Hot Corrosion

section epub:type=”chapter” role=”doc-chapter”>

Chapter 15
Hot Corrosion

15.1 Introduction

Alloys for high temperature application usually contain appreciable quantities of nickel: in particular, alloys based on the 80–20 nickel–chromium system have proved extremely successful because they can combine strength with good oxidation resistance at high temperatures. However, in gas turbine applications, where the products of combustion of the fuel used condition the atmospheres, oxidation resistance alone is insufficient, and the resistance of the alloys to carburization, sulfur attack, and corrosion by ash must also be considered.

Three main types of gas turbine may be listed: (i) aeroengine, (ii) industrial, and (iii) marine. Aeroengines burn clean fuel but ingest impurities in the air intake. This is aggravated in the marine gas turbines since the air intake includes sea salt. Industrial engines burn inferior fuel containing vanadium and sodium, but the air intake can be filtered. Turbine components are made of Ni‐ or Co‐based alloys. The blades and vanes are subject to cyclic temperature corrosion in alkali sulfate oxide ±chloride ±vanadic environment. Advanced turbines operate at high inlet temperatures (1350 °C) and the industrial turbines at 1150 °C. The blades are cooled to 850–900 °C. Catastrophic corrosion occurs at lower temperatures: 610–750 °C for Ni‐based alloys and 750–850 °C for Co‐based alloys. Coatings have been tested extensively in the field of gas turbines, and these include multicomponent alloy compositions and alloy‐ceramic formulas as well as thermal barrier coatings (TBCs). Superalloys form 50 wt% of gas turbine engines, the rest being about equal parts of Ti alloys, steels, and composites. Development trends in the superalloy systems tend toward directional solidification (DS), powder metallurgy (PM), and mechanical alloy processing (MAP).

In gas turbines used for aircraft applications, where highly refined fuels are used, corrosion, although still a problem, is usually secondary to stress considerations, but where turbines are used for marine and chemical plant or for production of cheap electrical power, it is often necessary for economic reasons to run on lower‐grade fuels, and these can introduce severe hot corrosion problems. Table 15.1 provides the relevance and importance of hot corrosion in various energy fields and the associated contaminants, which are responsible for the degradation of components during service.

Table 15.1 Relevance and importance of hot corrosion in energy systems

Energy systems Mixed oxidant reactions Hot corrosion
Batteries and fuel cells SO2, SO3, O2, H2, H2S, and H2O attack of alloys Fused halides and carbonate cells
Coal conversion and combustion systems CO, CO2, H2, H2S, and H2O
attack of alloys
Slag films, residue films, fly ash films in MHDa, salt film‐assisted coal combustion
Solar energy and energy storage Salt spillage or leakage from thermal storage tanks
Nuclear energy H–O–C in HTGR
Steam reactions with zircaloy
Fission product, salt condensation on cladding
Gas turbines SO2, SO3, and O2
attack of alloys
Na2SO4, NaVxOy, and
NaCl attack of alloys
Gas and oil recovery and magma energy H2, H2S, and H2O
attack of alloys


Thus, on studying corrosion under turbine conditions, the effects may be broadly divided into two categories, the first being the effect of the gaseous products of combustion and the second the effects of deposits. Sulfur present in the fuel and salt in the marine environment would lead to the formation of sodium sulfate. So it is sodium sulfate with or without sodium chloride (or nitrate) that has been most extensively investigated. However, internal sulfidation called green rot that occurs in a carburizing environment was reported by Hancock (1968) and Strochi et al. (1969), while black plague corrosion found on high‐strength nickel‐based turbine alloys was produced when a very dilute salt solution was injected into downstream air in minute quantities. Fuel quality does not cause “black plague,” which is primarily metallurgical in origin, and is regarded as an oxidation phenomenon as Belcher et al. (1967) report.

The object of the present chapter is to review some fundamental aspects of hot corrosion and the behavior of gas turbine materials in the presence of sodium sulfate (with or without sodium chloride), oxygen, and sulfur‐containing gases, which are the contaminants most likely to be commonly encountered in gas turbine environments.

15.2 Engine Description and Materials

The principle of the gas turbine is as follows. Air is drawn into the engine and compressed. Fuel is mixed with the compressed air and burned in a combustion chamber with the object of heating the gas. The hot compressed gas is then expanded through a turbine that extracts the energy from it; part of this energy is used to drive the compressor and the remainder is available for useful work. The maximum temperature in the cycle is the temperature of the gas entering the turbine, and this temperature is determined by the capability of the turbine to accept it. The hottest part of the turbine is the first stage, which consists of a ring of stationary airfoils, called the inlet nozzle guide vanes, and a ring of rotating blades called the first‐stage rotor blades or buckets.

The turbine inlet temperature in aircraft gas turbines rose progressively to a maximum of the order of 900 °C. Further increases in turbine inlet temperature were achieved as a result of the introduction of cooling of the vanes and blades using air drawn from the high‐pressure compressor.

The earliest alloys used for gas turbines, developed in the 1940s, were based on Nichrome (Ni–20%Cr) heater alloys and the Stellite (Co–10%Ni–35%Cr) supercharger alloys. The strength of the nickel‐based alloys was increased as a result of the introduction of the Ni3(Al,Ti) ordered phase, gamma prime. Increasing the volume fraction of gamma prime in the alloy required that the chromium content be reduced. No phase equivalent to gamma prime exists for cobalt‐based alloys, and the modern cobalt‐based superalloy has changed relatively little from its origins.

These early alloys owed their oxidation resistance to the formation of a protective layer of chromia, Cr2O3, on the metal surface. However, a number of the nickel‐based alloys with lower chromium contents also exhibited excellent oxidation resistance due to the formation of protective alumina scales. The components of the first stage of modern gas turbines are now generally coated with an alumina‐rich layer designed to develop an Al2O3 scale.

A major corrosion problem with the lower chromium nickel‐based gas turbine alloys was encountered in the early 1960s in aircraft that were being used in missions at low altitudes close to the sea. Under these conditions, a rapid corrosion of the blades and vanes was sometimes encountered, leading to severe destruction of the involved materials.

The increasing cost of fuel has resulted in a renewed interest in improving the efficiency of gas turbine systems. In most cases, this involves the modification of the properties of the existing alloys and the production of new materials. A partial listing of companies producing and supplying high temperature alloys is given below. The alloys referred to in this chapter that are produced by a specific company are given in parentheses. Each company produces a more extensive range of high temperature alloys, and a complete listing can be found at the web addresses provided:

15.3 Early Studies

15.3.1 General

It has been known for many years that at high temperatures in the presence of salt, heat‐resisting alloys undergo rapid corrosion. This problem was greatly intensified in the 1950s by development in jet engines and the application of nuclear energy. The extremely high temperatures (800 °C and above) encountered in jet engines and gas turbines required the development of new alloys with sufficient resistance to corrosion by the hot combustion gases as well as adequate mechanical properties at the high temperatures encountered. Alloys based on nickel and chromium are well known for their good resistance to oxidation at temperatures in the range of 800–1200 °C. With the demand for improvement in creep properties at elevated temperatures (≈900 °C), alloy compositions have been modified in order to attain these properties. This has been achieved by reducing the chromium content of the alloys, increasing the amounts of some of the elements already present, and introducing new elements. The alloy series – Inconel, Nimonic, Hastelloy, MELCO, and DISCO – are a result of this. Betteridge and Heslop (1974) and Sims and Hagel (1972) have discussed the Nimonic alloys in detail. Rather unexpectedly, these changes in composition have resulted in a susceptibility to a form of high temperature oxidation in which sulfur acts as a promoter and a propagator. This new phenomenon, most severe than the normal sulfur attack and generally referred to as hot corrosion (1967), has been encountered in gas turbines operating in marine environments. It is acknowledged that this form of attack is associated with the unfavorable operating environment.

Much research has been carried out on the chemical nature of the deposits and the mechanism of attack that results, but it is apparent from the diversity of opinions held by various workers in the field that the phenomenon was not adequately understood in 1970s (Jaffee and Stringer 1971; Stringer 1977; Seybolt et al. 1967).

However, there was general agreement that the corrosion proceeds in three steps:

  1. Deposition of a molten sulfur‐bearing slag on the surface of the superalloy engine components.
  2. Breakdown of the protective oxide film.
  3. Reaction between the sulfur from the deposit and the base alloy.

The mechanisms that have been postulated mainly differ in the explanation for the processes occurring in stages 2 and 3.

The slag that promotes sulfidation corrosion in a marine gas turbine is formed by the reaction of sulfur in the fuel and sodium from the sea salt ingested through the intake air (DeCrescente and Bornstein 1968). Under the oxidizing or de‐electronizing operating conditions, sulfur in the presence of a sodium salt such as sodium chloride (NaCl) is then capable of forming Na2SO4 via Na2O and SO2/SO3. The ability of NaCl to form sodium monoxide (Na2O) is well demonstrated by Quets and Dresher (1969). As it is known that sulfur dioxide is more effective in converting sodium chloride to sodium sulfate than sulfur trioxide, several workers (Birks et al. 2006) suggested that the transformation of NaCl to Na2SO4 involves the Hargreaves type of reaction, which is the basis of an old process for the manufacture of sodium sulfate (salt cake). The equations for the formation of Na2SO4 might then be




or combining these equations


As it has been shown by DeCrescente and Bornstein (1968), this reaction, besides being thermodynamically favorable in the temperature range of interest, is quite rapid. Sodium sulfate produced by this means, together with that already present in the sea spray (Schirmer and Quigg 1967), can condense as a slag on the hot engine parts. This concept is supported by the detection of sodium sulfate on parts corroded in service.

NaCl itself is almost never found in the deposits, due to its high dew point, but this does not rule out the possibility that impacting NaCl particles onto the alloy surface may play a significant part in the corrosion process. Therefore, as discussed by Condé (1972), NaCl may be one of the principal agents responsible for Na2SO4 hot corrosion.

15.3.2 Alloy–Na2SO4 Reaction

Reaction of a metal surface with its environment can occur only after the protective film is penetrated. In the case of the nickel‐based superalloys, this film is normally a mixture of Al–Ni–O and Cr–Ni–O phases composed mainly of chromium and/or aluminum oxides (Cr2O3 and Al2O3), and, in addition, a pure NiO type of oxide or NiCr2O4 spinel oxide may form (Bergman 1967). Whether Al2O3, Cr2O3, or NiCr2O4 is responsible for protection is the subject of much controversy (see above), but it is widely believed that the major contribution comes from Cr2O3 (Viswanathan 1968).

In the presence of a sodium sulfate slag, the oxide film is penetrated, bringing into direct contact the metal surface and the molten salt.

The mechanism of film penetration is uncertain, but several possibilities exist (Farrel et al. 1970). One explanation that has enjoyed wide acceptance suggests that chromium oxide (Cr2O3) is dissolved by reactions of the type


The events that occur after film destruction are not well understood. Earlier studies on this aspect have been mainly empirical. But, in the 1970s, Condé (1972) suggested that once the protective oxide film is penetrated, sodium sulfate may be reduced, or electronated, according to the following equations:






where R is an electronizing agent and M, in the case of complex superalloys, is composed of several elements including chromium. NaCl acts as the reducing agent.

Early in the 1950s, it was also proposed that Na2SO4 had to be chloride‐contaminated to promote rapid sulfidation attack. Waddams et al. (1969) proposed that Na2SO4 reacts with the base metal at cracks and pores where micro‐electronizing conditions may exist. Many other workers (Birks et al. 2006) report that neither the chloride ion nor externally introduced reducing conditions are prerequisites for hot corrosion. Bergman (1967) and Seybolt (1968) suggest that the depletion of chromium in surface zones through the formation of oxides and sulfides reduces the corrosion resistance of Cr‐depleted zones (nickel‐rich zones), thereby promoting gross oxidation and/or sulfidation. Quets and Dresher (1969) suggest that in order for hot corrosion to occur, there must be a simultaneous reaction between the nickel of the alloy and the oxide of one or more of the alloy ingredients, with sodium sulfate. The equations for the corrosion of a nickel–chromium alloy might then be



Bornstein and DeCrescente (1969, 1971) and Bornstein et al. (1972) show that the accelerated rates of oxidation associated with sulfidation attack are not related to preferential oxidation by sulfur in the alloy‐depleted zone and relate the rapid corrosion to the presence of oxide ions in the Na2SO4 melt.

The most reasonable models available in the seventies for the Na2SO4 hot corrosion of nickel‐based alloys were essentially those due to Goebel et al. (1973). In order to make progress toward discovery of the mechanism of hot corrosion, they constructed thermodynamic diagrams describing the stability and composition of pure Na2SO4 as well as the phases that may be formed during the exposure of nickel or aluminum to the sulfate environment. It was then necessary to consider reactions between oxides that are formed during the exposure of nickel or aluminum to the sulfate environment. It was then necessary to consider reactions between oxides that are formed on the surfaces of alloys and Na2SO4 and in addition reactions between the protective oxides and other components that are introduced into the modified Na2SO4 as a result of the oxide contaminant. Such studies indicated two types of hot corrosion.

The less severe type, referred to as Na2SO4‐induced accelerated oxidation, runs along the following lines:

  1. An oxide layer is formed on the alloy surface, the oxygen required for its formation coming from dissociation of Na2SO4 (see below).
  2. As a consequence of this, the sulfur activity at the oxide scale/Na2SO4 interface increases to a level at which metal sulfides can be formed in the alloy beneath the oxide scale, providing that sulfur can migrate through the oxide.
  3. Sulfide formation increases the oxide ion activity of the Na2SO4, eventually allowing the dissolution of the oxide as higher ions, for example, images in the case of aluminum.
  4. Accelerated oxidation of the alloy occurs until the oxidation activity of the melt is no longer capable of fluxing the oxide scale.

This model assumes that when Na2SO4 comes into contact with nickel, the equilibrium oxygen pressure in Na2SO4 is sufficient to oxidize Ni to NiO. The oxygen consumed because of NiO formation will displace the following dissociation equilibrium reaction existing in sodium sulfate to the right:


Thus, the Na2O and S activities rise at the NiO–Na2SO4 interface. This sulfur then migrates through the NiO layer to the Ni–NiO interface where, due to the lower oxygen activities, nickel sulfide phases can be formed. Removal of sulfur from the sulfate increases the Na2O activity to Na2SO4 in the vicinity of the alloy surface. In this Na2O‐rich layer, formation of nickelate (images) ions following the reaction


is likely to take place. Since the oxide ion activity of the unmodified Na2SO4 is not sufficient for nickelate formation, images decomposes as soon as it moves out into the Na2SO4. Consequently, a porous non‐protective nickel oxide scale is formed, and accelerated oxidation proceeds.


Figure 15.1 Model for the Na2SO4‐induced catastrophic oxidation of a Ni–31Al–Mo alloy. (a) Oxygen moves from the gas through Na2SO4 to the alloy surface where metal oxides are developed. (b) MoO3 reacts with Na2SO4 that decreases the oxide ion activity of Na2SO4. (c) The modified Na2SO4 layer reacts with the protective Al2O3 scale that results in its destruction. However, vaporization of MoO3 causes the Al2O3 to precipitate as a porous network at the Na2SO4/gas interface. (d) Rapid oxidation ensues and aluminum is preferentially removed from the alloy, and the alloy–scale interface becomes irregular. (e) The nickel‐rich islands developed because of preferential oxidation of aluminum are converted to NiO. The rapid oxidation is self‐sustaining because MoO3 is continually added to Na2SO4 by oxidation of molybdenum in the alloy (Goebel et al. 1973).

For the other type of hot corrosion, referred to as Na2SO4‐induced catastrophic oxidation, the following features are believed to be applicable:

  1. Oxide phases are developed on the alloy surface.
  2. Specific oxides (for example, MoO3, WO3, or V2O5) of the alloy scale form higher oxides by taking oxide ions from the Na2SO4, reducing its oxide ion content to a low enough level so that the protective oxide layer is thermodynamically unstable in the melt.
  3. Consequently, decomposition by ionic dissolution into the melt occurs, thereby allowing the salt to come into contact and react with the metal alloy substrate.
  4. Rapid oxidation ensues, being self‐sustained because the metal oxides responsible for this type of attack are continually formed in the Na2SO4 adjacent to the alloy surface.

Goebel et al. (1973) have schematically illustrated this model for a Ni–31Al–Mo alloy, as reproduced in Figure 15.1 where the abovementioned basic features are well documented.

In summary, the essential point to be emphasized is the crucial role of Na2O on the Na2SO4 hot corrosion of nickel‐based alloys: if the oxide ion activity of Na2SO4 increases to the point where oxide scales can partially dissolve in the melt, induced accelerated oxidation is to be expected; if the Na2O activity of the melt is so small that they are efficient dissociative fluxes for oxide scales, catastrophic oxidation occurs. In addition, the formation of sulfides beneath the oxide scales in some alloys can cause the specimens to swell, which results in increased oxidation (Birks et al. 2006).

15.3.3 Effect of Sodium Chloride

The previous discussion has been confined to molten sodium sulfate as a cause for metal destruction in gas turbines. Sodium chloride was shown to be a source of sodium for reaction with fuel sulfur to form Na2SO4. It is also possible that NaCl by itself could be responsible for corrosion by flue gas deposits. It is widely reported, as a result of laboratory studies, that mixed sulfate–chloride slags produce a higher corrosion rate than sulfate alone (El‐Dashan et al. 1972).

Although a condensate of Na2SO4 containing NaCl is unlikely to be formed above the dew point (the temperature at which salt species will condense on an inert solid surface) for NaCl (∼730 °C) under normal situations, its formation in the case of nonequilibrium conditions may occur (Farrel et al. 1970).

Corrosion by additive NaCl may be due in part to the formation of a low‐melting‐point mixture between these two sodium salts and also to the action of the chloride in causing a breakdown of the film. The Na2SO4–NaCl phase diagram illustrating the influence of NaCl on the melting characteristics of the mixture was constructed by Danek (1965). He concludes that since a molten slag is required for sulfidation, the melting point of the salt mixture and hence the NaCl content exerts a serious influence on turbine operating temperature and thus efficiency and fuel economy. Later, Halstead (1970) reviewed the general behavior of oxides in the presence of sodium chloride. This review emphasizes that in gas turbines the nickel oxide/NaCl reaction may be better described by a general equation:


Since the NiCl2 equilibrium pressure for the preceding reaction is significantly lower than the vapor pressure of NiCl2 at temperatures of about 900 °C, such a reaction could lead to a loss of nickel by vaporization. This loss would lead to an appreciably less protective oxide film than one produced directly by de‐electronation. Destruction of the protective chromium oxide film can also occur by the following reaction:


Figure 15.2 Potentiodynamic polarization curves for nickel in Na2SO4, Na2SO4–0.5% NaCl, Na2SO4–1% NaCl, and Na2SO4–5% NaCl at 900 °C in air.

The chloride may also have a catalytic effect on the following reaction:


to cause the internal sulfur penetration observed. In an attempt to develop a picture of the role of chloride in the corrosion caused by flue gases and their deposits, Cutler et al. (1971) carried out some studies and concluded that the chloride may be especially harmful in those cases in which the deposit environment fluctuates between electronizing and de‐electronizing conditions because in these situations the alloy is never capable of building up a protective layer. Other workers have considered that NaCl may possibly accelerate the Na2SO4‐induced hot corrosion, but its influence has not been examined. Sequeira and Hocking (1978a,b, 1981) studied the corrosion behavior of nickel and Nimonic 105 in molten Na2SO4, NaCl, and mixtures of these two salts, at 900 °C, in laboratory air and under O2 + SO2/SO3 atmospheres. The following conclusions were drawn from their electrochemical results, which were further supported by topochemical examinations of the corroded samples (Figure 15.2):

  1. Three different regions characterize the electrokinetic behavior of nickel electrodes in molten sodium sulfate at 900 °C in air. The first region corresponds to the onset of passivity and consists of simultaneous nickel dissolution and passivation, the second is related to the dissolution of passive nickel, and the third corresponds to the local oxidation of passive nickel and sulfate discharge. The passive film on Ni is mainly NiO. It should be noted that these three regions are only observed if the anodic polarization of Ni takes place immediately after the maximum negative potential is reached, i.e. at about five minutes after insertion of the Ni electrode into the melt. If the anodic polarization of Ni is recorded after 20‐60 minutes of its immersion into the melt, then auto‐passivation occurs, and the polarization curves are characteristic of a passive metal electrode (i.e. Ei=0 > Ep and only passive–transpassive behavior is observed).
  2. The E/I curves for nickel in Na2SO4–NaCl melts, in air, in the range of 0–80% NaCl, which are similar to the anodic polarization curve for Ni in pure Na2SO4, show that halide additions (especially those up to 25%) strongly affect the anodic behavior of Ni in the active region and have practically no effect on the passive region. The higher dissolution rates are represented by the equation

    which also suggested being a critical factor in the Ni passivation. Careful analysis of the transpassive part of the polarization curves also showed that the overpassive dissolution current is smaller in Cl melts than in images alone. This was further evidenced by particular experiments that have been concerned with determining the inhibiting role of NaCl on Ni in the passive–transpassive area.


    Figure 15.3 Diagram of a typical Ni section.


    Figure 15.4 Ni, 900 °C, Na2SO4–15% NaCl in air after potentiodynamic sweep AEI (×1200): close view of one end of a corrosion pit.


    Figure 15.5 Conditions as in Figure 15.4, X‐ray image (×1200): distribution of S.


    Figure 15.6 Conditions as in Figure 15.4, AEI (×300): complete view of corrosion pit.

    Micrography studies and electron probe microanalyses of potentiodynamic specimens polarized in Na2SO4–15%NaCl at 900 °C, in air, are shown in Figures 15.315.6. Uniform field and the presence of scratches indicating more dissolution of the surface corrosion products were observed; along the scratch lines, the growth of oxides could also be observed; moreover, there was heavy internal grain boundary attack, plus a few pits. Figure 15.3 shows a typical Ni section. No S, Cl, or Na was detected either in the voids or in the external NiO porous fragile layer (which is almost always green). Ni3S2 was identified (see Figure 15.3). NiO and Ni3S2 are, therefore, the main corrosion products. Absorbed electron images (AEI) of a corrosion pit are shown in Figures 15.415.6; a sulfur X‐ray image of the region shown in Figure 15.4 is shown in Figure 15.5. The Ni, S, and O contents of this region were determined. The oxygen content was below the detection limit of the electron probe. S and Ni countings were made using CdS and pure Ni as standards. The results showed that the material was not homogeneous, being higher in sulfur near the outer surface (Figure 15.3), so the measurements were split into two (outer and inner) before calculating the corrected values. The outer zone consists, therefore, of a network of Ni3S2, and the inner zone consists of a mixture of Ni and Ni3S2; it is probably this Ni–Ni3S2 liquid eutectic structure at 900 °C that undermines the exposed metal surface, making the corrosive process more extensive and, consequently, originating deep perforations.

  3. Further examination of the E/I curves for Ni in sulfate–chloride melts at 900 °C leads to the conclusion that the rate‐determining step of nickel dissolution in the melts is mainly the diffusion process into the melt.
  4. SO3 atmospheres destroy the passive capability found for Ni in Na2SO4–NaCl melts in air. The high rates of corrosion observed at high partial pressures of SO3 are shown to be mainly due to the effect of SO3 on the solubility of the corrosion products. More specifically, this corrosion stimulation is due to the SO3 role in promoting dissolution attack of species like NiS with the appearance of Ni2+ until these ions precipitate or the conjugate oxygen cathodic process (O2/O2−) takes place.
  5. Nimonic 105 does not exhibit the active–passive transition characteristic of nickel in sulfate media. In the presence of pure Na2SO4 at 900 °C, in air, two different regions characterize the E/I curve. The first region corresponds to the onset of passivity and dissolution of passive Nimonic; it is characterized by a small dissolution rate and is really an “active–passive” zone, as suggested by the physical appearance of the electrode surface polarized within its limits. The second region is a transpassive area and corresponds essentially to the sulfate discharge.
  6. Additions of NaCl, up to 15%, to the sulfate melt hardly affected the forward sweep for Nimonic 105 in pure Na2SO4, but severely affected the reverse one. This non‐repair capability of reforming a passive corrosion layer is due to the Al and Cr depletion at the Nimonic surface by NaCl attack, which is observed at high anode potential values and may be explained by reactions of the form

    Only gold members can continue reading. Log In or Register to continue

Aug 11, 2021 | Posted by in Fluid Flow and Transfer Proccesses | Comments Off on Hot Corrosion
Premium Wordpress Themes by UFO Themes