Functional Materials || Corrosion-Resistant Materials

12Corrosion-Resistant Materials

Vivekanand Kain

Corrosion Science Section, Materials Science Division, Bhabha AtomicResearch Centre, Trombay, Mumbai, Maharashtra, India

12.1 Introduction

Corrosion is termed as the ‘chemical or electrochemical reaction between a

material and its environment that leads to deterioration of the material and/or its

properties’ [1]. Each material has to experience an environment. It could be the

process fluid along with temperature, radiation and/or operating/residual stress.

A given material under stress would behave in a manner that is dictated by its

mechanical properties and may fail by ductile (overload) or brittle fracture includ-

ing fatigue fracture. Interaction with temperature may lead to additional degrada-

tion modes like creep, and interaction with radiation may lead to radiation damage

(including radiation hardening, radiation embrittlement, radiation-induced segrega-

tion). In this chapter, the scope will remain confined to the interaction of materials

with their process fluids, excluding the individual influence of radiation, tempera-

ture or stress. Also, it will be confined to metallic materials. It does not indicate

that non-metallic materials (e.g. composites, organic material, ceramic materials)

do not degrade because of interaction with environment.

The driving force for a material to corrode is its tendency to lower its (free)

energy. Materials exist in nature in the form of compounds (ores/minerals). This is

because they have lower energy in the form of compounds compared with the high-

er energy state if they were to exist in metallic form. Therefore, we have to provide

energy to the ores (lowest energy state) to extract metals (i.e. higher energy state).

Each of these metals tries to lower its energy by changing back to a compound

state (i.e. by corrosion). This is clear from the example of reactive metals (K, Mg,

Be, Al, Zn, and so on) for which we have to spend enormous amounts of energy to

extract them from their ore. These are the metals that corrode very quickly (i.e.

they have a very high tendency to lower their energy by forming compounds).

Corrosion has been traditionally classified as (a) uniform corrosion and (b)

localized corrosion. Uniform corrosion (or general corrosion) is the form of corro-

sion in which the anodic (and cathodic) reaction proceeds uniformly over the entire

exposed surface. Uniform corrosion leads to thinning of materials. Rusting of iron

is the most common example of uniform corrosion. Uniform corrosion rates of

most of the materials in a variety of environments are already established and are

Functional Materials. DOI: 10.1016/B978-0-12-385142-0.00012-X

© 2012 Elsevier Inc. All rights reserved.

http://dx.doi.org/10.1016/B978-0-12-385142-0.00012-X

available in corrosion handbooks [2�5]. On a tonnage basis, uniform corrosion

accounts for the maximum destruction of materials. However, it is also easily tack-

led in plant design by providing a corrosion allowance (with an additional safety

margin) that takes into account the uniform corrosion rate and design life of the

plant component.

When anodic reactions get concentrated at certain specific regions, they result in

localized corrosion [6�10]. Some examples of localized corrosion are galvanic cor-

rosion, crevice corrosion, pitting corrosion, intergranular corrosion, stress corrosion

cracking, selective leaching, erosion corrosion and so on. Oxidation of materials is

another form of corrosion that occurs on exposed surfaces (therefore close to gen-

eral corrosion) but need not occur at a uniform rate all the time. It is important to

glance through the mechanism of degradation through these forms of corrosion

[6�10] before looking at materials resistant to each form of corrosion.

Galvanic corrosion occurs when an alloy is electrically coupled to another alloy

in the same electrolyte. Upon coupling, the less corrosion-resistant material

becomes anodic and corrodes faster while the more corrosion-resistant material

becomes cathodic and is protected against corrosion (its corrosion rate is lowered).

Galvanic series in a given environment gives an indication of the relative resistance

of materials to corrosion. For example, steel is placed on the active (lower/anodic)

side of the galvanic series of materials in seawater compared with copper, which is

placed on the noble (upper/cathodic) side. Individually, when immersed in seawa-

ter, steel being more anodic corrodes faster than copper. Upon coupling in seawa-

ter, steel would corrode at an even faster rate and copper would get protected (its

corrosion rate would further come down). Galvanizing (coating Zn over steel) is a

classic example of corrosion protection using the cathodic nature of steel compared

with Zn.

Crevice corrosion occurs in openings too narrow to cause convective sweeps but

allowing ionic diffusion to take place. Crevice corrosion is associated with a small

volume of stagnant solution, caused by holes, gasket surfaces, lap joints, surface

(even corrosion product) deposits and crevices left by design—under bolt and rivet

heads. The progress of anodic reactions inside the crevice geometry takes place by

an autocatalytic process. This also leads to lowering of pH (down to 2) and concen-

tration of chloride (up to thousands of parts per million) inside the crevices.

Outside the crevices the pH may be neutral or alkaline and the chloride concentra-

tion may be as low as a few parts per million. However, crevice corrosion is associ-

ated with an incubation time. Stagnancy of solution (or its flow rate) and crevice

solution pH are important factors for its propagation. Increasing the flow rate of

the process solution, to dilute or wipe out the incipient crevice solution chemistry,

is a sure way to delay or avoid crevice corrosion damage. All materials, especially

those showing active�passive behaviour, are prone to crevice corrosion. The attack

takes the form of irregular surfaces within crevices, leading to shallow but widely

attacked regions.

Pitting corrosion propagates by a mechanism similar to crevice corrosion. It is

called the self-initiating form of crevice corrosion and occurs on open surfaces.

Materials showing passivity (stainless steels, nickel-based alloys, Al-based alloys,

508 Functional Materials

Ti and its alloys) are especially prone to pitting corrosion. Halides are the most

aggressive species that cause pitting corrosion. Pits preferentially initiate at the het-

erogeneities in the material (grain boundaries, precipitates, inclusions, segregation

location and so on) over which the passive film is weak or disturbed.

Selective leaching (also called parting/dealloying) takes place by selective

removal of one element from an alloy leaving an altered residual matrix. Removal

of zinc from brasses is a common example. The yellow colour of 70-30 brass turns

to red (copper) colour after selective removal of zinc from the surfaces. The resid-

ual matrix is also porous (hence weak) structure.

Intergranular corrosion (IGC) is the preferential corrosion of the grain bound-

aries with negligible corrosion of the grain matrix [11�15]. It may take the form of

selective corrosion at grain boundaries (grooving) to grain dropping (weight loss

and ultimately strength loss). The most common cause is due to depletion of an

alloying element at the grain boundary. Sensitization [5�15] is a term describing

depletion of chromium at grain boundaries to a level below 12 wt.% in stainless

steels. As the passivity at the surfaces over the chromium-depleted grain boundaries

is weak (compared with the passive film over the grain matrix where chromium

content is 12 wt.% or above), the passive film over the grain boundaries breaks eas-

ily when exposed to corrosive process fluids in applications. In stainless steels, it is

the formation of chromium-rich carbides that leads to the formation of chromium

depletion regions at grain boundaries. It is to be noted that in stainless steels them-

selves, one may encounter IGC due to segregation of certain alloying elements at

grain boundaries, for example segregation of Si or P in austenitic stainless steels

that leads to IGC of even annealed stainless steels [7,8] in highly oxidizing envir-

onments (boiling nitric acid containing hexavalent chromium ions). Similarly, non-

ferrous alloys also are prone to IGC [6,7,10]. Segregation of Zn in brasses, and Si

in aluminium alloys, also lead to IGC. Similarly, segregation of Fe impurity ele-

ments in Al alloys also leads to IGC.

Stress corrosion cracking (SCC) is a brittle failure at a low constant stress of an

alloy exposed to a corrosive environment. It requires a synergistic action of the cor-

rosive environment and tensile stress with the susceptible material. Classically, par-

ticular material�environment combinations are known to be prone to SCC [5�10].

Season cracking of brass (brass in ammoniacal solutions) and caustic cracking of

steels (steels in hot caustic solutions) have been known for decades to lead to SCC.

Tensile stresses and dislocation movement (therefore stresses in excess of yield

stress, even at localized regions in a material) are required for SCC to take place.

However, residual stresses (even lower than the yield stress, e.g. fit up or welding

stresses) are known to make a component undergo SCC although at localized

regions, for example at a discontinuity, the stresses will be high enough for disloca-

tion movement to occur.

The term ‘hydrogen damage’ refers [6�10], in general, to (a) hydrogen blister-

ing, (b) hydrogen embrittlement, (c) decarburization or hydrogen attack and (d)

hydrogen-induced SCC. Corrosion (cathodic) reaction releases nascent hydrogen on

the surface of the material. There may be other sources of atomic hydrogen, for

example welding or heat treatment in a moist and high-temperature environment.


A part of it may be adsorbed into the material. Molecular hydrogen does not diffuse

through metallic materials. Once nascent hydrogen enters materials in concentra-

tions higher than the solubility limit of hydrogen in that alloy, it may form volumi-

nous hydrides in certain alloys (e.g. zirconium alloys) and form hydride blisters. In

other alloys, the nascent hydrogen will diffuse through the thickness of the compo-

nent and may accumulate at discontinuities, for example void. At the discontinuity,

nascent hydrogen will combine to form molecular hydrogen. The equilibrium pres-

sure of molecular hydrogen with atomic hydrogen is so high that it can cause rup-

ture of any material. Hydrogen embrittlement is also caused by penetration of H

into a metal, which results in the loss of ductility and tensile strength, mainly

because of the weakening effect of hydrogen on metal�metal bond strength. In Zr

and Ti alloys, higher solubility of H at higher temperatures leads to higher H pick-up

and precipitation of brittle hydrides at lower temperatures. These brittle hydrides

lower the strength of the alloys. In other cases, atomic hydrogen in the material

reacts with carbon to form methane or with oxygen to form moisture. These high-

temperature processes lower the strength of the alloy (removing carbon) or cause

build-up of very high pressure (formation of moisture at high temperature). For caus-

ing SCC, atomic hydrogen (H) accumulates within the metal at the crack tip, leading

to localized weakening, either by void formation or by lowering the cohesive

strength. Cracks propagate by mechanical fracture of the weakened region (ahead of

the crack tip). Species like sulfur, which act as recombination poison for hydrogen,

tend to promote hydrogen damage.

Erosion corrosion [5�10] is the acceleration in rate of attack on a metal surface

owing to relative movement between the corrosive fluid and the metal surface.

Though all metals and alloys are prone to erosion corrosion, there are two distinct

mechanisms: alloys that depend on surface films for corrosion resistance undergo

erosion corrosion by a mechanism dominated by corrosion whereas soft alloys

undergo erosion mainly by mechanical damage. The erosion attack can be irregular

in geometry, ranging from corrosion pits to gullies. Inlet-end corrosion (for tubes)

and impingement attack at elbows/sharp bends are some common examples. Two-

phase process fluid results in severe erosion attack. Flow-accelerated corrosion

(FAC) is a subset of erosion corrosion in which the attack is mainly by electro-

chemical dissolution of the material surfaces [16�18]. FAC is dictated by the diffu-

sion rate of ionic species across the boundary layer and the solubility limit of the

dissolving ionic species in the process fluid. The velocity of the fluid affects both

the boundary layer and the solubility of dissolving ions reached in a given time in

the process fluid. Water chemistry parameters (pH, temperature and so on) also

affect the solubility limit, and other parameters like dissolved oxygen affect the

nature of surface oxides that form on the material.

Oxidation [5�10] is the reaction between a material and air/oxygen (in the

absence of water) or an aqueous medium. This is reflected in weight gain. The oxi-

dation process is treated as an electrochemical process that progresses with diffu-

sion of ions across the oxide layer. A broad indicator of oxidation corrosion is

taken to be the Pilling�Bedworth (PB) ratio, which is the ratio of the volume of

the oxide per gram atom of the metal to the volume of one gram atom of the metal.


A PB ratio close to one is taken as a sign of protective oxide. In addition to

meeting this criterion, the oxide should also have good adhesion, a high melting

point, low vapour pressure, high-temperature plasticity, low electrical conductivity

and low diffusion coefficient for metal ions/oxygen.

An understanding of the mechanism of the corrosion degradation process leads

us to develop and use materials that are resistant to corrosion degradation. This

brief description of the mechanisms of various forms of corrosion is intended to

make the reader broadly aware of these mechanisms. A detailed understanding of

the mechanism of corrosion degradation and the various terms used requires study

of the related references.

12.2 Materials Resistant to Uniform Corrosion

The uniform corrosion rate of most of the materials in commonly used environ-

ments is already established and is available in numerous handbooks. Iso-corrosion

charts [2�10] are a simple way to illustrate the regime of environmental variables

at which a given corrosion rate would be encountered. Depending on the applica-

tion, one can determine the corrosion rate for a candidate material from such data-

bases/iso-corrosion charts. At the design stage of components/plants, the uniform

corrosion rate of structural materials is considered in determining the thickness to

be used for fabrication. This is based on the corrosion rate and the design life of

the plant. In addition, a margin of safety is provided.

Materials resistant to uniform corrosion for specific applications are available.

Like steels are prone to uniform corrosion in atmospheric exposure conditions but

weathering steels (typical composition (all in wt.%): C, 0.10; Mn, 0.35; Si, 0.50;

Cu, 0.40; P, 0.10; Ni, 0.40 and Cr, 0.80) are generally resistant to atmospheric

exposures and usually do not need painting for corrosion protection [19]. These

weathering steels have a ferrite�pearlitic microstructure and derive strength mainly

from solid solution strengthening. These steels, containing Cu and P, do not require

painting but develop a ‘patina’ after a certain period of exposure to the atmosphere.

The Delhi iron pillar has not corroded for the past 16 centuries. It has been shown

[19,20] that the high slag content in the iron causes a high corrosion rate. This

leads to enrichment of phosphorus on the surface. The high P content catalyses for-

mation of the protective amorphous layer of δ-FeOOH. Conversion of FeOOH to

magnetite causes further protection from corrosion. Formation of iron phosphates

also help in lowering the corrosion rate.

When materials with known rates of corrosion have to be used for a long design

life (e.g. ship hulls and underground pipelines) generally a cathodic protection sys-

tem (impressed current system) is used [21] to bring down the corrosion rate

drastically.

Some materials like stainless steels are prone to uniform corrosion in solution-

annealed conditions but are prone to IGC in sensitized conditions in nitric acid

environments [22]. Therefore, one has to be careful in considering not only the


uniform corrosion rate for the solution-annealed plates to be used for fabrication

but also the IGC aspect for the weldments.

12.2.1 Additional Requirements from Corrosion-Resistant Materials

One has to consider that corrosion products can enter the process stream, which in

some cases might not be acceptable. For example, any contamination in the food

processing industry is unhealthy. Therefore, in such cases, one has to choose mate-

rials that do not release significant corrosion products in the process stream.

Materials with passivity (stainless steels, nickel-based alloys, Ti and its alloys, Al

and its alloys) are protected by the oxide film and exhibit very low rates of uniform

corrosion.

One example illustrates this point. In nuclear reactors (pressurized heavy-water

reactors), the carbon steel pipes are used in the primary heat transport circuit to

carry the hot water to (and back from) the steam generators. The water chemistry

is controlled and the uniform corrosion rate of the carbon steel is low in this sys-

tem. However, any crud (corrosion products) in the system would be activated

when passing over the nuclear fuel and would transport activity. To reduce this

problem of activity build up, a pre-commissioning procedure, known as hot condi-

tioning [23,24] is performed. In this procedure, the carbon steel surfaces are

exposed to a particular water chemistry (pH 10�10.4, dissolved oxygen , 5 ppb)

at temperatures around 250�C for a few days. This helps to create a surface oxide

(magnetite) that covers all the exposed surfaces of carbon steel and further brings

down the uniform corrosion rate.

Mechanical properties, weldability, fabricability and cost are some of the other

factors that are to be considered in selection of materials for a given application.

One should also ensure that it is the uniform corrosion that would be the degrada-

tion mechanism in a given application and not localized corrosion. One may select

stainless steels for an application against uniform corrosion but that ends up in

severe localized (pitting and crevice) corrosion if the process stream has chloride

ions.

Specific applications may necessitate other considerations. When selecting mate-

rials for use in radiation environments, one has to consider additional aspects of

activation and radiation-induced embrittlement. As stainless steels suffer from radi-

ation-induced segregation, which also leads to irradiation-assisted SCC in light-

water-cooled nuclear power reactors [25�26], the reactor pressure vessel is made

from a material that is much more resistant to irradiation damage. Low alloy steel

(SA 508 class 3) is typically used for reactor pressure vessel applications, and the

concentration of Cu, Ni and P is carefully controlled in this material to manage the

increase in ductile to brittle transition temperature due to irradiation from fast neu-

trons [27]. However, this material does not have good compatibility with the pri-

mary (high-temperature) water used in the reactors. Therefore, a lining of stainless

steel is laid over the inside surfaces of the reactor pressure vessel as stainless steel

has very good compatibility (and low uniform corrosion rate) with the reactor pri-

mary water at high temperature.


The material selection also has to take into account the economic implication of

the decision. Certain noble metals like gold and silver, and costly materials like

titanium would meet the process requirement for many applications. However, cost

and fabricability may stop their use in several applications.

12.3 Materials Resistant to Localized Corrosion

As previously mentioned, localized corrosion may take different forms: crevice,

pitting, selective leaching, IGC, SCC, erosion and oxidation corrosion. To identify

materials resistant to one of these forms of corrosion, one should understand the

mechanism of the specific form of corrosion first and then understand how the cor-

rosion resistance is developed.

12.3.1 Materials Resistant to Crevice and Pitting Corrosion

Materials that exhibit passivity are especially prone to pitting corrosion, whereas

crevice corrosion can affect materials that exhibit passivity as well as those that do

not. Therefore, a steel and a stainless steel may both undergo crevice corrosion but

stainless steel would be extremely prone to pitting corrosion. The resistance to pit-

ting corrosion improves when the passive film becomes stronger (more passive/

protective).

Alloying Additions to Resist Pitting Corrosion

In stainless steels, higher amounts of chromium and molybdenum in the material

make the passive film more protective [28�30]. Although a higher amount of chro-

mium gets incorporated in the surface oxide film to make it more protective, alloy-

ing with molybdenum does not result in its incorporation in the oxide film but

increases the chromium activity and repassivation kinetics. Molybdenum additions are

known [30] to increase resistance of stainless steels to pitting corrosion by one of the

following mechanisms: (a) the passive properties of the film improve, possibly by

enrichment of chromium in the film (Mo by itself does not get incorporated in the

film); (b) there is inhibition of pit growth by the adsorption of molybdate ions

(MoO42�), these molybdate ions get formed due to dissolution of the alloyed Mo dur-

ing the initial stages of pitting (but play no role in a rapidly propagating pit); and (c)

hexavalent Mo ions form complexes with negatively charged cation vacancies in the

passive film (Eq. (12.1)), at the interface of metal-oxide

Mo61 1 vV1nM -ðMoUvV1n

M Þð62 vnÞ1 ð12:1Þ

The complex formation not only decreases the concentration of free cation

vacancies in the passive film but also reduces the diffusivity of these species in the

film [6]. Therefore, it requires higher applied potential to cause breakdown of pas-

sive films and increase in induction time for breakdown.


It should be noted [28�30] that the beneficial effects of Mo addition decrease

with increasing temperature. The alloyed Mo is ineffective to inhibit pitting by bro-

mides. Also, there are variable effects of Mo alloying against pitting corrosion in

austenitic and ferritic stainless steels.

Alloying with nitrogen is also found to be beneficial against pitting corrosion.

Nitrogen is known [29,31,32] to (a) form ammonium ions in incipient pits that

raise the pH above the threshold level for pit propagation, (b) form nitrates that are

incorporated in the oxide film and are enriched under a passive film (such an

enrichment hinders rapid dissolution of the substrate after destruction of the passive

film), and (c) segregate and block the actively dissolving sites on metal surface.

There are several pitting resistance equivalent (PRE/PREN) parameters used [33]

for ferritic stainless steels (Eq. (12.2)), for austenitic stainless steels (Eq. (12.3))

and for the duplex stainless steel (Eq. (12.4)); these are given below:

PRE5%Cr1 3:3%Mo ð12:2ÞPREN5%Cr1 3:3%Mo1 30%N ð12:3Þ

or

PREN5%Cr1 3:3%Mo1 16%N ð12:4Þ

or that taking into account the synergy of Mo and N

PREN5%Cr1 3:3%Mo1 51%N1 6%MoUN� 1:6%ðNÞ2 ð12:5Þ

PRE/PREN is taken as a measure of resistance to pitting corrosion of stainless

steels. The maximum PREN value (as in Eq. (12.3)) of SS 304L is around 21, for

SS 316L it is around 31 and for SS 316LN around 35. A PREN value of over 40 is

taken as indicative of very high resistance against pitting corrosion. Generally, it is

shown the higher the PRE/PREN value, the higher is the resistance to pitting (as

reflected by pitting potential/critical pitting temperature). Typically, austenitic

stainless steels with PREN values higher than 40 are categorized as super-austenitic

stainless steels.Superferritic stainless steels typically have PRE values lower than 40. A typical

example is S44660, with a PREN value of 33.3�38.6. Some grades like S44635

with high Cr (24.5�26%) and 3.5�4.5% Mo and 3.5�4.5% Ni and Ti and Nb

additions to fix interstitials are known to have good resistance to pitting in seawater

applications. However, even this grade of superferritic stainless steel is prone to

crevice corrosion. Superferritic stainless steels have also reportedly suffered hydro-

gen embrittlement by cathodic protection systems used in seawater condensers.

Super-austenitic grades of stainless steels have PREN values over 40 and contain

higher Mo (6.0�6.5%) for 20% Cr to attain high resistance to pitting and crevice

corrosion. Addition of nitrogen in these stainless steels has also contributed

to improvement of resistance to localized corrosion. S31254 is a typical example,

with a PREN value of 44.7�48.6. Still, crevice corrosion resistance is not very

high and care should be taken with allowable crevice gaps. Compared with super-

austenitic grades, superduplex stainless steels have very good resistance against


pitting and crevice corrosion. These stainless steels are appropriate for use in

seawater-cooled condenser tubing applications. However, welding-related problems, high

operation temperature above 250�C or temperatures below �50�C remain a concern.

Another parameter (Eq. (12.6)) proposed to reflect the resistance of alloys to pit-

ting corrosion is the measure of alloying to resistance against corrosion (MARC).

This parameter [34] has been found to be useful in evaluating resistance to both

pitting and crevice corrosion.

MARC5%Cr1 3:3%Mo1 20%N1 20%C� 0:5%Mn� 0:25%Ni ð12:6Þ

Nickel-based alloys are inherently more resistant to pitting/crevice corrosion

than stainless steels. In fact, some nickel-based alloys (alloy 33/alloy 22) are

known not to pit even when exposed to a boiling solution of 25% NaCl. The criti-

cal pitting temperature, as determined by the ASTM G48/G150 test, is another indi-

cator of the susceptibility of materials to pitting corrosion. Higher values of the

critical pitting temperature indicate lower susceptibility to pitting corrosion. The

PREN/MARC and critical pitting temperature values of typical alloys have been

compared in some studies [34] and are shown in Figure 12.1 [8]. Figure 12.2 shows

the pitting and passivation behaviour of some commonly used austenitic, duplex

stainless steels and nickel-based alloys. Figures 12.1 and 12.2 clearly bring out the

beneficial effect of nitrogen addition in these alloys (e.g. 316L and alloy 33, with

250

50

Crit

ical

tem

pera

ture

(°C

)

100

Pitting corrosionCrevice corrosion

Alloy 33 without N

Alloy690

316L316LN

2205

317LMN 904L

1.4529

Alloy 625

Alloy 33

1.4565 S

30 35 40PREN = % Cr + 3.3%Mo + 30% N

45 50 55

Figure 12.1 Effect of alloy composition (PREN) on critical pitting and critical crevice

temperatures as determined by G48 test of ASTM in 6% FeCl3 for various stainless steels

and nickel-based alloys (8, reproduced with permission from John Wiley and Sons Ltd.),

with addition of data points for alloy 690, alloy 33 (with and without N) and 316LN.

Alloy 1.4565 S is Fe�24Cr�17Ni�4.5Mo�0.5N and alloy 1.4529 is

Fe�20Cr�25Ni�6Mo�1Cu�0.2N.


and without N addition) against pitting corrosion. MARC is an improvement over

the PRE/PREN parameter as it takes into account the contribution of Ni (and Mn),

which was left out from the previously mentioned parameters. Figure 12.1 shows

that using the PREN parameter, nickel-based alloys do not fit the otherwise clear

trend of correlation between pitting resistance and chemical-composition-based

parameters. Thus PRE/PREN may be used to compare austenitic stainless steels,

and MARC may be used to compare all austenitic alloys including austenitic stain-

less steels and nickel-based alloys.

Resistance to Pitting in Acids

Hydrochloric acid is an oxidizing acid and highly aggressive as it has low pH and

high levels of chloride. Small impurities of iron would be present as ferric (Fe31)

and make it even more aggressive environment that can cause pitting. In fact, for

hot hydrochloric acid containing ferric ions as impurity, there are no

suitable materials for use. Tantalum may be the only material used in such cases

but it is highly expensive. Other nickel-based alloys are also suitable under certain

concentration ranges. Hastelloy B is a chromium-free alloy (28Mo, 1Fe, 0.02C and

69Ni) that is good in all concentration ranges up to its boiling temperature. Even it

1E–8

–400

–200

0

200

Pot

entia

l (m

V(S

CE

))

400

600

800

1000

1200

–400

–200

0

200

400

600

800

1000

1200

1E–7

1N HCIAlloy 33

Alloy360

316L

316LN

Expt.Alloy

DuplexSS

1E–6 1E–5 1E–4 1E–3

Current density (A/cm2)

49.67

38.85

34

30.68

26.1

31.74

Figure 12.2 Effect of alloy composition (PREN) on pitting potential at room temperature in

deaerated HCl solution for various stainless steels and nickel-based alloys. The experimental

alloy is alloy 33 without addition of N. The numbers mentioned next to the alloys are the

PREN (as in Eq. (12.3)) values.


is attacked if aeration or strong oxidizing ions are present. Hastelloy C (65Ni,

16Cr, 115Mo, 2Fe and 0.02C) is good in moderate ranges of hydrochloric acid and

temperatures, even in the presence of oxidizing impurity ions.

Hydrofluoric acid is another acid that is capable of causing pitting corrosion.

At room temperature, magnesium shows very high resistance to it (if the concentra-

tion is above 5%) as it readily forms a fluoride film that protects it from corrosion

attack. However, rapid corrosion attack occurs if the concentration is below 1%.

Wrought Monel (70Ni, 30Cu) is an outstanding material for all concentrations and

temperatures. However, aeration and presence of oxidizing ions increase the corro-

sion rate. Similarly, Pb is a good material for storage of hydrofluoric acid of con-

centrations below 60%. Stronger acids, particularly anhydrous acids, attack lead.

Steel, on the other hand, is suitable for use in concentration ranges of 60�100%.

As concentration falls below 60%, corrosion rate increases rapidly. Anhydrous HF

is not particularly corrosive. Monels have been used in critical applications, for

example valve seats, valve trims and pump shafts. Cast irons suffer from graphiti-

zation and are therefore not used for handling anhydrous HF. Above a temperature

of 150�C, carbon steels are preferred to Monels or cupronickels. At higher tempera-

tures, stainless steels are rapidly attacked. Fluorine gas (anhydrous) is non-corro-

sive and even welded carbon steels are used for handling it. However, moist

fluorine gas (and its aqueous solutions) is extremely corrosive to all materials

except gold and platinum. It causes rapid pitting of stainless steels. Nickel, Monel

and aluminium are good for handling moist fluorine gas (and its aqueous solutions)

up to a temperature of 400�C.

Resistance to Crevice Corrosion

The resistance to crevice corrosion is indicated [7�10,35] by a parameter

Epit2Eprot. This is derived from the electrochemical polarization experiment in a

given solution. After start of pitting in the test (indicated by pitting potential, Epit,

with increase in current density from the low values in the passive regime), the cur-

rent density is allowed to increase up to a fixed value (typically 53 1023 A/cm2).

At this current density, the potential is reversed. Eprot is the potential where the cur-

rent density again reaches the current density corresponding to passivity. Above

Eprot, the pits do not initiate but an initiated pit keeps growing. A higher value of

Epit�Eprot is indicative of higher susceptibility of the material to crevice corrosion.

Another parameter indicating susceptibility to crevice corrosion is the critical crev-

ice temperature (i.e. the minimum temperature at which crevice corrosion starts to

occur in standard ASTM test G48). Alloy 33 is also highly resistant to crevice cor-

rosion. The critical crevice temperatures, as determined in ASTM G48/G150 tests,

are compared in some studies [34] and shown in Figure 12.1.

It is clear from Figures 12.1 and 12.2 that most of the stainless steels

become prone to crevice corrosion and pitting corrosion at room temperature or up

to 50�C in chloride environments [32�35]. In seawater applications, most of these

stainless steels would develop pitting (and crevice corrosion if crevice geometries

are present or if these alloys develop deposits and under-deposit corrosion starts).


One needs to use alloys with a PRE/PREN higher than 40 to use its resistance

in chloride environments against pitting corrosion. However, the presence of

inclusions may lead to pitting irrespective of the resistant base matrix. It has been

shown [35] that in a ferritic stainless steel with more than 26% Cr, inclusions can

induce crevices by a preferential dissolution at the matrix/inclusion interface if

they are less than submicrometre in size. If there is no crevice formation, the pitting

tendency is determined by chemical stability of the inclusions; hence, inclusions

with a less stable passive film (e.g. Al2O3) reduce the initial pitting corrosion

resistance.

Other Non-Ferrous Alloys

Apart from stainless steels and nickel-based alloys, there are other alloys that are

resistant to pitting and crevice corrosion. Copper and its alloys (especially cupro-

nickel) have been traditionally used in seawater applications [36,37] primarily for

their resistance to localized corrosion. The protective films of cuprous oxides and

hydroxychlorides grow slowly on the exposed surfaces. Dissolved oxygen aids

in formation of the protective films. However, it increases corrosion by functioning

as a cathodic depolarizer and by oxidizing cuprous to cupric ions. Crevice corro-

sion may also lead to pitting corrosion in the affected regions (even under-deposit

corrosion regions). Carbon contamination (from organic contamination in water

or leftover lubricant used during fabrication) on the cupronickel tubes is also

known to have resulted in pitting corrosion due to galvanic action. Sulfide pitting

in polluted seawaters is a common problem for 90-10 cupronickel. Otherwise,

90-10 cupronickel tubes in cold worked condition (for improved strength) are used

in those seawater-cooled applications where these are compatible with the process-

side environment.

Aluminium is another alloy that is especially prone to pitting and crevice corro-

sion. A pitting rate index (PRI, Eq. (12.7)) is used to predict the number of weeks

to achieve a maximum pit depth of 1 mm in aluminium and its alloys [38].

log PRI5 2:5� 0:28 log ðsulfateÞ1 0:18 log ðchlorideÞ

2 0:20 log ðpH27Þ2 3 100� �� 0:42 log

30; 000

R

� �

2 0:064 log ðcopper3 103Þ

ð12:7Þ

where R5 1/C, C is the conductivity in μS/cm and a pitting rate index of less than

25 weeks indicates aggressive water.

Water with conductivity approximately equal to 1 μS/cm usually shows a PRI of

over 10 years, and water with conductivity 200 μS/cm shows a PRI of a few

months. It was shown in separate studies [39] that aluminium 1100 (spent-fuel

clad) stored in a storage pool with water of specific conductivity of 2.5 μS/cmdeveloped pits of up to 0.8 mm depth and up to 1.1 mm width in areas near the

crevice regions. Pitting was observed in 550 days of exposure at ambient


temperature. Aluminium alloys (6067, 6061) were found to be much more resistant

to pitting corrosion. Galvanic contact with stainless steels/dissimilar aluminium

alloys was found to accelerate the crevice and pitting corrosion [40].

Titanium and zirconium alloys are also used in applications requiring resistance

to pitting and crevice corrosion. Titanium is by far the most common material for

seawater-cooled condenser tubing applications. Zirconium, being an expensive

material, is used when severe process corrosion conditions require its use. Grades

2, 7 and 12 of titanium (typical compositions given in Table 12.1) are commonly

used [41,41A,41B,42,42A] for heat-exchanger tubing applications.

Grade 2 titanium resists seawater up to 120�C; grades 7 and 12 can be used up

to 260�C in many heat-exchanger applications. Marine fouling is a possibility with

these alloys but concurrent pitting is not an issue. As titanium is inherently resistant

to wet chlorine and hypochlorites, there are no potential corrosion problems

encountered with chlorination to kill marine organisms. Crevice corrosion is also

not a problem unless process-side temperature is unusually high. Titanium in sea-

water and chloride service becomes prone to crevice and pitting corrosion under

specific concentrations of chloride ions and temperature, as shown in Figure 12.3

[42A]. Zirconium has the same resistance to seawater as titanium but is much more

prone to hydrogen pick-up and hydriding embrittlement [43]. Zirconium is more

resistant to crevice corrosion at temperatures above boiling point (e.g. 101�C) andis less affected by changes in pH and temperature. It is, however, far less resistant

to highly oxidizing conditions, such as might arise from chlorination procedures.

Zirconium is less resistant than titanium grade 12 in the event of salt-plugging

(crevice formation) of condenser tubes. Oxidizing contamination of ferric chloride

or hypochlorites may cause pitting and IGC of the weld heat-affected zone or even

SCC in zirconium and its alloys.

Table 12.1 Chemical Composition of Ti Alloys Commonly Used in Industries for

Condenser and Heat-Exchanger Tubing Applications [41A�42A]

Element

(wt.%)

Ti�50A

(ASTM Grade 2)

Ti�0.2 Pd

(ASTM Grade 7)

Ti�Code 12

(ASTM Grade 12)

N (max) 0.03 0.03 0.03

C (max) 0.10 0.10 0.08

H (max) 0.015 0.015 0.015

Fe (max) 0.30 0.30 0.30

O (max) 0.25 0.25 0.25

Al � � �V � � �Pd � 0.12�0.25 �Mo � � 0.2�0.4

Ni � � 0.6�0.9

Ti Remainder Remainder Remainder


12.3.2 Materials Resistant to Selective Leaching

Selective leaching, also known as ‘parting’ or ‘dealloying’, is the selective removal

of one element from an alloy leaving an altered residual structure [6,7,10].

Selective removal of Zn from brasses is one of the most common examples. The

70-30 brass, with its yellow colour, loses Zn selectively and its colour turns to red

(the colour of copper). Graphitization of cast iron in soil is another common exam-

ple where the iron matrix selectively corrodes, leaving behind a network of graph-

ite flakes (a porous and weak structure). Selective removal of Al from Al bronze in

HF/acidic medium with chloride, Si from silicon bronze, and Co from Co�W�Cr

alloy are other examples of selective leaching. Although selective leaching could

be layer type (high Zn brasses in acidic medium) or plug type (low Zn brasses in

neutral/alkaline medium), both lead to loss of strength but no apparent change in

dimensions (porous structures). Although it is necessary to overcome this problem

by switching over to cupronickels (70-90 Cu, 30-10 Ni) from brasses in severe cases,

it may be sufficient in many cases to use lower Zn brasses (red brass, 15 wt.% Zn

rarely shows failures due to selective leaching), or reduce the aggressiveness of the

environment by de-aeration, or by cathodic protection. In 70-30 brass, addition of 1%

Sn and small amounts of As, Sb or P protects it from selective leaching. Addition of

2% Al to brass also helps against selective leaching, and it is commonly used as a

material for tubing in seawater-cooled condensers.

In many instances, selective leaching may be limited to a thin layer on the sur-

face of the material only and may not affect performance. Removal of Ni and Mn

from surfaces of stainless steels in molten sodium in nuclear fast-breeder reactors

[44] and selective removal of chromium [45] from alloy 690 in molten glass

50

80

Met

al te

mpe

ratu

re (

°C)

130

180

240

Pitting and crevice corrosion

Crevice corrosion possible

Complete immunity

Sea

wat

er

pH 9

pH 3.5

pHIn

crea

sing

10 15 20Concentration, % NaCl

Figure 12.3 Susceptibility of

Ti to crevice and pitting

corrosion as a function of

temperature and chloride ion

concentration [42A].

Reproduced with permission

from The Institute of Materials/

Maney Publishing.


(vitrification of nuclear waste) are examples where depths of only a few tens

of micrometres from the surfaces are affected. In other industrial applications,

selective removal of Cr and Fe from nickel-based alloys in molten salt baths and

selective oxidation of Cr in low oxygen atmosphere at high temperature (leading to

Cr depletion on SS 430 trim on automobiles during bright annealing) are some of

the common examples [6].

12.3.3 Materials Resistant to IGC

Austenitic Stainless Steels

Stainless steels and nickel-based alloys are the categories of alloy that show sus-

ceptibility to IGC. Among austenitic stainless steels, in general a higher amount of

chromium, lower nickel and carbon content improves its resistance to sensitization

and IGC. The main reason for sensitization is the chromium depletion arising from

precipitation of chromium carbides at grain boundaries. A higher chromium content

does not allow (or makes it difficult to allow attainment of) depletion of chromium

to levels below 12 wt.%. Nickel reduces the solubility of carbon in the alloy matrix,

therefore lower nickel content would benefit against IGC. However, reducing

nickel and carbon contents makes the alloys depleted of elements that are austenite

stabilizers. Therefore, one cannot reduce these two elements drastically. On this

principle, an IGC-resistant alloy SS 304L (nitric acid grade (NAG)) that is actually

a subset of SS 304L [22,46,47] has been developed. The composition of these two

alloys is given in Table 12.2. It is clear that stricter control over Cr, Ni and C

allows control over IGC. This alloy is now commonly used in nuclear spent-fuel

reprocessing and waste management applications, and in plants that handle boiling

nitric acid against IGC.

A chemical-composition-based parameter, Creffective (5%Cr�0.18 (%Ni)�100

(%C)), has been correlated [48,49] to the susceptibility to IGC in nitric acid envir-

onments. SS 304L, which has Creffective greater than 14.0, is shown to be resistant

to IGC in practice C test of ASTM A262. This correlation was arrived at after eval-

uating the IGC resistance of 20 heats of SS 304L after a heat treatment at 675�Cfor 1 h (which simulates the worst microstructure produced in welding) in a nitric

acid environment (as per practice C, A262, ASTM). Therefore, one can quickly

Table 12.2 Typical Chemical Composition of SS 304L and 310L NAG Stainless

Steels (All in wt.%)

Grade of SS Cr Ni C Si Mn S P Others

Commercial Purity

304L

18�20 8�12 ,0.03 ,1.0 ,2.0 ,0.030 ,0.045

304L NAG 18�20 10�11 ,0.02 ,0.40 ,1.8 ,0.010 ,0.025 N,0.07

O,0.01

310L NAG 24�26 19�22 ,0.02 ,0.15 ,1.8 ,0.005 ,0.020


assess the resistance to IGC in nitric acid environments from the chemical

composition of SS 304L in its welded condition. However, it should also be noted

that ‘active’ inclusions (mainly sulfides/oxysulfides of Mn and sometimes of Mn

and Fe) accelerate IGC of stainless steels in oxidizing acids like nitric acid

[45,50,51]. This attack proceeds after the elongated (stringer type) inclusions are

dissolved in the strongly oxidizing acid, and the exposed stainless steel surfaces

beneath the inclusions dissolve at a fast rate. The corrosion rate is preferentially

along the grain boundaries as these are high-energy regions. This leads to change-

over of the initial attack along the elongated stringers of inclusions (from the

exposed end faces, for example cross-sectional surfaces of tubular products—tubes,

pipes and bars) to intergranular attack. This intergranular attack proceeds at a fast

rate even in a solution-annealed stainless steel. This attack is termed end grain cor-

rosion. This Creffective criterion is not applicable to stainless steels containing

‘active’ stringer types of inclusion. It should also be noted that nitric acids contain-

ing oxidizing ions (Cr61 /Fe31 and so on) are highly aggressive and when the con-

centration of these ions reaches a value that takes the operating potential to the

transpassive region of the stainless steels, severe IGC occurs, even in solution-

annealed stainless steels.

SS 316L is more resistant to IGC than SS 304L. This effect is attributed to the

presence of Mo in SS 316L. Mo binds carbon atoms and slows down their participa-

tion in the chromium carbide reaction [11�15]. Mo also improves the surface oxide

protection properties, necessitating a deeper chromium depletion to be present in a

given environment to cause IGC. Stainless steels with higher Cr content (e.g. SS

310L) are more resistant to IGC as these require more depletion of Cr to cause the

minimum Cr levels to fall below 12 wt.%. It is difficult to introduce delta ferrite in

the weld pool of SS 310L as it contains a high amount of nickel and chromium. It is,

therefore, important to control minor elements like silicon, phosphorus and sulfur to

as low levels as possible (see Table 12.2). With these precautions, SS 310L has been

used over the past five decades in welded condition without difficulties [52].

IGC is also caused in certain environments by segregation of specific elements

to grain boundaries (even though no Cr depletion may be present). Si and P segre-

gation causes susceptibility to IGC in annealed stainless steels when used in highly

oxidizing environments like boiling nitric acid with hexavalent Cr [8,11,53]. In

such cases [53], reducing the amount of P (below 0.01 wt.%) and Si (below

0.25 wt.%) imparts IGC resistance. A variant of SS 316L made by the containerless

process in Japan was able to reduce P content to 2 ppm from the typically 250 ppm

levels in commercial grades of stainless steel. It was shown [53] that the corrosion

rate of this material was 100 times lower than that of a corresponding commercial

grade S 316L in nitric acid tests as per practice C, A262, ASTM. Because IGC sus-

ceptibility shows a maximum as the Si content is increased, increasing Si content

above 4�5% also improves the IGC resistance of stainless steels. However, weld-

ability aspects need to be considered in such cases.

Many alternative alloys are considered for resistance against IGC in nitric acid

environments. Titanium-based alloys (especially Ti�Ta�Nb alloy) show excellent

resistance to IGC and exhibit low general corrosion rate in nitric acid [54,55]


medium (especially at room temperature). However, recently it has been reported

that these alloys too are prone to corrosion attack at the liquid�vapour interface in

nitric acid service. Above 90% nitric acid, there are reports [56] of SCC.

Zirconium-based alloys are reported to be highly resistant to corrosion in nitric

acid. However, these alloys are prone to heavy corrosion if even traces (ppm) of

fluoride impurity are present in the nitric acid [57]. These are also reported to be

prone to SCC in nitric acid service when the acid concentration is above 20%.

Aluminium and its alloys are considered highly compatible for fuming (86�100%)

nitric acid at ambient temperature [58]. However, dilution of the acid causes

increase in corrosion rates.

Grain Size

A stainless steel that has a very small grain size will have a large grain boundary

surface area over which the chromium carbide has to precipitate. For a given

amount of precipitation of carbide, a smaller grain size would be better if it is not

able to cover the grain boundary surfaces fully. However, if the amount of carbide

precipitating is large (e.g. high carbon alloy), the weight loss due to grain dropping

would start to increase early in the service life of the component. For a large-

grained material, the weight loss due to grain dropping would start later, when the

full depth of the grain is gradually corroded by IGC of the large grain surfaces.

Once started, the increase in weight loss would be fast [3,8]. Therefore, grain size

in the medium range is preferred for low carbon materials.

Stabilized Grades

Stabilization of the carbon with stabilizing elements (e.g. Ti and Nb) leads to use

of SS 321 and SS 347, respectively. Because these elements have higher affinity

with carbon to form carbides than chromium, all the carbon is fixed with these ele-

ments. This precludes formation of chromium carbides, hence resistance against

sensitization and IGC is imparted. However, controlled welding practices have to

be used to avoid knife line attack at regions immediately next to the weld pool

when such welded alloys are used in a service environment that is known to cause

IGC. A Ti/C ratio of over 8 or an Nb/C ratio of over 12 have been shown [59] to

lead to production of ferrite (upon solution annealing) and its subsequent transfor-

mation to brittle sigma phase. Excess Ti/Nb is known to lead to a higher tendency

for formation of delta ferrite [59] that easily transforms to sigma phase (a hard and

brittle phase). Therefore, recent versions of these alloys have specified Ti/C (or

Nb/C) ratios lower than the values mentioned above.

Ferritic Stainless Steels

The ferritic stainless steels are more prone to IGC than austenitic stainless steels.

The solubility of interstitials in the ferritic matrix is much lower than in the austen-

itic matrix. Therefore, the amount of carbon1 nitrogen should be of the order of

0.01 wt.% to resist IGC. In fact, carbon solubility is high in the temperature range


around 925�C (where austenite starts forming) and heating to these temperature

ranges and water quenching or cooling (either air cooling or furnace cooling)

causes rapid sensitization. The diffusion rate of chromium in the ferritic matrix is

very fast. Therefore, either annealing at temperatures where the chromium diffusion

rates are high (e.g. 800�C) or cooling very slowly through the temperature range

900�700�C leads to rediffusion of chromium back to the chromium-depleted

regions and desensitizes the material. Another method to avoid sensitization and

resist IGC in ferritic stainless steels is to alloy it with stabilizing elements like Ti

or Nb. Typically, solution annealing at 1200�C and then heat treating at

800�900�C to form carbides of these elements (and causing fast chromium diffu-

sion) induces resistance against IGC.

Duplex Stainless Steels

Duplex stainless steels are much more resistant to IGC than ferritic or austenitic

steel. There is a wide difference in chemical composition of the austenite and the

ferrite phase. Although austenite has more solubility for interstitial elements, it is

also rich in Ni, Mn and so on. The ferrite phase has more ferrite formers, for exam-

ple Cr, Mo, W, V. However, upon heating (welding or heating at temperatures

above 500�C), carbides form at the austenite�ferrite phase boundaries or at the fer-

rite�ferrite boundaries, leading to formation of chromium depletion regions around

these carbides, especially on the austenite side of the boundary. The ferrite phase

also transforms to brittle phases. Even short exposures to these temperatures/weld-

ing are enough to cause severe sensitization, proneness to IGC and formation of

brittle phases [60]. Therefore, thick section welds of these stainless steels need to

be fully solution annealed, wherever practical, if resistance against IGC is required

in a given application. Thus, in applications where welding is not required, for

example tubing in condensers/heat exchangers, these alloys remain highly resistant

against intergranular attack.

Nickel-Based Alloys

Nickel-based alloys, on the other hand, must have very low levels of carbon to

be resistant to IGC. Alloy 600 is known to develop chromium carbides at grain

boundaries and a large Cr depletion zone around it even in (low temperature) mill-

annealed condition [61�63]. Figure 12.4 shows heavy attack at the chromium-

depleted regions at grain boundaries in alloy 600 that had been sensitized and

tested by an electrochemical potentiokinetic reactivation (EPR) test. However,

these alloys are desensitized very quickly as the diffusion rate of Cr in the matrix

of nickel-based alloys is very fast. Alloy 600TT is the heat-treated condition of

alloy 600 in which it is heated at 700�C for 7 h and all the Cr depletion at the grain

boundaries is erased owing to diffusion of Cr from the grain matrix. In this condi-

tion, the alloy has chromium carbides at grain boundaries but the minimum Cr

levels in the surrounding depletion zones go up above 12 wt.%. Austenitic stainless

steels, on the other hand, take much longer times at the sensitization heat treatment


to become desensitized [64]. For example, SS 304 takes 15 days at 750�C or 5

days at 800�C for desensitization to take place (and for resistance against IGC in

nitric acid environments). Similarly, alloy 690TT also has a high resistance to IGC.

Owing to its high chromium content (28%), it is not easy to have alloy 690 in a

sensitized condition (i.e. at least one grain completely surrounded by chromium

depletion, below 12%).

Similarly, use of alloy 800 has now tended to use low carbon variety (0.03 wt.%

maximum) and to keep the Ti/C ratio greater than 12 (or Ti/(C1N). 8). An

annealing temperature of less than 1050�C is reported [65] to result in a signifi-

cantly higher resistance to sensitization of stabilized stainless steels compared with

that when the temperature is above 1050�C. The lower temperatures of annealing

favour formation of carbonitrides and reduce the carbon activity. Therefore, modi-

fied alloy 800 is annealed at 1000�C.

Grain Boundary Engineering

Grain boundary engineering has also been used to improve the resistance of stain-

less steels (and nickel-based alloys) to sensitization and IGC. One method is to

improve the fraction of special (low-angle) grain boundaries [66]. This is done by

thermo-mechanical processing. Typically, low levels of cold working (e.g. 5%

reduction in thickness) followed by low-temperature annealing (e.g. 927�C, 72 h, waterquenching) are shown to result in formation of a high fraction of special-angle grain

boundaries. Another method is to improve the fraction of high-angle grain boundaries

by a different combination of thermo-mechanical processing [67]. A key to improving

the resistance to IGC is attainment of the fraction of specific types of grain boundary

(special or random) to greater than 70�75%. This ensures connectivity of the same

type of grain boundary to provide the resistance to sensitization and hence to IGC

[68]. It was shown that chromium depletions at twin, special and random boundaries

were quite different for sensitized alloy 600, confirming the role of grain boundary

engineering on resistance to sensitization [69].

(A) (B)

Figure 12.4 (A) IGC after EPR test for alloy 600 showing width of attacked regions.

Unattacked carbides can be seen at higher magnification (after the EPR test) in (B) inside

the attacked grain boundaries.


Non-Ferrous Alloys

Non-ferrous alloys, for example cupronickels, are also susceptible to IGC. Use of

lower nickel content (or pure copper) imparts resistance to IGC. A failure analysis

[70] of condenser tubes made from 90-10 cupronickel (ASTM B111, grade C706)

showed that IGC (and cracking) of the tubes was caused by sulfur-containing media

(wet H2S) condensed on the shell side of a condenser having significant amounts of

H2S and ammonia (1�2 wt.% each). The hard-drawn (H105) condition of the tubes

was attributed to its higher susceptibility to IGC. Ammonia and sulfur-based com-

pounds are reported [71] to be the worst pollutants that greatly affect the corrosion

resistance of cupronickels. Even sulfate-reducing bacteria leading to formation of

sulfide species on the tubes is shown to be responsible for IGC in the river water/

cooling waterside of cupronickel tubes [71�73]. An example of intergranular

attack in cupronickel condenser tubes (at the site of under-deposit corrosion attack)

is shown in Figure 12.5. One explanation for preferential corrosion at grain bound-

aries of cupronickel (in this case 70-30 cupronickel) is that the sulfides produced by

bacterial action (or from other sources) present in occluded cells (or in general

media) react with nickel in the grain boundaries. Nickel was shown by polarization

studies to have increased corrosion rates with increasing levels of sulfide in synthetic

seawater [74]. Brasses are prone to IGC because Zn tends to segregate to grain

boundaries. Using brasses with lower Zn content ensures more resistance to IGC.

Figure 12.5 Intergranular attack in the under-deposit regions on the inner diameter (ID) side

of the cupronickel condenser tube.


Lowering Fe content in Al alloys makes them more resistant to IGC as Fe has

very low solubility in these alloys and the excess amount tends to segregate to

grain boundaries and makes it prone to IGC. Similarly, silicon tends to segregate to

grain boundaries in Al and its alloys. Reducing Si content in these alloys also helps

in making these alloys more resistant to IGC. IGC of Al�Zn�In alloy in 3.5%

NaCl solution has been reported [75]. The effect of additional Cu and Mg in

Al�Mn�Si alloy on IGC susceptibility has been investigated after heating at

600�C, followed by reheating at 200�C. IGC occurred in Al�Mn�Si�Cu alloy

after the heating for 1�10 h at 600�C, while pitting corrosion was observed in

Al�Mn�Si�Cu�Mg alloy. IGC in the former alloy was caused by selective disso-

lution of Cu-depleted zone. In the latter alloy, precipitates including Cu were

observed not only in grain boundaries but also in grain, even after a shorter heating

time at 200�C. Therefore, decreasing IGC susceptibility of the alloy was caused

by a smaller potential difference between the Cu-depleted zone and grain matrix.

In conclusion, addition of Mg in Al�Mn�Si�Cu alloy improves corrosion resis-

tance after heating at higher temperature [76]. The ASTM G110 test in an

NaCl1H2O2 environment is used for Al alloys as a test for checking its suscepti-

bility to IGC.

12.3.4 Materials Resistant to SCC

One of the material parameters related to the susceptibility to SCC is the stacking

fault energy (SFE). The lower the SFE, the wider will be the stacking fault and the

greater the probability of its presence on the material surface. The stacking fault

makes the breakage of the surface oxide film easier by easier planar slip and emer-

gence of slip steps on the surface, hence SCC initiation becomes easier. Stainless

steels have in general lower SFE than nickel-based alloys [7,33]. SS 304L (SFE:

24 mJ/m2) is more prone to SCC than pure nickel (SFE: 75 mJ/m2). Addition of the

certain elements (e.g. Ni, Cu, Mo, to some extent Cr and possibly C and N)

increases SFE; hence, increasing the amount of these elements in stainless steels

improves their resistance to SCC (in chloride-bearing environments). Therefore, SS

310L (21% Ni) is more resistant to SCC than SS 304L (12% Ni). Addition of Nb

significantly reduces SFE, and addition of N has shown a variable effect [33].

Austenitic Stainless Steels

In fact SS 304 has been known to be a material that is especially prone to SCC in

chloride environments. A plot of the susceptibility to SCC (time to failure) with the

nickel content in stainless steel (or in iron118% Cr alloy) shows [6�10] that the

minimum time to cracking is at a nickel content of 10 wt.% [77]. If nickel is either

increased or decreased, the time to cracking increases, indicating resistance to

SCC. This also shows that alloys with high nickel content will be highly resistant

to SCC in chloride environments.

A sensitized stainless steel is much more prone to SCC (in intergranular SCC,

IGSCC mode) than an annealed stainless steel. The sensitized condition provides


Cr-depleted regions at grain boundaries that allow for easier breakage of the oxide

film by aggressive ionic species (chlorides/sulfides) and have sluggish repassiva-

tion kinetics due to lower amounts of Cr in the depleted regions. Therefore, SS

304L would be more resistant to SCC than SS 304. Chloride ions are known to pro-

duce the branching type of SCC in austenitic stainless steels. A typical example of

the branching type of SCC in 304 type of stainless steel in sodium chloride solution

is shown in Figure 12.6.

SS 316L is more resistant to SCC than SS 304L. The presence of Mo does not

play a direct role in it. Owing to the presence of Mo (a strong ferrite stabilizer) in

SS 316L, the Cr (a ferrite stabilizer) content is reduced and the nickel (austenite

stabilizer) content is increased to obtain austenite stability. Therefore, it is this

altered chemistry of the alloy (higher nickel content) that imparts more SCC resis-

tance. Also, Mo increases the repassivation kinetics of the alloy, further improving

the resistance to SCC. Mo oxidizes to molybdate (MoO422 ) ions in the solutions

(e.g. chloride environment in an incipient crack). These ions then adsorb on the

crack surfaces or onto the active sites present in the passive film and inhibit attack

by chloride ions. The inhibiting effect of MoO422 on active dissolution is due to

formation of an insoluble lower-valence molybdenum oxide (with iron or chro-

mium ions or in its free form) that also acts as chloride getter. The presence of salt

films on crack fronts supports this hypothesis.

SS 316LN is even more resistant to SCC than SS 316L. This is due to the bene-

ficial effect of nitrogen alloying. Nitrogen, in some reports, has been shown [33] to

increase the SFE. It also gets incorporated in the surface oxide film as nitrates and

delays SCC initiation. It blocks active (dissolving) sites and prevents SCC initia-

tion. Formation of ammonium ions inside an incipient crack delays/prevents its

propagation. SS 316NG (nuclear grade) is a variant of SS 316L with nitrogen addi-

tion up to 0.12 wt.%, carbon restricted to 0.02 wt.% (and C1N# 0.13 wt.%) and

is used in most of the applications in boiling water reactors. Similarly, SS 304LN

is more resistant than SS 304L. However, higher amounts of nitrogen (up to

Figure 12.6 Chloride SCC of austenitic stainless steel showing typical branching of cracks

for the 304 type of stainless steel in boiling sodium chloride solution.


0.15 wt.%) are required in this case to impart resistance against SCC in BWR

simulated environment [78].

Nickel-Based Alloys

Nickel-based alloys are, in general, more resistant to SCC in chloride-bearing

environments [62]. However, addition of Mo and more Cr in these alloys also

makes them more resistant to SCC in chloride environments at higher temperatures.

For condensers cooled by seawater, Mo-bearing alloy N06625 (22Cr, 9Mo, 4Fe,

4Cb and 61Ni) is the least expensive alloy [43]. N10276 (16Cr, 16Mo, 6Fe, 4W

and 48Ni) is the next preferred nickel-based material [43]. N06110 (31Cr, 10Mo,

2W, 57Ni) and is used for high resistance in seawater and when its greater general

corrosion resistance is needed from the process side [43].

Nickel-based alloys also remain resistant in high-temperature aqueous environ-

ments [62]. For example, these alloys are used for steam generator tubing applica-

tions in nuclear reactors and operate at temperatures of 280�320�C. Currently

alloy 690, alloy 800, and alloy 600TT are commonly used in such environments,

largely because of their resistance to SCC [79]. Alloy 800 has to have a Ti/C ratio

of 12 and low carbon levels for resistance to sensitization and SCC [65]. In pressur-

ized water reactors, alloy 690 has to be resistant to primary water stress corrosion

cracking (PWSCC) in lithiated and borated water chemistry at temperatures of

approximately 320�C. On the other hand, the same alloy is in contact with the sec-

ondary water chemistry that is alkaline but may contain some levels of chloride

and operates at temperatures of approximately 280�310�C. However, a variant of

this reactor (the VVER) uses SS 321 as steam generator tubing material. Control of

the secondary water chemistry to avoid chloride ingress is a prerequisite for the use

of stainless steels as SG tubing material. For this, titanium and its alloys (Grades 7

and 12) are used as condenser tubing materials in seawater-cooled plants. In river/

lake water cooled plants, where the chloride levels are quite low, the condenser

tubing material used is SS 316L.

Grain Boundary Engineering

Grain boundary engineering has been shown to improve the resistance of stainless

steels (and nickel-based alloys) to SCC [66�69, 80�83, 83B�83D]. One method

used is to improve the fraction of special (low-angle) grain boundaries (as detailed

in Section 3.2). A growing SCC crack finds it difficult to go along a special (low-

angle, low-energy) grain boundary. Therefore, a large fraction of special grain

boundaries would increase the resistance of the material to IGSCC.

The grain boundary engineering concept to improve resistance to IGSCC was

first applied to nickel-based alloys [80]. Earlier applications [81] showed that ther-

mo-mechanical processing of alloy 600 (three cycles of cold working, each with

over 30% cold work, followed by annealing at 900�1050�C for 2�10 min) leads

to an increase in the fraction of special grain boundaries to over 60%. These studies

showed [81] the improvement in resistance to sensitization/IGC by the G 28A test


of ASTM for 72 h and improvement in resistance to SCC by exposure of UNS

N06600 rings stressed to a maximum of 0.4% strain to 10% NaOH solution at

350�C for 3000 h.

It has been shown [83B] for Ni�16Cr�9Fe�xC alloys that the high-angle (ran-

dom) boundaries resisted SCC in simulated primary water of pressurized water

reactors at 360�C more than the low-angle (special) boundaries. Also, the total

intergranular cracked fraction decreased with the increased fraction of special grain

boundaries, irrespective of the changes in alloy chemistry, microstructure or strain.

They also showed [83B] that discrete intergranular carbides reduce IGSCC in these

nickel-based alloys. It has been shown [83C] that twin boundaries (Σ3, Σ being

the reciprocal of coincident sites) resist SCC much better than other variants of

twins (Σ9, Σ27) in stainless alloys. Further, [83D] the frequency of IGSCC for a

low-energy special grain boundary in SS 316L and alloy 690 in a supercritical

environment at 500�C was 9�18 times lower than that for a random high-angle

grain boundary, at a given strain level of 25%.

In an in situ high-resolution X-ray tomography study [68] to observe SCC, it

was shown that there is a build-up of unfractured ligaments within the cracking

path for alloy 600. It was shown using a two-dimensional scale model that a special

boundary (at a triple point) ahead of the crack tip will not crack and acts as a

bridge, possibly because it remains as a non-sensitized ligament or because it has

an inherently higher resistance to cracking. It was further shown using a three-

dimensional scale model that crack will propagate through other connecting non-

special boundaries. Therefore, the crack path has unfractured ligaments.

Connectivity from one grain to another of random to random or of special to spe-

cial grain boundaries requires [12,68,83,83A] over 23% (in a two-dimensional

scale) grain boundaries of the same nature (hence mixed mode of SCC). When

over 89% (in a two-dimensional scale) of susceptible grain boundaries are present,

it will lead to pure IGSCC. Therefore, grain boundary engineering leading to

improvement in the resistance against IGSCC of nickel-based alloys and stainless

steels has been shown to be an effective tool by modelling as well as by experi-

mental evidence.

The susceptibility of copper-based alloys to SCC is well documented [70,71,74].

Ammoniacal cracking of brass is a common example of SCC. The α-brass has

been shown [84] to be prone to SCC in NaNO2 solution in concentrations from

0.01 to 1 M and pH values of 6.2�12 at room temperature. It was shown that as

long as the applied potential was higher than the repassivation potential (Erp),

transgranular SCC occurred in α-brass. This indicated that passivity breakdown at

localized regions was the main cause of SCC. If ammonia was added to these solu-

tions, intergranular SCC occurred. Adding Al (B2%) to brass makes it resistant to

selective leaching as well as to SCC.

12.3.5 Materials Resistant to Hydrogen Damage

High-strength materials are more prone to hydrogen embrittlement. Alloying steels

with nickel or molybdenum reduces its proneness to hydrogen embrittlement


[6�10,84A,84B]. ‘Clean’ steels (with lower inclusion/void content) are more resis-

tant to hydrogen attack. Also, steels with lower levels of ‘poison’ (for hydrogen

recombination reaction, e.g. sulfur) are more resistant [6,9�10]. In addition,

nickel-containing steels (or nickel-based alloys) have very low hydrogen diffusion

rates and generally do not suffer from hydrogen blistering [6,84A].

Hydrogen damage (or hydrogen embrittlement) in steels is often observed in the

oil and gas industry (e.g. in sour or H2S-containing environments). The corrosion

damage in the form of surface blisters and/or internal cracks in the absence of

applied stress is termed [84B�84D] hydrogen-induced cracking. In the presence of

stress, the damage is known [84B�84D] as sulfide stress cracking (SSC). The cor-

rosion reactions occurring in the presence of hydrogen sulfide are as follows:

Anodic reaction : Fe-Fe21 1 2e2 ð12:8Þ

Dissociation reactions : H2S-H1 1HS2 ð12:9Þ

HS2-H1 1S22 ð12:10Þ

Cathodic reaction : 2H1 1 2e2-2H ðatomic hydrogenÞ-H2ðgasÞm ð12:11Þ

The presence of H2S gas in acidic solution or hydrogen sulfide ions (HS�) inneutral and alkali solution reduces the rate of hydrogen gas formation on steel sur-

faces. Thus, a greater amount of atomic hydrogen diffuses into the steel. Sulfur

plays a catalytically active (otherwise passive) role in the embrittlement. The reac-

tion proceeds in the following manner:

Fe ðsurfaceÞ1H2S ðgasÞ-ðFeSÞðsurfaceÞ1 2ðHÞðsolutionÞ ð12:12Þ

The H atoms liberated in the surface reaction are prevented from recombina-

tion and its subsequent desorption by the catalytic poisoning effect of sulfur. The

hydrogen-induced microvoids or microcracks have been generally observed

[84A,84B] to nucleate heterogeneously at microstructural imperfections such as

grain boundaries, second-phase particle�matrix interphase boundaries, pre-exis-

tent internal voids or cracks and non-metallic inclusions (both with and without

externally applied stress). Therefore, reducing the sulfur content (or the sulfide

inclusion content) in steels is helpful in improving the resistance against hydro-

gen damage.

Materials that are known [6�10,84C,84D] to resist SSC are carbon and alloy

steels (in general, hardness ,HRC 22) in as-received (annealed) condition. If fab-

rication involves cold working/bending, a maximum permanent outer fibre strain of

5% is allowed [84C]. If the strain exceeds this value and hardness exceeds a speci-

fied value for a respective grade of steel, stress-relieving heat treatment (minimum


temperature 593�C) needs to be performed to reduce the hardness [84C]. However,

cold-worked line pipe fittings of A 106 grade B are acceptable, even with a cold-

worked strain of 15%, provided the hardness remains less than 190 HBW [84C].

Standard conditions of partial pressure of H2S gas and temperature of operation are

available [6�10,84D] for each alloy, defining the operation limits that do not lead

to cracking in H2S environments. Austenitic, ferritic or martensitic stainless steels

with a maximum hardness of HRC 22 are generally taken as resistant to SSC. In

general, the ductile austenite matrix is more resistant to hydrogen embrittlement

than the stronger ferritic matrix [84E]. Even the diffusion rates of hydrogen in aus-

tenite matrix are much slower and the solubility limit much smaller than the ferritic

matrix [84E]. Duplex stainless steels should have [84C] a maximum hardness of

HRC 28 (if PREN# 40) or HRC 32 (if PREN. 40), and precipitation hardenable

(austenitic base) stainless steels a maximum hardness of HRC 35 for resistance

against SSC. For duplex stainless steels, the PREN limits are determined by calcu-

lation as in Eq. (12.4) (but with the addition of 0.5 times weight percentage of W

to the Mo content). Nickel-based alloys (generally with hardness less than HRC

35) are taken as acceptable for resistance against SSC. Titanium alloys are prone to

hydrogen embrittlement [84C] if in galvanic contact with other active materials, for

example carbon steels in H2S-containing aqueous media at temperatures above

80�C. Specific hardness limits for each titanium alloy need to be followed while

making the selection of alloy for resistance to hydrogen damage. Aluminium- and

copper-based alloys are taken [84C] as resistant to hydrogen damage (SSC).

Aluminium alloys can suffer corrosion damage if exposed beyond a pH range of

4.0�8.5. Chloride ions can also cause pitting. Copper alloys are prone to SCC in

ammonia environments.

12.3.6 Materials Resistant to FAC

FAC is a subset of erosion corrosion [16�18,85]. The acceleration in the rate of

corrosion attack on a metal surface due to relative movement of the metal surface

and the corrosive fluid that happens in the form of dissolution of metal ions

through the process streams is termed FAC. Therefore, FAC is essentially an elec-

trochemical process. The material surface in any process fluid will have a boundary

layer and the extent of the boundary layer depends on the geometry of the compo-

nent, turbulence, flow velocity and viscosity of the process fluid. Any dissolving

ion from the material has to diffuse through this boundary layer [16�18,85,86].

This, essentially, is one of the two most important factors determining FAC. In

stagnant/low-flow conditions, the solubility limit can be approached and then fur-

ther dissolution would be very slow [16�18,85,86]. The solubility limit for any

metallic ion in a given process fluid is the second determining factor for FAC.

Flow of the process fluid results in provision of new liquid over the same metallic

surface, allowing it to take into solution even more quantities of ions as its solubil-

ity limit has not been reached. Two-phase flow (e.g. steam containing water) is

much more aggressive in causing FAC, as is the case in erosion corrosion. One

explanation of this is the known fact that soluble species partition to the vapour


phase from the liquid phase. This depletes the liquid of the amount of ionic species

(of iron) dissolved in it. Hence an acceleration of corrosion takes place. In addition,

other factors, like reduction in the boundary layer thickness and increased contribu-

tion of mechanical erosion mechanism, may also be taking place.

Single-phase FAC in thermal and nuclear power plants is encountered

[16�18,85,86] mainly at locations where the local flow velocity is very high either

because of some obstruction to fluid flow (e.g. flow measurement device/choked

valves) or the geometry of the components (pipe bends). Two (dual)-phase FAC is

also encountered [16�18,85,86] at locations where steam containing moisture is

available, for example steam extraction lines in the secondary circuits of reactors.

Single-phase FAC is characterized by a typical signature on the affected surfaces

by scallops (or orange peels). The surface film on carbon steel has been shown to

be magnetite in most plants. The overlaying scallop texture is a result of the micro-

scopic oxide structure that develops on underlying metal grains, especially from

pearlitic colonies. A typical scallop structure on FAC affected surfaces of carbon

steel from a nuclear power plant is shown in Figure 12.7A. ‘Tiger striping’ charac-

terizes dual-phase FAC on surfaces that are exposed to steam and water mixtures.

A typical example of a carbon steel component’s surface pattern that worked in

steam and water mixture is shown in Figure 12.7B.

The FAC rates for a 90� elbow predict that the FAC rate reduces to one-tenth

of the FAC rates for a carbon steel pipeline when Cr content is increased

[16�18,85,86] from 0.03 to 0.5%. A better grade of carbon steel (ASTM A335

grade 22 containing 1.9�2.6% Cr and 0.87�1.13% Mo, or grade 11 with 1.25% Cr

and 0.25% Mo) has been considered (and used on trial basis) in many plants for

better resistance to FAC. The surface oxide changes from Fe3O4 to FeCr2O4 with

addition of Cr in carbon steel. The impact strength of the welded Cr�Mo steels has

to be considered before a decision on their use in critical applications is taken.

Another important aspect is to establish if replacement with the same grade SA

106, grade B (but containing 0.25�0.40% Cr) would be effective. This has been

shown [17] to improve resistance to FAC and result in about 50% increase in

(A) (B)

100 µm

Figure 12.7 Typical signatures of (A) single-phase FAC indicated by ‘scallops’; and (B)

dual-phase FAC indicated by ‘tiger striping’ obtained from the FAC affected carbon steel

surfaces from the secondary circuit of a nuclear power plant.


service life. Even allowing a small amount of Mo (up to 0.15%) in the same grade

of steel would be more helpful.

Stainless steel and alloy steel were found [17,87] to be alternative replacement

materials for carbon steel as these are significantly less susceptible to FAC (with

SS grade 316L being the most suitable material). However, substitution with aus-

tenitic stainless steels will require some additional engineering aspects. These

materials have a high thermal expansion rate, on average 1.4 times greater than

that for plain carbon steel. Susceptibility to chloride stress corrosion of austenitic

stainless steels is another concern related with chloride contaminants in thermal

insulation. In a test on SS 316L pipelines exposed to steam of quality 26 wt.% at

292�C, 7.2 MPa pressure and water chemistry of pH 9.7, specific conductivity of

11 μS/cm at 25�C and dissolved oxygen less than 2 ppb (using all volatile treat-

ment), it [87] has been shown that the oxide formed on the stainless steel surfaces

was composed of α-Fe2O3 and FeO, total thickness was 0.60�2.40 μm and the den-

sity of the oxide was 1 g/cm3. For a steam velocity of 53 m/s and two-phase veloc-

ity of 29 m/s, the FAC rates were measured to be 1.3 and 1.7 μm/year in straight

and curved pipes, respectively, in a test lasting 2057 h. These rates confirm the

high resistance of stainless steels to FAC. Carbon steels (A106 grade B) in feeder

applications in the primary heat transport circuit of pressurized heavy-water reac-

tors that operate at 310�C at a pH of 10�10.5, dissolved oxygen less than 1 ppb,

show typical FAC rates of 175 μm/year. Local disturbances due to weld pool geom-

etry that lead to turbulence at the local spots in the feeder contours are shown to

add to the FAC rates.

Another method to have a better resistance against FAC is to alloy steels with

copper [16�18,85,86]. The steel corrodes with time, leaving behind copper (as

oxide) on the surfaces. This copper oxide film on the surface acts as a physical bar-

rier and helps in bringing down the FAC rate. However, certain applications do not

tolerate the presence of copper because of other complications, hence this approach

has limited use in plants.

Apart from corrosion-resistant materials, control of water chemistry [16�18,85]

too helps in minimizing FAC. Controlling pH between 9 and 9.5 for the secondary

circuit of nuclear reactors and keeping dissolved oxygen at a slightly higher level

(a few tens of parts per billion) are some of the measures that add to the resistance

against FAC.

12.3.7 Materials Resistant to Erosion Corrosion

Copper and its alloys (including cupronickels) and aluminium and its alloys are

prone to erosion corrosion [6,7,9,10,36,43]. For tubular products, like in condenser

tubing applications, inlet-end erosion corrosion is a common feature. When the

flowing water from the water box has to accommodate itself for ingress into tubes,

there is turbulence at the inlet ends of tubes. This is further aggravated by ingress

of air bubbles. This inlet-end erosion corrosion problem is addressed by use of

plastic ferrules at the inlet ends of tubes (or by barrier coatings). Even within the

safe range of velocity in a given application, deposition of scales/hard-shelled


organisms within the tubes become sites of local velocity disturbance/localized

turbulence and cause erosion corrosion at adjacent locations [43]. Impingement due

to particulate matter is another common reason for erosion of such materials.

Stainless steels, nickel-based alloys, titanium and its alloys, many more are

resistant to erosion corrosion [6,43]. Carbon steels are severely damaged by steam

containing entrained water droplets. This is even more of a problem for locations

that involve a change of flow direction/disturbance of flow, for example pipe

bends, elbows, pumps, valves especially flow control and pressure let-down valves,

centrifuges, impellers, turbine blades. In contrast, austenitic stainless steels (with

comparable hardness and strength) remain highly resistant to flowing wet steam.

Surface films play an important role in protection from erosion. Titanium readily

forms a strongly protective titanium oxide film that protects it from liquid erosion

corrosion. Similarly, stainless steels are protected by their stable passive film.

Carbon steels are generally more prone to erosion corrosion [88]. Figure 12.8

shows the surface of a carbon steel component that was exposed to an ammoniacal

environment containing precipitates (two-phase flow) and suffered severe erosion

corrosion. These eroded surfaces are devoid of a typical surface pattern that charac-

terizes FAC. However, carbon steels and lead have relatively good resistance to

certain concentrations of sulfuric acid under low to moderate flow rates [6]. Both

depend on a metal sulfate corrosion product film for erosion resistance. However,

both fail rapidly at higher flow rates when the metal sulfate film is removed by

liquid flow.

Figure 12.8 Erosion corrosion of carbon steel components in a fertilizer plant under

two-phase flow showing absence of typical ‘surface patterns’ that are the characteristic of

FAC [88].


Many times, a higher flow rate would also help in providing a greater

concentration of oxygen/inhibitors to metal surfaces, thereby offering better protec-

tion from corrosion. In a flowing seawater environment, austenitic stainless steels

behave better as the flow rates keep deposits from settling and avoid the start of

under-deposit corrosion and pitting corrosion.

12.3.8 Materials Resistant to Oxidation Corrosion

As mentioned earlier, oxidation is an electrochemical phenomenon. The metal

oxide acts as an electrolyte through which diffusion of cations (Mn1 ) and anions

(O22) occurs [6]. Depending on which (cation/anion) diffusion dominates, growth

of the oxide takes place at the oxide-gas or metal-oxide interface.

Both oxide and sulfide scales form on commonly used base metals (Fe, Ni, Cu,

Cr, Co and so on) and grow principally at the scale�gas interface by outward diffu-

sion of cations. However, owing to condensation of cation vacancies at the

metal�scale interface, a significant volume of voids forms at the metal�scale

interface [6].

For Ta, Cb, Hf, Ti and Zr, O2� predominates over cation diffusion [6]. So, sim-

ple diffusion control results in scale formation at the metal�scale interface. In

some of these materials, after a given time of oxidation, large volumes of oxide

form at the metal�scale interface. This leads to formation of tensile stresses and

hence cracking in the oxide, making it porous. From this stage the diffusion of O2�

does not control growth and it is a non-protective scale. However, before the stage

in oxidation when tensile stresses are not high enough to cause cracking, the oxides

remain protective.

In both the cases there are complications [6] due to (i) significant dissolution of

O and S in some metals, (ii) high volatality of oxides, halides and metals, (iii) low

temperature of some oxides/sulfides and (iv) grain boundaries in oxides/metals.

All the metal oxides are non-stoichiometric oxides. The n-type semiconductors

are metal excess oxides and have two vacant anion sites (VO2�) compensated by

four electrons for electroneutrality. The electronic current is carried by electrons

and ionic transport by oxide ions. Some common examples [6] of metal excess oxides

are ZrO2, CdO, TiO2, Al2O3, SiO2 and PbO2. In addition, Ta2O5 and Cb2O5

are metal-deficient oxides but have VO2� as predominant defects. In these metals/

alloys, addition of lower-valence metal ions increases the oxygen hole and the diffu-

sion-controlled oxidation rate increases. On the other hand, alloying/doping with high-

er-valence ions reduces the oxygen ion vacancies and the diffusion-controlled

oxidation rate reduces. Therefore, alloying Zr (valence 4) with Ta (valence 5)

improves its oxidation resistance.

The oxidation of Zr alloys is basically dictated by the matching of the lattice of

the monoclinic oxide with that of the base metal [89�92]. Matching of the oxide

and base alloy lattice takes place by contraction of the oxide lattice or by dilation

of the base alloy lattice. If an alloying element ion dilates the oxide lattice (when

replacing Zr41 ), the possibility of matching with the lattice of the metal reduces,

and the oxidation resistance of the alloy is reduced. In other words, if the alloying


element ion addition in the alloy matrix reduces the lattice size, the chances of

lattice matching of the oxide and the base alloy matrix reduce and its oxidation

resistance reduces.

Nitrogen ion (N32 ) has an ionic radius of 1.48 A, Zr41 has 0.82 A, Sn41

0.67 A and O22 has 1.36 A. Therefore, when an Sn41 replaces Zr41 , it produces

contraction of the crystal lattice of ZrO2. When N32 replaces O22 , it dilates the

oxide lattice. When both are present in the oxide, the dilating influence of nitro-

gen is compensated by the contracting influence of Sn41 . Up to about 0.05% N

in Zircaloy does not influence its oxidation behaviour. Above this amount, nulli-

fying the deleterious effect of nitrogen in the material requires about five times

the addition of tin. In Zircaloy-2 or Zircaloy-4 [91], the Sn composition is

1.20�1.70% to take care of the amount of nitrogen typically present. As N32

replaces O22 and Sn41 replaces Zr41 , both elements (N and Sn) individually

have a deleterious effect on oxidation resistance. However, when present

together, they nullify their deleterious influences. Other than Sn, there are other

elements, for example Sb and Nb, that have smaller ionic radii and thus counter

the deleterious effect of nitrogen.

Iron has a body-centred cubic (bcc) lattice and can exist as Fe21 or Fe31 with

ionic radius of 0.80 or 0.67A, respectively. So, replacing Zr41 by either Fe21 or

Fe31 in the monoclinic oxide lattice causes it to contract and allows better match-

ing of the oxide lattice with that of the base alloy. This improves the oxidation

resistance. Chromium also has a bcc lattice and has 0.83, 0.64 and 0.35 A ionic

radius for Cr21 , Cr31 and Cr61 ions, respectively. Replacing Zr41 in the oxide

with Cr31 causes beneficial effects on its oxidation resistance. Cr61 would not

form a solid solution and Cr21 somewhat worsens the oxidation resistance as it has

a slightly higher ionic radius than Zr41 . Similarly Ni21 has a lower ionic radius

(0.74 A) than Zr41 and causes improvement in oxidation resistance. Addition of Ni

to Zircaloy causes higher pick-up of hydrogen from water and steam at high tem-

peratures. This causes premature embrittlement of the alloy due to hydriding.

Therefore, an Ni-free alloy (Zircaloy-4) is used for applications where oxidation

resistance and resistance to hydrogen pick-up are required. Zircaloy-4 is the nickel-

free version of Zircaloy-2 and contains only 70 ppm Ni compared with 0.1% Ni in

Zircaloy-2. This helps in reducing the hydrogen uptake rate of the alloy. The

nascent hydrogen is produced in the corrosion reaction and a fraction of it is picked

up by the alloy. The pick-up rate is dependent on the alloy, with Zircaloy-4 picking

much lower than Zircaloy-2. Zr�Nb alloys too picks up a very small fraction of

the nascent hydrogen produced during the corrosion reaction.

Nb exists as Nb41 or Nb51 , having ionic radii of 0.67 and 0.66 A, respectively.

These are smaller than the radius of Zr41 , thus Nb reduces the dilation in the oxide

and is a beneficial element against aqueous oxidation. Zr�1Nb and Zr�2.5Nb

alloys are commonly used as materials in nuclear power plants for oxidation resis-

tance. Addition of more than 2.5% Nb poses fabrication problems as the alloy is

harder and less ductile than Zr-2.

The oxidation kinetics of alloys can therefore be linear, parabolic, cubic or loga-

rithmic (as shown in Figure 12.9A). In the initial stages of oxidation, Zircaloys


show an oxidation kinetics that is closer to the cubic rate law [89A,90,91,91A].

The general oxidation behaviour is shown [6,91,92] by Eq. (12.13):

ΔWn 5 kct ð12:13Þ

where ΔW is the weight gain in mg/dm2, kc is a constant and t is the time in hours.

The value of n for Zircaloy-2 in steam at 400�C is shown to be 2.84 and Kc to

be 1784.32 [90�92]. After exposure for a certain time, the oxidation kinetics of

Zircaloys change [90�92] to a linear rate law, as shown by Eq. (12.14):

XW 5 kLt ð12:14Þ

where kL is a constant.

For Zircaloy-2 at 400�C, the post-transition kinetics have been shown [90,91] by

Eq. (12.15):

ΔW 5 0:519t1 31:59 ð12:15Þ

At 400�C steam, the transition from a near-cubic to linear oxidation rate takes

place [90] after about 41 days and at 40 mg/dm2. In 360�C water, the transition

takes place after 112 days and at 34 mg/dm2. The corrosion rates are [90�92] only

23 1024 to 73 1024 in. of metal corroded per year, corresponding to an oxidation

rate of 0.37 and 1.27 mg/dm2/day (mdd) at the transition point at 360 and 400�Cfor Zircaloy-2. These rates are too low to be of consequence in applications

[90�92]. However, the transition to linear rates leads to spalling off of oxides

owing to stresses generated at the metal-oxide interface. This spalling off leads to

porous oxides that allow direct contact of water/steam with the metal surface,

(A) (B)

O

200 µm

Linea

r

Parabolic

Cubic

Logarithmic

Time (h)

Zircaloy

Oxide

Mount

Wei

ght g

ain

(mg/

dm2 )

Figure 12.9 Typical (A) oxidation kinetics of Zr alloy and (B) oxide morphology at a stage

where it starts to spall off.


leading to faster oxidation and hydrogen pick-up. A typical morphology of oxide

that has grown to a stage where oxides start to spall off is shown in Figure 12.9B.

Zr�Nb alloys, on the other hand, show an oxidation rate closer to parabolic and

do not show any transition in oxidation behaviour. The oxidation rate laws deter-

mined for these alloys are shown in Eqs. (12.16) and (12.17) [90�92]

For Zr�2.5 wt.% Nb,

log ΔW 5 0:37 log t1 1:149 or ΔW2:7 5 1278:88 t ð12:16ÞFor Zr�1 wt.% Nb,

log ΔW 5 0:38 log t1 1:029 or ΔW2:63 5 512:83 t ð12:17Þ

The p-type oxides, on the other hand, are metal-deficient oxides in which con-

duction is by diffusion of electron holes [6,90]. Addition of lower-valence cations

decreases the cation vacancy or increases electron holes and results in slower

growth rates than in pure metal oxides. Ni is an example that has p-type oxide.

Addition of Li11 to it would increase the oxidation resistance, and addition of

Cr31 would lead to faster oxide growth rates than for NiO.

High-Temperature Corrosion

Corrosion also occurs by direct reaction with an oxidizing gas at elevated tempera-

ture without the need for an aqueous environment [93,94]. This type of corrosion is

referred to as tarnishing/high-temperature oxidation/scaling. The mechanism of cor-

rosion remains the same as described for oxidation corrosion in an aqueous

medium. The rate of attack increases substantially with increasing temperature, and

the growth of the scale is either at the metal�scale or scale�gas interface, depend-

ing on whether it is a p- or n-type oxide/scale [6]. The initial rate of corrosion is

quite high, and subsequently, when the scale is non-porous and completely covers

the surface, the corrosion rate comes down with time and is ultimately controlled

by diffusion rate of the reactive species through the scale. High-temperature scales

are usually considered as oxides, but may be sulfides, carbides or a mixture of

these species. Sulfides typically show a greater diffusion rate of transport of anions

and cations than oxides of the same metal, hence they are inherently less protective

than oxides. For resistance against oxidation, the surface oxides of Cr2O3, Al2O3

and SiO2 are considered good. In pure state, Al2O3 shows the slowest rate of diffu-

sion of metal and oxygen ions, hence it is considered the best resistant against oxi-

dation. Alloys that contain sufficient amounts of elements (e.g. Cr, Al, Si) in

excess of the minimum level required for scale formation are more resistant to

high-temperature oxidation. Most oxide scales lose protectiveness owing to spall-

ation of the scale as a result of thermal cycling or stress generation at the metal�s-

cale interface. In addition, Cr2O3 is prone to being volatile at temperatures above

1010�C and its degradation is accelerated by high gas-flow rates.

When sulfur activity of the gaseous environment is sufficiently high, sulfides,

instead of oxide, form. In most alloys, Al2O3 or Cr2O3 should form in preference to


sulfides. Destructive sulfidation attack occurs only at sites where protective oxides

have broken down. The role of sulfur appears to be to tied up with Cr or Al as sul-

fide, effectively redistributing the protective scale-forming elements near the alloy

surface and thus interfering with the process of formation and reformation of the

scales. If most of the Cr/Al in the alloy surface is tied as sulfides, the other ele-

ments form subsequent sulfides/oxides and their rate of formation and growth is

quite higher. In many cases their melting point is lower and these molten sulfides

are present on the surfaces. Sulfur can diffuse through protective oxides and form

sulfides beneath them, in the metal. These sulfide phases may be preferentially oxi-

dized by the moving oxide front, so the sulfur is further pushed inwards, forming

sulfides deeper inside the material and causing faster oxidation. This has been

shown to lead to finger-type protrusions of oxide/sulfide from the alloy surface that

acts as spot of localized stress or to reduce the load-bearing capacity of the alloy.

Stainless steel is prone to such attack in the product gas from a coal gassifier.

Carburization of high-temperature alloys is likely only at very low oxygen par-

tial pressure, as protective oxides of Cr, Al and Si, and so on are more likely to

form than carbides. In carbon-containing environments, carburization may occur.

Carbon transport across continuous non-porous scales of Al2O3 or Cr2O3 is very

slow. Alloy pretreatments, for example smooth surfaces or pre-oxidation (that pro-

mote such scales), are also effective in decreasing carburization attack. In general,

the scales formed on high-temperature alloys often consist of multiple layers of

oxides, with the protectiveness being derived from the innermost layer, which usu-

ally is the richest in Cr/Al. Gaseous species like CO in the outer porous oxide

layers lead to sufficiently high carbon activity for carburization to occur. Creation

of microenvironments is also possible under deposits that create stagnant conditions

which are not permeable by the ambient gas. Once inside the alloy, the detrimental

effects of the carbon depend on the location, composition and morphology of the

carbides formed. Ferritic steels have more resistance to carburization than austen-

itic steels as ferrite has much lower carbon solubility. Similarly, Fe�Cr alloys con-

taining less than 12% Cr contain various amounts of austenite, depending on

temperature, and are prone to carburization. Such alloys with 13�20% Cr will

form austenite as a result of absorption of small amounts of carbon, but alloys with

more than 20% Cr can absorb considerable amounts of carbon before austenite

forms, so they are resistant to carburization. Minor elements, especially Si, Nb, W,

Ti and rare-earth elements, are known to promote resistance to carburization.

Hot corrosion is generally used to describe the form of accelerated attack experi-

enced by hot gas path components of gas turbine engines [93,94]. Type I hot corro-

sion occurs in the temperature range 850�950�C. It is a sulfidation attack on the

hot gas path parts involving formation of condensed salts (Na2SO4/K2SO4), which

are often molten at the operating temperature of the turbine. The source of sulfur

and sodium is from the fuel/ingested air. Sodium in air above just 0.008 wt.% is

sufficient to cause such an attack in turbine applications. Impurities in fuel, for

example V, P, Pb and Cl, may combine with Na2SO4 to form mixed salts with

even lower melting points and further broaden the range of conditions over which

this form of attack occurs. It has been shown that a high Cr concentration in the


alloy is required to resist type I hot corrosion in turbine applications. The use of

alloy 738 and alloy 939 for the first-stage blades/buckets and FSX 414 for the first-

stage vanes/nozzles in nearly all the applications shows that these are the materials

with the best combination of high-temperature strength and resistance to hot corro-

sion. Type II or low-temperature corrosion occurs in the metal temperature range

650�700�C, which is well below the melting temperature of Na2SO4 (884�C). Itproduces characteristic pitting attack that results from formation of low melting

mixtures of essentially Na2SO4 and CoSO4, corrosion products resulting from the

reaction between the blade/bucket surface and SO3 in the combustion gases. The

melting point of this mixture is 540�C. Here, the partial pressure of SO3 in the

gases is critical for the reaction to occur. Co-free nickel-based alloys are therefore

free from type II hot corrosion. The resistance to this type of hot corrosion

increases with increasing Cr content of the alloy (or the coating).

Acknowledgements

I thank my colleagues in the Corrosion Science Section, Materials Science

Division, Bhabha Atomic Research Centre, who have worked extensively on vari-

ous forms of corrosion and have provided corrosion- and metallurgy-related support

to various units of the Department of Atomic Energy, India. This chapter has

brought together in one place the expertise in the Corrosion Science Section on var-

ious forms of corrosion and the vast experience on combating these in different

plants.

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