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

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the core consisted of pearlite and ferrite. Upon

rapid cooling, the carburized case had a small

amount of martensite with a correspondingly

lower amount of pearlite, and in the core, a lower

amount of ferrite with the remainder being pear-

lite was obtained. However, for the 12KhN3A

alloy, air cooling produced predominantly mar-

tensite and retained austenite in the case and

correspondingly less pearlite. The core con-

tained bainite (Ref 76).

In practice, it was shown that gears manu-

factured from 20KhGB steel exhibited an aver-

age reduction in diameter (0.7 to 1.4 mm) from

463+0.3 after carburizing. This was attributed

to reduced stability of austenite in the pearlite

region. Conversely, for the 12KhN3A steel, the

gears increased in size (0.2 to 0.5 mm) after

carburizing because austenite was more stable in

the pearlite region. These data suggest that dis-

tortion can be controlled by alloy and cooling

rate selection, or improper selection of cooling

rate or alloy can result in unacceptable compo-

nent distortion (Ref 76).

The effect of phase transformation behavior

in the case and the core due to thermal gradients

such as those occurring during heating and

cooling will affect shape and size distortion of

small parts, as illustrated by Fig. 32 (Ref 19). If

the thermal and transformational stresses exceed

the yield strength of the steel, corresponding

distortion will occur.

Quenching and Grinding Cracks (Ref 77)

Quenching cracks occur when tensile

stresses of the ﬁrst kind are greater than the

material strength. Quenching cracks typically

occur during, or in some cases after, quenching at

temperatures less than the M

temperature. Sus-

ceptibility to cracking increases with the carbon

content of the steel, increasing austenitizing

temperature, and cooling rate, especially in the

M

transformation temperature range.

The probability of cracking during quenching

increases with the presence of stress raisers

Fig. 32 Resulting distortion after heat treatment of different steels after quenching in oil and water

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such as notches, abrupt changes in section size,

keyways, and holes. Quenching cracks are irre-

versible but can be minimized by appropriate

design modiﬁcations, reduction in hardening

temperature, and use of lower quench severity.

The reasons for the occurrence of quenching

defects include:

 Poor surface cleanliness, such as residual

forging and metalworking lubricants, and so

on, which leads to nonuniform quenching,

increased thermal gradients, and soft spots

 Incorrect loading and arrangement of parts

in the furnace, which leads to nonuniform

heating and related distortion

 Excessive heating rates, which may lead to

warping and cracking

 Lack of a protective atmosphere to eliminate

oxidation and decarburization of the steel

surface, which will lead to reduction in

mechanical properties after hardening and

a decrease of hardness of the superﬁcial

layer

 Excessive cooling rates and incorrect im-

mersion into the quenching bath, causing

cracking, warping, and twisting

 Insufﬁcient cooling rates or undersized

quench tanks, which will inhibit the desired

martensitic transformation

Quenching cracks, which are characterized by

relatively large depth and short length, rarely

occur in the case of carburized or carbonitrided

components. This can be explained by more

beneﬁcial compressive surface stresses than

those typically formed in higher-carbon-

containing through-hardened parts. However,

internal cracks may form in the core below the

carburized surface or in the transition zone, that

is, in the places with the largest tensile stresses in

carburized parts. Cracks may also form on the

surfaces and corners of carburized parts, which

is related to triaxial tensile stresses in these

locations. Therefore, to prevent cracking, the

case must be sufﬁciently deep so that stresses

developed at any point below the surface are less

than the fatigue limit of the material at that point

(Ref 1).

When steel contains greater than 0.5% C in a

martensite matrix, such as in the carburized case,

intergranular fracture along prior-austenite grain

boundaries may occur. In this situation, the

intergranular fracture is due to the presence of

both phosphorus and cementite formation on the

austenite grain boundaries during austenitizing

or cooling from austenitizing temperatures.

Krauss has referred to this fracture mechanism

as quench embrittlement and has suggested

that the mechanism for this to occur is analogous

to quench cracking in through-hardened steels,

which is due to the formation of tensile sur-

face stresses during quenching, as described

previously (Fig. 33) (Ref 78). However, since

relatively high surface compressive stresses are

present in properly carburized steels, quench

cracking should not exist if conditions for po-

tential intergranular cracking are present.

Fig. 33

Scanning electron micrographs of overload case fracture surfaces in carburized SAE 8620 steel. (a) Quenched directly after

carburizing at 927



C (1700



F). (b) Reheated to 788



C (1450



F). Both specimens were tempered at 145



C (300



F).

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Krauss has reported that carburized steels will

fail by intergranular cracking if sufﬁcent bend-

ing or tensile stresses are applied to offset the

compressive stresses in carburized cases. The

fracture map shown in Fig. 34 illustrates three

conditions where susceptibility to intergranular

cracking can be minimized: carburizing, inter-

critical austenitizing, and applications where

loading is Hertzian or compressive (Ref 79).

In some cases, microcracking can occur with

higher-carbon lath martensite matrices, which

are present in the hardened carburized case.

Microcracking is due to contact of martensitic

plates with each other or with the austenitic grain

boundaries. The potential for microcracking

increases with the austenization temperature.

Prior-austenite grain size also affects micro-

crack density, which decreases with decreasing

prior-austenite grain size. However, microcrack

density is not affected by quench severity

(Ref 78, 80). Figure 35 illustrates examples of

microcrack formation in the carburized case of

SAE 8620 steel (Ref 78).

The presence of microcracks can further lead

to a surface defect called ﬂaking, which refers to

ﬂakelike fracture and subsequent peeloff. (Note:

Flaking is initiated at microcracks that may also

be caused by surface damage due to lubricant

contamination by chips, shavings, burrs, or

abrasive powder ingression into the lubricating

system.)

Krauss reported that fatigue resistance

decreased with increasing microcrack forma-

tion. Crack initiation occurred at the site of the

microcrack that acts as a stress concentrator (Ref

78). Although microcracks can be removed by

surface grinding and polishing, fatigue failure

may still be initiated at prior-austenite grain

boundaries, intergranular surface oxides, or

surface defects such as scratches, machining

marks, and surface asperities due to roughness.

Aksenova et al. showed that contributing

factors to cracking of case-hardened gear wheels

included residual tensile stresses in the case-core

interface, grain growth and overheating, super-

saturation of the case with carbon, excessively

high cooling rates during quenching, too low

a quenching temperature, and insufﬁcient resi-

dence time in the quench tank after immersion

(Ref 81).

Fig. 34

Krauss fracture map illustrating conditions where susceptibility to intergranular cracking can be minimized: carburizing,

intercritical austenitizing, and Hertzian or compressive contact loading

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McEvily et al., reported the use of fracto-

graphic analysis to explain cracking of a

carburized AISI 9310 gear at the case-core

transition zone. One of the factors was the dif-

ferential in the Poisson ratio (n) between the case

and the core that develops during a monotonic

bending test when the core deforms plastically

while the case deforms elastically. McEvily

used as an example to illustrate this point a

9 mm diameter round bar with a 1 mm case.

When the bar is ﬁrst loaded initially and the

deformation is elastic, the Poisson ratio for the

case and the core is equal. Upon further bending,

the core will deform plastically while the case is

deforming elastically. However, because of the

difference in the Poisson ratio between the case

(n = 0.3) and the core (n = 0.5) at this point,

radial tensile stresses will develop. For a tensile

strain of 1%, the radial strain in the case (relative

to the centerline) would be 0.003, and the radial

strain in the core would be 0.005. To remain

compatible, the difference in the strain and

corresponding displacements must be accom-

modated by the formation of a tensile stress.

However, this radial stress will result in a state of

triaxial stress that will promote brittle behavior

by inhibiting plastic deformation. The total

tensile stress may then be developed sufﬁcient

to result in rupture at the case-core interface

(Ref 82).

Grinding Cracks. Grinding may be used for

postprocessing of components to remove growth

and distortion that may have resulted from car-

burizing and carbonitriding. Grinding is also

performed to remove such metallurgical features

as carbide ﬁlms, internal oxidation, and high-

temperature transformation products that may

impart deleterious performance properties. In

addition, grinding processes are commonly used

to create the desired surface ﬁnish to improve

bending and contact fatigue and lubrication

properties.

Surface cracks in the carburized case may

occur during the grinding process, which can be

attributed to microstructural transformations

and thermal stresses producing tensile forces. If

the tensile forces in the case exceed the material

strength, then surface cracks will result. This is

due to the difference in the speciﬁc volume of

the transformational phases present in the case

structure, primarily martensite and austenite.

Structural defects, such as those caused by

inclusions, will also inﬂuence the susceptibility

of the steel to cracking.

Generally, crack creation during grinding is

inﬂuenced by thermoelastic tensile stresses that

are created in the surface cooling zone during

grinding. They are dependent on thermal and

mechanical properties of the material, maximum

contact temperature, grinding feed depth, and

cooling rate.

One study conducted on the generation of

grinding cracks showed that microstructural

heterogeneity, such as the presence of carbide

inclusions (particularly those with a mean dia-

meter of 6 to 10 mm), which were shown to be

associated with large internal residual stresses,

were a predominant cause of grinding cracks

(Ref 83).

Grinding burns arise when excessive heat is

generated during the grinding process, and this is

characterized by surface discoloration. The term

grinding burn refers to localized surface tem-

perature increases at least sufﬁcient to cause

tempering of the martensitic surface, resulting in

localized soft spots. Furthermore, since carbide

precipitation volume contraction accompanies

tempering, the burnt areas are in tension and, if

the resulting tensile stresses are sufﬁcient, sub-

ject to tranverse cracking. However, in other

cases, the increase may be in excess of the Ac

temperature, producing an austenitic surface

that, upon rapid cooling, may produce a hard,

light-etching, martensitic thin layer at the sur-

face. This induced defect is known as a rehar-

dening burn, which is characteristically

surrounded by a layer of tempered steel (Fig. 36)

(Ref 84). In this case, the rehardened zone is in

compression due to the martensitic volume

expansion, and the surrounding areas of tem-

pered martensite are in relative tension. Crack-

ing may occur in the area surrounding the

rehardened material or in the interface between

the two (Ref 84).

Fig. 35

Microcracks in the martensitic case of a coarse-

grained SAE 8620 steel

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There are two possibilities to prevent the

occurrence of thermal defects in the surface

resulting from the grinding process:

 Reduction of heating due to the rotation of a

grinding wheel by reducing the speed of the

grinding wheel will reduce heating and

friction. Conversely, increasing the speed of

the grinding wheel will increase heat pro-

duction and the potential for the formation of

grinding burns due to overheating.

 Increasing heat abstraction due to the

grinding process and reducing the contact

time between the grinding wheel and part

will reduce the tendency for grinding burns.

Grinding cracks may exhibit characteristic

short, parallel cracks (Fig. 37), or they may

exhibit a “chicken-wire” pattern and are typi-

cally between 0.076 and 0.13 mm (0.003 and

0.005 in.) deep. The parallel cracks are typically

deeper than the chicken-wire pattern. Grinding

cracks form perpendicular to the grinding

direction. As indicated previously, the potential

for grinding cracks is affected by improper heat

treatment or a metallurgical structure that is

prone to cracking. For example, if the surface

temperature exceeds the Ac

temperature, the

steel in this region may transform to austenite,

then upon rapid cooling, a hard martensite layer

may form. This effect is called a rehardening

burn. Grinding cracks may be detected by a

magnetic particle test.

Some carburizing steels such as chromium

and chromium-manganese steels, which include

SAE 5120 and 20MnCr5, may undergo over-

carburizing with subsequent cracking of the case

upon cooling, which then renders the part more

susceptible to the formation of grinding cracks

(Ref 1). Overcarburizing leads to the formation

of a complex carbide network that is excessively

brittle, which causes greater susceptibility to

grinding cracks.

Severe grinding may lead to the development

of residual tensile stress, which can be the

initiation point for crack formation. Since cracks

will not propagate into layers of compressed

stress, it therefore may be advantageous to shot

peen the part prior to grinding to prevent the

formation of grinding cracks (Ref 86).

Parrish reported that the potential for forma-

tion of grinding cracks may be minimized by

(Ref 84):

 The thermal conductivity of the steel is an

important design variable, and free carbides

and retained austenite have an adverse effect

on thermal conductivity.

 The surface carbon concentration should be

between 0.7 and 0.9%.

 Parts should be tempered immediately after

quenching.

 The tempering temperature should be as

high as possible while still achieving the

necessary surface hardness.

Improper Case Depth (Ref 77). In a recent

study conducted by Bahnsen et al. on carburized

SAE 5120 test specimens to rate the relative

inﬂuence of surface carbon content, case depth,

and carburizing temperature on distortion, it was

Fig. 37

Example of grinding cracks on the ﬂank of a worm

gear. Source: Ref 85

Fig. 36

Microstructure of a section through a rehardening

burn. Original magniﬁcation: 500·

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reported that of these variables, the most domi-

nant effect was observed for case depth (Ref 87).

Case depth of carburized steel is determined by

the carburizing time and the available carbon

potential at the surface. One of the most com-

mon defects of carburized and carbonitrided

materials is an insufﬁcient or excessive case

depth. For example, when prolonged carburiz-

ing times are used to produce a deep case depth,

a high carbon potential will produce a high

carbon content on the surface and the possible

corresponding formation of excessive retained

austenite or free carbides, which may lead to an

improper residual-stress distribution in the case-

hardened part. Therefore, a high carbon poten-

tial may be suitable for short carburizing times

and shallow case depths but not for prolonged

carburizing times and deep case depths. Fur-

thermore, the fatigue limit of carburized SAE

8620 steel was related to case depth and also

microstructure, distribution of retained auste-

nite, depth of internal oxidation, and near-

surface compressive residual stresses (Ref 10,

21). Bending-fatigue strength decreased with

the increasing case depths due to the presence of

increasing internal oxidation and nonmartensitic

transformation products at the surface. Wear

behavior of carburized 8620 steel is also related

to case depth (Ref 88).

To assure optimal quality during the produc-

tion of case-carburized parts, the following are

essential:

 The processing temperature should be

accurately controlled.

 Maintain temperature uniformity throughout

the load.

 Rack the parts to assure uniform gas ﬂow

throughout the load.

 Use uniform circulation of atmosphere

throughout the load in the furnace

 To properly control case depth, either use

shim stock or sample the parts periodically

during the carburizing cycle.

 Conduct the process at the lowest acceptable

temperature and time.

 The components constituting a load should

possess uniform size and surface area with

respect to each other, using empirically

established carburizing conditions.

 Avoid carburizing and carbonitriding in the

same furnace. Use a separate furnace for

each process, if possible.

The optimal case depth for a speciﬁc com-

ponent and steel alloy is based on the design and

service conditions of the component. Typically,

the case depth is designed to provide the

necessary residual-stress distribution for the

wear requirements for the part (Ref 89). Typi-

cally, the greater the case depth, the greater the

fatigue strength (Ref 23). For carburized steel

gears used in the automotive industry, for

example, SAE 8620, hardened case depths are

generally 0.8 to 1.4 mm. Improper case depth

may be caused by establishing an unnecessarily

restrictive case thickness speciﬁcation that is not

appropriate for the process or the particular

furnace in use, which leads to decreased and

nonuniform hardness and unacceptable material

properties.

Insufﬁcient Case Hardness and

Improper Core Hardness (Ref 77)

One reason for insufﬁcient case hardness is

the presence of incorrect microstructure, such as

bainite. The appearance of bainite in a carbur-

ized case in even small amounts will sig-

niﬁcantly decrease fatigue strength, especially

contact fatigue strength. It is an important

microstructural defect to be avoided. The pre-

sence of bainite in the carburized case is

particularly problematic because it cannot be

detected by hardness measurements and by the

severity of the quenchant used to harden the steel

after carburizing. Figure 38(a) illustrates a bai-

nitic case microstructure of carburized SAE

8620. Figure 38(b) shows the core structure. In

the case of carbonitrided components, this is less

important because nitrogen increases harden-

ability of steel more than carbon.

Insufﬁcient as-quenched case hardness is

caused by:

 Insufﬁcient carbon content in the entire case

or in the superﬁcial zones

 Increased retained austenite content

 Insufﬁcient case hardenability

 Insufﬁcient case depth

On the other hand, core hardness is dependent

on carbon content, steel alloy hardenability, and

section size.

Case carbon content less than 0.4 to 0.5% C is

easily detected by conducting a spark test.

Insufﬁcient case carbon content resulting from

the gas carburizing process occurs when the

process is conducted with a low carbon poten-

tial, inadequate furnace pressure, or cooling the

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load from the diffusion temperature to the

quenching temperature without a protective

atmosphere or improper atmosphere composi-

tion. These problems may be rectiﬁed by addi-

tional soaking in a carburizing atmosphere with

the proper carbon potential.

Hardness reduction due to an increase in

retained austenite content may occur with a

direct quenching after the carburizing process.

This is even more critical with steels containing

increased chromium or nickel, which may pro-

duce retained austenite levels as high as 80 to

100% in the surface after quenching. As the case

thickness increases, the zone containing retained

austenite may be sufﬁcient to cause a signiﬁcant

reduction in hardness. To reduce the retained

austenite content from a direct quenching pro-

cess, it is important to select the proper steel

alloy for carburizing and to use the appropriate

quenchant and quenchant conditions, including

the use of a subzero treatment, if necessary.

Reduced hardenability usually does not occur

throughout the entire case but only in the 0.001

to 0.01 mm depth from the surface. Common

problems leading to reduced hardenability in-

clude internal oxidation and overcarburization

of surface zones. To improve case hardenability,

internal oxidation and overcarburizing should be

prevented.

When the case depth is too shallow, the

observed hardness is dependent on the load

applied. For example, for thicknesses of 0.3

to 0.4 mm, if the surface hardness is deter-

mined using a 60 kg load, the value will be

approximately 12 HRC units higher than if the

hardness is determined using a 150 kg load. In

practice, it is possible for the hardness to differ

by 1 to 2 HRC units from published values for

the material. Hardnesses that are i2 HRC units

less than published values may be due to the use

of an incorrect steel grade or an insufﬁcient

austenitizing temperature for the steel harden-

ability and section size in use. Hardness values

iHRC units higher than published values may

be due to the use of an incorrect steel grade,

excessive case depths, or puncturing a copper

layer or paste on surfaces that are protected from

carburization.

Core Microstructure. The design material

properties of case-hardened steels are not only

dependent on a martensitic case but also on the

microstructural composition of the core. An

important design criterion is the ultimate tensile

strength, which is dependent on the micro-

structure of the core. For example, soft cores

(5770 N/mm

, or 50 ton/in.

) are suggestive

of a core with high ferrite content, as shown in

Fig. 39(c) (Ref 19), and a hard core (41240 N/

, or 80 ton/in.

) would be expected for a

predominantly martensitic structure, as shown in

Fig. 39(a) (Ref 19). Intermediate structures

would be bainitic structures, such as those illu-

strated in Fig. 39(b) (Ref 19). The effect of core

microstructure on ultimate tensile strength is

illustrated by Fig. 40 (Ref 90). The approximate

relationship between the core microstructure

and hardness for a Ni-Cr-Mo steel is illustrated

in Fig. 41 (Ref 19).

Fig. 38

The case and core microstructure of carburized SAE 8620 test specimens (0.95% C potential); carburized at 955



(1750



F), quenched into a 50:50 mixture of sodium nitrate and potassium nitrate at 250



C (480



F), held 120 min, then air

cooled and tempered at 250



C (480



F) for 240 min. The case is lower bainite (56 HRC), and the core (42 HRC) is lath martensite. These

images were made by etching with 10% Na

(sodium metabisulﬁte). The magniﬁcation bar is 20 mm. Courtesy of G. Vander Voort,

Buehler Ltd., Lake Bluff, IL

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Inﬂuence of the Transition Zone. There

is a transition zone between the case and the

core, and the thickness of a transition zone is

dependent on carburizing time (transition zone

thickness increases with carburizing time),

carburizing medium, and carburizing tempera-

ture (excessive carburizing temperature will

lead to increased pearlite content in the core).

After subsequent hardening, if there is too great

a transition between the case and the core, there

is increased potential for peeling and chipping,

resulting from the presence of a martensitic

carburized case and an uncarburized core con-

taining sorbite.

Fatigue cracks occur most often in the tran-

sition zone, which subsequently propagate into

the core and into the carburized case. When the

gradient between the case and the core changes

too rapidly, operations such as grinding or

when a component is subjected to bending due

to heavy loadings, could lead to peeling and

chipping failures. For example, deep grinding

could not be performed. Increasing the depth

of the transition zone will increase the strength

of adhesion of the case to the core. Then, if the

transition zone is sufﬁciently large, deep grind-

ing operations may be performed.

Figure 42 provides the microstructure of

two transition zones. Figure 42(a) illustrates a

gradual transition between the case-core

microstructure. A more rapid transition between

the case-core microstructure is illustrated

in Fig. 42(b). Typically, a transition zone

such as that illustrated in Fig. 42(a) is desired,

since it will exhibit a lesser tendency for chip-

ping.

Inﬂuence of Surface Carbon Content

Overcarburizing or Overcarbonitriding.

Important microstructural defects related to

carburized or carbonitrided case structure

include overcarburization or overcarbonitriding

of the case and coarse grain structure. Excessive

carbon content (carburizing) or carbon and

nitrogen (carbonitriding) is typiﬁed by the pre-

sence of carbides or carbonitrides in the case,

which creates an almost continuous nonetching

area. Case-hardened steels with these micro-

structures characteristically exhibit increased

brittleness, a propensity for chipping during

grinding and use, and decreased fatigue strength

and pitting resistance.

Fig. 39

Microstructures obtained by cooling a 0.16%C-

3%Ni-Cr steel from 920



C. (a) Fast cool (920–

200



C in 30 s), giving low-carbon martensitic structure of

1590 MPa ultimate tensile strength (UTS). Original magniﬁca-

tion: 800· . (b) Intermediate cooling (920–250



C in 200 s),

giving bainitic structure of 1360 MPa UTS. Original magniﬁca-

tion: 800· . (c) Slow cool (920–250



C in 104 s), giving a ferrite/

pearlite structure of 740 MPa UTS. Original magniﬁcation:

800·

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Fig. 40 Approximate effect of microstructure on the ultimate tensile strength of low-carbon, low-alloy steels

Fig. 41

Approximate relationship between core microstructure and hardness of a Ni-Cr-Mo carburizing steel (approximately 4%

alloy content) with approximately 0.16% C. The alloy content/carbon content extension (upper right corner of the ﬁgure)

permits phase percentage plots to be adjusted in relation to the ﬁxed hardness scale to approximate core strength for other steels. Below

250 HV represents slow-cooled (normalized) and annealed steels, and bainite can be read as bainite, pearlite, or spheroidized carbides.

Above 250 HV refers to quenched steels. For the 180



C tempered condition, there will be zero change at 360 HV and below, but there

will be a 20 HV loss at 100% martensite.

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Figure 43 illustrates the effect of carbon

content on the hardness of martensite in carbon

and alloy steels. Increasing carbon content to

0.5% increases hardness from 20 to 65 HRC.

However, increases in carbon content to ap-

proximately 1% do not produce a corresponding

increase in hardness above 65 HRC.

Hardness and mechanical properties are

related not only to carbon content but also the

composition of carburized steels, which is illu-

strated in Fig. 44. This ﬁgure shows the corre-

sponding carbon gradients of the cases of three

carburized steels: chromium-molybdenum, car-

bon, and nickel (Ref 25). The data in Fig. 44

show that the presence of chrome and molyb-

denum increases the case carbon content, and

nickel decreases the case carbon content. The

case carbon content is increased due to the pre-

sence of carbide-forming elements; their struc-

tures in the carburized case inﬂuence the

mechanical properties of the steel. Over-

carburization may also lead to quench cracking

(Ref 91).

The inﬂuence of steel composition on the

microstructure of a carburized and quenched

20H steel is shown in Fig. 45. Figure 45(a) shows

a martensite-retained austenite microstructure

with some carbides dispersed throughout the

structure. Fig. 45(b) shows a martensitic-

retained austenite structure with network car-

bides. This structure is due to the chrome content

in the steel and leads to varying carbon content

during carburizing. Additional examples of the

case microstructure of SAE 8620 steel and core

microstructures of SAE 1524 and 8115 steel are

shown in Fig. 46 to 48 respectively.

Decarburization is the opposite of carbur-

izing. While carburization is performed to

increase carbon content in the surface of steel,

decarburization is the process by which carbon

is lost from the surface of steel. Decarburization

can lead to catastrophic failures of components

(Ref 92) and must be minimized because of

fatigue failure such as bending and contact

fatigue (Ref 93). Figure 49 illustrates decarbur-

ization of a poorly carburized SAE 8620 steel,

and Fig. 50 shows the microstructure at higher

magniﬁcation.

Decarburization occurs at temperatures in

excess of 700



C in the presence of gases that act

as decarburization agents, which include carbon

dioxide (CO

), water vapor (H

O), hydrogen

), and oxygen (O

). The decarburization

process involves the following chemical

reactions of molecules with carbon in the steel

surface (C

) until there is an equilibrium

Fig. 42

Micrographs illustrating transition zones between the carburized case and the uncarburized core. (a) Illustrates a gradual

transition. (b) Illustrates a rapid transition between the case and the core microstructures. Original magniﬁcation: 500·

Sources of Failures in Carburized and Carbonitrided Components / 213

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