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

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This is common for hot work tools due to the

low as-quenched hardness and the high hard-

enability of this class of tool steels. For example,

the use of improper temperatures may lead to the

speciﬁed hardness, which can even be homo-

geneous, but the mechanical properties, such

as toughness or tempering resistance, can be

deeply affected. A poor practice during quench-

ing can lead to high brittleness with no observed

changes in hardness or strength. These two as-

pects are described in this section.

Before continuing, a useful recommendation

should be made. Due to the high value of die-

casting tools and their high productivity, a strong

effort has been made for improvements in this

ﬁeld. An important recommendation for H13

tool steel was written by the North American

Die Casting Association (NADCA) (Ref 32).

This recommendation provides important

Fig. 28

Impact toug hness and hardness as a function of

tempering temperature. Retained austenite content is

also shown. Notice the hash-marked area, indicated as a temper

embrittlement region, where very low toughness is observed; this

region coincides with the peak hardness. Source: Ref 24

Fig. 29

Example of the importance of tempering resistance

instead of initial working hardness. (a) Hot forging

punch showing wear and cracks as the normal failure condition.

For maximum wear resistance, initial hardness was established at

56 HRC for H13 tool steel. However, the end-life mechanism was

related to hardness reduction in the working (heated) areas, as

shown in (b). This grade was substituted by a higher-molybdenum

grade, brand name VHSUPER, with higher tempering resistance,

as shown in (c) by the longer times necessary for hardness

decrease. The substitution lead to 50% longer tool life. Courtesy

of Villares Metals

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information for quality assurance of tool steel

and also for applied heat treating. It is very

helpful for die-casting applications as well as for

quality analysis of AISI H-series steels applied

to other processes.

Incorrect Hardening and/or Tempering

Temperatures. To analyze the effect of incor-

rect temperatures on hot work steel properties,

several studies were conducted in the labora-

tory (Ref 12). Toughness was evaluated by

unnotched impact testing, according to the

NADCA procedure (Ref 32). Three conditions

were simulated. To show the effect of low

hardening temperature on tempering resistance,

conditions 1 and 2 were evaluated:

 Condition 1: standard condition for H13

grade, with hardening temperature at

1020



C and double tempering at 610



(for 2 h at temperature), leading to 45 HRC

 Condition 2: Low hardening temperature (at

890



C), simulating a condition for furnaces

that reach up to 900



C. For this condition,

tempering should be reduced to 250



 Condition 3: higher hardening temperature,

increasing from 1020



C to 1150



C. Such

temperature is currently used for high-speed

steels heat treated to lower hardness. For 45

HRC, tempering was slightly increased, to

640



C, also twice for 2 h.

Figure 32 evaluates the reduction in toughness

promoted by inadequate heat treating conditions,

as well as the respective microstructures. The

low-tempering situation, condition 2, causes a

substantial reduction in toughness (40%) as well

as a loss in tempering resistance (the hardness

decrease was six times higher than expected).

Toughness reduction in this condition, is caused

by incomplete austenite transformation, causing

a heterogeneous microstructure (Fig. 32c), as

well as by the low tempering temperature, which

does not promote adequate martensite stress

relief. On the other hand, the decrease in

Fig. 30

Examples of failures caused by excessive hardness. (a) Tool made of DIN 1.2714 tool steel (similar composition to AISI 6F3

and L6) that fractured after a short life. For this tool, the hardness was expected to be approximately 40 HRC, but a sample was

analyzed and found to be 50 HRC. Arrows indicate cracking location and cracking initiation site. (b) Microstructure showing light areas,

indicating excess retained austenite and untempered martensite, another indication that low tempering temperature was employed and/

or only one tempering treatment. This led to high hardness as well as a brittle microstructure. (c) H13 forging die that cracked prematurely

(arrows). Hardness was measured at 52 HRC but expected values were approximately 44 HRC. The excessive hardness was caused by

short tempering times and low temperatures. (d) Typical H13 microstructure tempered at low temperatures. Etched with 4% nital. It is

lighter than usual (compare to Fig. 32b) due to low-temperature tempering, which causes poor pr ecipitation of alloy carbides and thus

enhances corrosion resistance during etching. Courtesy of Villares Metals

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tempering resistance is caused by the insufﬁcient

dissolution and reprecipitation of alloy carbides

(poor secondary hardening), which is the main

strengthening mechanism in hot work tool steels

at high temperatures.

The most intense embrittlement was pro-

duced by condition 3—too high hardening

temperature. In this situation, very coarse aus-

tenite grains are produced (Fig. 32d), leading to

increased grain-boundary embrittlement. When

high hardening temperatures are used, precipi-

tation of proeutectoid carbides on grain bound-

aries intensiﬁes, causing a great reduction in

toughness, as observed (90% lower).

Some sources studied the use of higher hard-

ening temperature as one way to improve ther-

mal fatigue (Ref 22, 33). In fact, the increase

in hardening temperature leads to more dis-

solution of alloy carbides, rich in vanadium and

molybdenum, which increases the content of

alloy elements in solid solution and enhances

secondary hardening. A simple observation of

this is shown in condition 3, where tempering

should be increased 30



C to attain the same

hardness as the usual hardening condition.

In terms of tempering resistance, 30



C (54



the dislocation in tempering curve indicates a

substantial increase in tempering resistance,

because temperature effect is exponential to

time effect in tempering conditions. This phe-

nomenon also explains some advantages found

in specimens austenitized at 1100



C compared

to 1020



C (Ref 22). However, modiﬁcations in

hardening temperatures are rarely possible in

practical (industrial) conditions. Increasing

the hardening temperature deeply affects the

precipitation behavior on grain boundaries dur-

ing quenching, causing intense embrittlement.

Figure 33 shows this effect (note the dashed

lines), but a full explanation is given in the fol-

lowing section (especially regarding Fig. 36).

Before continuing, it is interesting to show a

case of failure caused by excessive hardening

temperature. Figure 34 presents such a case—a

tool that cracked after low performance. It is

easy to see the grain-boundary crack propaga-

tion, caused by the coarse grain size as well as by

precipitation of carbides on grain boundaries.

This is denoted by the preferential and strong

etching of austenite grain boundaries.

Slow Cooling during Quenching. Reach-

ing the ﬁnal hardness in tool steels is quite a

Fig. 31

Heat checking resistance (lower readings indicate higher resistance) as a function of unnotched impact toughness and

hardness of H13 steel. Heat checking is evaluated by the photographs on the left; the rating is calculated by adding the

column representing the largest cracks (leading) and the column representing the severity of the cracks (network). See text for discussion of

these results. Source: Ref 31

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simple task for quenching. Although there are

some exceptions, tool steels normally have

high hardenability, with the as-quenched hard-

ness attained even if improper procedures are

applied. However, hardness, although necessary,

is not sufﬁcient for the high required performance

of tool steels, as already shown. In this context,

quenching should be considered as a process for

promoting the required mechanical properties,

not only for attaining a speciﬁed hardness.

For this reason, consider Fig. 33 again. For

both hardening temperatures, hardness higher

than 500 HV (~49 HRC) is obtained within a

wide range of the continuous cooling tranfor-

mation (CCT) curve. Considering the hardness

reduction after tempering, it is therefore quite

simple to obtain the ﬁnal hardness with different

quenching practices, even using air quenching.

However, two further important aspects should

be considered. First, the dashed lines in the CCT

diagram (Fig. 33) indicate formation of proeu-

tectic carbides. As mentioned in the previous

section, if austenitizing temperature increases,

more alloy elements go into solid solution by

carbide dissolution. In cooling, the process is

reversed, and such carbides tend to form again.

This happens by precipitation in high-energy

areas, the most important being the grain

boundaries. The result is a ﬁlm of carbides

between grains, which weakens the interface

and promotes failure (Ref 35). Such a phenom-

enon is marked by two characteristics: a strong

etching at austenite grain boundaries (because

the interfaces of carbides and steel are regions

(b) condition 1 (c) condition 2 (d) condition 3

Fig. 32

Laboratory simulation of adequate and inadequate heat treating conditions for AISI H13. The ﬁrst situation (condition 1) is the

recommended heat treatment: hardening at 1020



C, followed by two tempering treatments at high temperature. In this case ,

45 HRC was desired, and thus tempering was performed at 610



C. Condition 2 involves a very low hardening temperature , where

austenitizing was done at 890



C. To reach 45 HRC, specimens were heat tempered at 250



C. In addition to toughness reduction, the

heat treating condition caused reduction of tempering resistance. Condition 3 describes a situation with excessively high hardening

temperature (1150



C), with tempering done at 640



C to attain 45 HRC. In (a), the impact toughness is presented, and in (b) to (d), the

microstructure relative to each condition is shown (same magniﬁcation; etched with 4% nital). Source: Ref 12

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more prone to corrosion) and, in stronger cases,

intergranular failure. As shown in Fig. 33, an

increase in austenitizing temperature causes

dislocation of the dashed lines to the left,

indicating stronger precipitation, even when

high cooling rates are applied. Secondly, a slow

cooling also affects the previous microstructure,

forming bainite instead of martensite. Although

the precipitation necessary for secondary hard-

ening is practically not affected (Ref 33),

modiﬁcation of the initial microstructure, from

martensite to bainite, also reduces toughness

Fig. 33

Continuous cooling transformation diagrams for H13 tool steel austenitized at 1030



C (1885



F) (top) and 1100



(2010



F) (bottom). Note the dislocation of the dashed line, indicating more pronounced proeutectic carbide precipitation on

grain boundaries for the high austenitizing temperature condition. Source: Ref 34

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(ref 29, 33 and 36). The effect of decreasing the

cooling rate is thus embrittlement for both

mechanisms.

Therefore, both bainite and carbide pre-

cipitation on grain boundaries must be avoided

by preventing slow cooling in quenching. This is

important advice for tools sensitive to failures

caused or assisted by cracks. One example

is presented in Fig. 35 for hot forging dies

that failed after short service time. The

microstructural analysis showed coarse grain

sizes and strong precipitation on grain bound-

aries (Fig. 35b, c), illustrating the interaction

between the two effects. As a ﬁnal result, strong

embrittlement occurs (Fig. 35d) as well as a

clearly intergranular fracture (Fig. 35e).

Typically, hot work tool steels were oil

quenched, but today (2008), vacuum heat treat-

ing with pressurized nitrogen quenching has

become very popular. In this treatment, cooling

rate control is rather critical, since it is related

to nitrogen pressure and gas circulation as well.

If too strong and heterogeneously applied,

cooling may lead to strong distortion or even

quenching cracks. On the other hand, grain

boundary embrittlement occurs easily if the

cooling rate is too slow. A guide for evaluating

tool heat treating quality is described by the

NADCA (Ref 32), including the use of coupons

for destructive testing after heat treating as well

as advice for vacuum hardening.

Another important issue is the step in which

heat treatment should be applied. With the

advance in machining technology, the feasibility

for machining in higher hardness has increased;

machining hot work dies up to 50 HRC is rather

common by means of high-speed machining

technologies (high cutting speeds with low

feed). Consequently, it is common, mainly in

forging dies, to machine from prehardened

blocks. However, the probability of embrittle-

ment increases as the section size of the tool

increases. Figure 36 shows the effect of section

size and austenitizing temperature on the tough-

ness of H13 steel. The tendency for toughness

loss is evident when larger sizes or higher

austenitizing temperatures are used, because

they are directly related to the grain-boundary

embrittlement effect and are also affected by

bainite formation. Even if the quenching process

uses a strong cooling medium, large tools are

unavoidably sensitive to embrittlement in core

regions. Therefore, heavy-section tools with

deep engravings should be heat treated only after

rough machining to avoid embrittlement of tool

working regions. The most important example

in this ﬁeld is die-casting dies. They are usually

heat treated only after machining to improve

Fig. 34

Die failure caused by excessive hardening tem-

perature. Two tools were analyzed: one made of

VHSUPER steel (commercial brand name) and the other of AISI

H13. (a) One of the tools cracked in the position denoted by the

arrow, where a sample was cut for analysis. (b) Typical micro-

structure from H13 tool and (c) from VHSU PER tool with 54 HRC

(45 HRC was expected). Note the coarse grain size, approxi-

mately ASTM 3 to 4. For these grades, grain size is expected to be

approximately ASTM 7 to 10. Courtesy of Villares Metals

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(a)

Fig. 35

Example of die failure in a hot forging die caused by coarse grain size and strong precipitation of proeutectoid carbides on

austenite grain boundaries. (a) Aspect of the tool. (b) and (c) Microstructure showing the coarse grain size (approximately

ASTM 4; expecte d ASTM 8 to 10), marked by preferred etching on carbides present at grain boundaries and the coarse martensite laths.

Samples were taken from the tool midradius and analyzed regarding (d) impact toughness in the as-received condition and after new heat

treating to the same hardness and (e) fracture of impact-tested specimens (for the initial condition—as-received) by scanning electron

microscopy. Note the strong increase in toughness after new heat treating, indicating the deleterious effect of carbide precipitation on

grain boundaries, producing intergranular failure in impact specimens. Courtesy of Villares Metals

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toughness, because die resistance to heat check-

ing is directly related to this property (Fig. 33).

Even if the quench is applied after machining,

it is still important to control the Quench cooling

rate (avoiding too slow conditions), due to the

possibility of embrittlement of surface regions,

which are the working areas of die-casting dies

(Ref 32).

Failures of Nitrided Tools. Nitriding is

commonly used for several tools and dies. For

example, hot extrusion dies are typically nitrided

for all uses. They normally work for production

of parts in aluminum or other nonferrous alloys,

mainly for construction applications. Different

from other hot working processes, extrusion is

continuous and involves constant ﬂow between

the conformed alloy, and the tool steel. This

enhances the wear condition, which is the usual

end-life mechanism. Nitriding considerably

improves surface hardness and, consequently,

wear resistance, this being the reason for its

application in virtually all extrusion dies. The

same approach is applied in hot forging tools.

Nitriding is currently used in several tools and

dies to improve wear resistance. However, as

hardness increases, surface toughness decreases,

and for deep tools, especially those prone to

cracks, nitriding is prevented. In die casting,

nitriding is used in some cases, claiming that the

increase in surface hardness tends to increase the

initiation of heat checking cracks; however, as

cracks cross the nitriding layer, it has no effect at

all. Depending on the nitriding layer condition, it

can be harmful to crack initiation.

Normally, the nitride layer in tool steels is 0.1

to 0.3 mm thick. Hardness is higher than in

carbon steels, due to the formation of alloy

nitrides. In hot work steels from the H series,

chromium is very important for this effect, and

the maximum surface hardness approaches 1100

HV. Typical nitrided microstructures present a

diffusion layer and a ﬁne white layer on the tool

surface (Fig. 37a). In H-series tool steels, the

Fig. 36

Charpy V-notch (CVN) results for different heat

treating conditions of H13 tool steel, carried out at

room temperature and at 425



C. Specimens cooled at various

rates, simulating the core of 150 and 300 mm round bars, as well

as an air-cooled 25 mm specimen. Results were tested for dif-

ferent austenitizing temperatures. Toughness reduction is evident

at higher austenitizing temperatures and larger sizes, both related

to lower cooling rates during quenching. Source: Ref 33

Fig. 37

Tool steel surface after nitriding. (a) White and dif-

fusion layers (b) Coarse nitrides precipitated on grain

boundaries. See text for discussion. Courtesy of Villares Metals

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nitriding layer is darker after etching, due to the

depletion of chromium content from the steel

matrix, thus causing a reduction in corrosion

resistance. In carbon or low-alloy steels, the

nitride layer appears lighter in the micro-

structure and can increase corrosion resistance.

In hot work tool steels, it is rather common to

observe coarse nitrides precipitated on grain

boundaries inside the diffusion layer, as shown

in Fig. 37(b). Depending on the intensity, this

embrittles the tool and forms an exceptional

route for crack propagation (Fig. 42). Preci-

pitation of coarse nitrides on grain boundaries

can also be accompanied by a thick white layer

on the surface. This layer is extremely hard

and brittle, because it is uniquely composed of

nitrides. It is therefore a common region for

crack initiation. The combination of a coarse

white layer and grain boundary precipitation of

coarse nitrides commonly leads to premature

crack initiation, with damage to several tools,

especially forging tools prone to cracking or

for die-casting dies. Some examples of failures

assisted by this phenomenon are shown in

Fig. 38. It is thus recommended, for these

applications and in general, that the nitriding

process be controlled in order to avoid both a

thick white layer and grain-boundary nitrides.

Today (2008), this control is usually performed

in computer-controlled gas nitriding and plasma

nitriding processes.

Excessive Heating Causing Tool Fail-

ure. As discussed previously, hardening of hot

work tool steel mainly results from precipitation

of alloy carbides during the tempering treatment.

This phenomenon is also known as the fourth

stage of tempering and occurs after the mod-

iﬁcation of martensite and the formation

Fig. 38

Examples of undesirable microstructures encountered on the surface of nitrided tools. For both cases, the core micro-

structures are correct, indicating proper hardening and tempering procedures. (a) Surface and (b) core microstructure of

a nitrided forging tool, showing (in a) the problems of a coarse white layer and nitrides on grain boundaries. (c) and (d) Extracted from a

die-casting die failure analysis, also for surface and core respectively. Note the strong precipitation on grain boundaries in (c), whereas

core regions are quite well heat treated (to approximately 44 HRC), leading to 300 J of unnotched (NADCA) impact strength. Never-

theless, an unexpected failure occurred (after less than 100 shots), caused primarily by improper tool use but also assisted by the nitriding

layer condition.

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of cementite. At low temperatures, typically

below 500



C, the thermodynamic driving force

(free energy) is insufﬁcient for such carbide

precipitation (mainly molybdenum or vanadium

carbides), and no hardness peak is observed

(Fig. 27). As the temperature increases, carbide

formation starts and very small carbides (a few

nanometers) precipitate within martensite laths.

Because they are ﬁne and large in number, these

carbides are proper for obstructing dislocation

movement, causing substantial improvement in

strength. This occurs only at high temperatures,

with precipitation of alloy carbides pertaining

to materials used in hot working, in this case,

hot work tool steels. For higher temperatures,

carbides tend to transform to more stable types or

to coarsen, which increases the size and thus re-

duces the total number of precipitated carbides.

This results in loss of the dislocation blocking

effect, leading to strength reduction. The total

effect is schematically shown in Fig. 26.

Temperature and time affect the precipitation

behavior. From a thermodynamic point of view,

the lowest energy is obtained after alloy carbide

precipitation (reducing the free energy of ele-

ments in solid solution) and if their sizes are

as large as possible (reducing the surface free

energy). Consequently, the higher the time

or temperature, the easier it is for carbides to

become larger, and strength tends to decrease.

In tempering curves, such as those shown in

Fig. 24 and 28, this can be observed for a ﬁxed

time (twice for 2 h) according to temperature.

However, a complete view is given by using

the tempering parameter, as shown in Fig. 39.

This combines the effects of both time and

temperature in only one variable, the parameter,

proportional to time and to a logarithm of the

temperature. Through this mathematical rela-

tion, it can be seen that temperature actually

has the highest effect; however, if excessively

long times are used at appreciably high tempera-

tures, the same effect may occur. For example,

4 h (twice 2 h tempering) at 600



C (1110



parameter = 32,700) leads to ~46 HRC for H11,

which is equivalent to 15 h at 577



C (1070



F),

or 3 days at 550



C (1020



F), or approximately

1 month at 500



C (930



F). That is why

below 500



C, heating has practically no effect.

However, a second situation should also be

considered. Low times at excessively high

temperatures can produce rather important

effects. This is the situation for several tool

applications, where tools are in contact at very

high temperatures but for very short times. For

example, consider a forging tool that reaches

750



C(~1380



F) for a half-second during each

stroke. After producing 1000 parts, this tool will

be exposed to 750



C for a total time of 8.3 min.

Such heating is equivalent to a tempering para-

meter of 35,260 or 2 h at 691



C (1275



F). In

H11, such heating is capable of reducing the

hardness from 46 to less than 30 HRC.

It is in this context that the failures described

in this section should be considered. In some

situations, tools are correctly heat treated, but,

during hot working operation, heating is so

high that strength is reduced and failure is

accelerated. One example, shown in Fig. 40, is

a hot forging punch that usually exhibits low

life if compared to the whole set of tools. It is

made of DIN 1.2885, a highly temper-resistant

hot work steel that has high amounts of

molybdenum and cobalt for improvement of

this property. Even though the analysis showed

a continuous reduction of hardness approaching

the punch tip, which contacted the hot infor-

ging part. Besides the hardness values, it is

interesting to note the microstructural behavior.

Regions far from the hot working areas are

typically tempered martensite, becoming dar-

ker closer to the punch tip. As previously

explained, secondary carbides are not visible

Fig. 39

Tempering curve as a function of the time-tempera-

ture parameter for H11 steel containing 0.40% C,

0.92% Si, 5.09% Cr, 1.34% Mo and 0.52% V. For this curve,

t = time in hours, and T = temperature in



F+460. Source:

Ref 37

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