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

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region of martensite plates along the prior-

austenite grain boundaries (Ref 19).

There are basically three different modes of

fracture through the tempered martensite of

specimens tempered in the range of 260 to

370



C. First is brittle transgranular fracture,

which results from the decomposition of

retained austenite in the second stage of tem-

pering, as mentioned previously. Films of

retained austenite between laths of martensite in

quenched medium-carbon steels transform into

thin plates of cementite on tempering. The sec-

ond mode of fracture associated with TME is

intergranular. This kind of fracture is quite

common and has been related to phosphorus

segregation to the austenite grain boundary

Fig. 14 Effects of titanium, vanadium, chromium, and molybdenum on tempering hardness behavior. Source: Ref 5

Fig. 15

Hardness and toughness of a tool steel as a function

of tempering temperature. Charpy V-test performed

at room temperature in the short-transverse direction

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during austenitization. However, data on the

effect of phosphorus on the fracture surfaces in

the untempered condition indicate only that the

presence of phosphorus at the prior-austenite

grain boundary is not sufﬁcient for the devel-

opment of TME. So, the interaction between

phosphorus and cementite is necessary for the

intergranular mode of TME (Ref 20, 21).

It is important to understand that the fracture

is occurring along a prior-austenite grain

boundary, which is now a high-angle ferrite

grain boundary (Ref 22). Krauss (Ref 12)

describes another type of transgranular fracture

mode associated with TME, which is observed

in 4340-type steels. It is interlath cleavage

induced by cracking parallel to the cementite

formed from the retained austenite transforma-

tion. Transgranular fractures in TME may be

related to the interlath carbide thickness; thinner

carbides cause interlath fracture, and thicker

carbides promote translath cleavage. In some

low-carbon steels, embrittlement is associated

with peculiar carbide morphology that provides

numerous sites for microcrack initiation, grow-

ing by microvoid coalescence and then fracture,

with little gross plastic deformation.

Silicon additions to carbon steels raise the

temperature range in which TME occurs (as

mentioned earlier), because silicon delays both

the conversion of the transition carbide (e)to

cementite within the martensite laths as well

as cementite coarsening at boundaries at

higher tempering temperatures (Ref 16, 17). The

effect of silicon content on the impact prop-

erties of 0.6%C-0.47%Mn-0.52%Cr-1.77%Ni-

0.19%V-1.0%-2.5%Si steels can be seen in

Fig. 17 (Ref 16).

Investigations using AISI 4140 steels, have

shown that austenitizing temperature has an

inﬂuence on the TME phenomenon. High aus-

tenitizing temperature was found to be asso-

ciated with more pronounced TME, favoring

brittle failure modes, even in specimens showing

virtual absence of phosphorus segregation.

These investigations found that high austenitiz-

ing temperatures increase carbide dissolution in

austenite, apparently due to more intensive car-

bide precipitation and growth during tempering

(Ref 23).

When TME appears, there is no heat treat-

ment to reverse the effect, other than to reaus-

tenitize and quench the steel, then temper

in a temperature range where TME does not

occur. TME is also called one-step embrittle-

ment.

Temper Embrittlement. The second type

of embrittlement from tempering at high tem-

peratures is known as temper embrittlement

(TE). Temper embrittlement occurs when tem-

pering in the high-temperature range of 450 to

600



C. It is not a major problem because it may

be avoided simply by quenching from the tem-

pering temperature. This embrittlement can be

reversible under high-temperature tempering.

When steel that has undergone TE is heated

to a temperature above 600



C and then cooled

very quickly, its impact is restored. Therefore,

such brittleness is termed reversible. Temper

embrittlement is also called two-step embrittle-

ment, because two tempering treatments or a

heating step and a cooling step are sometimes

required to induce embrittlement. Figure 18

shows different situations where embrittlement

may or may not appear.

Fig. 16 Illustration of toughness loss after tempering in the embrittlement range. Source: Ref 17

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Temper embrittlement of this category of

steel takes place due to impurity segregation at

the grain boundaries and, ﬁnally, decohesion of

the grain boundary. This leads to intergranular

fracture morphology. It is signaled by a material

toughness loss. It is pointed out that grain-

boundary segregation depends on the alloying

elements of the steel.

Carbon steels with less than 0.5% Mn are not

prone to reversible TE. The phenomenon can

only appear in alloy steels. Alloying elements

may have a different effect on steel after tem-

pering at the steel propensity to TE. Unfortu-

nately, the most widely used alloying elements,

such as chromium, nickel, and manganese,

promote TE. When taken separately, they pro-

duce a weaker effect than in the case of

combined alloying. The highest embrittlement

effect is observed in chromium-nickel and

chromium-manganese steels.

A fundamental fact is that alloy steels of very

high purity are utterly unsusceptible to TE,

which is caused by the presence of various

impurities, such as phosphorus, antimony, and

arsenic, in commercial steels. Relatively small

amounts of these elements, on the order

of 0.01% or less, have been related to TE (Ref

24–27).

Steels made of pure elements do not become

brittle after tempering as can be seen in Fig. 19.

Molybdenum is one of the main alloying

elements in many low-alloy steels and is an

effective method of alleviating TE. Small addi-

tions of molybdenum (0.2 to 0.3%) can diminish

TE, while greater additions enhance the effect.

The mechanism of actuation is related to

molybdenum segregation (equilibrium and non-

equilibrium) during quenching, and the quen-

ched-in vacancies play a role in the temper

process (Ref 28, 29).

As mentioned earlier, in chromium-

molybdenum steel, phosphorus is the major

embrittling element. However, the kinetics

Fig. 17

Charpy tests at room temperature as a function of tempering temperature for 0.6% C steel with different silicon contents.

Source: Adapted from Ref 16

Fig. 18

Schematic diagram showing thermal cycles leading

to the reversal of temper embrittlement and the de-

embrittlement heat treatment

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of phosphorus segregation in this steel are dif-

ferent from that in steels due to molybdenum.

Molybdenum has a strong effect in lowering the

solubility of phosphorus in iron. This is inter-

preted as due to a molybdenum-phosphorus

scavenging reaction; so, if molybdenum is free

in solution, it appears to either prevent phos-

phorus from segregating to the grain boundaries

or to reduce the brittleness potency of phos-

phorus at the grain boundary. Although molyb-

denum is an effective element to reduce the

susceptibility for TE, the precipitation of

molybdenum as carbide must be taken care of.

To avoid that, it is observed that vanadium is

added in this grade of steel. Vanadium is a strong

carbide former compared to molybdenum and

chromium. Vanadium initially forms MC-type

carbides; this changes the molybdenum-to-

carbon and chromium-to-carbon ratios. The

increase in molybdenum-to-carbon ratio is

favorable for Mo

C-type carbides, and that of

the chromium-to-carbon ratio is favorable for

carbides in this grade of steel. These

changes in carbide formation sequences basi-

cally slow down the precipitation of molybde-

num as carbides. When the molybdenum in solid

solution in the ferrite matrix is fully removed,

phosphorus is free to segregate, and the material

thus becomes embrittled (Ref 28, 29).

Failure analysis of high-temperature studs

concluded that the failure related to TE could be

delayed due to the presence of molybdenum and

vanadium. However, reductions of phosphorus

contents in the steel and a grain size of

approximately 10 mm could reduce the tendency

for brittle fracture (Ref 30).

Table 4 shows elements that can segregate to

the former austenite boundaries that are now

ferrite boundaries (Ref 2).

This segregation was shown in a conclusive

form through Auger electron microscopy per-

formed on intergranular fracture surfaces. This

technique allowed the exact determination of

atomic specie concentration segregated in con-

ﬁned boundaries with approximately 10 A

depth. This fraction varies from 0.3 to 2.0 for

steels when these matrix elements are lower than

0.1%.

Although there are similarities in the effects

of the two types of embrittlement from a prac-

tical standpoint, TME and TE are separable into

two different phenomena because they occur in

two different ranges of temperature and also

because TME is a much more rapid process than

TE. The former develops during the ﬁrst hour

of the normal tempering period, and it is inde-

pendent of section size and/or cooling rate after

tempering. Second, TE needs many hours to

develop, and it is an important concern, mainly

for heavy sections that are tempered at higher

temperatures (out of embrittlement range) and

cooled very slowly over a period of many hours

through the critical range of embrittlement.

Temper embrittlement develops in relatively

soft structures because of the high tempering

temperature required to develop it (Ref 12).

Temper embrittlement is clearly associated

with the high-angle ferrite grain boundaries

(former austenite grain boundaries). Thus, it is

expected that austenite, which has the necessary

concentration of segregants to allow the devel-

opment of TE when aging tempered martensite,

Fig. 19

Inﬂuence of phosphorus and antimony on room-

temperature impact energy as a function of temper-

ing temperature in a Ni-Cr-Mo steel. Arrow shows the laboratory

alloy. Source: Ref 19

Table 4 Chemical elements that can segregate

to the grain boundary

Group

IV B V B VI B

CNO

Si P S

Ge As Se

Sn Sb Te

... Bi ...

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will develop TE upon aging a bainitic micro-

structure (Ref 22).

In fact, bainitic microstructures, present in

many kinds of low- and medium-alloy structural

carbon steels, can also develop embrittlement

when tempered (TBE). Tempering at low

temperature has only a small effect on the

cementite size and morphology. Consequently,

the low-temperature embrittlement phenomena

are not found in conventional bainitic micro-

structure (Ref 31).

However, for higher temperatures, embrittle-

ment can be developed. The tempering tem-

perature range for TBE depends on the chemical

composition and microstructure of the steel

(Ref 32).

Mechanisms that provoke TBE are similar

to TME, that is, precipitation of cementite

formed by the decomposition of retained auste-

nite ﬁlm at boundaries of bainitic ferrite laths or

even around the martensite-austenite islands

during soaking at the tempering temperature

(Ref 33, 34).

Secondary quenching (formation of marten-

site during cooling from the tempering tem-

perature) and transformation of retained

austenite into martensite because of plastic

deformation after tempering are other possible

reasons for the occurrence of TBE. The inﬂu-

ence of impurity segregation as well as the

as-quenched microstructure on the TE phe-

nomenon in low-alloy steel was veriﬁed. The

susceptibility of lower bainite or martensite to

embrittlement was investigated. It was shown

that martensitic microstructures are more sus-

ceptible to intergranular fracture than bainitic

microstructures (Ref 35).

Mechanical Tests for TE Determination.

As previously mentioned, TE is characterized

by decreasing the impact resistance by heating

and maintenance in the critical temperature

where the phenomenon occurs (450 to 600



or by slow cooling through this temperature

range. The impact resistance can be recovered

by heating up the embrittlement temperature

range (4600



C), followed by a quick cooling.

The TE apparently does not have an inﬂuence on

hardness, yield strength, and elongation mea-

sured in conventional tension tests. The same

can be said about fatigue resistance, although

this phenomenon produces a drastic reduction in

toughness fracture and also increases the tran-

sition temperature, mainly for steels with nickel,

chromium, and magnesium additions.

Experimental methods that are more adequate

for TE studies consider the kind of load, the test

temperature, and also the velocity of the

imposed deformation. Figure 20 shows the test

results of AISI 1340 steel, performed in different

tempering temperatures.

Figure 21 indicates that the embrittlement

phenomenon of tempered martensite is just

sensitive to determinate test conditions. The

room-temperature tensile test, using unnotched

proof tests, does not present changes in terms of

reduction of section and tension rupture in the

Fig. 20

AISI 1340 steel quenched and tempered (at different temperatures). A, unnotched sample at room-temperature test.

B, unnotched sample at very low-temperature test. C, notched sample at room temperature. D, Charpy impact test at room

temperature. Source: Ref 19

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embrittlement temperature range. It is possible

to make an analogy with the embrittlement

phenomenon that occurs in ferritic stainless steel

having chromium above 15% when exposed

to temperatures at approximately 475



C. In

some cases, this embrittlement is not detected

in tension tests, and almost no modiﬁcations

occur in elongation during the test, but a strong

variation in impact resistance occurs, decreasing

the value compared to the same unembrittled

alloy steel, as seen in Table 5 (Ref 36).

In examining Fig. 21, it is possible to once

again realize the importance of the kind of test

used to verify the embrittlement phenomenon.

As seen in Fig. 20, a common method of

detecting loss of toughness during tempering is

the impact test, which measures the energy

absorbed in fracturing a sample in a speciﬁed

impact loading for a speciﬁed specimen size and

geometry. Results can be reported as impact

energy (for fracture) at a speciﬁc temperature or

impact energy as a function of the test tem-

perature. It is possible also to give the impact

transition temperature, which is the temperature

where the impact energy-temperature curve

changes from a relatively high value to a rela-

tively low value. This temperature is often taken

as the temperature at the inﬂection point into the

impact energy-temperature curve, as shown in

Fig. 21 (Ref 22).

In Fe-Ni-C steel, the TME associated with

the formation of grain-boundary cementite was

observed only when the impact test was per-

formed below the critical test temperature

(approximately 40



C). Impact toughness is

controlled by the intergranular fracture below

this critical temperature, while it is controlled

by transgranular fracture (i.e., the matrix

toughness) above this critical temperature.

Temperature is an important test parameter to

determine embrittlement during tempering

(Ref 37).

Temper embrittlement is not detected by

simple plots of impact energy versus tempera-

ture. However, it is detected by more extensive

impact testing that measures variation of the

impact transition temperature with the temper-

ing temperature. In order to determine the

maximum embrittlement temperature and the

boundaries of the brittleness to reversible TE,

the secondary tempering method is recom-

mended (Ref 38).

It is often possible to tell if a steel has failed

because of one of these embrittlement problems

by examining the fracture surface. A grain-

boundary fracture mode is characteristic of both

types of embrittlement. However, in steels with

Fig. 21

Transition of the fracture behavior of two hypothetical steels in two situations. A, tough. B, embrittled. Source: Adapted from

Ref 22

Table 5 Mechanical properties of ferritic

stainless steel at room temperature under two

different heat treatment conditions

Mechanical

properties

Annealed

at 800 °C

Annealed at 800 °C

and tempered at

450 °C for 4 h

, Kgf/mm

58.4 59.7

, Kgf/mm

33.0 40.0

A, % 22.8 23.6

Impact,

Kgm/cm

(J)

12 (94.5) 1.4 (11.0)

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extremely low levels of phosphorus and sulfur,

TME can display a cleavage surface.

Although traditionally, TME is usually

detected by Charpy tests, a series of experiments

were carried out on three commercial steels

to explore the possibility of characterizing TME

by macro- and microhardness tests. Results

indicate distinct hardness peaks in two steels and

an inﬂexion in the other at approximately the

TME temperature. These experiments are based

on the fact that TME is associated with im-

purity segregation, and microhardness measure-

ments have elucidated such segregation effects

(Ref 39).

Interaction of the TE phenomenon with

Hydrogen Embrittlement. Hydrogen embrit-

tlement (HEM) occurs when high-strength steels

absorb an excessive amount of hydrogen in a

variety of environments. Hydrogen can be

introduced into the material in service or during

materials processing, for example, in the pre-

sence of hydrocarbons or hydrogen sulﬁde or

during pickling in acids, plating, welding, and

heat treatment (Ref 40).

When tensile stresses are applied to a hydro-

gen-embrittled component, it may fail prema-

turely. Hydrogen embrittlement failures are

frequently unexpected and sometimes cata-

strophic. An externally applied load is not

required, because the tensile stresses may be due

to residual stresses in the material. The threshold

stresses that cause cracking are commonly

below the yield stress of the material.

High-strength steel, such as quenched and

tempered steels or precipitation-hardened steels,

are particularly susceptible to HEM.

Tensile stresses, susceptible material, and the

presence of hydrogen are necessary to cause

HEM. Residual stresses or externally applied

loads resulting in stresses signiﬁcantly below

yield stresses can cause cracking. Thus, cata-

strophic failure can occur without signiﬁcant

deformation or obvious deterioration of the

component. Very small amounts of hydrogen

can cause HEM in high-strength steels.

It is believed that hydrogen is likely to be

trapped by structural defects due to its extremely

low solubility in the iron lattice. The most

commonly recognized structural defects that

have a strong hydrogen trapping effect are grain

boundaries, dislocations, carbides, and micro-

voids. In the case of carbides, their size is related

to the hydrogen trapping effect. For the boron-

bearing steel tempered to 1050 and 1300 MPa,

ﬁne cementites having an effect on hydrogen

trapping were found. Lower susceptibility to

HEM was found when this steel was tempered

to a lower strength level, due to delayed onset of

brittle intergranular fracture (Ref 41).

The fracture mechanism produced by this

embrittlement is not simple, because it is related

to the imposed load as well as the hydrogen

amount in the sample. Tests to determine this

embrittlement process are frequently static tests,

taking many hours or even many days to obtain

the results. This kind of test is sometimes called

static fatigue. Figure 22 shows the relation

between the crack propagation mode for differ-

ent load levels.

Hydrogen embrittlement susceptibility is

related to the hardness of the ferrous matrix,

and a very small amount of hydrogen can be

deleterious in the case of high hardness of the

matrix. Under the inﬂuence of hydrogen, the

resistance to fracture or crack growth of steels

is greatly reduced in gaseous hydrogen or

hydrogen-containing environments. Thus, some

experiments with 4140 steel under different heat

Fig. 22

Illustration of a crack growing at different load levels.

Source: Adapted from Ref 19

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treatments were carried out. Experiment results

indicated that the acceleration of the crack

growth in hydrogen was more pronounced for

quenched and tempered conditions when the

tempered temperature range coincided with the

TME range of the 4140 steel (230 to 370



C).

At higher tempered temperature (approximately

550



C), the inﬂuence of hydrogen became

insigniﬁcant. Specimens in the austempered

condition have the best performance on the tests,

as seen in Fig. 23 (Ref 42).

The effect of the microstructure on HEM was

investigated for a low-carbon (Mn-Si-Cr) steel.

Microstructure formed by bainite and martensite

has better behavior than tempered martensite in

intermediate-temperature tempering. At the

same strength level, the impact energy of the

mixed microstructure is 17% higher than only

tempered martensite. It is attributed to the fact

that bainite-martensite delays the TME onset

(Ref 43).

Similar work was performed by Lantsman

et al. Experiments with cadmium-plated

65S2VA steel springs were carried out. In this

process, the steel absorbs hydrogen, which leads

to HEM. Results show that the susceptibility of

HEM has a strong dependence on the pre-

liminary heat treatment and structure. With an

identical hydrogen content, the austempered

steel, with lower internal stresses, will fracture

under a higher load than the steel subjected to

standard quenching and tempering (Ref 44).

Interaction of the TE Phenomenon with

Liquid-Metal Embrittlement. The presence of

select metallic specimens on the surface of

various alloys can provide for a very detrimental

reaction under load, known as liquid metal

embrittlement (LME), causing brittle fracture by

intergranular cracking. LME is the reduction in

elongation to failure that can occur when nor-

mally ductile metals are stressed while in contact

with liquid metals. Failure of components rela-

ted to LME is less common than failures caused

by other processes, such as fatigue, HEM and

stress-corrosion cracking, but a signiﬁcant

number of industrial failures related to LME do

occur (Ref 45–47).

LME should depend on time of contact with

liquid metal while the solid is stressed. When

this occurs, the solid metal fails instantly,

because the ﬂow of liquid metal into the crack tip

during crack growth plays a signiﬁcant role (Ref

48). The liquid metal affects the fracture beha-

vior at the tip of the crack, reducing the critical

stress intensity for fracture and altering the

micromechanism of fracture at the crack tip

(Ref 49).

Sources of the aggressive elements vary,

including unintentional or accidental exposures,

for example, during fabrication or service when

there is intimate contact between the structural

and embrittling metals and also when there are

tensile stresses above a threshold value. Even-

tually, small amounts (0.1 g) of embrittling

metals can produce extensive cracking by LME

(Ref 45).

Plain carbon and low-alloy steels may be

embrittled by exposure to liquid lead, cadmium,

brass, aluminum bronze, copper, zinc, lead-tin

solders, and lithium (Ref 50, 51). Other metal

alloys also have susceptibility to LME. Table 6

lists embrittling environments for some com-

mon structural materials (Ref 45).

Fig. 23

Impact values of 4140 specimens after various

types of tempering treatments. Source: Adapted from

Ref 42

Table 6 Examples of embrittling metal

environments for some structural materials

Structural materials Embrittling environments

High-strength martensitic steel Hg, In, Sn, Pb, Cd, Zn, Li, Cu

Austenitic stainless steel Zn, Cu, Li

Titanium alloys Hg, Cd, Ag, Au

Aluminum alloys Hg, Ga, In, Sn, Pb, Cd, Zn, Na

Copper alloys Hg, Ga, Bi, Zn, Li, Sn, Pb, In

Zirconium alloys Hg, Cd, Cd-Cs, Zn

Nickel alloys Hg, In, Li, Zn, Ag

Magnesium alloys Na, K, Rb, Cs, Zn

Source: Ref 45

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Industrial environments provide diverse

sources of embrittling metals, for example,

during application of coatings or during service

if temperatures and stress are sufﬁciently high

and there is close contact with the substrate, as

mentioned previously.

During soldering, brazing, or welding, LME

can be developed if the residual or assembly

stresses are sufﬁciently high. Metal lubricants

and overheated bearings are other examples that

can provoke LME (Ref 45). Literature presents

various cases of failures associated with LME,

such as failures in gas turbines, presented by

D.W. Cameron (Ref 50). In this case, high

temperatures inherent in the gas turbine aggra-

vated the phenomenon.

Research has shown that the martensitic steel

91 (9% Cr, 1% Mo) is prone to LME by liquid

lead when some conditions are fulﬁlled (Ref 52–

54). Quenched and tempered steels are suscep-

tible to lead embrittlement, and the conditions to

develop this phenomenon are:

 Presence of either external or internal lead in

the steel

 Tensile loading

 Temperature between 200 and 480



If any one of these conditions is not observed,

LME is avoided (Ref 12).

There are a great number of possible

mechanisms for this embrittlement. In a general

way, it occurs in metals when they are in contact

with low-melting metals and can happen even

when the stresses occur at lower temperatures

than the melting point. LME is not typical, but it

is an important failure mechanism, and indus-

trial awareness of potential problems is still

limited.

Case Studies

Case Study I: Grinding Cracks. The origin

of grinding cracks may be related to low efﬁ-

ciency of the cooling system, microstructure and

material cleanliness, and also excessive rate of

material removal per pass. The heat generated

due to these parameters can produce visual

characteristics on the worn surface, as seen in the

cam shaft surface shown in Fig. 24. This com-

ponent (AISI 5160) was induction hardened and

tempered.

The surface cam shaft heat effect can be

studied by metallography and microhardness

determinations (100 g load) in a cross section

containing the worn surface. Figure 25 presents

this microstructure with a low-depth white layer

(approximately 50 mm) containing the lower

indenter marks, which means fresh martensite

(indentations 1 to 3). The fresh martensite has

different chemical reactivity than tempered

martensite or even other austenite decomposi-

tion products. Because of this the overheated

surface can be identiﬁed by special chemical

etching methods, for example, cooper sulﬁde.

Table 7 presents the hardness values.

Fig. 24 Worn surface with excessive generation of heat (arrow)

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Going deeper in the case (right side of Fig. 25)

close to the white layer, a darker etched zone

(indentations 4 to 9) has lower hardness (bigger

indentations), which means that the temperature

was sufﬁcient to cause tempering and reduce the

local hardness. Indentations 13 to 21 (400 to

1000 mm) show higher hardness (Table 7) rela-

ted to the process tempering operation, where

hardness values are close to the martensite as-

quenched condition.

This condition is deleterious because worked

surfaces are more prone to grinding cracks.

One example can be seen in Fig. 26. The

observed grinding cracks are small and difﬁcult

to detect by nondestructive tests as well as

optical techniques classiﬁed as macrographic

techniques. In this case, they could be seen only

through microscopy techniques applied directly

on the worn surface or using replica methods.

These defects must be avoided, because they

may increase the precore failure probability.

Machining operations must be carefully per-

formed, but surface compressive residual stres-

ses may contribute to avoiding the grinding

cracks.

Case Study 2: Transgranular and Inter-

granular Crack Path. Intergranular cracks are

frequently related to the TE phenomenon.

However, during low-temperature tempering,

very close to the embrittlement range, a pre-

dominantly intergranular crack path was found

in an AISI 5160 commercial steel after a tem-

pering operation at 2000



C (Ref 55).

Fig. 25 Microstructure close to the worn surface. Etched with 2% nital

Table 7 Microhardness values of the worn cam

shaft cross section

Identation Depth, mm Hardness, HV

0.1

Average

1 40 782 ...

2 35 803 789

3 35 782 ...

4 95 433 ...

5 92 433 440.3

6 90 455 ...

7 155 520 ...

8 150 493 525.3

9 160 563 ...

10 220 592 ...

11 220 606 613.3

12 225 642 ...

13 405 642 ...

14 405 782 704.7

15 405 690 ...

16 610 858 ...

17 605 690 770

18 625 762 ...

19 1000 724 ...

20 1000 772 766.3

21 1000 803 ...

22 Base material 256 ...

23 Base material 230 245.3

24 Base material 251 ...

Fig. 26

Grinding cracks on AISI 5160 steel cam shaft after

induction hardening and low tempering (high-hard-

ness tempering). Original magniﬁcation: 200·

304 / Failure Analysis of Heat Treated Steel Components

Name ///sr-nova/Dclabs_wip/Failure_Analysis/5113_285-309.pdf/Chap_09/ 18/8/2008 3:35PM Plate # 0 pg 304