1

MATERIALS SELECTION FOR FUTURE MARINE GAS TURBINES

David A. Shifler, Ph.D, P.E. Office of Naval Research Naval Materials Division

Arlington, VA 22203-1995

Keywords: gas turbines, hot corrosion, superalloys, fatigue, creep, overlay coatings

ABSTRACT

Selection of materials for marine gas turbines demand that such materials have a variety of properties

such as high strength, toughness, creep resistance, fatigue resistance, as well as resistance to degradation

by their interaction with the environment. High temperature alloys derive their resistance to degradation

by forming and maintaining a continuous protective oxide surface layer that is slow-growing, very stable,

and adherent. In aggressive environments, superalloy oxidation and corrosion resistance needs to be aug-

mented by coatings that are compatible with the alloy substrate. Propulsion materials for Naval shipboard

gas turbine engines are subjected to the corrosive environment of the sea to differing degrees as a result

of several external factors. Increasing engine fuel efficiency and platform capabilities require higher op-

erating temperatures that may lead to new corrosion mechanism that degrades coatings and/or substrate

materials. Contaminants from several sources can influence corrosion rates which, in turn, can adversely

affect the life in these propulsion or auxiliary gas turbine engines.

This paper will dwell on some past and current results of materials testing and offer some views on

future directions into materials research in high temperature materials in aggressive environments that will

lead to new advanced propulsion materials for shipboard applications.

INTRODUCTION

The gas turbine engine generally was developed for main propulsion and auxiliary power for aircraft,

ships or other military and commercial platforms that require a higher power density than diesel engines

can generate. Achieving enhanced efficiency for marine gas turbines is a major challenge as the surround-

ing environment is highly aggressive. The successful performance and life of a marine gas turbine engine

depends not only on the design but also on the selection of appropriate materials of construction. Materials

for U.S. Navy shipboard gas turbines were largely developed between 1965 and 1995 to withstand the

harsh salt-laden environment [1]. Marine gas turbine engines serve as primary and auxiliary power

sources for several current classes of ships in the U.S. Navy. Service experiences in the 1960’s and

2019 Department of Defense – Allied Nations Technical Corrosion Conference

Paper No. 2019-0218-0314-000039

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1970’s led to the discoveries of severe high temperature corrosion in shipboard gas turbine engines that

largely absent in aircraft turbine engines or most commercial gas turbine installations. Although corrosion

is one consideration in the design process, the engine design must include materials that will perform

economically and possess adequate resistance to mechanical and severe thermal service conditions when

the engine is exposed to the operational marine environment.

Although various thermodynamic cycles such as the Atkinson and Humphrey cycles are being con-

sidered for gas turbine engine applications, they are still in various stages of research and development.

Current gas turbine engines employ the Brayton cycle. The Brayton cycle is a thermodynamic cycle that

describes the workings of a constant pressure heat engine which is an isentropic and isobaric process.

Ambient air is drawn into a piston compressor, where it is compressed. The compressed air then runs

through a mixing chamber where fuel is added, and pressurized air and fuel mixture are then ignited in an

expansion cylinder and energy is released, causing the heated air and combustion products to expand

through a piston/cylinder. A portion of the work extracted by the piston/cylinder is used to drive the com-

pressor through a crankshaft arrangement. The efficiency, , of the ideal Brayton cycle is:

= 1- 𝑇1

𝑇2 = 1- (

P2

P1)(−1)/

(eq. 1)

Where is the heat capacity ratio, T1 and P1 is temperature and pressure at atmospheric, T2 and P2 is the

temperature and pressure at the compressor exit [2]. Thus, raising the temperature, T2, increases the effi-

ciency of the gas turbine engine. Thus in the current effort to increase gas turbine efficiency or increase

power of a given gas turbine engine, efforts to increase engine temperatures are ongoing.

The U.S. Navy employs gas turbine engines for aero-propulsion and for propulsion and auxiliary

power for select surface ships. In this paper, only marine gas turbine engines will be discussed. All marine

gas turbines currently in the Fleet were derived (marinized) from aircraft engines. For the U.S. Navy, only

GE and Rolls Royce are currently manufacturers of marine gas turbine engines for propulsion and auxil-

iary power. The 501-K17 and 501-K34 were derived from Rolls Royce T56 Series III and T56-427 aero

engines [3]. The GE LM2500 was derived from CF6 and CF39 aero engines [3]. Using aero-derived en-

gines for developing marine gas turbines is due in part to minimize costs for engine development.

DEGRADATION FACTORS AFFECTING ENGINE PERFORMANCE

Materials for marine gas turbines, particularly in the engine hot sections, demand that such materials

have a variety of properties such as high strength, toughness, creep resistance, fatigue resistance, as well

as corrosion resistance. Nickel-base and cobalt-base superalloys have been the materials of choice in the

hottest areas of a marine gas turbine engine. Nickel-base superalloys are an extraordinary class of mate-

rials that have been crucial to the continued development of high performance turbine engines. These

superalloys enable load-bearing structures because they have the highest homologous temperature of any

common alloy system (Tm = 0.9, or 90% of their melting point). The widespread use of nickel-based

superalloys in turbine engines coupled with the fact that the thermodynamic efficiency of turbine engines

is increased with increasing turbine inlet temperatures has, in part, provided the motivation for increasing

the maximum-use temperature of superalloys. This paper will focus on alloy stability, creep resistance,

fatigue, and oxidation/corrosion resistance.

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Failure Modes of Industrial Gas Turbines Engines

Both GE and Rolls Royce utilize their industrial gas turbine line of gas turbines to support the U.S.

Navy’s requirements for marine gas turbines for propulsion and auxiliary power needs. Though marine

gas turbines tend to be very reliable, degradation and failures do occur. Blades account for 42% of failures

observed in gas turbines [4]. Factors leading to component failures need to understood, designs altered,

if possible, and advanced materials chosen to lessen material degradation, and failure occurrences. Meher-

Homji identified problems found in compressor airfoils and in the engine hot section [5]. Problems ob-

served in the compressor include; high cycle fatigue, axial compressor surge, rotating stall and flutter,

blade rubs, foreign and domestic object damage (FOD/DOD), migration of blading, root attachment prob-

lems, and erosion. Vibration from several sources is generally the root cause for high cycle fatigue of

compressor blading, while FOD or excessive inlet airflow distortion causes compressor surges leading to

blade failure. Air inlet design and corresponding air inlet velocities into an axial compressor should avoid

harmonics and strong resonances that could lead to blade vibration and surges that also cause distortions.

Compressor corrosion may lead to blade pitting and erosion by sand or fly ash can impact blade aerody-

namic performance and mechanical strength [5]. Decreasing a cross-section of a blade can expose a

component to premature creep as a result of increasing the resultant tensile stress imposed on the compo-

nent.

In the engine hot section, component degradation and failure may be caused by high cycle fatigue, low

cycle fatigue, thermal fatigue, creep, stress rupture, erosion, and oxidation/high temperature corrosion [5,

6]. Temperature fluctuations, temperature cycling, or large temperature gradients can lead to a severe

vibratory environment that can influence the fatigue resistance of the superalloy blade. Uneven blade or

nozzle clearances along with nozzle wakes, or mechanical forces can excite blade resonance that can im-

pose additional stresses onto the blades. Under the platform of a blade is an area that is subjected to high

stresses during operation and is also prone to corrosion that serves as initiation sites for fatigue damage.

Fretting damage leading stress concentrations and high cycle fatigue (HCF) damage may also be experi-

enced [5]. Although properties of nickel-base superalloys (thermal and microstructural stability, tensile

strength, fatigue strength, creep resistance, and hot corrosion resistance) have proven optimal for the hot

sections of marine gas turbines [7, 8, 9], degradation and failures do occur, often caused by multiple

mechanism acting conjointly. Rao et. al. observed that low and high temperature hot corrosion, erosion

and fatigue caused turbine blade failure for a 100MW turbine engine used in marine applications [10]. If

one is to choose a nickel base superalloy for a marine gas turbine engine, we must know what the engine

requirements will be: know will the engine will be operated, what is the temperature range for engine

operations, what is the expected engine power rating, what is the prevailing operating environment (ma-

rine, aero, land-based), what are the physical, mechanical, and corrosion resistant properties of these sup-

eralloys for marine uses. This will be discussed in the following sections.

Chemistry, Microstructure, and Thermal Stability

The microstructure of a typical superalloy consists of several different phases and may contain up to

approximately ten constitutive elements and over ten trace elements [11]:

(i) The gamma phase, denoted γ. This exhibits a FCC structure, and in nearly all cases it forms a con-

tinuous, matrix phase in which the other phases reside. It contains significant concentrations of ele-

ments such as cobalt, chromium, molybdenum, ruthenium and rhenium, where these are present,

since these prefer to reside in this phase.

(ii) The gamma prime precipitate, denoted γ(Ni3Al). This forms as a precipitate phase, which is often

coherent with the γ-matrix, and rich in elements such as aluminum, titanium and tantalum. In nickel–

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iron superalloys and those alloys rich in niobium, a related gamma double prime ordered phase,

γ(Ni3Nb), is preferred instead of γ.

(iii) Carbides and borides. Carbon, often present at concentrations up to 0. 2wt%, combines with reac-

tive elements such as titanium, tantalum and hafnium to form MC carbides. During processing or

service, these can decompose to other species, such as M23C6 and M6C, which prefer to reside on

the γ–grain boundaries, and which are rich in chromium, molybdenum and tungsten. Boron can

combine with elements such as chromium or molybdenum to borides which are present on the -

gran boundaries [12].

Improvements in engine thrust, fuel consumption and durability continue to be largely dependent upon

advances in superalloy technology and cooling systems [13]. Advances in casting-alloy development as

well as the improvement in melting practices have significantly contributed to increasing the temperature

capabilities of Ni-base superalloys strengthened by the ' phase [14-17]. Directional solidification [18, 19]

and single-crystal castings [19] are among the most important methods developed to control the micro-

structure. Extensive reviews have dealt with the mechanical properties of the ' phase based upon the

Ni3Al composition as well as '-containing alloys [20,21]. Typically, the mechanical properties of '-

strengthened alloys are evaluated after solution annealing, and a standard ageing treatment designed to

precipitate the strengthening ' phase. Subsequent prolonged exposure to higher temperatures, however,

can cause significant changes in the volume fraction and morphology of the ' phase [22-28]. Misfit dis-

locations between the ' and phases can form readily either during prolonged exposure to elevated tem-

peratures [22] or during plastic deformation [27, 28], resulting in loss of coherency between the two

phases. Other important microstructural changes include precipitation of topologically close-packed

phases such as sigma phase [29, 30] and decomposition of primary carbides [28, 31, 32].

Marine gas turbines run cooler than aero engines; engine temperatures have operated in the 500-

940°C range with the upper temperature operations occur less than 10 percent of the engine’s operating

life. For these temperatures, cobalt- or nickel based superalloys have been utilized as substrate materials

for hot section blades and vanes such as shown in Table 1. It is projected that future marine gas turbine

engines may operate at 600 to 1050°C with possible excursions to 1100°C. Marine engines are also pro-

jected to run near full power. For higher temperature engine operations marine superalloys must remain

stable over time that can extend beyond 20,000 to 30,000 hours. Most marine grade superalloys today are

used in the multicrystalline, equiaxed mode, but with the drive towards higher operating temperatures

future superalloy engine component alloys will be required to be directionally-solidified or in single crys-

talline form to endure greater threats to creep. New coatings will be utilized to protect the nickel-based or

cobalt-based superalloys from the aggressive marine environment.

Creep

The yield stress of nickel-based superalloys do not strongly decrease with increasing temperatures

like other alloys systems. In many cases, the yield stress of superalloys increases with increasing temper-

ature, until about 800°C is reached. Above 800°C the yield stress decreases rapidly [12]. Creep is a time-

dependent deformation that is a function of the combined effect of stress and temperature. There are

typically stages of creep: primary, secondary, and tertiary creep. Orr et al. was first to describe creep and

diffusion was initiated by a thermal activation energy [33,34]. There are several mechanisms that have

been proposed for creep; the rate-controlling mechanism depends on the stress level and on the tempera-

ture. For temperatures above T > 0.5Tm(°K), creep is a function of applied stress. There are four major

creep mechanisms proposed: (1) diffusion creep, either diffusion of vacancies (Nabarro [35] or Herring

[36] or diffusion in the grain boundaries and sliding (Coble [37]), (2) dislocation creep where dislocations

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Table 1- Nominal Compositions of Cast Superalloys

Nominal Composition, Weight %

Alloy

Designation C Ni Cr Co Mo Re Al Ti Ta W Fe Zr Other

CMSX-3 -- bal 8.0 4.8 0.6 -- 5.6 1.0 6.3 8.0 -- -- 0.1Hf

CMSX-4 -- bal 6.5 9.6 0.6 3.0 5.6 1.0 6.5 8.0 -- -- 0.1Hf

CMSX-6 -- bal 10.0 5.0 3.0 -- 4.8 4.7 6.0 -- -- -- 0.1Hf

Inconel 713C 0.12 74 12.5 --- 4.2 6 0.8 1.75 --- --- 0.1 0.012B 0.9Nb

Inconel 625 0.05 bal 21.5 -- 9.0 -- 0.2 0.2 -- -- 2.5 -- 3.6Nb

Inconel 718 0.04 53 19 --- 3 0.5 0.9 --- --- 18 -- 0.1Cu 5Nb

Inconel 738 0.17 61.5 16 8.5 1.75 3.4 3.4 --- 2.6 --- 0.1 0.01B

Inconel 792 0.2 60 13 9 2.0 3.2 4.2 --- 4 --- 0.1 0.02B 2Nb

MAR-M-246 0.15 60 9 10 2.5 5.5 2 --- 12.5 --- 0.05 0.015B

MAR-M-247 0.15 59 8.25 10 0.7 5.5 1 3 10 0.5 0.05 0.015B

MAR-M- 509 0.6 10 23.5 54.5 --- --- 0.2 3.5 7 --- 0.5 ---

Rene 80 0.17 60 14 9.5 4 3 5 --- 4 --- 0.03 0.015B

TMS 162 -- bal 2.9 5.8 3.9 4.9 5.8 -- 5.6 5.8 -- -- 6.0Ru. 0.09Hf

Rene N4 0.06 62 9.8 7.5 1.5 4.2 3.5 4.8 6 --- --- 0.004B, 0.5Nb, 0.15Hf

Rene N5 -- bal 7.0 8.0 2.0 3.0 6.2 -- 7.0 5.0 -- -- 0.2Hf

Waspaloy 0.07 57.5 19.5 13.5 4.2 1.2 3 --- --- 1 0.09 0.005B

pinned by various obstacles are overcome by dislocation climb by either vacancy or interstitial generation

or destruction [38,39], (3) dislocation glide occurs at high stresses and thermal activation, does not depend

on diffusion, and (4) grain boundary sliding with little grain deformation and limited diffusion.

The creep strengthening of polycrystalline nickel alloys is promoted by solid-solution strengthening

from the presence of solute atoms and from precipitation hardening due the presence of phases such as

[12]. Elemental additions to superalloys to support creep strengthening are [40]:

a) Elements that form substitutional solid solutions within the austenitic matrix: cobalt, iron,

chromium, vanadium, molybdenum, tungsten.8Elements that form precipitates: aluminum, ti-

tanium, niobium, tantalum (ex. Ni3Al, Ni3Ti).

b) Carbide-forming elements: chromium, Molybdenum, Tungsten, vanadium, niobium, tanta-

lum, titanium.

c) Elements that segregate along the grain boundaries: magnesium, boron, carbon, zirconium

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d) Elements forming protective and adherent oxides: chromium and aluminum.

e) Rare earths: yttrium, hafnium, zirconium, etc.

The mechanical properties of nickel alloys depend strongly on the condition of the microstructure

which is controlled by the chemical composition and the applied processing conditions. The yield stress

is related to the distribution of the phase. In a study by Piearcey et. al. displayed results which indicated

that at or beyond peak stress, the behavior of a nickel alloy is determined by the strength of the phase

[41]; above the peak stress temperature, the yield stress of a + two-phase alloy obeys the rule of

mixtures. The phase has a profound effect on the creep resistance of nickel alloys: as an increase of

volume fraction occurs, the creep resistance also increases.

Marine grade alloys used in gas turbines have tended to be multicrystalline, but to increase power

density and efficiency, higher operating temperatures are required. Thus, to avoid creep, marine gas tur-

bine alloys may be need to be in the form of single crystals or be directionally solidified to improve creep

resistance at higher operating temperatures. The larger the grain size, the greater the degree of strength-

ening [42]. Larger grain sizes, in general, possess higher creep resistance than multicrystalline grains.

For alloys of similar chemical composition, single crystal alloys can sustain and operate under higher

stresses without sign of creep damage than equiaxed, multicrystalline alloys since single alloy microstruc-

ture minimize the surface grain boundaries that enables grain boundary sliding. Figure 1 shows the com-

parative strength of an alloy in equiaxed, directionally-solidified, and single crystal alloy structures [43].

For the same general chemical compositions, the single crystalline structure has the highest creep strength

due to the loss of grain boundaries.

If single crystal alloys are chosen for marine gas turbine engines several guidelines to maximize creep

resistance and maintain alloy stability should be followed [44 ]. The first guideline specifies that the pro-

portions of -forming elements (ca. Al, Ti, Ta) should be sufficient to form fraction around 70%. Creep

deformation has been observed to be restricted to channels that lie between precipitates since disloca-

tions cannot penetrate these precipitates. To improve creep resistance the volume fraction of should be

minimized. Further studies have shown that the fraction vs creep performance is optimized at about

70% due to / interfaces also imparting resistance to creep deformation.

Figure 1 – Comparative creep strengths of an equiaxed, directionally- solidi-

fied, and single crystal alloy [43].

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The second guideline declares that the composition of the alloy needs to be controlled so that the /

lattice misfit is small, which minimizes the / interfacial energy so that coarsening is restricted. There

is a strong temperature dependence of the lattice misfit and the expansion coefficient of is much less

than the expansion coefficient of [45]. The resultant misfit, , between and coefficients may be

either positive or negative. There is some evidence that a negative misfit leads to an alloy with greater

creep resistance than an alloy with a positive misfit. This is due to the additive effects of a negative misfit

stress with a positive external applied stresses from dislocation activity in the channels [45].

Concentrations of creep-strengthening elements such as Cr, Re, W, T, Mo, and Ru must be significant,

but need to be insufficient to cause precipitation of topological closed packed (TCP) phases. Re has a

strongest creep-strengthening effect on alloys. However, the concentrations of refractory elements in

superalloys is limited. Excessive concentrations of Cr, Mo, W, and Re promote precipitation of TCP

phases which denotes a lack of alloy stability, which negatively impacts creep resistance and rupture life.

TCP precipitation destroys the / structure which is responsible for creep resistance.

The final guideline for superalloy design is that the alloy composition must avoid surface degradation

from exposure of hot working gases. Corrosion and/or oxidation needs to be avoided to maintain the wall

thickness and the load bearing capacity of the engine component. The operating temperure needs to be

known since oxidation kinetics and phase changes among NiO, Cr2O3, and Al2O3, can destroy passivity.

Small rare earths additions such as Hf, Y, and La, strongly bind to sulfur impurities in the alloy. The

“gettering” effect improves oxidation resistance by removing sulfur from the metal/oxide interface thus

strengthening the Van der Waals bond between the oxide and metal substrate [46].

Fatigue

Superalloys in equiaxed, directionally solidified, or in single crystal form are susceptible to fatigue

failures due to cyclic, oscillating stresses that are usually in the elastic range. Initiation sites for fatigue

occur at stress concentrations sites, often at surface pits, casting pores, machining marks, or cracks in the

oxide scale. At these initiation sites, plastic deformation proceeds as intense localized slip with dislocation

activity focused on a limited number of lattice planes [47]. At high temperatures (~ 700°C) and high

stresses, persistent slip band (PCBs) can shear precipitates. Dislocation activity is restricted to chan-

nels at higher temperatures and low stresses. Two distinct forms of fatigue occur: (1) low-cycle fatigue

(LCF) and high-cycle fatigue (HCF).

LCF occurs when the stress amplitude is at or above the elastic limit of the alloy failure generally is

experienced with 105 cycles, and most of the fatigue life is expended in the propagation stage (as opposed

to the initiation stage). Events attributable to LCF damage originates from abrupt lading changes such as

during engines startup or stopping. Under LCF test conditions, generally under total strain control, fatigue

performance depends strongly on the orientation of the loading axis relative to the alloy crystallographic

orientation (assuming single crystal usage). LCF life decrease rapidly with increasing temperature.

HCF originates from smaller stress amplitudes in the elastic range, typical as imposed on turbine

blades by vibrations or resonance. Most of the fatigue life under HCF is consumed in the initiation stage

and failure occurs in excess of 105 cycles. The natural resonant vibration frequencies of single crystal

turbine blades range from a few to several thousand hertz. Excitation frequencies in this range from

turbulent flow around the airfoil will promote HCF. HCF loads are superimposed on the steady-state

mean stresses caused by centrifugal loading of the blade and thermal loads from the thermal gradients

imposed through the blade by passing hot gas and cooling air and failure typically occurs in the range of

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107 cycles. Creep and fatigue interactions can alter the appearance of the fracture surfaces, the fracture

appearance on the mean stress vs. the magnitude of the cyclic stress [44].

Corrosion

Hot Corrosion

Though creep and fatigue may be operative and impact the design life of a turbine component, expo-

sure to the very aggressive marine environment often leads to corrosion that participates in initiating and

accelerating degradation by both mechanisms.

Hot corrosion is a complex process involving both sulfidation and oxidation [48-51]. This form of

accelerated oxidation affects alloys and coatings exposed to high-temperature gases contaminated with

sulfur and alkali metal salts [52,53]. Hot corrosion proceeds via two stages: (1) an incubation stage and

(2) a propagation stage. The characteristics of hot corrosion are determined by exposure time, the nature

of the deposit (gas, liquid, or solid), amount of deposit, deposit chemical composition, temperature, gas

composition, alloy or coating composition, alloy process (i.e. cast or wrought), coating application

method, gas velocity, cyclic versus isothermal conditions, erosion, and/or specimen geometry. Contami-

nants derived from the gas phase, fuel sulfur, and salts from the air intake form alkali metal sulfates. If the

temperature of the alloy is below the dew point of the alkali sulfate vapors and above the sulfate melting

points, molten sulfate deposits are formed [53]. In the open literature, molten sodium sulfate is credited

to be the principal agent in causing hot corrosion [54,55]. Sulfur compounds come from three sources in

a marine gas turbine engine: (1) sulfur from the combustion fuel that often ranges from 15 ppm to 1 wt.%

(or more) depending on the grade of fuel, (2) sulfate salts and other sulfur species contained in the marine

air that are ingested into the hot section of the turbine engine, and (3) seawater contamination of the fuel

from seawater-compensated fuel tanks.

Salt aerosols may contain 0.01 ppm in good weather and as much as 3 ppm in bad weather [56]. Air

intakes of marine gas turbine are protected by filters may remove up to 70 percent of the ingested salt

[57]. Salt in fuel may be about 1 ppm [20]. The residence time of gases in the combustion chamber is

very short (5-10 ms). The sulphation rate of sodium chloride was measured to determine whether gaseous

sodium sulfate is formed at low salt loadings and short residence times in the turbine engine. The results

indicated that significant gaseous Na2SO4 formed too slowly [56] to contribute to sulfate formation.

Two distinct forms of sulfate hot corrosion exist, where fluxing with corroding salts overcomes the

protective oxide scales that form on superalloys or coatings. As shown in Figure 2 both Type I and Type

II hot corrosion can be viewed as departures from the Arrhenius oxidation rate vs. temperature.

The thermodynamics of oxyanion melts is in equilibrium with the oxide anion and the appropriate gas

species. For sulfate melts, the acid-base reaction is:

Na2SO4 Na2O + SO3 (g) (eq. 2)

where Na2O is the base and SO3 is the acid. Rapp indicated that by measuring Na+ cation and O2- anion

concentrations, the Na and O2 activities in the salt could be measured independently to yield log𝑎𝑁𝑎2𝑂,

the melt basicity, and with an oxygen sensing electrode, its oxidation potential [58]. The melt acidity is

log𝑎𝑆𝑂3.

Bornstein and DeCrescente [59,60] performed experiments on nickel-based superalloys and con-

cluded that the hot corrosion attack of this nickel-based alloy occurred due to basic Na2O from these nitrate

or sulfate salts, which destroyed the alloy NiO passive film by basic fluxing according to eq. (3):

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Figure 2 – Profile of shipboard engine corrosion rates of Type I and Type II hot corrosion versus tem-

perature in a marine environment compared to the Arrhenius oxidation rate vs. temperature. Courtesy

of U.S. Navy

NiO + Na2SO4 2Na+ + NiO22− + SO3 (eq. 3)

Although the role of SO3 is not fully settled in Type I hot corrosion, it is the general consensus that the

reaction of the Na2SO4 with the nickel overrides any influence of the gaseous environment [61].

Type I high temperature hot corrosion/sulfidation occurs through basic fluxing and subsequent disso-

lution of the normally protective oxide scales by molten sulfate deposits that accumulate on high temper-

ature component surfaces and is characterized by a smooth base alloy oxide interface and a continuous,

uniform precipitate depleted zone containing discrete sulfide particles beneath the oxide scale [62]. High

temperature hot corrosion usually occurs at metal temperatures ranging from 850 to 950C (1560–

1740F). Type I hot corrosion involves general broad attack caused by internal sulfidation above 800C;

alloy depletion is generally associated with the corrosion front. Molten sodium sulfate is the principal

agent in promoting hot corrosion [54,55]. Basic fluxing attack involves raising the Na2O activity in the

molten sulfate by formation of metal sulfides [60,62,63]. Very small amounts of sulfur and sodium or

potassium can produce sufficient Na2SO4 or K2SO4 to initiate the attack. Unfiltered air quality may contain

up to 2600 ppm Na2SO4, 19,000 ppm NaCl, and other seawater-derived species that further contributes to

the environmental corrosive impact on materials in shipboard marine gas turbines [64]. The incubation

stage involves the time for conditions to develop before dissolution of the protective scale occurs, allowing

direct access of the melt to the underlying alloy. Once the protective oxide scales are penetrated, the

substrate alloy or coating becomes vulnerable to accelerated oxidation and sulfidation. Above approxi-

mately 950C (1740F), oxidation becomes the primary material degradation mechanism.

A major characteristic of Type II, low-temperature hot corrosion, is pitting attack [65-71]. At the

corrosion product-gas interface the pit was enriched in cobalt or nickel depending upon the base element

of the alloy, but cobalt/nickel concentration decreases as the pit becomes deeper as the corrosion pro-

gresses into the alloy [61]. Under conditions where sodium salts ingested into the gas turbine causes

deposition of sodium sulfate (Na2SO4) on hot section airfoils, hot corrosion of coatings and alloys with

moderate (10-20%) chromium content occurs at significantly greater rates at 700°C than at 800° to

1000°C. This is viewed as Type II, low temperature hot corrosion that occurs in the temperature range of

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650–750C (1202–1382F), where layer type corrosion is characterized by pitting attack, an uneven base

alloy/oxide interface, and the absence of subscale sulfides. Here the partial pressure pSO3 is relatively

high for melts deficient in the oxide ion concentration leading to acidic fluxing of metal oxides that cause

pitting from the formation of low melting mixtures of Na2SO4 and NiSO4 or Na2SO4 and CoSO4 in Ni–

Cr, Co–Cr, and Co–Cr–Al alloys [72-78]. A number of mechanisms have been proposed to account for

the Type II hot corrosion [71,72,79,80], all of which propose that at temperatures of 650C to 750C solid

Na2SO4 is converted to a liquid solution of either Na2SO4–CoSO4 or Na2SO4–NiSO4 [60]. Smeggil was

able to demonstrate, however, that low temperature hot corrosion could occur as low as 530°C with mixed

sulfate salts [81].

Jones showed that SO3 was the critical component in forming low temperature hot corrosion deposits

[82]. Type II hot corrosion is initiated when the effective SO3 pressure (the sum of gas stream and locally

produced SO3) is high enough to form CoSO4-Na2SO4 compositions within the liquid phase region of the

CoSO4-Na2SO4 system [83]. It is also known that a partial pressure of SO3 of about 10-5 atm is required

for this melt to form [82]. This SO3 pressure exists over a range of oxygen pressures up to the CoO/Co3O4

equilibrium [61]. Above this oxygen pressure the following reaction must be used to calculate SO3 pres-

sure, (eq. 4):

𝟏

𝟑 CO3O4 + SO3 COSO4 +

𝟏

𝟔 O2 (eq. 4)

SO3 pressure increases as the oxygen pressure increases, but the total pressure cannot exceed 1 atm for

tests conducted at atmospheric pressure. Reaction with the liquid sulfate and the alloy surface causes

decomposition of the mixed sulfate and SO3 is generated at the alloy-molten sulfate interface near 1 at-

mosphere pressure. High SO3 pressure causes pitting. Very small changes in PSO3 can cause substantial

increases in corrosion if a transition from a solid to a liquid sulfate occurs. Though high SO3 concentra-

tions favor corrosion, Type II hot corrosion is possible even when the gas stream SO3/SO2 ratio is very

low [83].

When nickel-base superalloys contain certain refractory constituents such as tungsten, molybdenum,

or vanadium, catastrophic hot corrosion due to Mo, W, or V in the alloy reacting with Na2SO4 or NaNO3,

forms Na2MoO4 Na2WO4 or Na3WO4 with a high activity of MoO3, WO3, or V2O5. Alumina protective

scales undergo alloy-induced acidic fluxing involving the following reaction (eq. 5):

Na2SO4 + MoO3 NaMoO4 + SO3 (eq. 5)

This form of hot corrosion, which has been called alloy-induced acidic fluxing, is observed at temper-

atures between 900°C and 1000°C. Mishra has shown that the alloy-induced acidic fluxing process is

influenced by the SO3 pressure in the gas phase [84]. When coated with Na2SO4 those alloys with 2 wt%

or more Mo showed degradation products similar to those observed previously in Mo-containing alloys,

which undergo alloy-induced acidic fluxing Type I hot corrosion even though the temperatures were in

the Type II hot corrosion range [85]. The Na2SO4 deposit can become acidic due to alloy-induced acidic

fluxing to form the very acidic refractory alloy oxides.

The incubation stage involves the time for conditions to develop before dissolution of the protective

scale occurs, allowing direct access of the melt to the underlying alloy. Once the protective oxide scales

are penetrated, the substrate alloy or coating becomes vulnerable to accelerated oxidation and sulfidation.

11

New Insights into Hot Corrosion

It was generally accepted that if the fuel sulfur was removed from fuels, turbine engine materials

would cease to experience hot corrosion. Recent Navy standards for Navy F-76 diesel distillate have

dropped the maximum fuel sulfur content down to 15 ppm [86]. Unlike previous fuel specifications, this

Navy specification allows up to maximum of 50 volume percent of synthetic fuel from a number of spec-

ified synthetic fuel processes to be mixed with petroleum-based fuel. Military and commercial specifica-

tions for aviation fuels have also dropped the maximum fuel sulfur content to15 ppm, yet instances of hot

corrosion still persist. Alkali and alkaline earth metals such as Na, Ca, Mg, Mg, K, must be tightly con-

trolled to contain no more than 1ppm maximum in NATO F-76 during synthetic fuel blending. Alkali

contaminants from synthetic fuel production are several orders of magnitude higher and would otherwise

accelerate corrosion is not tightly controlled [87]; aviation grade fuels have no specified limits on alkali

and alkaline earth metals. Impurities in fuel may combine with Na2O to form mixed salts with reduced

melting temperatures. Copper affects fuel stability, lead forms low-temperature eutectics, sodium and

potassium increases corrosion, and calcium acts as surface glue that enhances corrosion activity [88].

High carbon content in the fuel can also increase corrosion rates because char particles can trap sulfur-

containing species on turbine airfoils. Reducing environments have a far more detrimental effect of hot

corrosion of turbine materials than viewed in oxidizing environments [54,88,89].

Although tight controls on fuel quality may mitigate the effects of many potential species, elements,

and compounds that can influence hot corrosion, other sources can impact the degree of hot corrosion in

marine and aero gas turbines. Sulfur particulate matter (PM) may still enter the combustion chamber via

the ship air intake or from seawater contamination of the fuel. Seawater contains chlorides and sulfates

of magnesium, calcium, and potassium salts, in addition to sodium sulfate and sodium chloride as shown

in Table 2. Sodium sulfate traditionally was viewed as the sole reactant responsible for hot corrosion, yet

magnesium, potassium, and calcium sulfates are also contained in seawater in significant amounts [90].

The sulfate speciation and total solids in seawater and the content of atmospheric particulate matter (PM)

varies considerably around the world and in some places, such as China and India, the high PM content

seems to cause an increased risk of deposit-induced hot corrosion due to atmospheric pollutants and PM

chemical content (i.e. sulfates, nitrates, chlorides, ammonium cations, Si, Ca, Al, Fe, Na, K, and Mg). The

increasing operating temperatures of turbine engines and activities within regions of the world that have

relatively high pollutant (PM, SO2, etc.) levels appear to be working conjointly to cause previously unob-

served forms of high-temperature corrosion.

Table 2--Major Seawater Constituents

at (35 Parts per Thousand (o/oo) Salinity) and 25 C. Reference [90]

Ion or Molecule Concentration

mmol/kg g/kg

Na+ 468.5 10.77

K+ 10.21 0.399

Mg+2 53.08 1.290

Ca+2 10.28 0.4121

Cl- 545.9 19.354

SO4-2 28.23 2.712

Source: S.J. Dexter, C. Culberson, "Global Variability of Natural Sea Water," Materials Performance, v. 19 (9),

16 (1980).

12

While sodium sulfate itself is molten at 884C, mixtures of Na, K, Ca, and Mg sulfates increase hot

corrosion range because these salt mixtures melt at 695-698C and thus are partly to fully molten at 700°C.

The mixture of MgSO4 and Na2SO4 promotes accelerated Type II hot corrosion at 700C because the

eutectic salt mixture melts at 666C [91]. Figure 3 shows that the average metal loss increases as the

liquid volume fraction increases at 700C. Other impurities such as phosphorus, lead, chlorides, sand, and

unburned carbon could also lower salt melting temperatures, altering the sulfate activity, changing the

solution chemistry, or influencing the salt acidity/basicity that contributes to accelerating hot corrosion.

Figure 3 - Hot Corrosion at 700°C in Air: Effect of Salt. Salt mixtures have lower melting

temperatures. The corrosion rate increases as the volume fraction of liquid increases. Unpublished, courtesy of Z. Tang and B. Gleeson, University of Pittsburgh

However, the deposit chemistry influencing hot corrosion is much more complex, consisting of mul-

tiple sulfates and silicates [92-94]. The variability of SO2 concentrations and the prevailing sulfate salt

composition of seawater mentioned earlier at different cities around the world are shown in Table 3. Sul-

fate speciation and content of atmospheric PM varies considerably around the world [65]. The designa-

tion PM2.5 designates particulate particles less than 2.5μm in size. The increasing operating temperatures

of turbine engines and activities within regions of the world that have relatively high pollutant (PM, SO2,

etc.) levels are working conjointly to cause previously unobserved forms of high-temperature corrosion

[92]. Dust and similar matter tend to be hydroscopic, picking up water. The average SO2 concentration

(0.055 ppm SO2) in Beijing is about 8 times greater than the SO2 levels in New Delhi, India, and 27 times

greater than Los Angeles [94]. Seasonal variations can raise Beijing sulfur dioxide concentrations to 0.14

ppm during the winter months [96,97]. In addition, the sulfur concentration within a gas turbine engine

would be amplified by the pressure ratio; for instance, a 30:1 pressure ratio would concentrate the sulfur

Table 3 - Regional Variability of Environments

*Assumes that all measured Na+, K+, Mg2+ and Ca2+ in PM are present in the form of sulfate.

Source: Unpublished, courtesy of Z. Tang and B. Gleeson, University of Pittsburgh

13

content 30 times the observed atmospheric values of 1.65 ppm to 4.2 ppm maximum, respectively, in

Beijing. Gleeson reported that 2 ppm sulfur in the atmosphere (such as SO2) could promote the transition

from internal to external Al2O3-scale formation [98] and form a thin continuous alumina layer.

The work of Lawson et al. [99] showed that at 700°C and 1,000C in either acidic (high pSO3 ) or

basic (high Na2O activity) sulfate melts, the hot corrosion of bulk Al2O3 is minimal and is mainly con-

trolled by impurity content, with high levels of impurities (SO2) above 2 ppm leading to hot corrosion

attack. This is in contrast to a statement by Pettit who indicated that Type II hot corrosion is not be

observed unless the fuel sulfur content was 0.25 wt% (~ 25,000 ppm) or above [100]. Figure 4 shows that

only 2.5 ppm SO2 is needed to accelerate hot corrosion.

Figure 4 – Effect of pSO2 on Na2SO4- induced Corrosion at 700C. 2.5 ppm atmospheric SO2 can greatly

accelerate hot corrosion over clean air conditions. Source: Unpublished, courtesy of Z. Tang and B.

Gleeson, University of Pittsburgh

Hot Corrosion involving CaSO4

The drive for operating gas turbines at higher temperatures has revealed another mechanism for hot

corrosion involving CaSO4 in equilibrium with CaO as depicted in eq. (6):

CaSO4 CaO + SO3 (g) (eq. 6)

The reactivity of -NiAl (Ni–20Co–16Cr–23Al–0.1Y (at.%)) + -Ni-based NiCoCrAlY (Ni–30Co–

33Cr–12Al–0.1Y) alloys with and without CaO deposits was studied by means of isothermal exposures

in air for durations ranging between 5 and 250 h [101]. In the absence of a deposit, both alloys formed

adherent Al2O3 scales, with no or very little oxides of Ni, Co or Cr. The layer sequence was Al-rich CA2

nearest to the substrate, followed by equimolar CA, and occasionally the Ca-rich C12A7 and C3A toward

the CaO deposit [101]. A thin, continuous Al-rich oxide layer, identified as -Al2O3 was always present

between the metal and the first aluminate layer.

Reaction with CaO at 1100C produced multi-layer scales of Al2O3 and calcium aluminates (xCaO–

yAl2O3), and a mixture of liquid calcium chromate and nickel–cobalt oxide particles. Calcium chromate

formation was a rapid, transient process, and the transition to a steady-state of slower Al2O3 growth was

favored by increasing the alloy fraction. The thermally-growing Al2O3 reacted with the deposit to form

calcium aluminates in a solid-state diffusion process, which led to an increased oxidation rate. Calcium

oxide is known to react with the chromium contained in MCrAlY alloys to form a low-melting calcium

chromate, which is liquid at 1100C [102]. The calcium chromate was found to develop rapidly on the -

rich Ni–30Co–33Cr–12Al–0.1Y alloy within 5 h exposure, and to not significantly thicken thereafter.

14

Unlike increasing resistance to sodium sulfate or vanadate-induced hot corrosion by increasing the chro-

mium content of an alloy, increasing the chromium content of a superalloy subjected to CaO-induced

corrosion actually increases the severity of degradation.

The reaction of cast NiCoCrAlY alloys with oxide–sulfate deposits in CO2–H2O–O2 was studied at

1100C. Applying sulfate–oxide deposits to Al2O3-forming NiCoCrAlY alloys led to enhanced degrada-

tion due to the presence of CaO, SiO2, and Na2SO4 [103]. The minimum Al concentration needed to form

an external Al2O3 scale was increased compared with deposit-free exposures, as Al2O3-forming composi-

tions transitioned to internal Al2O3 and external Cr2O3 growth in the presence of certain deposits. Model

deposits were used to investigate the role of each constituent in the complex reaction morphology observed

with an industrial fly-ash. Two main modes of degradation were identified, which involved Al2O3 disso-

lution in molten Na silicate and solid-state Al2O3 reaction with CaO. Both led to enhanced Al consump-

tion and promoted non-selective oxidation. Additions of Al2O3 or SiO2 decreased the CaO reactivity due

to the formation of aluminates or silicates, while Na2SO4, on the contrary, enhanced the degradation by

providing rapid mass transport in the molten state, and reduced alloy/scale adherence.

The Cr-rich -phase initially reacted with CaO to form a molten Ca chromate, Ca aluminates, and Ni

and Co oxides. Once an external Al2O3 scale was established, it either reacted with CaO to form Ca alu-

minates, or dissolved in molten Na silicate. In both cases, this increased Al consumption, promoting Al2O3

failure for -rich alloys. These primary reaction modes were affected by the deposit compositions: SiO2

and Al2O3 were found to reduce the reactivity of CaO, due to the formation of Ca aluminates and silicates

in the deposit. Sodium sulfate, on the other hand, provided faster Ca transport and enhanced its reactivity

toward the alloys. It also produced a locally high sulfur activity at the metal/oxide interface, and reduced

scale adherence [103].

CMAS in Ship Gas-Turbine Engines

Recent inspection of a shipboard gas turbine component under the platform has indicated the apparent

presence of calcium-magnesium-alumino-silicate (CMAS) [104]. This type of attack has been observed

in aero gas turbine engines when sand and similar siliceous matter melts; however, this phenomenon has

not been known to occur under temperatures below about 1130C. Metal temperatures in a shipboard

engine rarely exceed 950C, and temperatures are even lower at ‘under the platform’ dovetail sites. NaCl

additions to CMAS turbine deposits have depressed the melting temperature of CMAS deposits in labor-

atory tests observed in chloride-free deposits [105]. A typical example of a high pressure turbine (HPT)

blade from a marine gas turbine engine showing apparent CMAS ingestion is depicted in Figure 5 [106].

Figure 5. A typical example of a HPT blade in the hot section of a marine gas turbine

engine showing significant ingestion of apparent CMAS deposits.

15

The maximum temperature capability of thermal barrier systems used in gas turbines is often limited

by deposits of CMAS [105,107,108]. It is not completely known at this time the corrosive effect of CMAS

on shipboard materials. The CMAS material melts over a range TM ~ 1150-1240oC, depending upon the

specific composition of the material. CMAS penetration within the thermal barrier coating (TBC) pore

network is arrested at a depth dictated by the rate of the capillary-driven flow, the crystallization kinetics

of CMAS, and the imposed thermal gradient [109-111]. When the TBC surface exceeds the particular

CMAS melt temperature, the excellent wetting characteristics of the CMAS cause it to penetrate into

pores/cracks and between the columns of the TBC to a depth where the temperature of the TBC equals

that of the CMAS melt temperature. Upon cooling, it solidifies, modifying the local TBC compliance

which increases the TBC’s susceptibility to spalling. Release of the elastic strain energy stored in the

coating, caused by the thermal expansion mismatch between the coating and the substrate upon cooling,

supplies the underlying driving force for the TBC delamination and spalling [112]. It is not completely

known at this time the corrosive effect of CMAS will have on ship engine materials.

COATINGS THAT IMPROVE PERFORMANCE

Materials for marine gas turbines, particularly in the engine hot sections, demand that such materials

have a variety of properties such as high strength, toughness, creep resistance, fatigue resistance, as well

as corrosion resistance. Nickel-base and cobalt-base superalloys have been the materials of choice in the

hottest sections of a marine gas turbine engine and the task to design these alloys to resist the various

modes of degradation is a difficult. Superalloys such as MAR-M-509, Inconel 738, or Waspaloy (Table

1) exposed to the corrosive marine environment have higher levels of chromium that will better resist hot

corrosion than alloys such TMS 162, Rene 5, or CMSX-4. The chromium increase does limit the strength

and temperature range that the superalloy may operate due to the potential precipitation of TCP phases

which negatively impacts creep resistance and rupture life. TCP precipitation destroys the / structure

which is responsible for creep resistance.

To overcome limitations of superalloy chemistry to optimize engine performance exposed to the tem-

perature and environmental variations, coatings are usually applied to superalloys. Overlay and diffusion

coatings are metallic coatings that are generally applied on marine hot section gas turbine components.

Diffusion coatings consist of a substrate alloy surface layer enriched with oxide formers such a Cr,

Al, or Si [113,114]. Diffusion aluminide coatings are typically processed via slurry-fusion or pack cemen-

tation. Subsequent heat treatments help develop the proper mechanical properties and allow further coat-

ing diffusion. The principal protective oxide which causes improved oxidation and corrosion resistance

of aluminide coatings is Al2O3, although less protective oxides can form if other alloying elements are

present in the substrate, either in solution or as precipitated phases [115]. Simple aluminides on nickel

superalloys may contain up to 30 wt% Al although diffusion coatings containing hyper- and hypo-stoichi-

ometric Al compositions exist. Simple NiAl coatings are resistant against oxidation, but for conditions

where hot corrosion can occur, incorporation of Pt, Cr, and or Si are more resistant.

Degradation of aluminide coatings may be initiated by basic fluxing of Al2O3, and proceed by both

basic fluxing and sulfide oxidation [115]. If molybdenum or tungsten is present in the coating derived

from their presence in the substrate, acidic fluxing of the protective oxides can occur. One engine manu-

facturer has indicated that certain aluminide coatings perform well at temperatures in the Type 1 hot cor-

rosion range (800-950C) but may have poor resistance to Type II hot corrosion. Cracking of aluminide

coatings may also occur by creep if the service temperature is below the ductile-to-brittle transition tem-

perature [116].

16

Overlay coatings are applied with minimal direct contribution from the substrate alloy. Typical over-

lays have a composition of MCrAlX where M is Fe, Co, Ni, or a combination of these elements and X are

oxygen-reactive elements such as Zr, Hf, Si, and Y [115,117,118]. Overlay coatings are applied by several

processes including electron beam physical vapor deposition (EB-PVD) or by thermal or plasma spraying

where Typically these coatings initially contain cobalt or nickel, with 15-wt.%Cr, 12-wt.% Al, and 0.1-

0.9-wt.%Y. To improve its resistance to Type II hot corrosion, the cobalt was partially replaced with

nickel, the aluminum content was reduced somewhat, and the chromium content was increased from 15-

20-wt.% to almost 32-wt.%. Hafnium is used as a replacement for yttrium in some MCrAlX coatings,

where X (Y or Hf) promotes coating adhesion and oxidation resistance under thermal cycling conditions.

In recent years, MCrAlY coatings have been air plasma sprayed (APS) as a cost-savings procedure and

for reportedly better control of chemical composition. There are various procedures and methods by which

plasma spraying may be done, many of which are proprietary or patent-protected. The principal oxide in

the outer portion of a MCrAlY overlay coating is Al2O3. Yttrium is more prevalent in the outer scale and

promotes alumina formation for improved oxidation resistance and better scale adherence. Hafnium also

promotes oxidation resistance and like yttrium, prevents sulfur from segregating to the coating/substrate

interface, thus improving the overlay coating adhesion. Ni, Co, or both and chromium are found in the

inner half of the coating nearer to the coating/substrate interface. Increases in chromium content improve

both the hot corrosion resistance and the oxidation resistance of the overlay coating, but tend to make the

coating less ductile. If a single coating can operate successfully over a range of temperatures with differ-

ent forms of corrosion attack like type I and II hot corrosion and high temperature oxidation, the coating

essentially respond to local temperature in such a way that it will form either alumina or chromia protective

scale as appropriate. High purity alumina scales offer best protection against high temperature oxidation

and type I hot corrosion and chromia scales against type II hot corrosion.

The date, there is some limited use of thermal barrier coatings (TBCs), but as general ship engine temperatures rise, increased usage of TBCs is expected. TBCs often consist of a porous zirconia layer, ZrO2, that is stabilized in the metastable tetragonal phase by approximately 8 wt% yttria, Y2O3. The yttria-stabilized zirconia (YSZ) is applied using either plasma spraying or by EB-PVD over a bond coat which is usually MCrAlY (where M is Co and/or Ni) or a diffusion aluminide. The bond layer is typically 75-125 m and the TBC is 125-375 m thick. Application of the TBC system to hot section components in a gas turbine has been able to decrease metal surface temperature by about 150C. TBCs also smooth out hot spots thus reducing thermal fatigue stresses. A thermally grown oxide (TGO) layer is formed between the bond coat and the TBC during fabrication and tends to grow thicker upon high-temperature exposure of the TBC.

Electron beam physical vapor deposition (EB-PVD)-applied TBCs tend to produce a columnar struc-ture with fine ribbon-like voids and porosity within the columns aligned normal to the plane of the coating that results in superior compliance to residual stresses. Plasma sprayed TBCs have disk-like voids aligned parallel to the coating plane that promote superior insulating efficiency.

During service, coatings degrade at two fronts: the coating/gas-path interface and the coating/substrate

interface. Deterioration at the coating/gas path interface is the result of interaction of the coating materials

with the prevailing environment; a major contributor to coating degradation in marine gas turbine engines

is hot corrosion. However, on the second front at the coating/substrate interface, interdiffusion contributes

to coating degradation as a function of exposure time which is often less noticed, but is vitally important

to the service life of high temperature coatings.

Coatings and superalloy substrates generally have widely diverse compositions due to the functional

requirements of balancing mechanical properties and corrosion resistance. Solid-state diffusion at the

17

coating/substrate interface occurs. This interdiffusion causes compositional changes on both sides of this

interface as a function of time and temperature; this can compromise the substrate properties and deplete

the coating of critical species. Interdiffusion between the coating and the substrate [119] can modify the

oxidation and corrosion resistance of the coating and the mechanical properties of the coating-substrate

system. Interdiffusion slowly changes the chemical composition and subsequently, may alter the micro-

structure of the coating and/or the substrate. The presence of sulfide species in the coating may also lead

to diffusion via grain boundaries or other relatively easy diffusion pathways into the substrate and/or the

coating. This may alter the long-term performance of a coating/alloy system in a hot corrosion environ-

ment in a way that would not occur in an oxidation environment.

The fundamental importance of surfaces, interfaces, grain boundaries, coating structures and associ-ated structural defects, enrichment-segregation phenomena, and transport properties are essential towards understanding corrosion degradation mechanisms [120]. The phase transformations and surface reactions of the coating during hot corrosion are basically the same as those that take place during oxidation, except that the overall rates during hot corrosion are faster [121]; quality control is critical to ensure coatings are resistant to these rapid, and often severe, reactions.

It must be recognized that the performance of a high-temperature coating is dependent on the substrate and on the coating application process [122]. Identical coatings applied on different substrates do not behave in the same way. In addition, identical substrates may not behave equally because casting param-eters and structural segregation and processing may differ from one sample lot to another [122]. Coatings cannot be successfully developed without considering substrate effects and the variables in coating fabri-cation processes. Coatings applied by different techniques do not perform identically in terms of perfor-mance/lifetime because the process and heat treatment variables may vary in different lots or from differ-ent coating suppliers [122].

High-temperature coatings provide resistance to corrosion and degradation by the formation of a pro-

tective oxide film. Currently, alumina and chromia, and possibly silica have been the oxides of choice.

The nature of the oxide is dependent upon many variables, not all of which have been identified or under-

stood [122]. It has been stated that there is a need to develop improved physics-based models linking the

thermochemical and thermomechanical behavior of thermal barrier coatings to their combustion environ-

ments [123]. Bornstein acknowledged that there needs to be an increased understanding of the coatings

manufacturing methods, coating structure, corrosion mechanisms, and coating failure mechanisms [123].

The use of thermodynamically-based modeling and development of multi-component phase diagrams will

help in predicting possible harmful phases in the development of new high-temperature coatings.

A better understanding of all application and processing steps for overlay coating systems is required

and coating technology should move toward non-line-of-sight application processes to avoid weak loca-

tions where the coating is non-existent or inadequately thin. Potentially improved coating systems are

limited in their application by the scope of the present understanding and control of processing [122].

Coating compositions must be tailored for the deposition process (EB-PVD, plasma spray, electrodeposi-

tion, electrophoresis, etc.) and for the substrate. The basis for potentially detrimental substrate interactions

with coatings from interdiffusion must be thoroughly understood and considered in the early phases of

substrate and of coating development [122].

The effects of sulfur on coating adherence and the addition of reactive elements of coatings to improve

coating adherence to the substrate and combat the sulfur effect needs to be better understood before these

coating additions can be fully exploited in advanced alloys and coatings [121]. Alloy development must

be optimized for coatability, hot corrosion resistance, and thermomechanical fatigue resistance rather than

just for high-temperature creep [122]. Unknown tradeoffs still exist.

18

CONCLUSIONS

Selection of materials for marine gas turbines demand that such materials have a variety of properties

such as high strength, toughness, creep resistance, fatigue resistance, as well as resistance to degradation

by their interaction with the environment. Nickel-base and cobalt-base superalloys have been the materials

of choice in the hottest areas of a marine gas turbine engine. Nickel-base superalloys are an extraordinary

class of materials that have been crucial to the continued development of high performance turbine en-

gines. Major failure modes of gas turbine engine components are creep and fatigue. However, in marine

gas turbines, a major factor influencing these mechanical failure modes is corrosion. Hot corrosion is a

major drive degrading superalloys and causing accelerate failures by creep and/or fatigue by decreasing

the component thickness. Understanding high temperature corrosion and then using that knowledge to

design and select corrosion resistant materials or apply protective measures will permit performance, ca-

pability, and efficiency improvements.

Coatings are often used when the alloy chemistry and mechanical properties of the alloy far short of

engine performance goals. There are several factors that can influence the performance of high-tempera-

ture coatings in terms of their resistance to corrosion, spallation, and mechanical wear. The complexity

of the environment, where there can be an almost infinite mixture combinations of sulfur, carbon, oxygen,

halogens, water vapor, molten salts, and other contaminants which can interact with a specific coating can

make prediction of coating performance difficult. Coating performance prediction is further modified and

complicated by: (1) application process; (2) processing variables; (3) and substrate structure and compo-

sition. Different substrates may cause vastly difference performance of the same coating. Different coat-

ing application processes may also greatly alter the performance of the same coating and of the substrate

interactions with identical coatings.

The interactions of creep, fatigue, and corrosion need to be identified and understood in order to derive

solution to minimize turbine component failures. For optimal performance for marine gas turbine engines,

the selection of a nickel-base superalloy likely will need to be paired with a compatible coating that best

resists corrosion by the marine environment, is chemically and metallurgically stable, experiences mini-

mal interdiffusion with the alloy substrate, does not negatively impact creep or fatigue itself, and can be

reliably applied. Using the overall systems approach, the goal will be to create multiscale, multivariable

physics and chemical based models and tools to establish the fundamental science base for ultrahigh tem-

perature material systems that can survive shipboard environments over a wide range of temperature-

stress–environment–time variable fields that may be encountered.

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