Canale L.C.F., Mesquita R.A., Totten G.E. Failure Analysis of Heat Treated Steel Components

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carbon-rich carburizing atmosphere. As the

temperature increases further to Ac

, additional

ferrite will transform to austenite by combining

with carbon from the furnace atmosphere, pre-

ferentially over carbon transfer from the spher-

oidal carbides already present until the austenitic

transformation is complete, at which point there

will be an equilibrium between the austenite and

carbides (Ref 112).

Those steels containing alloying elements

such as chromium and manganese that will

reduce the eutectoid carbon content and Ac

temperature are more likely to develop globular

carbides. These elements will increase the

solubility limit of carbon (A

), thus shifting the

equilibrium diagram to the left, as shown in

Fig. 64 (Ref 112).

It should be noted that although the terms

spheroidal carbides and globular carbides are

often used interchangeably, these carbides may,

in fact, possess a round, angular, or even nee-

dlelike appearance. Also, although the structures

are typically designated as M

C, their actual

ratio of the element/carbon composing the car-

bide being observed is not only dependent on the

time and temperature of heat treatment but pri-

marily on the elemental availability in the steel

during formation.

There is a three-step process of high-density

carburizing that is conducted to increase case

hardness by aggressive precipitation of cemen-

tite (Fe

C) to improve surface fatigue strength

(Ref 113). Generally, grain-boundary carbide

precipitation will form network carbide struc-

tures that are susceptible to quench cracking and

result in reduced fatigue strength. Interestingly,

it is possible to obtain desirable surface fatigue

strength by vacuum or plasma carburizing a

chromium-containing steel such as SCr420H

or a chromium-molybdenum steel such as

SCM420H (Ref 114). This process involves the

formation of a M

-type globular carbide that

is approximately 1 mm in diameter. These

microstructures are reported to exhibit excellent

surface fatigue strength and rolling fatigue

strength under high bearing loads (43 GPa) at

relatively high temperature (100 to 300



C),

which are unachievable in the presence of larger

(i3 mm) M

C carbides.

Reheating a carburized steel to a temperature

below A

causes spheroidized carbide particles

to form. (Both austenite and cementite are stable

at this temperature.) Since these carbides bind

some of the carbon in the case, there will be an

increase in the M

temperature. In addition,

grain-boundary migration is reduced (Ref 115).

Carbides are typically very hard. For exam-

ple, the microhardness of globular carbides in a

carburized case of a plain carbon steel has been

measured to be 41000 HV, and carbides in a 2%

Ni-Cr steel have been measured to be approxi-

mately 800 HV (Ref 112). Therefore, the pre-

sence of free globular carbides is often assumed

to improve component wear, abrasion, and

scufﬁng resistance. However, in a study reported

by Parrish on the effect of globular carbides on

contact fatigue with carburized 2% Cr-Mn steel,

it was shown that the presence of massive

Fig. 63

Bad globular carbide formation in the case of a car-

burized 9310 steel. Etchant: boil in alkaline sodium

picrate solution (45 s). Scale = 10 mm. Courtesy of G. Vander

Voort, Buehler Ltd., Lake Bluff, IL

Fig. 64

The saturation surface carbon content (A

) of var-

ious carburizing steels as a function of carburizing

temperature as related to the iron-carbon equilibrium diagram

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carbides was detrimental, while a beneﬁcial

effect was obtained if the carbides were ﬁner and

better distributed, as shown in Fig. 65 (Ref 112).

The presence of globular carbides may also

lead to grinding cracks, which are also related to

surface residual stresses.

Network Carbides. Globular carbides are

most typically formed when the carbon content

is less than the A

. However, under some

conditions, it is possible for globular carbides to

still be formed when the carbon content is

greater than the solubility limit (in excess of the

eutectoid composition) and austenite is super-

saturated with respect to carbon. Most typically,

these are conditions where carbon will pre-

cipitate in the grain boundaries as cementite

(Fe

C) during slow cooling from the carburizing

temperature, leading to the formation of network

carbides. Figure 66 illustrates network carbide

structure observed in a broken carburized AISI

P5 steel tool.

The equilibrium diagram shown in Fig. 67

illustrates the conditions for the formation of

network carbides when excess carbon is pre-

cipitated from austenite as Fe

C (Ref 112).

When the carbon steel supersaturated with car-

bon content of C

is cooled from t

to t

,Fe

will begin to precipitate. As cooling continues

to t

, additional carbon will have precipitated

as Fe

C until the carbon content is C

. At the

eutectoid temperature t

, carbon precipitation as

C will have stopped, and the austenite carbon

content C

will transform to a pearlitic eutectoid

microstructure. The relative proportion of Fe

to austenite can be determined from Fig. 67

using the lever rule (C

B).

Although austenite is supersaturated with

respect to carbon during the carburizing process,

and carbide precipitation at austenitic grain

boundaries will occur during cooling, if the steel

is quenched from the carburizing temperature,

the excess carbon can be retained by the result-

ing as-quenched martensitic/retained-austenite

microstructure. Typically, during carburizing,

the load is cooled in the furnace from the car-

burizing temperature to the temperature from

which the steel will be hardened. Since the Fe

migrates to the grain boundaries during slow

cooling, faster cooling by quenching will allow

retention of the carbon in the martensitic/auste-

nitic structure. Thus, cooling rate can be used to

control the amount of network carbide forma-

tion. Alternatively, carburized steel can be

cooled to ambient temperature and reheated to

820 to 860



C and quenched. However, it has

been shown that traces of the network carbides

remained in the microstructure even after heat-

ing to 900



C (Ref 112). Parrish also has

reported that if steel containing excess carbon is

subcritically annealed prior to quenching, a ﬁne

Fig. 65 Contact fatigue strength of carburized 25Kh2GHTA steel (tempered at 180 to 200



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dispersion of relatively unharmful spheroidized

carbides will be obtained instead of network

carbides (Ref 112).

Karpov studied the effect of cooling from the

carburizing temperature on the nature of net-

work carbides formed when quenching gas-

carburized 11 by 11 by 56 mm 07Kh16N6 steel

test specimens. The test specimens were car-

burized, cooled, then reheated to austenitize to

1020



C, and then cooled to the quenching

temperature at 0.036



C/s, which is less than the

critical quenching rate that leads to network

carbide segregation. Quenching was conducted

in water at room temperature. From Table 11, it

is evident that network carbide formation begins

at 900 to 850



C for this alloy and is completed

at 600



C. From the phase diagram for this alloy,

the A

temperature is 860



C (Ref 116). Dur-

ing the course of this work, Karpov found that a

nondestructive electromagnetic ﬂaw detector

could be used to rank the network carbide size

(Ref 116).

Network carbides have been reported to

reduce surface fatigue (pitting) resistance of

carburized steel used for bearing applications

(Ref 114). Like globular carbides, complex

carburized networks in an overcarburized case

will reduce the potential for carbide redissolu-

tion during reheating, which will lead to in-

creased brittleness and grinding cracks (Ref 1).

The presence of network carbides in a Kaplan

turbine blade constructed from improperly car-

burized 17CrNiMo6 steel was reported to be a

major contributor to failure by an intergranular

microcracking mechanism (Ref 117). The

cracks seemed to follow the path of the network

carbide structure. Parrish has summarized var-

ious studies and concluded that continuous net-

work carbides do reduce fatigue properties,

leading to premature cracking failures by a

stiffening mechanism (Fig. 68) However, other

studies with partial nertwork carbides showed

no deleterious effects (Ref 112). The presence of

network carbides also is not expected to produce

an adverse effect with respect to wear under

heavily loaded conditions or scufﬁng (Ref 112).

The three most common failure modes of

carburized steels are ductile fracture, cleavage,

and intergranular fracture (Ref 23). Ductile

fracture is caused by nucleation growth and

coalescence of voids that are initiated at inclu-

sion sites and second-phase particles. Cleavage

fracture occurs by separation at crystallographic

planes by a transgranular pathway. Intergranular

fracture, such as that involving cementite

Fig. 66

Micrographs of a broken carburized P5 tool steel die.

Note the layer of cementite along the surface and the

heavy grain-boundary network in (a). The case was 59.5 HRC,

and the core was 22 GRC. (a) The case is shown at an original

magniﬁcation of 100· (100 mm bar). (b) The case at an original

magniﬁcation of 1000· (10 mm bar). (c) The core at an origi-

nal magniﬁcation of 400· (25 mm bar) in Nomarski differential

interference contrast (note the “old” and “new” ferrite, as in dual-

phase steel). The austenitizing temperature for the case is

approximately 1475



F, which is in the two-phase ﬁeld for the

core. New ferrite formed due to limited hardenability in the

quench. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

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deposition at grain boundaries, involves crack-

ing on grain boundaries and is due to (Ref 23):

 Precipitation of a brittle phase (such as net-

work carbides) on the grain boundary

 Hydrogen embrittlement

 Environmental-assisted cracking

 Intergranular corrosion

 Grain-boundary cavitation and high-

temperature cracking

Film or Flake Carbides. Surface-ﬁlm or

ﬂake carbides are composed of a continuous or

discontinuous carbide ﬁlm with typically little or

no penetration into the case structure, which is

caused by cooling of the carburized steel in the

furnace with high carbon potential. Parrish

summarized various previously published re-

ports that stated such carbide ﬁlms contain

approximately 19% Fe

C, 16% austenite, with

the balance being martensite. These ﬁlms cover

a nonmartensitic layer of approximately 30 mm

thickness, with carbides penetrating into the

grain boundaries (Ref 112). Koistinen showed

that such ﬁlms have high tensile surface

(0.025 mm depth) residual stress, as shown in

Fig. 69 (Ref 44). Until now, no further infor-

mation relating to the presence of these carbide

types on carburized steel properties has been

reported.

Fig. 67 The use of the iron-carbon equilibrium diagram to illustrate network carbide formation

Table 11 Effect of initial cool-down period on

network carbide formation of carburized

07Kh16N6 steel

Cool-down temperature(a), °C

Network carbide

severity rating(b)

1020 (initial temperature) 1

950 1

900 1–2

850 1–3

800 1–3

700 3–4

600 5–5

(a) The cooldown occurs from 1020 to the temperature shown at a rate of

0.036



C/s. (b) 1, very ﬁne carbide network; 3–4, failure rating; 5, largest

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Noncarbide Inclusions

Two of the primary causes of fatigue failure

are inclusions and surface defects. Inclusions

may be metallic impurities or metallic oxides

(Ref 118). Metallic element inclusions (impu-

rities), although typically in trace quantities,

may be traced to the scrap used in the steel-

making process. These elements cause inter-

granular segregation, which may lead to crack

Fig. 68 Comparison of bending fatigue of carburized 12Khn3 gears showing adverse effect of network carbides

Fig. 69

Residual-stress distribution of carburized SAE 1018 steel with a ﬁlm-carbide layer formed due to a high carburizing potential.

The surface layer consisted of 16% Fe

C, 16% retained austenite, and the balance was as-quenched martensite.

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formation, detrimental precipitate formation,

and are often observed as “slivers” in the ﬁnal

product.

Metallic oxide inclusions vary in morphology

and composition. Some sources of these oxides

include:

 Deoxidation products, such as alumina

inclusions, that are formed by the reaction of

dissolved oxygen and the added deoxidant,

such as aluminum. Alumina inclusions are

dendritic when formed in the presence of

high oxygen concentration.

 Reoxidation products are generated when

the aluminum remaining in the liquid steel is

oxidized by FeO, MnO, or SiO

and other

oxides in the slag or refractory materials or

by exposure to the atmosphere.

 Slag entrapment occurs in metallurgical

ﬂuxes entrained in the steel. This occurs

during transfer between steelmaking vessels.

These inclusions are typically spherical.

 Exogenous inclusions from other sources

include dirt, broken refractory brickwork,

and ceramic lining materials. They act as

sites for heterogeneous nucleation of alu-

mina and may include a central particle.

 Chemical reactions may produce oxides

from inclusion modiﬁcation when calcium

treatment is imperfectly performed. Inclu-

sions containing CaO may also originate

from entrained slag.

All steels contain various noncarbide in-

clusions. It is well known that inclusions, in

addition to surface defects and inhomogenities

such as retained austenite, nonmetallic inclu-

sions, and inhomogeneities, can reduce fatigue

strength and negatively inﬂuence ductility and

toughness (Ref 119). Although harder inclusions

such as metallic inclusions are more harmful

than softer inclusions, nonmetallic inclusions

are still generally harmful. In addition, the de-

leterious effects of inclusions increase with size.

In one study conducted by Bomas and

Schleicher on the effect of inclusions on the

fatigue strength of carburized 16MnCrS5 (SAE

5115) steel, it was found that subsurface fatigue

crack initiation was initiated by nonmetallic

inclusions up to depths of 1.4 mm (Ref 120).

Similar results were obtained for a study of

MnS-induced bending fatigue failure of car-

burized EN39B steel (Ref 121).

Even in clean steels, oxide and sulﬁde inclu-

sions exist. From a study of the effect of defects

such as inclusions on the fatigue bending

strength of carburized SCM20 steel, a model of

fatigue crack initiation during fatigue bending,

such as would occur in a gear tooth, was devel-

oped (Ref 122). This model is illustrated sche-

matically in Fig. 70. Figure 70(a) illustrates the

chromium and manganese oxides and grain-

boundary segregation of silicon oxides present

in a decarburized surface layer. When the sur-

face is loaded, the grain boundary or oxides act

as a stress raiser for microcrack formation, as

shown in Fig. 70(b). Although most of the cracks

remain in the decarburized layer, the most cri-

tical crack penetrates deeper into the steel, as

shown in Fig. 70(c). The crack that initiates fati-

gue failure was thought to be due to a combi-

nation of the microcracks shown in Fig. 70(b). In

this case, the decarburized layer in combination

with the presence of inclusions was projected to

be the source of fatigue failure (Ref 122).

Fig. 70

Model of fatigue crack initiation due to the presence of inclusions in a nonmartensitic (decarburized) steel layer. Source:

Ref 122

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Micropitting

Micropitting fatigue usually occurs on heav-

ily-loaded surface-hardened components and

is characterized by a frosted or gray-stained

appearance under thin-ﬁlm lubrication condi-

tions (Ref 87, 123). Numerous small cracks in

the surface may exceed the depth of the micro-

pits (Ref 124). Under magniﬁcation, small pits

approximately 10 mm deep will be observed.

The surface will appear etched with a pattern

that sometimes follows the slightly higher ridges

left by cutter marks or other surface irregula-

rities on the ﬁnished component. Micropitting is

inﬂuenced by high surface loads, frictional heat

generation due to poor lubrication caused by

insufﬁcient ﬁlm thicknesses in the partial elas-

tohydrodynamic lubrication regime, excessive

retained austenite, tangential speed, and lubri-

cant additives and oxidation (Ref 123).

Micropitting is strongly inﬂuenced by the

relative lubricating oil ﬁlm thickness and can

be quantitatively related to both the surface

condition (roughness) and the thickness of

the lubricating ﬁlm by the lambda (L) value

(Ref 124, 125):

L=h=s=

ﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃ

)

where s

and s

are the root mean square surface

roughness of the two opposing wear surfaces,

and h is the lubricating oil ﬁlm thickness. When

Li3, there is full ﬁlm lubrication with no

asperity contact. When 0.85L53, there is

partial elastohydrodynamic lubrication. When

L50.8, there is a boundary lubrication condi-

tion. When L51, micropitting will occur, and

once micropitting occurs, pitting fatigue (mac-

ropitting) will be accelerated (Ref 124). How-

ever, if macropitting does occur, it is often

characterized by an arrowhead or fan shape (Ref

123). There have been reports of substantial

improvements in fatigue lives, such as with

carburized 9310 steel, with corresponding reduc-

tions in surface roughness (Ref 125).

For gears, there is a critical temperature where

pitting fatigue and scufﬁng are likely to occur.

This is called the critical scufﬁng temperature

), which is calculated from:

where T

is the equilibrium temperature of

the gears before meshing, and T

is the ﬂash

temperature, which is the instantaneous tem-

perature rise due to localized friction heat at the

point where the gear teeth mesh. The value of T

is controlled by gear geometry design, and the

value of T

is controlled by the lubricant visc-

osity and surface roughness.

To minimize micropitting:

 Use higher operational speeds and smooth

material surfaces.

 Use the recommended amount of clean, dry

lubricant with the highest viscosity permis-

sible.

 Reduce the lubricating oil temperature and

surface loading.

 Use the optimal amount of case carbon

content in carburized gear materials.

Contact Fatigue Pitting (Macropitting)

Pitting failures occur when fatigue cracks are

initiated on the tooth surface or just below the

surface. Usually, fatigue pits are the result of

surface cracks caused by metal-to-metal contact

of asperities or defects due to insufﬁcient lubri-

cant ﬁlm thickness. They are dependent on the

Hertzian contact surface stress and the number

of stress cycles (Ref 124). Surface asperities of

the harder material of a wear contact will lead to

damage of the softer surface, sometimes by a

work-hardening mechanism, leading to the

creation of microcracks that then become fatigue

pits as the wear process continues (Ref 124).

Pitting damage is commonly encountered with

rolling element bearings, gears, and machine

components subject to cyclic rolling-sliding

motion under a load. Initially, fatigue pits may

occur in localized areas and may range in size

from 0.38 to 0.76 mm (0.015 to 0.030 in.) in

diameter (Ref 85).

Vinokur et al. examined the effect of case

carbon content of carburized 18KhGNMFL steel

(1.3% Mn, 1.3% Cr, 0.8% Ni, 0.25% Mo, 0.1%

V) on the contact endurance of fatigue tests

conducted with a wear test with an applied stress

of 3500 MPa using a lubricating oil. Contact

fatigue was the average of 10 tests to determine

the number of cycles until pitting was observed.

The case depth was approximately 1.8 mm. The

carburized steel was hardened from the inter-

critical range and tempered at 170



C. The

results of this study are summarized in Fig. 71

(Ref 126). The contact endurance increases with

carbon content up to approximately 0.9% C and

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then decreases. The optimal hardening tem-

perature was at approximately 850



C, which is

in the intercritical temperature range, just below

the upper critical temperature for the steel.

Although fatigue pits are usually initiated at

the surface, subsurface initiation is relatively

common in case-hardened rolling element

bearings with serious inclusion problems. In

these situations, failures usually do not follow

the case-core interface (Ref 125, 127). The

potential for micropitting, pitting, and spalling

phenomena may be assessed from the lambda

(L) value (see the section “Micropitting” in this

chapter); however, even relatively smooth sur-

faces and lubricating ﬁlm thickness, such as

those used for high-speed gears, may exhibit

pitting failures due to the presence of subsurface

cracks. The subsurface cracks may be caused by

the presence of inclusions that act as stress

concentrators, causing the crack to propagate

parallel to the surface and subsequently break

through the surface.

Spalling failure occurs when the wear process

causes several pits to join together. These larger

craters are usually caused by more severe over-

loading conditions. As the number of stress

cycles increases, the pitting process will con-

tinue in an effort to relieve stresses. The rule is

that spalling cracks initiate where the ratio of

shear stress to Vicker’s hardness is maximum

(Ref 124). However, this relationship is not

correct when there are excessive amounts of

retained austenite. Generally, it is assumed that

spalling will occur when L41.

To prevent pitting fatigue, either the surface

loading must be decreased to a level below the

endurance limit of the material or the hardness

must be increased to increase the endurance

limit (Ref 28). Pitting may also be reduced by

instituting a break-in period at reduced loads and

speeds to improve gear tooth contact (Ref 85,

124].

Other potential causes of fatigue pitting

include hydrogen embrittlement due to water

contamination of the lubricant and particle

contamination of the lubricant, which act as

surface stress-concentration points that lead to

pitting failure.

Case Crushing

If the case depth is too deep, case-core

separation may occur due to the tips of the gear

teeth becoming too brittle and possibly break-

ing, if the case depth is too thin, the strength

of the gear teeth will be reduced, causing pre-

mature pitting, or it may lead to a condition

called case crushing (Ref 131). Case crushing

occurs in heavily loaded case-hardened com-

ponents such as gears. Case crushing occurs by a

subsurface fatigue process where the high-cycle

contact stress exceeds the endurance limit. This

will occur when subsurface stresses exceed the

strength of the core. Case crushing failures may

have a similar appearance to pitting, although it

often occurs as longitudinal cracks on the sur-

face of only one or two gear teeth, where sec-

tions of the tooth surface may subsequently

break away. However, the case material may

appear to have chipped away from the core in

large ﬂakes (Ref 85). The observed cracks will

move toward the case-to-core boundary and then

to the gear surface (Ref 129).

Adequate case support is provided by proper

core structure to not only prevent case crushing

but also to transmit torque, support bending

loads, and provide adequate toughness to prevent

brittle fracture. The presence of any ferrite will

contribute to reducing the toughness of the core.

Case crushing may be prevented by increasing

the case depth and possibly the core hardness.

For general applications where core hardnesses

of 30 to 45 HRC are speciﬁed, the required case

depth can be estimated from (Ref 128):

Case depth to 50 HRC=½1:2 · 10

7

(W)=F

where W is the force in pounds pressing the

surfaces together, and F is the length of the line

Fig. 71

Effect of case carbon content and hardening tem-

perature on the contact endurance limit of carbur-

ized 18KhGNMFL steel (1.3% Mn, 1.3% Cr, 0.8% Ni, 0.25% Mo,

0.1% V). Source: Ref 126

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contact in inches. The strength of the core can be

determined from hardness. Since case crushing is

promoted by shear, the shear strength of the core

must be determined and can be estimated from

Fig. 72 (Ref 128). Generally, the subsurface

stress/strength ratio should not be greater than

0.55.

Eryu et al. have studied case-crushing fracture

mechanisms of carburized 20CrMnTi by

scanning electron microscopy (Ref 130). They

showed that the surface features of the primary

cracks exhibited scaly features and that there

were two features of the fracture surfaces of the

branching cracks. The fatigue steps and dimples

were analogous to materials of higher strength,

and the fragmentation pattern was analogous to

brittle material. In addition, spherical particles

were observed that were composed of a-iron,

which were proposed to be caused by the

movement of the faces of the primary crack due

to shearing and compressive stresses.

Pitting Corrosion

Pitting corrosion is a localized penetrating

corrosion attack of typically corrosion-resistant

steel resulting in a mass loss of the steel (Ref

131). Pitting corrosion is related to localized

discontinuities of a passive layer caused by

mechanical imperfections, inclusions, surface-

localized chemical attack of the passive layer by

salts such as chlorides, or by overaggressive

lubricant additives. After the corrosion pit is

created, the localized chemical surrounding is

much more aggressive than the surrounding area

of the uncorroded material.

Initiation of the pitting process is dependent on

temperature and on the steel surface, including

the presence of sediment. In some applications

where pitting corrosion is more prevalent, such

as steel in concrete structures, the pitting corro-

sion process is characterized by the temperature

or narrow range of temperatures above which

pitting will nucleate. The creation of corrosion

sediments will lead to a temperature decrease.

Pitting corrosion will only occur above this cri-

tical temperature. Therefore, to increase the

lifetime of steel used in reinforced concrete by

reducing the rate of pitting corrosion, frequent

sediment removal (cleaning) is recommended.

Most often, pitting corrosion is initiated by the

presence of chloride salts, and the rate of cor-

rosive attack is steel alloy dependent. The cri-

tical concentration of chlorides for different

steel alloys cannot be deﬁned, because corro-

sivity is dependent on other chemicals that may

be present, which will affect the rate of corrosion

attack. However, since pitting corrosion is

typically relatively fast, it should be prevented.

Resistance of steel to pitting corrosion is

dependent on the alloy composition (chromium,

molybdenum, tungsten, nitrogen). Relative cor-

rosion resistance of steel alloys may be empiri-

cally quantiﬁed by:

Relative corrosion resistance=%Cr

+3:3 ½%Mo+0:5 (%W)+16 (%N)

Chromium and molybdenum are also useful

alloy additions to minimize the potential for

stress-corrosion cracking.

Corrosion pitting may also be caused by

chemical attack of the steel surface by lubricant

additives such as extreme-pressure additives,

particularly in the presence of acid, water, or

contaminants. Also, during use, the oil itself will

oxidatively degrade, producing acidic by-pro-

ducts that may lead to corrosion pitting. In

addition to pitting, corrosive attack may occur at

the grain boundaries of the carburized case. It is

particularly important that components exposed

to saltwater, liquid chemicals, or other foreign

materials during use should be sealed from their

operating environment (Ref 85).

Pitting corrosion may also occur during heat

treatment, particularly those processes involving

salt baths (Ref 132). Heating of steel with scale

on a surface not only accelerates decarburizing,

Fig. 72

Shear strength of carburizing steels as a function of

hardness

232 / Failure Analysis of Heat Treated Steel Components

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but pitting corrosion is also accelerated. The

corrosive attack is increased with temperature

and holding time. If the scale is not uniform, then

pitting corrosion is localized to those areas

where scale is present. If residual salt from the

bath crystallizes on the surface of the steel,

violent boiling may occur during subsequent

quenching in oil, which may result in blister

formation on the surface. After cleaning, loca-

lized pitting corrosion will then occur. To avoid

pitting corrosion during furnace heating, parts

should be thoroughly cleaned.

In some cases, stainless steel machine parts

are carburized to reduce wear. However, car-

burizing a stainless steel (1Kh16N2AM) reduces

corrosion resistance where machine parts are

used in humid environments (Ref 133). The

potential for pitting corrosion increases as the

amount of d-ferrite increases. Corrosion resis-

tance is reduced as the quenching temperature

prior to carburizing is increased. Although

increasing the quenching temperature after car-

burizing to 950 to 1100



C does not affect cor-

rosion resistance, decreasing the quenching

temperature to 800 to 850



C reduces corrosion

resistance. Increasing the tempering tempera-

ture decreases corrosion resistance, as shown in

Fig. 73 (Ref 133). The corrosion resistance of

carburized steel is greatest after stress relieving

at 250



Partial Melting

Partial melting occurs when there is nonuni-

form heating of the surface of the steel, such that

some areas are heated to the liquation tempera-

ture (the partial melting temperature of an alloy)

(Ref 77, 131). Corners and edges are particularly

susceptible to partial melting. Microscopically,

the presence of partial melting is typically

observed as black spots containing retained

austenite in a large cluster of carbides. Macro-

scopically, partial melting is accompanied by

the formation of tiny surface cracks.

Partial melting occurs when the carburized

steel is heated to an excessively high tempera-

ture, resulting in incomplete or selective car-

burizing of the surface. For example, partial

melting may occur during stray current ﬂow into

the load from electrodes used to heat salt pot

furnaces or if a load is placed too close to the

furnace hearth, so that some areas of the load are

heated to an abnormally high temperature. To

avoid this defect, heating in salt baths with

appropriate composition and at appropriate aus-

tenitizing temperature should be conducted by

keeping the load at a recommended distance

from the heating electrodes. Similarly, when

heating in a conventional furnace, the load

should be properly placed to facilitate uniform

heating.

It is also important to be aware of the liquation

temperature (beginning of melting) of the alloy

being heated. Some typical examples of the

approximate liquation temperatures for different

steels are provided in Table 12, where the

soaking temperature is 2 to 3 min. Precise deﬁ-

nition of the partial melting temperature range is

typically a difﬁcult task, because of the rela-

tively large data scatter due to potential com-

positional variation within the alloy, variation of

carbon content, and, in some cases, relatively

large carbide segregation. These structural var-

iations favor the potential for partial melting to

Fig. 73

Effect of tempering temperature on corrosion resis-

tance of carburized stainless steel 1Kh16N2AM. The

corrosion test was conducted in a humidity cabinet. Source:

Ref 133

Table 12 Approximate liquation temperature

for various steel alloys

Steel alloy

Approximate liquation

temperature, °C

SW14 1320

SW18 1330

SW7Mo 1280

SK5 1350

SKC 1280

SK5V 1270

SW12C 1260

SK10V 1250

Sources of Failures in Carburized and Carbonitrided Components / 233

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