloy from oxidation/hot

corrosion/thermal degradation and allows the design of the superalloy to focus on resisting the

mechanical stresses of operation. This leads to the focus of this report, the role of surface coatings.

Figure 7: Cyclic oxidation tests at 1100˚C of 1hour intervals of various uncoated single-crystal superalloys (1)





Surface Coatings

In order to improve the hot corrosion resistance of gas turbine blades specialised protective surface

layers are applied. There are many types of surface coatings available however the main three types

are: thermal barrier coatings (TBC), diffusion coatings, and overlay coatings. TBC’s are used to

insulate the turbine blade from the high operating temperatures by a few hundred degrees

centigrade. As the melting point of Ni-base superalloys are in the vicinity of 1450˚C which is also

their operating temperature, without TBC’s their safe application would not be possible. Thermal

protection is not the concern of this report so will not be touched on any further.

Diffusion and overlay coatings are used to protect the superalloy from environmental oxidation and

other various corrosive attacks. Diffusion coatings involve the powder deposition (and subsequent

diffusion) of aluminium onto (and into) the surface of a superalloy. The aluminium reacts with the

nickel to form intermetallic compounds of Ni3Al, NiAl and Ni2Al3. Heat treating increases adhesion

and results in inter-diffusion with the substrate generating an aluminium rich layer on the surface of

the component. With the high surface concentration of aluminium the superalloy substrate is now

protected from oxidation as the aluminium rich layer forms an aluminium-oxide (Al2O3) scale, a thick

continuous protective layer to the harsh oxidising conditions (1) (6).

Overlay coatings are an improvement in oxidation protection of superalloys than diffusion coatings

due to a more chemically stable oxide scale generated. The oxide scale is more chemically stable as

the deposited powder is prealloyed to be independent of the substrate alloy (some interdiffusion

will occur). High temperature overlay coatings are of the form MCrAlX, where M denotes nickel (Ni)

or cobalt (Co), Cr is chromium, Al is aluminium and X denotes a minor proportion element added to

enhance adherence of the oxide layer to the substrate: yttrium, hafnium and silicon are used.

Similar to the diffusion coating the aluminium forms an aluminium-oxide scale to resist oxidation

whereas the chromium protects against other various hot corrosion mechanisms (1) (6).

Table 4: Composition of commonly used MCrAlY overlay coatings in wt% (6)

Figure 8 below shows the significant synergistic improvement in surface coating protection of

overlay coatings with the addition of silicon and hafnium.





Figure 8: Progressive degradation of EB-PVD and plasma sprayed overlay coatings, showing synergistic benefit of silicon

and hafnium additions (1)

The method of deposition greatly influences the effectiveness of the surface coating. Diffusion

coatings are deposited using pack-cementation or chemical vapour deposition (CVD). Pack-

cementation involves the component to be immersed in the powder mixture to be deposited and is

heated within a sealed vessel in a protective atmosphere. The pack powder consists of the coating

element source, an activator (such as NaF, NaCl, or NH4Cl), and an inert filler material often alumina

to prevent the source from sintering. The deposition requires the activator to react with the

aluminium source to for aluminium halides which deposit onto the surface of the substrate alloy to

react and release aluminium. A flaw of pack cementation is that it is used to coat surfaces which are

brought into direct contact with the pack, complex cooling channels in gas turbine blades would not

be able to be coated and thus left vulnerable to oxidation/hot corrosion attack, hence the use of

CVD. CVD is similar to pack-cementation however the aluminisation takes place in the gaseous

phase (hence vapour deposition), requiring the reaction of AlCl3 powder in argon with hydrogen gas

to take place to form aluminium monochloride which reacts with the superalloy substrate to release

aluminium at the surface (1).

Overlay coatings (and TBC’s) are deposited by means of electron beam physical vapour deposition

(EB-PVD) and plasma spraying. The process uses an electron beam to vaporise an ingot of coating

material, once the coating material has evaporated it is deposited onto the surface of the superalloy

substrate. The key difference of EB-PVD to CVD is that no chemical reactions take place and it is a

much more expensive process to operate (1). Plasma spraying uses the energy from a thermally

ionised gas to melt and propel fine metal or oxide particles from powder form onto a surface so they

adhere and agglomerate to a produce a coating. Plasma sprayed layers have a characteristic ‘splat’

surface morphology. The advantage of plasma spraying is that the powder composition closely

matches that of the deposited coating, this is not always the case with EB-PVD due to evaporation





rates. A disadvantage of plasma spraying is the difficulty to produce dense layers without some

layer porosity.

Component testing

Gas turbine blades of different aircraft will operate under varying conditions, when concerning hotcorrosion the important lifetime factors are the amount of fuel air contamination and the blade

operating temperature. A key example of how salt contaminants influence turbine blade life is

shown below in Figure 9. An exponential decrease in life is caused by a small increase in parts per

million (ppm) of sodium salt contaminant. In this example the coating used is a platinum-aluminium

coat, where platinum is uniformly electroplated followed by a CVD layer of aluminium.

Figure 9: Effect of sodium sulphate hot corrosion on turbine blade life at 870˚C (7)

The surface coating of platinum-aluminium intermetallic is highly corrosion resistant. The example

above is the conventional method of taking lifetime data directly. The blades were run within the

same test machine side by side in severe corrosive conditions. An interim evaluation after 11,300

hours of service (consisting of 289 starts) showed the uncoated blade had 0.005inch corrosion attack

over 50% of the airfoil with approximately 0.01inch penetration at the base of the airfoil. The

coated blade on the other hand had no visual evidence of attack other than two small roughened

regions. Another example is through the work of Sidhu et al (8) where superalloys of composition

shown in Table 5 were tested uncoated and coated with an high velocity oxy-fuel sprayed (HVOF aform of plasma spraying) Ni-15.3Cr-3.1B-4.8Si-4.2Fe-0.6 wt% coating in a molten salt environment

consisting of Na2SO4 and V2O5 in the ratio of 40:60 wt% at 900˚C. HVOF is a thermal spraying

technique producing coats with lower porosity, higher hardness, superior bond strength, and lower

decarburisation than other spray techniques. The given composition of salt contaminants provides a

eutectic with a low melting point of 550˚C and is a rather aggressive environment.

Table 5: Sidhu et al superalloy compositions





The coated and uncoated alloys were subject to a cyclic regime of 50 cycles, each cycle being 1 hour

of heating at 900˚C. A thickness of 3-5mg/cm2

of salt contaminant mixture were deposited with a

camel hair brush on the preheated samples. The micrographs (Figure 10) show clearly that the

uncoated samples suffered a significant amount of spallation and sputtering of the surface due to

molten salt corrosion and clearly indicate the presence of cracks. The use of x-ray diffraction (Figure

12 & Figure 13) outlines the main phases present on the surface before and after oxidation/hot

corrosion. By clearly identifying the corrosion products the attack mechanism can be identified for

the coated and uncoated superalloys.

Figure 10: SEM analysis for the uncoated and coated samples after 50 cycles; a)uncoated Superni 600, b) coated Superni

600, c) uncoated Superni 601, d) coated Superni 601, e) uncoated Superfer800H, f)coated Superfer800H





Figure 11: SEM analysis of the as sprayed NiCrBSi coating

Figure 12: X-ray diffraction patterns for uncoated superalloys after 50 cycles; a) uncoated Superni 601, b) uncoated

Superfer 800H, and c) uncoated Superni 600

The x-ray diffraction peaks show that the uncoated corroded Superni 600 has the main phases of

NiO, Fe2O3, NiCr2O4, Ni(VO3)2, FeV2O4, and FeV occurring. Corroded Superni 601 shows similar

behaviour but peak intensities of NiO and Ni(VO3)2 are low.





Figure 13: X-ray diffraction patterns for coated superalloys after 50 cycles; a) coated Superni 600, b) coated Superni 601,

c) coated Superfer 800H

Figure 14 below shows clearly the uncoated Ni-base superalloys performed better than the Fe-base

superalloy yet the same surface coating resulted in practically equal oxidation/hot corrosion

resistance. This is an excellent outline of the significant benefits of surface coatings and how the

superalloy composition influences oxidation/hot corrosion of the uncoated superalloy.

Figure 14: Weight gain of specimens for up to 50 cycles





Bibliography

1. Reed, Roger C. The Superalloys Fundamentals and Applications. s.l. : Cambridge University Press,

2006. 978-0-521-85904-2.

2. Hot corrosion of some superalloys and role of high-velocity oxy-fuel spray coatings- a review.

Sidhu, T S, Agrawal, R D and Prakash, S. s.l. : Elsevier, 2005, Surface & Coatings Technology, Vol.

198, pp. 441-446.

3. Lai, George Y. High-Temperature Corrosion of Engineering Alloys. s.l. : ASM International, 1997. 0-

87170-411-0.

4. An investigation of blades failures in combustion turbines. Viswanathan, R. s.l. : Pergamon, 2001,

Engineering Failure Analysis, Vol. 8, pp. 493-511.

5. Hot corrosion in gas turbine components. Eliaz, N, Shemesh, G and Latanision, R M. s.l. : Elsevier,

2002.

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8. Hot corrosion behaviour of HVOF-sprayed NiCrBSi coatings on Ni- and Fe-based superalloys in

Na2SO4- 60% V2O5 environment at 900C. Sidhu, T S, Prakash, S and Agrawal, R D. s.l. : Elsevier,

2006, Acta Materialia, Vol. 54, pp. 773-784.

9. Bernstein, Henry L. Materials Issues For Users Of Gas Turbines. San Antonio, Texas : Gas TurbineMaterials Associates.

10. On the surface preparation of nickel superalloys before CoNiCrAlY deposition by thermal spray.

Bardi, U, et al. s.l. : Elsevier, 2004, Vol. 184, pp. 156-162.

11. The effect of EB PVD coatings on structure and properties of nickel-base superalloy for gas

turbine blades. Tchizhik, A A, et al. St. Petersburg, Russia : Elsevier, 1996, Surface & Coatings

Technology, Vol. 78, pp. 113-123.

12. Vapour aluminide coating of internal cooling channels, in turbine blades and vanes. Smith, A B,

Kempster, A and Smith, J. s.l. : Elsevier, 1999, Vols. 120-121, pp. 112-117.

13. Hot corrosion of materials: a fluxing mechanism? Rapp, Robert A. s.l. : Pergamon, 2002,

Corrosion Science, Vol. 44, pp. 209-221.

14. Chemistry and Electrochemistry of Hot Corrosion of Metals. Rapp, Robert A. s.l. : Elsevier, 1987,

Materials Science and Engineering, Vol. 87, pp. 319-327.

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salt fluxing degradation final report- gary bywater

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