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

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under optical microscope; however, their effect

on corrosion resistance (and thus on etching

behavior) is evident. High tempering tempera-

tures cause an increase in precipitation and

carbides, leading to reduction of corrosion

resistance due to the depletion of alloy ele-

ments in the steel matrix and to the increased

sites (interfaces) for corrosion. This explains

why the microstructure becomes darker where

heating is more intense and hardness is lower

(this will always happen if the hardening

practice is not changed). In the punch in

Fig. 40, heating is so high that secondary

carbides have become large enough to be

observable by optical microscope. This is

typically the annealed state for this alloy, not

obtained through transformation annealing but

by temper annealing.

The effect shown in Fig. 40 occurs in several

hot working tools. A further example is shown in

Fig. 41 and 42 of a forging die for automotive

valves. After some time in production, wear of

the forging die produced grooves on the surface

of forged products (Fig. 41a), leading to die

substitution. The analysis revealed that wear

was deeply inﬂuenced by thermal fatigue crack-

ing (Fig. 41c, d). By comparing heated and

not heated areas, two further problems were

observed. First, some regions on the die surface

had been heated to high temperatures, leading to

reaustenitizing, rehardening, and the formation

of untempered martensite (Fig. 42a). Second, the

tool surface had serious indications of nitriding

problems, with coarse nitride formed on grain

boundaries and crack propagations through

these regions (Fig. 42c). The combination of

intense surface heating and nitriding led to a

“crazy pattern” of hardness, differing substan-

tially from that encountered in non-heat-affected

areas (Fig. 42b). This was caused by the com-

bination of rehardening and nitrogen diffusion to

the core regions. The end result, of course, was a

deep embrittlement of the die, as observed by its

intense cracking.

In cases similar to those in Fig. 40 to 42, two

possibilities are possible to solve or reduce the

problem. The ﬁrst is to employ one steel grade

with a higher tempering resistance, normally a

higher-alloyed grade; however, this approach

has some limitations. If the situation is only an

Fig. 40 Microstructures of hot forging punch made of DIN 1.2885. Etched with 4% nital for the same time in all microstructures

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increase in tool resistance to softening, changing

the steel grade may be a solution. For example,

this was the situation described in Fig. 29.

However, it should be noted that highly alloyed

grades can also be more brittle, leading to

problems if gross cracking or heat checking

are important issues for tool failure. Also,

the increased alloy is directly related to steel cost,

with an analysis of the cost and beneﬁts of steel

substitution being important. The best way is to

consider the total tool cost for each produced

part, which includes the increase in steel cost

but also its higher performance as well as the

reduction of setup times for tool substitution.

However, in some situations, heating is so

intense during tool application that it is im-

possible to solve the problem only by modifying

the tool steel employed; thus, there is a second

possibility concerning cooling applied during

hot work tooling. Coolants and lubricants are

both important in hot work tooling, in many cases

being the same product. In cases where appre-

ciable regions of tools are reaching more than

700



C or even reaustenitizing, the best practice

is to work on the lubricating/cooling involved in

the process, before considering changes to the

tool steel used. It would be helpful to contact the

lubricant manufacturer and obtain updated

Fig. 41

(a) Forging tools for production of automotive engine valves; the analyzed die was painted in blue. (b) Forged valve and, in

detail, the grooves produced when using worn dies. (c) Three analyzed regions, marked “A”, “B” and “C”. “A” is the main

wear region, but “B” is where a large amount of thermal fatigue cracking is observed. “C” is a region not affected by process heat and used

as a reference. (d) Thermal cracks of region “B”, under scanning electron microscopy. Courtesy of Villares Metals

Failure Analysis in Tool Steels / 347

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(b)

Fig. 42

(a) Microstructure of working regions of valve forging die in Fig. 41 (region A). Note the intense cracking and thick nitrided

layer (double the expected). Also note the surface white layer, which corresponds to brittle untempered martensite, obtained

by rehardening of the tool surface. (b) Hardness proﬁle of this region (A) and another region not affected by heat (C) in Fig. 41(c). (c) Two

pictures of region C, not affected by heat, that have a mechanical crack within the nitriding layer. Note the presence of a coarse white

layer and nitrides on the grain boundaries. The micrographs in (c) refer to the same region, but the one on the right is slightly underfocus for

better observation of nitrides on the grain boundaries. Courtesy of Villares Metals

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information on the possibilities for a speciﬁc

application. Besides the type of lubricant, a very

important issue is the way the lubricant is

applied. In modern processes, robots may be

used for homogeneous application of lubricants.

It is also important to evaluate the possible

changes in tool design to avoid heat during hot

working. The answer is complex and should

consider tool life, process productivity, and the

necessary investments, always aiming for the

minimum cost per produced part.

Conclusion

All theory and data discussed in this chapter

may be summarized by the following points:

 Heat treating of tools and dies may be con-

sidered crucial for their performance. Sev-

eral examples show that success cases as

well as premature failures are often related to

heat treating quality.

 Hardness measurements alone are usually

not the best indication of proper quality in

too steel heat treating. Other properties such

as strength, toughness, and wear resistance,

as well as the microstructural features, are

better indicators as to whether a heat treat-

ment was done correctly or not.

 Assessing these properties is usually difﬁcult

in tools and dies because it would be a

destructive (and expensive) test. Therefore,

guaranteeing the correct procedure is the

best way to assure that the heat treatment of a

tool was performed correctly and thus

enables adequate tool performance.

 In order to apply such correct conditions, the

best approach is to follow the speciﬁcations

for each grade which is normally provided

by the tool manufacturer. For several spe-

ciﬁc points such as design, tool use, and

surface treatments, the examples given with

this chapter may be helpful.

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Case Studies of Steel Component

Failures in Aerospace Applications

Scott MacKenzie, Houghton International, Inc.

STEEL has limited use in modern aircraft, but

it is used in landing gear systems, arrestment

systems for naval aircraft, and for support

equipment. Because of the high strength levels

required (1790 to 2070 MPa, or 260 to 300 ksi,

yield) to maintain justiﬁcation for the use of

steel, these very high-strength steels are prone to

environmental effects (stress-corrosion cracking

and hydrogen embrittlement), ﬂaws created

during manufacturing (laps, seams, machining

gouges, and grinding), and heat treatment (dis-

tortion, decarburization, and quench cracking).

Care is needed to protect and inspect these high-

value and safety-critical steel components. The

following case histories illustrate typical failures

experienced by these high-strength steels.

The case histories in this chapter illustrate

a variety of failure mechanisms. The causative

reasons vary from manufacturing to operational

to environmental. Because of the many different

types of root causes, the failure engineer must be

aware of the manufacturing process, the assem-

bly process, as well as the environment to which

component is exposed to effectively determine

the primary cause of failure.

Failure Analysis of a Catapult

Holdback Bar

This investigation analyzed the failure of a

repeatable-release holdback bar. This bar con-

sisted of a failed strain bar and a failed T-head

(rod end connector). These components failed

during characterization fatigue testing. The

strain bar and the T-head (rod end connector) had

been subjected to spectrum loading. This fatigue

spectrum consisted of two 365 MPa (53 ksi)

tensile loads followed by a single 469 MPa

(68 ksi) tensile load application. The strain bar

failed at 4875 cycles. The T-head failed at 3235

cycles; the desired lifetime for the strain bar and

the T-head was 4500 cycles.

The failed AISI 4340 steel strain bar and the

failed cadmium-plated AISI 4330V steel T-head

were submitted for analysis.

Figure 1 shows the as-received failed strain

bar. The fracture occurred at the aft radius of

the circumferential retainer ring groove located

near a slot in the strain bar. Figure 2 shows

the as-received failed T-head and the location

Fig. 1 As-received strain bar from fatigue testing of a hold-back bar assembly

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Failure Analysis of Heat Treated Steel Components

L.C.F. Canale, R.A. Mesquita, and G.E. Totten, editors, p 351-393

DOI: 10.1361/faht2008p351

www.asminternational.org

of the fracture. Fracture occurred in the run-out

area of internal threads at a distance of appro-

ximately 50 mm (2 in.) from the forward end.

All of the parts were subjected to a magnetic

particle inspection. In the strain bar, two cracks

were found approximately 180



apart in the re-

tainer ring groove location shown in Fig. 3. The

crack lengths were 22 and 25 mm (0.85 and

0.99 in.). No cracks were detected in the T-head.

The radius at the primary fracture of the strain

bar was determined to be 1.3 mm (0.050 in.) At

the secondary cracks, the radius was 1.14 mm

(0.045 in.). At present, there is no dimensional

requirement for these radii.

Fig. 2 As-received T-head connector from fatigue testing of hold-back bar assembly

Fig. 3

Magnetic particle inspection indications found in the retainer ring groove on the strain bar. (a) Crack indication 1 showing a

crack length of 25 mm (0.99 in.) between arrows. (b) Crack indication 2 showing a crack length of 22 mm (0.85 in.) between

arrows. Original magniﬁcation: 1.4·

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The fracture surfaces were examined visually

and at up to 50· magniﬁcation using a stereo-

microscope. Figure 4 shows macrographs of the

fracture surfaces of both the strain bar and the

T-head. The strain bar primary fracture surface

exhibited multiple fatigue origins around the

outside periphery of the fracture. The T-head

fracture exhibited fatigue origins primarily near

internal thread roots.

A scanning electron microscope (SEM) was

used to examine the fracture surfaces and to

document the mode of failure. Figure 5 shows

the fracture topographies of the primary fracture

in the strain bar. The secondary fractures in the

strain bar are shown in Fig. 6. Fatigue striations

emanated from fatigue origins located on the

outside periphery of the fractures. Fracture rid-

ges also gave evidence of fatigue origin loca-

tions on the outside periphery of the fracture.

Figure 7 shows topographic features of the

fracture surface of the T-head failures. Fatigue

striations emanated from multiple origins loca-

ted on the inside peripheries of the fractures.

Hardness measurements were made on the

strain bar and T-heads to verify heat treat con-

ditions. The hardness of the strain bar was

determined to be 41.6 HRC, and this met the

hardness requirement of 40 to 43 HRC. The

hardness of the T-head was determined to be

48.8 HRC. This hardness value exceeded the

requirement of 46 to 48 HRC.

Metallographic sections were prepared

through the strain bar and a typical T-head at

the locations shown in Fig. 4. The specimens

were prepared using standard metallographic

techniques. Figure 8 shows the microstructures

of the strain bar and a typical T-head.

The inclusion contents of the metallographic

specimens appeared to be relatively high, as

is observed in air-melted steels. An energy-

dispersive spectroscopy analysis identiﬁed man-

ganese sulﬁde inclusions as being present.

Manganese sulﬁde inclusions are typically pres-

ent in air-melted steels and can severely affect

fracture toughness (Ref 1). Inclusions can act as

initiation sites for fatigue failures; however, the

failure of these parts could not be associated

with the presence of inclusions.

Banding was also present in both the strain

bar and T-head metallographic specimens.

Microhardness measurements made on the light

and dark bands in the strain bar and T-head

yielded hardness readings of 44.1 HRC (light),

42.6 (light), 49.8 (dark), and 49.1 (dark), respec-

tively. This microstructure is typical of an AISI

4340 component that has been improperly

normalized (Ref 2). The light etching areas are

generally untempered or lightly tempered mar-

tensite, while the darker etching regions

are martensite that has been tempered more

thoroughly. The lighter etching regions contain

higher hardenability than the darker etching

regions. This is typically due to segregation of

chromium, generally as the result of inadequate

normalizing prior to heat treatment. This type of

nonuniform structure is not optimal. The harder

regions tend to have lower fracture toughness

and have higher notch sensitivity than the dar-

ker, more tempered regions. Because of this, it

was thought that the premature fracture occurred

Fig. 4

Macrographs of primary fracture surfaces. (a) Strain bar, with maximum depth of fatigue at arrow (4.24 mm, or 0.167 in.),

looking forward. Original magniﬁcation: 2.4·. (b) T-head, with maximum depth of fatigue at arrow (2.90 mm, or 0.114 in.),

looking aft. Original magniﬁcation: 2.2·

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because of improper heat treatment. The pres-

ence of a higher level of inclusions could also

contribute to notch sensitivity reduction.

Cracking in a Main Landing Gear

Attach Pin

Multiple pins exhibited cracks in the ﬂanges

during magnetic particle inspection as part of

the manufacturing process. Some parts were

returned to the vendor for rework, while others

were reworked in-house.

The pins were machined from 300M steel.

The pins were then heat treated to a strength

range of 1930 to 2070 MPa (280 to 300 ksi).

The pin was ground and chromium plated.

Figure 9 shows the main landing gear (MLG)

lever attach pin as received for examination. The

magnetic particle inspection indications in the

ﬂange are shown in Fig. 10.

After the chromium plating had been

removed, the ﬂange area was subjected to a

temper-etch operation of 1% nitric acid in

methanol. Numerous areas on the ﬂange etched

dark, indicating that these regions had been

overtempered (Ref 3). A micrograph of the

temper-etched inspected ﬂange is shown in

Fig. 11.

A typical crack was opened, and the fracture

surface was examined visually and with a stereo-

microscope. The fracture surface was discolored

(Fig. 12). The discoloration appeared to be simi-

lar to a temper color in steel (straw or tan). After

the chrome plating was ground to ﬁnal dimen-

sions, the attach pin had received a stress-relief

operation that could have produced this dis-

coloration.

An SEM was used to examine the fracture

surface. Figure 13 shows the SEM fractographs

of the surface. The fracture topology was pri-

marily intergranular.

Fig. 5

SEM fractographs documenting typical topographic features of the strain bar primary fracture surface. (a) Overall view of strain

bar origins (micron bar is 1.5 mm long). (b) Typical origins (arrows) on outside edge of strain bar fracture (micron bar is

1500 mm long). (c) Typical fatigue striations on strain bar fracture (micron bar is 10 mm). (d) Typical worn area on strain bar fracture

(micron bar is 10 mm)

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A hydrogen analysis was conducted on the

MLG lever attach pin before and after baking for

24 h at 190



C (375



F). The results of the

analysis indicated that before baking, the lever

attach pin contained 1.6 ppm of hydrogen, and

after baking, the attach pin contained 1.4 ppm

of hydrogen. These amounts of hydrogen are

considered to be low and will not produce a

hydrogen embrittlement failure (Ref 4). Also,

fractures caused by hydrogen embrittlement are

generally not discolored.

Metallographic sections were prepared

through the magnetic particle indications using

standard metallographic techniques. Shown in

Fig. 14 are areas where the chromium plating

had been reduced by possible grinding.

Figure 15 shows micrographs of a typical crack.

This crack extended through the chromium

plating and into the base metal to a depth of ap-

proximately 0.25 mm (0.010 in.). There was

no visual indication of untempered or over-

tempered martensite.

Based on the results of this investigation,

it was concluded that the MLG lever attach pin

contained numerous probable grinding cracks in

the ﬂange.

MLG Linear Actuating Rod and Cylinder

The part was removed from the aircraft and

was reported to have accumulated 50 ﬂight

hours. The failed part was from an early pro-

duction lot manufactured with electroless nickel

on the inner diameter surface rather than the

drawing requirement of electrolytic nickel. The

manufacturing sequences for the MLG linear

actuating piston rod cylinder were:

 A 300M steel bar was machined to the re-

quired outside diameter and was then bored,

honed, and nickel plated.

 The bar was swaged in an approximate

76 mm (3 in.) length on the end opposite

Fig. 6

SEM fractographs of the fracture surfaces of cracks 1 and 2 in the strain bar. (a) Fracture surface of crack 1, showing typical

fatigue zones at arrows (750 mm). (b) Typical fatigue striations in crack 1 fatigue zones (30 mm). (c) Fracture surface at crack 2 of

the strain bar, showing typical fatigue zones at arrows (830 mm). (d) Typical fatigue striations in crack 2 fatigue zones (1.5 mm)

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