en 353

92

Chapter 4

RESULTS AND DISCUSSION The results obtained by performing various experiments, as described in Chapter 3,

for different steel grades are presented in this chapter. On the basis of results obtained, an

effort has been made to understand the effect of WC/C PVD coating on the fatigue

behaviour of different low-alloy steels investigated in this study.

4.1 Mechanical properties

Tensile and hardness tests were conducted on various steel specimens according to

the procedures described in section 3.3. The results obtained are presented and discussed in

the following sub-sections.

4.1.1 Tensile properties

Standard tensile test specimens of various steel grades were tested on a universal

testing machine to obtain their stress-strain diagrams, the values of yield and ultimate

strengths. All the steels tested exhibited similar stress-strain behaviour but with slightly

different values of yield and ultimate strengths. A representative stress-strain diagram

obtained for SAE8620 steel in green state is depicted in Figure 4.1.

In order to determine the elastic coefficients for various grades of steels in the green

state, tensile tests were conducted on a tensile testing machine by gradually applying the

load in small installments and recording the strain with the help of a mechanical

extensometer, as explained in section 3.3.1. The results obtained by conducting tensile tests

on various grades of steels are summarized in Table 4.1.

93

Figure 4.1: Stress-Strain graph for SAE8620 steel in green state

Table 4.1: Results of tensile tests conducted on various steel grades in green state Steel Grades Mechanical Properties SAE8620 20MnCr5 EN353 SCM420

Elastic Constant (GPa) 218 268 190 225 Yield Strength (MPa) 390 375 320 365 Ultimate Strength (MPa) 659 682 580 665

4.1.2 Microhardness

Following the procedure outlined in section 3.3.2, micro-hardness tests were

conducted on transverse cut-sections of the case carburized (uncoated) and case carburized

- WC/C coated specimens, referred to as uncoated and coated specimens, respectively

hereafter. The micro-indentations were made on the cut cross-section under a load of 200gf

for a dwell time of 15s. The micro-hardness values were estimated by measuring diagonals

of the indentations. The variation of hardness, observed for different steel specimens in the

uncoated and coated states, on moving from the surface towards the core, is shown in

Figure 4.2.

94

Figure 4.2: Effect of coating on microhardness profiles of various steels

95

As ASTM E 384 [263] recommends availability of sufficient backup material at the

site of indentation, the hardness measurements for the outer surfaces of the uncoated and

coated specimens were taken on their flat ends, rather than transverse cut-sections.

Hardness of the coating is estimated through microindentation measurements performed on

the flat coated surface under a load of 10gf and a dwell time of 15s. The hardness of the

coating is found to be HV1504±174.

It is evident from Figure 4.2 that in all the steels, the coating leads to a loss in

hardness of around 35HV near the surface, which diminishes gradually on moving towards

the core. This is attributed to the processing of case carburized specimens during PVD

coating at a temperature of 220°C, which is on the order of temperatures encountered

during low temperature tempering of low-alloy steels [268]. In general, beyond a depth of

around 1100 µm from the surface, the loss of hardness becomes almost negligible. It

appears that owing to relatively low processing temperature during the coating process, the

loss in hardness of the steel specimens is marginal.

4.1.3 Characterization of WC/C coating through Raman Spectroscopy

The Raman spectra of the WC/C coating, obtained through procedure described in

section 3.8, are shown in Figure 4.3. The two graphs correspond to the spectra recorded at

two different locations on the flat end of a single specimen.

The Raman spectra of various allotrophic forms of carbon are well-documented in

literature [269,270]. The peak observed at 1580 cm-1 for both the locations corresponds to

the first order graphite peak (G) of sp2 bonded carbon atoms, while the peak at 1360 cm-1,

observed at location-2 refers to the disorder-induced peak (D). The peak observed at 2726

cm-1 (location-1) corresponds to the second order graphite peak (G′ ). The sharpness of G

and G′ peaks signifies the presence of highly oriented pyrolytic graphite (HOPG), which

undergoes significant broadening in case of disordered films [Refer location-2]. The

96

difference in spectra recorded at different locations in a single specimen is known to result

from the variations in local deposition and annealing conditions [108].

Figure 4.3: Laser Raman spectra of WC/C coating at two different locations.

Both the Raman spectra reveal the presence of sp2 bonding among carbon atoms

within the PVD layer, thereby confirming the presence of graphitic carbon in the deposit.

4.2 Testing and analysis of SAE8620 steel specimens

The effect of material processings, viz. case-carburization and deposition of WC/C

coating, on the fatigue performance of SAE8620 steel is presented in the following sub-

sections. The results obtained by conducting fatigue tests are correlated with the influence

of these processes on various properties, viz. hardness, residual stresses etc., and

fractographic features.

97

4.2.1 Metallographic observations

Metallographic observations of the case carburized and tempered SAE8620 steel

specimens reveal the presence of martensitic structure of plate morphology in the case, as

shown in Figure 4.4 (a). Internal oxidation along boundaries of prior austenitic grains can

be seen extending to a depth of about 8-10 µm in this figure. The carbon-rich case is also

dotted with carbide precipitates that appear bright in the dark anodic film deposited by

Klemm’s – I reagent, as indicated by arrow marks in Figure 4.4 (b). Tempering of the

carburized steel is responsible for the formation of these spherodized carbides, which in-

turn leads to a reduction in the carbon content of austenite. The carbide particles help in

maintaining fine grains by inhibiting their growth. Besides, lower amounts of carbon in

austenite facilitate greater martensitic transformation by lowering the martensitic start

temperature (Ms). At some locations, chunks of carbide precipitates, segregated along prior

austenitic grain boundaries, can also be observed in Figure 4.4 (c) and (d). The core was

dominated by lath martensite [Refer Figure 4.5].

Based on the composition of SAE8620 steel given in Table 3.1 the martensitic start

temperature (Ms) of this steel can be calculated from the following equation, as suggested

by Payson and Savage [271]:

( ) MoSiNiCrMnCFM S 2020305060570930 −−−−−−=° Equation 4.1

From this relationship, the Ms temperature of SAE8620 steel works out to be

728.7 °F, or 387 °C. The estimated value of Ms temperature corresponds to the core of

carburized specimens, where chemical composition remains essentially unaltered. Once Ms

temperature is known, the expected volume fraction of martensite ( Mv ) can be calculated

from the following relationship, which was put forth by Koistinen and Marburger [272]

for low-alloy steels:

( )TMM

sev −−−= 011.01 Equation 4.2

where T is the quenching temperature.

98

Figure 4.4: Metallographs of case carburized SAE8620 specimen’s cross-section:

(a): Martensitic structures in the case, revealed by etching in 3% Nital for 10s; (b): Chunks of carbides at a depth of 300µm, (c) and (d): Pictures reproduced in true colour to reveal carbide segregates along prior austenitic grain-boundaries. Specimens in (b), (c) and (d) etched face-up for 3 min in Klemm’s - I reagent;

Figure 4.5: Metallographs showing the presence of lath martensite in the core of SAE8620

specimen etched with 3% nital for 5 seconds, viewed using: (a): 10X objective and (b): 40X objective.

It is evident from Equation 4.2 that near complete martensitic transformation would

occur when SAE8620 steel having Ms temperature of 387°C is quenched from austenitic

99

range to 120°C. Further, it also appears from the above discussion that substantial amounts

of retained austenite would be present within the outermost layers of carburized case.

However, contrary to this, the metallographic observations made on the carburized case

reveal a high volume fraction of plate martensite within the outermost 80µm layer [Refer

Figure 4.4 (a)]. The low fraction of retained austenite observed in the outermost layers is

attributed to the following two factors. First, an increase in Ms temperature of the case,

resulting from loss in carbon from the outermost surface layers due to decarburization,

internal oxidation [273,274] and precipitation of carbides. The second factor responsible for

this phenomenon is the occurrence of macrostraining during quenching of SAE8620 steel

specimens from austenitic temperature. Due to negative carbon gradient within the

carburized case, the Ms temperature will increase with depth. As a consequence, the

martensitic transformation would first commence at some depth below the surface during

quenching, thereby leading to development of tensile residual stresses within the surface

layers, where the phase is still austenitic. These residual tensile stresses would in-turn

stimulate a more complete martensitic transformation near the surface [275,276].

The gradual variation in microstructure upon moving from the surface towards the

core is captured by means of a number of metallographs recorded at successive depths and

compiled in the form of a mosaic, shown in Figure 4.6. It can again be observed that up to a

depth of around 80µm, the microstructure is largely dominated by plate martensite,

accompanied by small amounts of retained austenite. Upon moving further inwards, we

encounter a mixed – plate and lath morphology of martensite. Beyond a depth of about

700µm from the outermost surface, which corresponds to the effective depth of case, well-

defined packets of lath martensite can be observed, which become more orderly upon

moving further towards the core. The high volume fraction of lath martensite observed in

the core of SAE8620 steel specimens is commensurate with the predictions made by using

Equation 4.2 [272]. Lath martensite is generally the preferred microstructure for the low-

carbon core, since it is known for its better properties of strength and toughness in

comparison to the ferrite-pearlite microstructure [277].

100

Figure 4.6: Mosaic of metallographs showing variation in microstructure with depth in case carburized and tempered specimens made of SAE8620 specimen, etched in 3% Nital for 10 seconds.

In some of the specimens, small amounts of lower bainite [278] were also observed in

some regions of the core [Refer Figure 4.7 (a)], as revealed from the micrographs of case

carburized and WC/C coated specimens shown in Figure 4.7 (b). During the etching of

101

polished specimen with 3% nital for 2 hours, the etchant attached the martensitic matrix at a

higher rate than the lower bainitic phase, thereby leading to formation of raised

microstructures [Figure 4.7 (b)], which are further highlighted by means of an additional

source of light placed towards lower left side of the specimen, during the capture of

micrograph. The presence of small blow-holes in the material is also revealed through

scanning electron micrograph shown in Figure 4.7 (c).

The EDAX elemental analysis of the specimen [Figure 4.7 (d)], recorded at the

position marked with cross-hair in Figure 4.7 (c), revealed the dominance of Fe, Mo and C,

with their respective wt. percentages as 75.87%, 13.7% and 7.87%.

Figure 4.7: (a): Cross-section of case carburized, tempered and coated specimen etched in 3% nital for 2 h, followed by light polishing; (b): Magnified view of region “A” in figure (a) showing the presence of lower bainite in martensitic matrix; (c): Magnified view at location “B” in figure (a); and (d): Result of EDAX elemental analysis performed at location marked with cross-hair in (c).

102

4.2.2 Estimation of residual stresses

The residual stress-field present in the material plays an important role in determining

its fatigue performance. The residual stresses in the specimens were estimated through X-

ray diffraction technique, following standard procedures described in sections 2.5.2 and 3.7.

On the basis of location and intensity, the peak corresponding to (211) plane (ref. Figure

3.26) was found to be suitable for the purpose of estimating lattice strain. The peak

exhibited broadening, which is a commonly observed phenomenon in most of the hardened

and fine-grained steels [234]. The X-ray diffraction pattern was recorded at an interval of

0.0098° with a total of 255 readings in the 2θ range from 80.99° to 83.50° in the tilt range

(ψ) from – 40° to +40°, while the divergent slit was kept at 2mm. Parabolic curve fitting

was employed for determining the position of peaks at individual tilts. The results of X-ray

diffraction analysis, performed on uncoated and coated specimens, are summarized in Table

4.2 and shown graphically in Figure 4.8. Psi-splitting has been observed in the d vs sin2ψ

plot [Refer Figure 4.8] for both the specimens, which signifies the presence of tri-axial

stresses near the surface of specimens. The shape of both these plots, however, rules out the

presence of texturing in the material [220,221].

Table 4.2: Determination of lattice strain in SAE8620 steel specimens through X-ray diffraction

Uncoated Specimen Coated Specimen S. No. Tilt, ψ sin2ψ d-spacing Tilt, ψ sin2ψ d-spacing

1 -40.00 0.4132 1.17060 -40.00 0.4132 1.17113 2 -35.09 0.3305 1.17079 -35.09 0.3305 1.17152 3 -29.86 0.2479 1.17098 -29.86 0.2479 1.17195 4 -22.99 0.1525 1.17122 -22.99 0.1525 1.17239 5 -16.71 0.0827 1.17148 -16.71 0.0827 1.17282 6 0.00 0 1.17184 0.00 0 1.17334 7 16.71 0.0827 1.17186 16.71 0.0827 1.17303 8 22.99 0.1525 1.17175 22.99 0.1525 1.17268 9 29.86 0.2479 1.17162 29.86 0.2479 1.17228 10 35.09 0.3305 1.17147 35.09 0.3305 1.17189 11 40.00 0.4132 1.17129 40.00 0.4132 1.17151

103

Residual stress (σΦ) corresponding to the recorded peak positions is estimated from

Equation 3.4, as reproduced below:

∂∂

+

=ψν

σ φψ

φφ 2

0)( sin1

1d

dE

hkl

To estimate the value of material constant, K ( )[ ]ν+= 1/E , of the martensitic case,

we have used E = 208 GPa and ν = 0.29 from the reported literature [266]. The estimated

value of K (= 161.2 GPa) is in fair agreement with the value (167.5 GPa) reported in the

literature for 4820 steel under carburized condition [220]. Using these constants, the value

of residual stress (σΦ) is estimated as -305.6±2.9 MPa for uncoated and -684.6±3.1 MPa for

coated specimens. The fairly high magnitude of residual compressive stresses observed in

the case-carburized condition is attributed to the high degree of martensitic transformation

near the surface, as reported in section 4.2.1. After coating the carburized specimen with

WC/C layer, the magnitude of residual compressive stresses becomes more than double.

Figure 4.8: d vs sin2ψ plot for uncoated and coated specimens made of SAE8620 steel

104

4.2.3 Fatigue testing

The results obtained by conducting 4-point rotating bending fatigue tests on standard

fatigue specimens made from SAE8620 steel in the green state are reported in Table 4.3.

The average number of cycles to failure corresponding to a given stress level is denoted by

Nf , which is estimated by taking the average of N1 and N2. The results obtained by

performing fatigue tests on case carburized specimens in uncoated and coated conditions

are reported in Table 4.4 and Table 4.5 respectively.

Table 4.3: Results of fatigue tests conducted on specimens made of SAE8620 steel in green state

Cycles to failure Stress (MPa) N1 N2

Average (Nf

Green) Std. Dev. (σ

Green) 425 10164 13472 11818 2339 396 14136 20163 17150 4262 365 40357 52341 46349 8474 327 109370 149867 129619 28636 300 236783 352208 294496 81618 279 356783 542760 449772 131506 262 583675 882430 733053 211252

Table 4.4: Results of fatigue tests conducted on case-carburized and tempered (uncoated) specimens made of SAE8620 steel


Average (Nf

Uncoated) Std. Dev. (σ

Uncoated) 1000 21341 26178 23760 3420 970 36756 51451 44104 10391 930 78824 112570 98697 23862 902 116280 162092 139186 32394 850 292365 412878 352622 85216 810 547443 813357 680400 188030 800 665382 946726 806054 198940

Table 4.5: Results of fatigue tests conducted on case-carburized, tempered and WC/C coated specimens made of SAE8620 steel


Average (Nf

WC/C) Std. Dev. (σ WC/C)

995 10502 15325 12913 3412 975 17783 25578 21681 5512 960 30451 36562 33507 4321 935 58749 78637 68693 14063 910 98834 152016 125425 37605 885 240679 336284 288482 67603 860 561146 838254 699700 195945

105

In order to quantify the change in fatigue life as a function of maximum alternating

stress, a linear model of the form: Y = A + BX [Refer Equation 2.5] has been fitted

separately to the fatigue data obtained for specimens in three different states – (i) green, (ii)

case-carburized and (iii) case carburized, followed by WC/C coating. The term Y in the

model equation represents logarithm of the number of cycles to failure (Nf) while X denotes

the maximum alternating stress (S) in MPa. The parameters A and B in the linear model are

estimated through regression analysis, as prescribed in ASTM E 739 standard [121]. The

estimated values of parameters are summarized in Table 4.6.

Table 4.6: Values of parameters in the linear model for fatigue life of SAE8620 steel

Material & State Parameter A Parameter B SAE8620 (Green) 8.839 -0.0114 SAE8620 (Case carburized) 11.923 -0.0075 SAE8620 (WC/C coated) 16.659 -0.0126

The experimental fatigue data, along with the fitted linear models and 95%

confidence bands, are plotted on semi-log scale, i.e. Log (Nf) vs Stress (S), in Figure 4.9 for

SAE8620 steel specimens in green, case-carburized (uncoated) and case carburized–WC/C

coated states. The parameters A and B appearing in the linear model are dependent on the

bulk and surface properties of the material apart from the testing conditions. These

parameters can be of prime importance while estimating the fatigue life of SAE8620 steel

in green, uncoated and coated states, processed in a fashion similar to that employed in the

present work.

The endurance limit of the specimens in various states was estimated through

staircase tests conducted for 2×106 cycles. The results obtained are summarized in Table

4.7. Taking the number of less frequent event (i.e., survival or failure) as the basis for

statistical analysis of these staircase tests, the data reduction is carried out according to the

Dixon-Mood method [Refer section 2.3.2.4]. The intermediate calculations and final results

thereof are given in Table 4.8.

106

Table 4.7: Results of staircase tests conducted on specimens made of SAE8620 steel Green Case Carburized WC/C Coated S. No.

Stress (MPa)

Cycles Stress (MPa)

Cycles Stress (MPa)

Cycles

1 279 356783 810 813357 885 240679 2 262 2000000* 800 665382 860 838254 3 279 542760 790 2000000* 835 2000000* 4 262 882430 800 2000000* 860 2000000* 5 245 2000000* 810 547443 885 336284 6 262 583675 800 946726 860 561146 7 245 2000000* 790 879452 835 2000000* 8 --- --- 780 2000000* --- ---

* No failure; Test suspended.

Table 4.8: Results of Dixon-Mood statistical analysis of staircase test data obtained for SAE8620 steel specimens

i Green Case Carburized WC/C Coated 0 S0 = 245 MPa nDM,0 = 2 S0 = 780 MPa nDM,0 = 1 S0 = 835 MPa nDM,0 = 2

1 S1 = 262 MPa nDM,1 = 1 S1 = 790 MPa nDM,1 = 1 S1 = 860 MPa nDM,1 = 1


3 --- --- S3 = 810 MPa nDM,3 = 0 --- ---

Stress interval (d)= 17 MPa

( )( )∑= iDMDM niA , = 1

( )( )∑= iDMDM niB ,2 = 1

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.22 (<0.3)

Mean stress (µs) = 259 MPa

Std. dev. (σs) = 9.01 MPa


( )( )∑= iDMDM niA , = 3

( )( )∑= iDMDM niB ,2 = 5

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.67 (≥0.3)

Mean stress (µs)= 795 MPa



( )( )∑= iDMDM niA , = 1

( )( )∑= iDMDM niB ,2 = 1

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.22 (<0.3)


Std. dev. (σs)= 13.25 MPa

107

An interesting phenomenon observed in Figure 4.9 is the cross-over in S-N curves of

the uncoated and coated specimens at a stress cycle of around 105, which implies that the

coated specimens show poor life in comparison to their uncoated counterparts in the low

cycle fatigue regime, but exhibit better performance in the high cycle regime. The lower

slope of S-N curve observed in Figure 4.9 for the coated specimens is also apparent from

the lower value of parameter B observed for WC/C coated specimens [Refer Table 4.6]. Su

and co-workers [10,11] in their studies on fatigue behaviour of various PVD coatings

applied to normalized AISI 1035 steel, also reported better performance of the coated

specimens in comparison to their uncoated counterparts for stress cycles higher than 5×105.

Figure 4.9: S-N graphs for SAE8620 steel specimens in green, case carburized (uncoated)

and case carburized - WC/C coated states.

108

The effect of case carburization as well as WC/C coating on the fatigue strength of

SAE8620 steel has been estimated by comparing the fatigue strength of case carburized

(uncoated) and case carburized – WC/C coated specimens with that of the specimens in

green state. The percentage change in fatigue strength at a given number of stress cycles is

computed from the relationship given below:

100% ×−

= Greenf

Greenff

SSS

strengthfatigueinChange

Equation 4.3

where fS refers to the fatigue strength of either uncoated ( UncoatedfS ) or coated ( CWC

fS / )

specimens. It denotes the magnitude of maximum alternating stress at which the specimens

would fail after the given number of cycles (Nf). The value of fS has been computed from

the linear fitted model, given in Equation 2.5, developed for steel specimens in different

states, viz. green, case carburized (uncoated) and case carburized – WC/C coated. To

estimate fS , Equation 2.5 has been re-arranged as:

( )B

ANS f

f

−=

log

Equation 4.4

The values of parameters A and B for different types of steel specimens are

substituted from Table 4.6. The percentage change in fatigue strength has been computed

for different number of stress-cycles, ranging from 103 to 106. The results obtained are

summarized in Table 4.9 and depicted graphically in Figure 4.10.

Table 4.9: Effect of case carburization and subsequent WC/C coating on the fatigue strength of SAE8620 steel.

SfGreen Sf

Uncoated SfWC/C %Change in Fatigue Strength

A=8.8392 A=11.9234 A=16.6587

No. of Cycles to failure, Nf

B=-0.0114 B=-0.0075 B=-0.0126 Case

Carburized vs Green

WC/C Coated

vs Green 1000 512 1187 1080 132 111

10000 424 1054 1001 148 136 100000 337 921 922 174 174

1000000 249 788 843 216 239

109

It can be observed from Figure 4.10 that both uncoated and coated specimens show

substantial improvement in fatigue performance as compared to specimens in green state.

This enhancement in fatigue performance is attributed to the presence of residual

compressive stresses within the surface layers of the uncoated as well as coated specimens.

Therefore, the sequence of martensitic transformation [Refer section 4.2.1], which affects

the magnitude of residual stresses in the outermost surface [250], in-turn, also influences

the resulting fatigue properties. The presence of residual compressive stress delays the

crack propagation within the uncoated and coated specimens, thereby increasing their

fatigue performance in comparison to steel specimens in green state. For both types of

specimens, the percentage gain in fatigue strength increases with increasing number of

cycles. The increase observed for the coated specimens is steeper than that observed for

uncoated specimens. For bending stresses higher than 922 MPa, which corresponds to a

fatigue life of 105 cycles, the case carburized (uncoated) specimens exhibit better fatigue

performance than their WC/C coated counterparts. The presence of WC/C PVD coating on

the case carburized specimens results in improved fatigue performance for stress levels

below 922 MPa (Nf > 105 cycles). The gain observed in fatigue strength of the WC/C

coated specimens increases with increasing number of cycles upto the endurance limit of

the coated specimens.

The beneficial effect of PVD coating on fatigue performance is also reported by

Puchi-Cabrera et al. [197]. They reported 3 to 20 times enhancement in finite fatigue life of

AISI 316L stainless steel specimens when they were coated with PVD layer of TiN. In their

study, the fatigue strength of the coated specimens was observed to be 22% higher than that

of the uncoated specimens.

110

Figure 4.10: Percentage change in fatigue strength of uncoated and coated SAE8620 steel specimens with respect to specimens in green state.

4.2.4 Fractographic observations

In order to identify the mode of fatigue failure and reveal the topography of fatigue

fracture surfaces, fractographic observations were made on the fatigue specimens in various

states, viz. green, case carburized and WC/C coated, tested at different stress levels. The

results obtained are presented and discussed in the following sub-sections.

4.2.4.1 Fractographic observations on specimens in green state

Fractographic observations on specimens in green state, failed at different stress

levels, reveal that different crack propagation mechanisms operate at different levels of

remote loading stress. The fracto-macrographs of the fracture-surfaces of the specimens,

arranged in increasing order of applied stress, are shown in Figure 4.11(a-g).

111

Figure 4.11 (a) and (b) show the fracture surfaces obtained at relatively low stress

levels of 262 and 279 MPa respectively. On a macroscopic scale, surfaces in both the

figures are oriented normal to the main principal stress (longitudinal axis of the specimen)

and exhibit planar topography with gradually increasing roughness as the crack propagates.

The region of crack propagation under stage-II [94] is characterized by smooth and bright

fractured surface, vividly captured in Figure 4.11 (b), where a bright, crescent-shaped

region opposite to the zone of final fracture has been demarcated with a dotted curve. The

portion between this dotted line and the region of final fracture exhibits higher roughness,

characteristic of rapid crack propagation (Stage-III). A few white patches appearing in this

zone are the result of rubbing that took place between the two halves of the specimen

during the course of failure. The area of final fracture is rather small in these cases owing to

relatively lower applied stress. The geometry and macroscopic orientation of the fractured

surface implies mode-I (tensile) crack propagation [279].

The examination of fractured specimen at low magnification revealed that the

outermost region opposite to the site of final failure exhibited near specular finish in a band

of approximately 0.5mm width, which is attributed to the initial, slow pace of crack

propagation. The same region also exhibited a number of ratchet marks [Refer Figure 4.12],

which indicate the occurrence of multiple fatigue crack nucleation sites [280]. The

examination of the specimen under SEM [Figure 4.13] reveals the formation of wedge-

shaped extrusion sliver (marked with an arrow), resulting from slip bands within the

crystallites, typical of stage-I crack propagation [93,281,282].

112

Figure 4.11: Fatigue fracture surfaces of SAE8620 steel specimens in green state, tested at

(a): 262 MPa (b): 279 MPa (c): 300 MPa (d): 327 MPa (e): 365 MPa (f): 396 MPa (g) and (h) 425 MPa.

113

Figure 4.12: Optical micrograph of specimen tested at 279MPa, showing ratchet marks and transgranular crack propagation.

Figure 4.13: SEM image of fatigue specimen tested at 279 MPa, showing transgranular crack propagation, along with ratchet mark and extrusion sliver on the outer surface (identified with arrow-mark).

114

Figure 4.14: SEM image of fractured specimen tested at 279 MPa, showing fatigue striations on multiple plateaus

At some places in the region of mode-I crack propagation, fatigue striations could be

observed on the fractured surface, as revealed in the scanning electron micrograph given in

Figure 4.14. The striations are found to be bowed out roughly in the direction of crack

propagation and their topography consists of multiple plateaus. The SEM image is taken in

the stage-II region at a place midway between the crack initiation site and the region of

rapid fracture. Unlike stage-I, where striation spacing can be much more than the actual

crack growth rate [98], the stage-II striations are generally formed during each loading

cycle. The observed striation spacing, which is on the order of 0.5µm, indicates a relatively

short duration of stage-II crack propagation phase (5,000-10,000 cycles) in comparison to

the total fatigue life of the specimen. Thus, it can be concluded that in case of specimens

tested under high cycle fatigue, stage-I crack propagation phase constitutes the major

proportion of fatigue life.

As the applied stress is increased beyond 279 MPa, the topography of fractured

surfaces transforms to that shown in Figure 4.11 (c) and (d), which were observed in

115

specimens tested at stress levels of 300 and 327 MPa respectively. The fractured surfaces

reveal formation of shear lips, inclined at an angle of approximately 45° with the transverse

section of the specimens. Therefore, with the increase in stress level beyond 279 MPa, the

mode of failure shifts from tensile (mode-I) to shear (mode-II), thereby leading to formation

of shear lips. Though, the formation of shear lips is not very well understood, but their

occurrence in metals having FCC or BCC lattice structure is already known and has been

extensively reported for flat, plate-shaped fatigue specimens made of aluminium alloys

[283,284] and to a lesser extent for steel plates [285]. However, the literature consulted so

far does not reveal the occurrence of shear lips in case of cylindrical steel specimens. In

general, the formation of shear lips is attributed to the fact that owing to lesser restraint on

the material at the free surface in comparison to that in the bulk, a large plastic deformation

occurs in the region where the crack front meets the free surface, which in turn promotes

the formation of shear lips [109]. The width of shear lips in the present study does not reach

an equilibrium state but continues to grow, which is in agreement with the model proposed

by Zuidema and Blaauw [286].

The advancing intersection of the crack front with the free surface is marked with

arrows in Figure 4.11 (c). As one moves in the direction of arrows, a very well defined start,

followed by a gradual increase in width of the shear lips is observed. A similar feature can

also be noticed in Figure 4.11 (d), where two mutually perpendicular shear lips are

observed at +45° and -45°, starting point of one of which is identified with an arrow.

Double shear lips are known to cause a deviation in the direction of crack propagation, as

described by Schijve and co-workers [109,287]. However, in the present study, owing to

cylindrical shape of the specimens, the ridge formed between the two shear lips naturally

follows a helical path, instead of the straight-line path, as reported for flat specimens by

Schijve and co-workers.

Figure 4.11 (e) represents a more complex fractured surface of the specimen tested at

365 MPa, where crack nucleation and macroscopic growth occurred on multiple planes.

Although, the cracks could not be observed during the course of tests, but the final fractured

116

surface reveals cracks originating on multiple planes, which subsequently merged with each

other. At various sites, the crack propagation has occurred through mixed modes – I and II,

while the region of stable crack growth has been reduced significantly in the wake of

relatively larger applied stresses.

When the stress level is further increased to 396MPa, which is on the order of yield

strength of the SAE8620 material [Refer Table 4.1], an interesting topography, resembling

a “helical staircase” emerges, as revealed in Figure 4.11 (f). As the material near the outer

surface of the specimen is subjected to severe plastic deformation, the transition of crack

propagation from mode–I to mode–II also initiates at the outer periphery. It appears that the

unstable crack front has led to the formation of a “helical staircase” topography on the

fractured surface. Observations under scanning electron microscope [Refer Figure 4.15]

reveal fibrous appearance of the fractured surface, which is attributed to mode–II

transgranular crack growth [288]. Fatigue striations could not be observed anywhere on the

fractured surface, though it exhibited tire-tracks at a few places, as shown in Figure 4.16.

A further increase in stress level to 425MPa, which is well above the static yield

strength of SAE8620 steel, causes the entire material within the gauge length of the

specimen to undergo plastic deformation. The extension of plastic zone from the surface

right up to the center of the cylindrical specimen is confirmed from the topography of the

fractured surface shown in Figure 4.11 (g). The region exhibiting unstable crack front,

which was restricted to the outer periphery in case of specimen tested at an applied stress of

396 MPa [Refer Figure 4.11 (f)], now gets extended up to the neutral axis. The specimen

tested at 425 MPa also exhibited the formation of multiple macroscopic cracks, some of

which coalesced and led to the final fracture, as revealed in the side view of the specimen in

Figure 4.11 (h). For the specimens tested at stress levels of 396 and 425MPa, the plastic

straining was also accompanied by hysteresis losses within the material, thereby leading to

a considerable rise in temperature of the specimens.

117

Figure 4.15: SEM image of crack geometry under mode-II in fatigue specimen tested at

396 MPa.

Figure 4.16: SEM image of fractured specimen tested at 396 MPa, showing presence of tire tracks

118

4.2.4.2 Fractographic observations on case-carburized specimens

The fracture surfaces of the case carburized (uncoated) and case carburized – WC/C

coated specimens failed at different levels of alternating stresses were examined optically as

well as under scanning electron microscope [Refer Figure 4.17 (a) - (h)]. The examinations

were carried out to investigate the crack propagation mechanisms operative in uncoated and

coated specimens and to develop an understanding of the effect of coating on fatigue

behaviour of the case carburized SAE8620 steel.

Figure 4.17 (a) shows the optical fractograph of uncoated fatigue specimen, tested at

810 MPa, where arrow points the region of crack initiation. The crack is observed to

propagate nearly normal to the direction of principal stress (which is also the direction of

specimen axis). Almost diametrically opposite to the crack initiation site, a small step is

formed on the fractured surface due to merging of the two crack-fronts, which were slightly

out-of-plane with respect to each other. Fractographic studies conducted at higher

magnification revealed that the cracks originated either at some defect on the outermost

surface [Refer Figure 4.17 (b)] or through intergranular decohesion along prior austenitic

grain boundaries [Refer Figure 4.18]. Crack initiation through intergranular decohesion is a

commonly observed fracture mode in the carbon rich case of the carburized steels and

occurs when the applied stress exceeds the sum of residual compressive macro-stresses and

cohesion strength of the material at prior austenitic grain boundaries [239].

The crack thus initiated, however, can not always continue to progress in the same

intergranular fashion and gets arrested as a result of small plastic zone (smaller than the

grain size) at the tip of these sharp intergranular cracks. The presence of plastic zone leads

to strain-induced transformation of retained austenite, present near the outer surface of the

carburized specimen, into martensite, which further induces compressive stresses at the

crack tip [289]. As a result, the crack now propagates through stable transgranular mode,

till it reaches a critical size, whereupon it propagates as unstable fracture.

119

The dominance of transgranular mode of crack propagation is clearly revelaed in

Figure 4.17 (a) and (b). In Figure 4.17 (b), the region of stable crack propagation appears in

distinct contrast to the remaining region of unstable crack propagation and can be employed

for estimating fracture toughness of the material [Refer section 2.6].

It is interesting to note that despite the presence of large residual compressive stresses

on the order of 306 MPa, the crack initiation still occured in the outermost case-carburized

layer and not at the root of the carburized case. Similar observations regarding crack

initiation are noticed by other researchers [151, 290]. The initiation of cracks at the surface

reveals the importance of surface conditions in deciding the fatigue life despite the presence

of large residual compressive stresses at the surface.

The fractographs recorded on uncoated fatigue specimen tested at a relatively higher

stress of 1000 MPa reveal that the fractured surface is again dominated by the region of

transgranular mode of crack growth, as is evident from Figure 4.17 (c), where the crack

initiation site is marked with a circle. The regions marked with rectangles exhibited fatigue

striations and tire tracks, as is evident from the scanning electron micrograph shown in

Figure 4.17 (d), which is recorded at some location within the regions marked with

rectangle.

The top and side views of the fractured surface of the coated fatigue specimen tested

at 860 MPa are shown in Figure 4.17 (e). The fracture surface of the coated specimen is

also observed to be dominated by transgranular mode of crack propagation, similar to that

observed in Figure 4.17 (c) for uncoated specimen tested at 1000 MPa. The magnified view

of the crack initiation region marked in Figure 4.17 (e) is shown in Figure 4.18, where the

dominance of transgranular crack propagation mechanism is quite apparent. The crack

originates through intergranular decohesion, but immediately switches to transgranular

mode, whereupon it propagates circumferentially in both the directions and finally meets at

the diametrically opposite side, where it reaches along two slightly different planes, marked

as 1 and 2 in Figure 4.17 (e).

120

Figure 4.17: Fractographs of specimens made of SAE8620 steel: (a): Optical fractograph of

uncoated fatigue specimen tested at 810 MPa, showing crack initiation site; (b): Scanning electron micrograph of the location marked by arrow in (a); (c): Optical fractograph of uncoated specimen tested at 1000 MPa, showing crack initiation site; (d): Scanning electron micrograph showing striations and tire tracks at a location within the regions marked by rectangles in (c); (e): Top and side views of coated fatigue specimen tested at 860 MPa; (f): Scanning electron micrograph showing adherence of coating at the failed section marked with rectangles in (e); (g): Optical fractrograph of coated fatigue specimen tested at 910 MPa, showing ratchet marks and crack initiation site; (h): Scanning electron micrograph at a location within the region marked with rectangle in (g), showing tire tracks.

121

Scanning electron micrograph of the region marked with a rectangle in Figure 4.17

(e), is shown in Figure 4.17 (f). The micrograph reveals that the coating remains bonded

even at the thin tip of the failed specimen. The examination of fractured surfaces of the

other specimens also indicates that no delamination of coating from the substrate wall

occurs. Earlier investigation reported by Carvalho [291] also confirms the absence of

delamination in case of WC/C coatings applied to steel substrates.

The fractured surface of the coated fatigue specimen tested at a relatively high stress

of 910 MPa is characterized by the regions of unstable crack growth, along with a large

number of ratchet marks, as identified by arrows 1 and 2 in Figure 4.17 (g). The arrow

marked 3 in Figure 4.17 (g) shows the crack initiation site. The presence of tire tracks is

clearly revealed in scanning electron micrograph shown in Figure 4.17 (h), which is taken

in the region marked with a rectangle in Figure 4.17 (g).

The comparison of crack-initiation sites observed for uncoated and WC/C coated

specimens, as shown respectively in Figure 4.17 (b) and Figure 4.18, reveals that similar

crack initiation and growth mechanisms operate in both the specimens. The overall crack

propagation is found to be transgranular in nature, while the region of final ductile fracture

exhibits microvoid coalescence morphology. Measurements made on the stable crack

propagation region [Refer section 2.6] shown in Figure 4.17 (b) provide depth, a = 322µm

and width, c = 911 µm, on the basis of which, fracture toughness [Refer Equation 2.21]

works out to be KIC = 25.59 mMPa , which is quite high for the steel in case carburized

condition [240].

An important observation is that the crack progresses largely through transgranular

fracture mode even through the high hardness case. As a result, there is no change in the

mode of propagation when the crack progresses to regions that are low in carbon content. In

materials, which are susceptible to intergranular cracking, a short-term, secondary stage of

fatigue crack propagation, characterized by transgranular mode of fracture, observable

fatigue striations and secondary cracking is often reported to occur prior to the final ductile

122

overload fracture [240,292]. However, no such transition from intergranular to

transgranular mode took place for case carburized specimens made of SAE8620 steel, as

the dominant mode of crack propagation was always transgranular.

Figure 4.18: Magnified view of crack initiation region of specimen shown in Figure 4.17

(e), depicting intergranular initiation and transgranular propagation.

For the coated specimen shown in Figure 4.17 (g), which was tested under low cycle

fatigue, the fracture surface exhibited crack-initiation on multiple planes, as shown in

Figure 4.19. Crack-initiation on multiple planes is also confirmed by the formation of

multiple plateus [Refer Figure 4.20] near the outermost edge of fractured surface.

These observations imply that formation of multiple cracks is encouraged under low

cycle fatigue when the specimen is in coated condition. This can be attributed to substantial

straining of the substrate under high applied stresses, which lead to formation of multiple

cracks within the coating, some of which grow to macroscopic scale. This phenomenon

seems to be responsible for relatively poor performance of the coated specimens under low

cycle fatigue. Otherwise, there is no significant difference in the surface topography as well

as associated crack propagation mechanisms for the uncoated and WC/C coated specimens,

particularly for specimens that failed at comparable number of cycles. This is in agreement

with the observation reported by Baragetti and Tordini [188]. The presence of coating

123

modifies the level of stress at which the failure would occur after a given number of load

cycles. In other words, the modes of failure remain more or less the same for both, uncoated

and WC/C coated specimens. However, the stress level for their occurrence gets modified

in the presence of WC/C coating.

Figure 4.19: Optical fractograph taken on side-wall of coated specimen shown in Figure

4.17 (g), revealing the formation of multiple cracks under low-cycle fatigue.

Figure 4.20: Close-up view of specimen shown in Figure 4.17 (g), revealing crack

formation on multiple planes.

124

4.2.5 Effect of WC/C coating on fatigue behaviour of SAE8620 steel

The observed change in the fatigue performance of the coated specimens, in

comparison to their uncoated counterparts, can be explained on the basis of residual stress

measurements reported in Section 4.2.2. The factor which seems to be responsible for

reducing the low-cycle fatigue strength of the coated specimens is the development of

multiple cracks on the hard WC/C layer, when the substrate undergoes considerable

deflection due to high applied loads under low cycle fatigue [Refer Figure 4.19 and Figure

4.20]. Some of these cracks propagate into the case-carburized layer of the substrate and

lead to early fatigue failure.

However, when it comes to high cycle fatigue, the coating seems to provide an

effective protection to the surface by suppressing the crack propagation, thereby leading to

considerable improvement in fatigue performance. The observed improvement in high

cycle fatigue performance of WC/C coated specimens is therefore attributed to higher

residual compressive stresses (685 MPa) in the coated specimens than those observed in the

case carburized (uncoated) specimens (306 MPa).

125

4.3 Testing and analysis of 20MnCr5 steel specimens

This segment of the study investigates the effect of case carburization and WC/C

coating on hardness, residual stress and fatigue performance of 20MnCr5 steel. For this

purpose, a comparative study has been performed on specimens made of 20MnCr5 steel in

green, case carburized (uncoated) and case carburized – WC/C coated states, in a manner

similar to that followed for SAE8620 steel. The results obtained are presented and

discussed in the following sub-sections.


The metallographic observations performed on 20MnCr5 steel in green state revealed

the presence of fine pearlitic microstructures [Refer Figure 4.21] with interlamellar spacing

on the order of 0.5 µm. The occurrence of fine pearlitic microstructure is attributed to the

presence of chromium and manganese in this steel [293], as revealed from Table 3.1.

Figure 4.21: Metallograph of 20MnCr5 steel (etched in 3% Nital for 4s) in green state revealing the presence of pearlitic microstructure.

The microstructure of case carburized and tempered specimen of 20MnCr5 steel is

depicted in Figure 4.22. The microstructures are recorded across depth of the specimen and

compiled in the form of a mosaic.

126

Figure 4.22: Mosaic of metallographs, showing variation of microstructure with depth in case carburized and tempered specimens made of 20MnCr5 steel, etched in 3% Nital for 8 seconds.

127

The microstructure within the outermost 400µm is dominated by plate martensite,

dotted with segregates of carbide particles, which appear as white dots, scattered upto a

depth of 300µm. These precipitates are vividly revealed upon etching with Klemm’s - I

reagent, as shown in Figure 4.23.

Figure 4.23: Metallograph of case carburized and tempered specimen made of 20MnCr5 steel, etched in Klemm’s - I reagent for 3 minutes, revealing the presence of carbide particles in the carbon-rich case.

Apart from plate martensite, some amount of retained austenite can also be observed

upto a depth of around 400µm in Figure 4.22. The amount of retained austenite reduces

significantly on moving from the outermost surface to a depth of around 400µm, beyond

which, its decrease becomes marginal. This reduction in retained austenite with increase in

depth is attributed to a rise in Ms temperature, as a consequence of negative gradient in

carbon content.

The amount of retained austenite observed in case carburized and tempered

specimens made of 20MnCr5 steel is relatively higher than that observed in SAE8620 steel

specimens processed in similar fashion [Refer Figure 4.6]. It may be attributed to relatively

lower Ms temperature of 368°C for 20MnCr5 steel, as estimated from Equation 4.1 [Refer

Table 3.1 for chemical composition] in comparison to that estimated for SAE8620 steel

128

(Ms = 387°C). Due to lower Ms temperature, the transformation in the core of specimens

made of 20MnCr5 specimens would have commenced at a later stage, whose time would be

overlapping with the start of martensitic transformation at the surface. As a result of this

alteration in the sequence of martensitic transformation, the advantage of strain-induced

martensitic transformation would be harnessed to a lesser extent for specimens made of

20MnCr5 steel, in comparison to those made of SAE8620 steel. The significantly higher

amounts of retained austenite in the outermost layers, in comparison to those observed for

SAE8620 steel, provide an explanation for the considerably lower values of hardness

observed for 20MnCr5 steel specimens [Refer Figure 4.2].

Similar to SAE8620 steel, the microstructure of 20MnCr5 specimens also reveals a

mixed morphology of plate and lath martensite at a depth of 400 µm to 600 µm below the

surface. Beyond a depth of around 600 µm, the microstructure is dominated by lath

martensite, which becomes more and more orderly with further increase in depth.


The residual stress in the uncoated and WC/C coated specimens made of 20MnCr5

steel was estimated through X-ray diffraction technique, as described earlier for SAE8620

steel (Ref. section 4.2.2). The results obtained are summarized in Table 4.10 and depicted

graphically in Figure 4.24. Psi-splitting is observed to occur to a greater extent than that

observed earlier for SAE8620 steel [Refer Figure 4.8], thereby implying higher magnitude

of triaxial stresses in the surface layers of 20MnCr5 steel specimens.

Residual stresses corresponding to the recorded peak positions are estimated from the

relationship given in section 3.7.2. From the recorded shift in peaks, the value of residual

stress (σΦ) is estimated as -184.8±3.0 MPa for the uncoated and -449.2±3.0 MPa for the

coated specimens. The magnitude of residual compressive stresses in the uncoated

specimens made of 20MnCr5 steel is about 40% lower than that observed in the

corresponding specimens made of SAE8620 steel [Refer section 4.2.2]. This is attributed to

the presence of significantly higher amounts of retained austenite in the surface layers of

129

20MnCr5 steel specimens, as compared to that observed in SAE8620 steel specimens. The

deposition of WC/C coating, again led to enhance the magnitude of residual compressive

stresses, but they still remained 35% lower than the compressive residual stresses observed

in WC/C coated SAE8620 steel specimens.

Table 4.10: Determination of lattice strain in 20MnCr5 steel specimens through X-ray diffraction

Uncoated Specimen Coated Specimen S. No. Tilt, ψ sin2ψ d-spacing Tilt, ψ sin2ψ d-spacing

1 -40.00 0.41318 1.17081 -40.00 0.41318 1.17089 2 -35.09 0.33047 1.17092 -35.09 0.33047 1.17117 3 -29.86 0.24789 1.17105 -29.86 0.24789 1.17148 4 -23.99 0.16530 1.17120 -23.99 0.16530 1.1718 5 -16.71 0.08267 1.17137 -16.71 0.08267 1.17213 6 0 0 1.17169 0 0 1.17260 7 16.71 0.08267 1.17177 16.71 0.08267 1.17256 8 23.99 0.16530 1.17174 23.99 0.16530 1.17238 9 29.86 0.24789 1.17166 29.86 0.24789 1.17217 10 35.09 0.33047 1.17158 35.09 0.33047 1.17194 11 40.00 0.41318 1.17148 40.00 0.41318 1.1717

Figure 4.24: d vs sin2ψ plot for uncoated and coated specimens made of 20MnCr5 steel

130



fatigue specimens made of 20MnCr5 steel in various states, viz. green, case carburized

(uncoated) and case carburized – WC/C coated, are reported in Table 4.11, Table 4.12 and

Table 4.13 respectively.

Table 4.11: Results of fatigue tests conducted on specimens made of 20MnCr5 steel in green state

Stress (MPa) Cycles to failure Average Std. Dev. N1 N2 Nf

Green σ Green

398 11783 17245 14514 3862 361 25498 38673 32086 9316 315 67672 118260 92966 35771 283 157356 273353 215355 82022 255 283679 622208 452944 239376

Table 4.12: Results of fatigue tests conducted on case-carburized and tempered (uncoated) specimens made of 20MnCr5 steel


Uncoated σ Uncoated

1000 7407 10164 8786 1949 935 11661 17140 14401 3874 920 18766 26673 22720 5591 900 324294 782463 553379 323974 880 640758 2000000* --- --- 840 775436 2000000* --- --- 772 2000000* --- --- ---


Table 4.13: Results of fatigue tests conducted on case-carburized, tempered and WC/C coated specimens made of 20MnCr5 steel


WC/C σ WC/C 941 5604 8807 7206 2265 888 6963 10315 8639 2370 842 11773 19618 15696 5547 765 26115 43525 34820 12311 720 59872 101734 80803 29601

131

The data obtained from the fatigue experiments is plotted on a semi-log scale, i.e.,

Log (Nf) vs Stress (S), in Figure 4.25 along with the results of statistical analysis. For

specimens in green and WC/C coated states, the data points are observed to lie on a straight

line. The extent of finite life portion of the S-N curve for uncoated specimens was, however,

too short to be confidently ascertained as linear on semi-log scale. Nevertheless, for the

sake of comparison with other types of steels, a linear curve based on the model Y = A +

BX, described in ASTM E 739 standard [121], has been fitted to the experimental data. The

estimated values of parameters A and B in the linear model, obtained for 20MnCr5 steel

specimens in various states, are summarized in Table 4.14. The values of parameters A and

B for 20MnCr5 steel in the case carburized (uncoated) state have been computed by taking

into account the data given in only first three rows of Table 4.12. However, it was not

possible to work-out the 95% confidence bands for the range of data available for uncoated

specimens.

Table 4.14: Values of parameters in the linear model for fatigue life of 20MnCr5 steel

Material & State Parameter A Parameter B 20MnCr5 (Green) 8.210 -0.0102 20MnCr5 (Uncoated) 8.504 -0.0046 20MnCr5 (WC/C coated) 8.235 -0.0048

Unlike SAE8620 steel, the S-N graph for case carburized (uncoated) specimens made

of 20MnCr5 steel appears to be indicative of double-knee formation. The first knee forms at

around 20,000 cycles, after which, the curve remains asymptotic almost up to 106 cycles.

But failure of some specimens in the vicinity of 106 cycles implies a further dip in the

fatigue performance. The behaviour beyond this point couldn’t be ascertained with

confidence, as testing for higher number of cycles, on the order of 107, was beyond the

scope of present study. Accordingly, this portion of the graph is represented by an

extrapolated, dashed line. For the case carburized – WC/C coated specimens, the knee-

formation was observed to occur at 105 cycles.

132

As expected, the fatigue performance of 20MnCr5 steel specimens in case carburized

and tempered (uncoated as well as WC/C coated) states is significantly superior to that

noticed for specimens tested in green state. But unlike SAE8620 steel, the coated specimens

made of 20MnCr5 steel exhibit inferior fatigue behavior than their uncoated counterparts in

both low as well as high cycle fatigue regimes.

The endurance limit of the specimens in various states was estimated through

staircase tests conducted for 2×106 cycles. The results of these tests are summarized in

Table 4.15. Taking the number of less frequent event (i.e. survival or failure) as the basis

for statistical analysis of staircase tests, the data reduction is carried out according to the

Dixon-Mood method [refer section 2.3.2.4]. The intermediate calculations of the statistical

analysis along with the final results obtained for endurance limits, pertaining to specimens

in various states, are presented in Table 4.16.

Table 4.15: Results of staircase tests conducted on specimens made of 20MnCr5 steel Green Case Carburized WC/C Coated S. No.

Stress (MPa)

Cycles Stress (MPa)

Cycles Stress (MPa)

Cycles

1 255 283679 920 26673 765 43525 2 230 564563 900 324294 720 59872 3 205 2000000* 880 2000000* 675 2000000* 4 230 2000000* 900 782463 720 2000000* 5 255 622208 880 640758 765 26115 6 230 2000000* 860 2000000* 720 101734 7 --- --- 880 2000000* 675 2000000* 8 --- --- 900 2000000* --- --- 9 --- --- 920 18766 --- ---


133

Table 4.16: Results of Dixon-Mood statistical analysis of staircase test data obtained for 20MnCr5 steel specimens

i Green Case Carburized WC/C Coated




3 --- --- S3 = 920 MPa nDM,3 = 0 --- ---


( )( )∑= iDMDM niA , = 5

( )( )∑= iDMDM niB ,2 = 9

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.22 (<0.3)

Mean stress, (µs) = 234 MPa



( )( )∑= iDMDM niA , = 4

( )( )∑= iDMDM niB ,2 = 6

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.5 (≥0.3)




( )( )∑= iDMDM niA , =

( )( )∑= iDMDM niB ,2 =

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.22 (<0.3)



134

Figure 4.25: S-N graphs for 20MnCr5 steel specimens in green, case carburized (uncoated) and case carburized - WC/C coated states.

135


Fractographic observations were recorded to study the topography of surfaces

fractured under fatigue and to identify the associated fracture mechanisms at different stress

levels. The topography of specimens failed in green state at different level of stress was

found to be similar to that observed for SAE8620 steel specimens in green state [Figure

4.11]. However, the examination of fractured surfaces of case carburized and tempered

specimens in uncoated and coated states revealed the fracture surfaces to exhibit

significantly different topography and associated crack propagation mechanisms, in

comparison to their counterparts made of SAE8620 steel.

Figure 4.26 shows macrographs of fractured surfaces of uncoated as well as WC/C

coated specimens made of 20MnCr5 steel. It can be observed, that on a macroscopic scale,

the fracture surfaces of the uncoated as well as WC/C coated specimens tested at various

stress levels generally exhibit three distinct regions: (i) Outermost region exhibiting grainy

appearance and having a number of ratchet marks, indicating crack propagation through

intergranular decohesion cracking and multiple crack origin sites respectively,

(ii) Intermediate, donut shaped region, exhibiting river-patterns, characteristic of secondary-

stage transgranular fracture, generally observed in low-carbon regions [240,292] and

(iii) Innermost region, featuring final – ductile fracture due to overload. The region of final

(ductile) fracture is almost concentric with the cross-section of the specimen, which is a

typical feature of failure under pure rotating-bending fatigue [280]. The extent of ductile

fracture is observed to shrink with the increasing magnitude of maximum alternating stress

encountered during the fatigue test.

Similar to SAE8620 steel specimens, the crack in 20MnCr5 specimens (in uncoated

as well as WC/C coated states) was always found to initiate at the surface. The fracture

surfaces of the uncoated and coated specimens exhibited similar features, except that the

coated specimens tested under low cycle fatigue regime exhibited a mild tendency to form

shear lips [Figure 4.26 (e) and (f)], which was not observed in the uncoated specimens.

136

Figure 4.26: Optical fractographs showing macroscopic features on uncoated / coated

specimens made of 20MnCr5 steel, fatigued at various stress levels: (a): Uncoated specimen tested at 1000 MPa, (b): Uncoated specimen tested at 935 MPa, (c): Uncoated specimen tested at 920 MPa (d): Uncoated specimen tested at 900 MPa (e): WC/C coated specimen tested at 941 MPa (f): WC/C coated specimen tested at 888 MPa (g): WC/C coated specimen tested at 842 MPa and (h): WC/C coated specimen tested at 765 MPa.

137

Figure 4.27 shows the magnified views of regions marked with rectangles in Figure

4.26. These images are formed as mosaics covering entire swaths on the cross-section,

constructed by digitally stitching together individual images captured at successive

locations on the fractured surface. It is revealed from these images that intergranular

decohesion along prior austenitic grain boundaries is the dominant mode of crack initiation

and propagation.

The regions of stable, transgranular crack propagation are too small and difficult to

observe in most of the specimens. One such region is identified with the help of a dashed

line in Figure 4.28. Quantitative measurements taken on the region of stable crack

propagation [Refer section 2.6] provide depth, a = 79µm and width, c = 415 µm, for which,

fracture toughness, KIC = 16.62 mMPa , as estimated from Equation 2.21. This value is

quite low in comparison to the value of KIC (= 25.59 mMPa ) estimated for SAE8620

steel [Refer section 4.2.4.2], where transgranular fracture was observed to be the dominant

mode of crack advancement.

Owing to the limitations on depth-of-field in optical microscope, the topography of

surface exhibiting intergranular decohesion is better revealed by scanning electron

micrograph shown in Figure 4.29. This type of fracture is generally known to occur within

the high hardness region [294] and is also promoted by grain-boundary precipitates as well

as environmental assisted corrosion [295]. The intermediate region showing river-pattern

topography was found to exhibit cleavage-like transgranular fracture [Refer Figure 4.30].

A closer examination of the edges of fractured surfaces under optical microscope

revealed the presence of small loops of fracture zones in the vicinity of outermost layers,

that could be observed in both the specimens (uncoated and coated), as shown in Figure

4.31 (a-e). At some places, these cracks led to chipping-off of some material from the edge,

as shown in Figure 4.31 (d). This is in stark contrast to the observations made on SAE8620

steel specimens and is attributed to the formation of complex crack geometry under

intergranular crack propagation. As multiple crack fronts continue to progress while

138

changing their direction of propagation, some portions of material near the outermost

surface gets chipped-off.

Figure 4.27: Magnified views of regions marked with rectangles in the corresponding

fractographs given in Figure 4.26.

139

Figure 4.28: Micrograph indicating intergranular crack initiation and small region of

stable transgranular growth (marked by dashed-line), followed by dominantly intergranular propagation in the specimen shown in Figure 4.26 (d).

Figure 4.29: Scanning electron micrograph showing intergranular cracking in the

specimen shown in Figure 4.26 (b).

140

Figure 4.30: Magnified view of specimen shown in Figure 4.26 (a), depicting three

different regions of crack propagation: (A): Region dominated by intergranular fracture, (B): Region of cleavage-like transgranular fracture, characterized by river pattern and (C): Region of ductile fracture.

As such, no significant difference was observed in the surface topography as well as

the associated crack propagation mechanisms among the uncoated and WC/C coated

specimens, particularly for specimens that had failed at comparable number of cycles. The

formation of shear lips near the outermost region of the coated specimens subjected to low

cycle fatigue [Figure 4.26 (e) and (f)] can be attributed to the increase in magnitude of

residual compressive stresses in the outermost layers of these specimens, which might have

hindered the occurrence of failure under mode – I.

141

Figure 4.31: Magnified views showing the presence of curved-cracks in various specimens whose fractographs are shown in Figure 4.26: (a): Non-radial cracks in uncoated specimen shown in Figure 4.26 (d); (b): Looping cracks in coated specimen shown in Figure 4.26 (e); (c): Completely looped cracks in coated specimen shown in Figure 4.26 (f); (d): Composite micrograph showing chipped-off material in specimen shown in Figure 4.26 (c), and (e): Curved-cracks in coated specimen shown in Figure 4.26 (h).

142

4.3.5 Investigation of double-knee behaviour in 20MnCr5 steel

It is revealed from Figure 4.25 that while the WC/C coated specimens made of

20MnCr5 steel exhibit a well-defined endurance limit of 713MPa, the behaviour of

uncoated specimens is suggestive of double-knee formation, with estimated mean

endurance limit of 890MPa. These observations are commensurate with those reported by

Naito [162] and Parrish [276], where such behaviour is attributed to the presence of either

high temperature transformation products (HTTP) or internal oxidation [Refer section

2.4.1].

With a view to investigate the possible reason behind the double-knee behaviour, a

detailed study was performed on the failed uncoated specimens. The influence of

microstructure in the wake of crack was studied by preparing the fracture surface for

metallographic analysis. When viewed at high magnification, deep-penetrating grain-

boundary oxides could be observed in the outermost layers of polished and etched

specimens, as shown in Figure 4.32 (a) and (b). In order to ascertain the depth of oxide

formation, the specimens were polished lightly after etching with Nital. The resulting

micrographs, given in Figure 4.32 (c) and (d) reveal the presence of extensive network of

grain-boundary oxides. The oxides were found to be present up to a depth of about 18-

25µm, which is substantially greater than the critical depth of 15µm [276,296], thereby

affecting fatigue behaviour of the specimens.

In order to further confirm the presence of oxides within the surface layers, the

specimens were viewed under scanning electron microscope and EDAX elemental analysis

was performed over the region exhibiting oxidation. Figure 4.33 (a) shows the appearance

of oxides. The EDAX plot of the region marked with rectangle in Figure 4.33 (a) is shown

in Figure 4.33 (b). The results obtained through EDAX quantitative analysis [Refer Table

4.17] also confirm the presence of oxygen, thereby implying grain boundary oxidation in

the outermost surface of the specimens.

143

Figure 4.32: Micrographs indicating the presence of internal oxidation in 20MnCr5

specimens: (a) and (b): Specimens etched in 3% Nital for 5 seconds; (c): Specimen polished after etching to reveal the depth of penetration of oxides; (d): Micrograph taken in the vicinity of chipped-off portion appearing towards left end in Figure 4.27 (b); the image is digitally processed for enhancing depth of field by stitching together portions of various photographs taken by shifting the focal plane of the microscope.

Figure 4.33: (a): Scanning electron micrograph showing the presence of oxide precipitates within a grain near the surface of a case-carburized 20MnCr5 steel specimen; (b): EDAX spectrum, confirming the presence of oxygen in the region identified with a rectangle in (a).

144

Table 4.17: Results of quantitative EDAX elemental analysis Element Weight % Atomic %

C 10.83 35.33 O 1.17 2.87 Cr 1.04 0.79 Mn 1.16 0.83 Fe 85.79 60.19

Dowling et al. [297] suggested that the presence of tensile residual stresses at the

surface of fatigue specimen may indicate the presence of high temperature transformation

products (HTTP), which is also responsible for double knee behaviour, as observed in this

study for uncoated 20MnCr5 steel specimens. However, in the present study, HTTP could

not be observed in the specimens. In order to rule out the absence of tensile residual stresses

in the outermost surface of the specimen, the residual stress measurements were made

across the depth of uncoated specimens.

In order to perform these X-ray diffraction residual stress measurements, a number of

specimens were cut from the ends of uncoated fatigue specimens. These specimens were

mounted face-down on a magnetic bed and their sectioned sides were ground flat, so as to

achieve equal thickness among all the specimens. The thickness of these specimens after

grinding was measured to high degree of accuracy using a digital micrometer.

Subsequently, the carburized faces of these specimens were polished using SiC papers of

various grades, so as to remove material upto a certain, pre-determined depth from each

specimen. Difference between the initial thickness (immediately after grinding the back

face) and final thickness after polishing with SiC paper served as a measure of depth below

surface, up to which the material had been removed by the SiC paper. After polishing, the

specimens were etched in 3% Nital and finally washed with ethanol. The final depth below

original surface was determined to an accuracy of 1 µm before characterizing the specimens

for estimation of residual stresses. The results of residual stress analysis are given in Table

4.18 and shown graphically in Figure 4.34.

145

Table 4.18: Determination of lattice strain and associated residual stresses at various depths in case carburized specimens made of 20MnCr5 steel.

d-values (Å) obtained at various depths below surface S. No. sin2ψ

1 µm 24 µm 85 µm 219 µm 320 µm

1 0.41318 1.17082 1.17065 1.17085 1.17081 1.17116

2 0.33047 1.17102 1.17082 1.17100 1.17092 1.17123

3 0.24789 1.17125 1.17098 1.17116 1.17105 1.17131

4 0.16530 1.17150 1.17118 1.17129 1.17120 1.17139

5 0.08267 1.17174 1.17136 1.17148 1.17137 1.17148

6 0 1.17218 1.17158 1.17175 1.17169 1.17163

7 0.08267 1.17219 1.17149 1.17178 1.17177 1.17162

8 0.16530 1.17209 1.17137 1.17169 1.17174 1.17157

9 0.24789 1.17195 1.17124 1.17159 1.17166 1.17152

10 0.33047 1.17180 1.17108 1.17148 1.17158 1.17146

11 0.41318 1.17165 1.17096 1.17135 1.17148 1.17138

Residual Stress (MPa) -319.6 -266.4 -221.9 -184.8 -118.4

Figure 4.34: d vs sin2ψ plots for case-carburized, uncoated specimens made of 20MnCr5

steel at various depths below surface.

146

The variation of residual stress across depth of the case carburized and tempered

(uncoated) specimens made of 20MnCr5 steel is plotted on semi-logarithmic scale in Figure

4.35. As expected, the magnitude of residual compressive stress decreases with increasing

depth. The absence of tensile residual stresses at the surface of the specimens rules out the

chances of formation of HTTP in the present study.

These observations thus imply that the observed anomalous double-knee behaviour of

case carburized and tempered, uncoated specimens made of 20MnCr5 steel is due the

formation of internal oxides during carburization. These oxides provide sites of

discontinuities, or material inhomogenieties, which increasingly govern the fatigue crack

initiation and propagation [298]. Similar observations concerning shift in the knee of S-N

curve for specimens having internal oxidation are also reported by Krauss [239].

Figure 4.35: Residual stress profile in case carburized 20MnCr5 steel showing variation of

residual stress with depth below surface.

With a view to further investigate the effect of internal oxidation on the observed

fatigue behavior of case-carburized and tempered (uncoated) specimens made of 20MnCr5

steel, the fatigue tests were conducted on a specimen after removing the surface layers

affected by internal oxidation through grinding with successive grades of SiC paper.

147

Though the resulting specimen had a neck diameter of 5.906mm, which is outside the

prescribed tolerance of 6±0.05 mm, the specimen exhibited significantly improved fatigue

performance. When subjected to a stress of 1014 MPa, it survived 5,29,574 cycles, which is

much higher in comparison to the average life of 8,786 cycles, exhibited by regular

uncoated specimens tested at 1000 MPa [Refer Table 4.12].

Another interesting observation was that the crack now had a sub-surface origin,

thereby forming a fish-eye, which appears as a bright spot towards the top of Figure 4.36

(a). Fish eye features appear bright because the freshly fractured surface does not come in

contact with air till the crack-front reaches the outermost surface. When viewed at higher

magnification, radial marks emanating from the crack initiation site can be observed [Refer

Figure 4.36 (b)]. Examination at still higher magnification reveals the optically dark area

(ODA) at the site of crack origin, as shown in Figure 4.37. ODAs are known to appear only

for the specimens fractured after large number of cycles and are absent in case of specimens

failed at small number of cycles [294]. This area appears dark because of different fracture

features, which give rise to a very rough morphology.

Figure 4.36: Optical fractographs of uncoated specimen, polished to remove surface layers

affected by internal oxidation, fatigue tested at 1014 MPa. (a): Fracture macrograph showing the formation of fish-eye; (b): Magnified view of the fish-eye appearing in (a).

148

Figure 4.37: Optical fractograph of ODA at the crack initiation site within the fish-eye.

All these observations pertaining to sub-surface crack origin imply that a defect-free

surface, obtained by complete removal of oxidation-affected layers, leads to a remarkable

improvement in fatigue performance of the case carburized and tempered specimens made

of 20MnCr5 steel.

4.3.6 Effect of WC/C coating on fatigue behaviour of 20MnCr5 steel

It is revealed from the investigation that the fatigue behaviour of case carburized –

WC/C coated specimens is quite inferior to that observed for case carburized (uncoated)

specimens made of 20MnCr5 steel. Though the magnitude of residual compressive stress

among the coated specimens is considerably higher than the residual stress in the uncoated

specimens, the dominance of crack propagation through intergranular decohesion cracking,

coupled with grain-boundary oxidation, seems to be responsible for inferior fatigue

behaviour of the coated specimens. It appears that since the hardness of steel gets

149

compromised to some extent during the coating process [Refer Figure 4.2], the cracks find

an easy passage along weak boundaries of prior austenitic grains, especially in the vicinity

of oxidized layer.

The observed fatigue limit of 713 MPa for coated specimens in the present study is

around 5% higher than the value of 680 MPa reported by Brugger [161] for 20MnCr5 steel

that was free from HTTP formation; except that this limit is found to occur at 105 cycles in

our case. For the uncoated specimens, the first knee is found to form at a considerably

higher stress level of around 890 MPa. However, after remaining horizontal for some

extent, the graph again exhibits decline in the vicinity of 106 cycles. Though the fatigue

behaviour beyond this point could not be ascertained with confidence, as tests beyond

2×106 cycles were not possible due to experimental limitations, still, the double-knee

formation is clearly implied by the results obtained.

150

4.4 Testing and analysis of EN353 steel specimens

The influence of case carburization, followed by subsequent coating with WC/C layer

on the fatigue performance of EN353 steel is presented in the following sub-sections.

Following the procedure similar to that described in sections 4.2 and 4.3 for other steels

investigated in this study, the fatigue performance of EN353 steel has also been evaluated

under various states, viz. green, case carburized (uncoated) and case carburized – WC/C

coated and the results obtained have been correlated with the microstructure, hardness and

residual stress to identify and explain the prevailing crack propagation mechanisms.


A mosaic of metallographs, showing variation of microstructure across depth in a

case-carburized specimen made of EN353 steel, is shown in Figure 4.38. In order to prepare

the metallograph, the specimen was first etched in 3% Nital for 5 seconds, followed by

immersion in potassium metabisulfite solution for 20 seconds. The difference in thickness

of the film deposited on various microstructural constituents led to formation of different

colours through interference of light. The plate martensite is imparted yellow to green tint,

while the matrix phase appears deep red in this mosaic of slightly over-etched

metallographic specimen.

Some internal oxidation near the outermost edge of the specimen can be observed in

these metallographs. The oxides appear as a network of dark bands along prior austenitic

grain boundaries. The depth of penetration (~ 15 – 20 µm) of these oxides is on the order of

austenitic grain size. Figure 4.39 (a) shows high-contrast metallograph, vividly revealing

the presence of internal oxidation along prior austenitic grain boundaries. The region

marked with rectangle in Figure 4.39 (a) is shown magnified in Figure 4.39 (b), where

inter-grain as well as intra-grain oxides can be observed separately. The formation of these

oxides generally results due to oxidation of chromium and manganese [157,158] present in

EN353 steel [Refer Table 3.1].

151

Figure 4.38: Mosaic of metallographs, showing variation in microstructure with depth in case carburized and tempered specimens made of EN353 steel, etched in 3% Nital for 5 seconds, followed by immersion in potassium metabisulfite solution for 20 seconds.

152

A careful examination of Figure 4.38 reveals that martensitic structures occupy

different volume fractions at different depths below the outermost surface. The outermost

region, extending to a depth of around 50 – 70 µm, contains high volume fraction of plate

martensite, exhibiting fine acicular morphology (etched yellow-green). This region is

followed by a sudden increase in the content of retained austenite (etched deep-red), which

extends over a band of around 100 µm width. At a depth of around 160 µm, the volume

fraction of plate martensite (etched green) again begins to increase and becomes highest at

depths ranging from 450 to 750 µm. Beyond a depth of 750 µm (which corresponds to the

effective depth of case), a mixture of plate and lath martensite can be observed, which

gradually gives way to pure lath martensite, extending right into the core of the specimen,

which is vividly revealed through nital-etch [Refer Figure 4.40 (a) and (b)].

At higher magnifications, segregates of carbide particles could be seen dispersed in

the case, as revealed in the colour metallograph shown in Figure 4.40 (c). In this

metallograph, the red matrix phase corresponds to retained austenite, while the carbide

precipitates can be seen as bright white spots, some of which have been identified with

arrow-marks.

The martensitic start (Ms) temperature, as estimated from Equation 4.1, corresponding

to the chemical composition of EN353 [Refer Table 3.1], works out to be 367°C, which

corresponds to 94% volume fraction of martensitic transformation [Refer Equation 4.2] at a

quenching temperature of 120°C [Refer section 3.2]. This is in agreement with the high

volume fraction of lath martensite observed in the core, where the chemical composition

essentially remains unaltered even after case carburization.

The value of Ms temperature obtained for this steel is quite comparable to that

calculated for 20MnCr5 steel in section 4.3.1, but is significantly lower than that obtained

for SAE8620 steel [Refer section 4.2.1]. Thus, during quenching of EN353 steel specimens,

the martensitic transformation in the core would have commenced at a relatively later stage

in comparison to that in the specimens made of SAE8620 steel. This altered timing would

153

have adversely affected the possibility of strain-induced martensitic transformation near the

surface of specimens made of EN353 steel. Further, the underlying band of retained

austenite (observed at a depth ranging from 60µm to 160µm) would have served to provide

a soft base – incapable of transmitting stresses to the layer above. These arguments explain

the relatively lower volume fraction of martensite in the outermost layers of EN353 steel

specimens, as compared to those made of SAE8620 steel.

Figure 4.39: (a): Metallograph indicating the presence of internal oxidation in case

carburized and tempered specimens made of EN353. The specimen was lightly polished after etching with nital; and (b): Magnified view of region marked with rectangle in (a).

154

Figure 4.40: Metallographs of the cross section of case carburized and tempered specimen made of EN353 steel. (a) and (b): Lath martensite in the core; and (c): Composite metallograph showing martensite, retained austenite and chunks of carbide precipitates (marked with arrows) within the carburized and tempered case. Specimens in (a) and (b) etched with 3% nital for 5s, specimen in (c) were etched with 3% nital for 4s, followed by immersion in potassium metabisulfite for 12s.


The X-ray diffraction residual stress estimation was performed according to the

procedures described earlier [Refer sections 2.5.2, 3.7.2 and 4.2.2]. The results obtained are

summarized in Table 4.19 and depicted graphically in Figure 4.41. Similar to the

observations made on SAE8620 and 20MnCr5 steels, slight ψ-splitting has also been

observed in the diffraction patterns recorded for EN353 steel.

155

Table 4.19: Determination of lattice strain in EN353 steel specimens through X-ray diffraction

Uncoated Specimen Coated Specimen S. No.

Tilt, ψ sin2ψ d-spacing Tilt, ψ sin2ψ d-spacing

1 -40.00 0.41318 1.17026 -40.00 0.41318 1.17043

2 -35.09 0.33047 1.17037 -35.09 0.33047 1.17069

3 -29.86 0.24789 1.17053 -29.86 0.24789 1.17091

4 -23.99 0.16530 1.17069 -23.99 0.16530 1.17118

5 -16.71 0.08267 1.17087 -16.71 0.08267 1.17144

6 0 0 1.17113 0 0 1.17173

7 16.71 0.08267 1.17111 16.71 0.08267 1.17159

8 23.99 0.16530 1.17102 23.99 0.16530 1.17137

9 29.86 0.24789 1.17092 29.86 0.24789 1.17115

10 35.09 0.33047 1.17079 35.09 0.33047 1.17090

11 40.00 0.41318 1.17068 40.00 0.41318 1.17069

Figure 4.41: d vs sin2ψ plot for uncoated and coated specimens made of EN353 steel

156

The residual stress is estimated to be -225.0 ± 3.6 MPa for uncoated and -399.9 ± 3.8

MPa for coated specimens. Thus, for this steel too, the case carburized surface contains

residual compressive stresses, which undergo considerable enhancement upon coating with

WC/C layer through physical vapour deposition. The surface residual stresses in the

uncoated and coated specimens made of EN353 steel are significantly lower than their

counterparts made of SAE8620 steel. This is attributed to relatively higher amount of

retained austenite in the vicinity of outermost surface of EN353 steel specimens [Figure

4.38] than that observed in specimens made of SAE8620 steel [Figure 4.4 (a), Figure 4.6].



fatigue specimens made of EN353 steel in various states, viz. green, case carburized

(uncoated) and case carburized – WC/C coated are reported in Table 4.20, Table 4.21 and


Table 4.20: Results of fatigue tests conducted on specimens made of EN353 steel specimens in green state


Green σ Green

555 6730 8775 7753 1446 537 9113 21938 15526 9069 480 15675 27846 21761 8606 426 33456 66488 49972 23357 357 105672 194063 149868 62502 317 171248 353672 262460 128993 270 316893 667382 492138 247833

Table 4.21: Results of fatigue tests conducted on case-carburized and tempered (uncoated) specimens made of EN353 steel



1000 8467 15264 11866 4806 925 20752 66754 43753 32528 860 69815 98762 84289 20469 765 176743 327483 252113 106589 686 289467 615463 452465 230514 663 556733 925362 741048 260660

157

Table 4.22: Results of fatigue tests conducted on case-carburized, tempered and WC/C coated specimens made of EN353 steel


WC/C σ WC/C 923 4563 8727 6645 2944 843 16220 22881 19551 4710 726 122589 239346 180968 82560 680 298356 445892 372124 104324 655 566229 896546 731388 233569

Similar to the procedure followed earlier for other steels [Refer sections 4.2.3 and

4.3.3], the change in fatigue life as a function of maximum alternating stress, has been

expressed in terms of linear model Y = A + BX, described in ASTM E 739 standard [121].

This linear model has been fitted separately to the fatigue data recorded for specimens in

green, case-carburized (uncoated) and case carburized – WC/C coated states. The estimated

values of parameters A and B in the linear model are summarized in Table 4.23.

Table 4.23: Values of parameters in the linear model for fatigue life of EN353 steel

Material & State Parameter A Parameter B EN353 (Green) 7.304 -0.0061 EN353 (Uncoated) 9.171 -0.0050 EN353 (WC/C coated) 10.787 -0.0076

The endurance limits of the specimens in various states were ascertained by means of

staircase tests, conducted for 2×106 cycles, the results of which are summarized in Table

4.24. The data obtained from the staircase tests is processed according to the Dixon-Mood

method [Refer section 2.3.2.4], by taking into account the number of less frequent event –

i.e., survival or failure. The outline of analysis, along with the results obtained, is presented

in Table 4.25.

158

Table 4.24: Results of staircase tests conducted on specimens made of EN353 steel Green Case Carburized WC/C Coated S. No.

Stress (MPa)

Cycles Stress (MPa)

Cycles Stress (MPa)

Cycles

1 270 667382 686 615463 680 445892 2 245 2000000* 663 925362 655 2000000* 3 270 316893 640 784560 680 298356 4 245 754038 617 2000000* 655 566229 5 220 1209843 640 2000000* 630 789473 6 195 2000000* 663 556733 605 2000000* 7 220 2000000* 640 2000000* 630 2000000* 8 --- --- 663 2000000* 655 896546 9 --- --- 686 289467 630 2000000*


Table 4.25: Results of Dixon-Mood statistical analysis of staircase test data obtained for

EN353 steel specimens i Green Case Carburized WC/C Coated 0 S0 = 195 MPa nDM,0 = 1 S0 = 617 MPa nDM,0 = 1 S0 = 605 MPa nDM,0 = 1





( )( )∑= iDMDM niA , = 1

( )( )∑= iDMDM niB ,2 = 1

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.22 (<0.3) Mean stress, (µs) = 233 MPa Std. dev. (σs) = 28.17 MPa


( )( )∑= iDMDM niA , = 4

( )( )∑= iDMDM niB ,2 = 6

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.5 (≥0.3) Mean stress (µs)= 652 MPa Std. dev. (σs) = 19.71 MPa


( )( )∑= iDMDM niA , = 4

( )( )∑= iDMDM niB ,2 = 6

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.5 (≥0.3) Mean stress (µs)= 643 MPa Std. dev. (σs)= 21.42 MPa

159

The experimental fatigue data, along with the results estimated from the fitted linear

models and 95% confidence bands, is plotted in Figure 4.42. It is observed from the figure

that Log (Nf) vs Stress (S) curves for EN353 steel specimens in various states follow

straight-line paths. The case carburized specimens in uncoated as well as WC/C coated

condition exhibit much better performance than the specimens in green state. The

endurance limits for case carburized (uncoated) as well as case carburized – WC/C coated

steels are nearly 180% higher than the endurance limit for steel in green state. The fatigue

performance of case carburized and tempered (uncoated) specimens deteriorates upon

application of WC/C coating. The loss observed in fatigue strength diminishes with

increasing number of cycles to failure (Nf).

The contribution of case carburization as well as application of WC/C coating on

carburized specimens, towards percentage change in fatigue strength of the specimens

subjected to a given number of stress cycles, as computed from Equation 4.3, is given in

Table 4.26 and depicted graphically in Figure 4.43.

These observations are somewhat similar to those for SAE8620 steel [Refer section

4.2.3] except that the coated specimens exhibit poor fatigue response than the uncoated

ones even at stress cycles beyond 105. However, the endurance strengths of uncoated and

WC/C coated specimens are observed to be quite identical.

Table 4.26: Effect of case carburization and subsequent WC/C coating on the fatigue

strength of EN353 steel. Sf

Green SfUncoated Sf

WC/C %Change in Fatigue Strength

A= 7.3040 A=9.1708 A=10.7869


B=-0.00608 B=-0.00503 B=-0.00763 Case

Carburized vs Green

WC/C Coated

vs Green 1000 708 1228 1021 73 44

10000 543 1029 889 89 64 100000 379 830 758 119 100

1000000 214 631 627 194 193

160

Figure 4.42: S-N graphs for EN353 steel specimens in green, case carburized (uncoated) and case carburized - WC/C coated states.

161

Figure 4.43: Percentage change in fatigue strength of uncoated and coated EN353 steel specimens with reference to specimens in green state.

162


Fractomacrographs recorded on case carburized, uncoated as well as WC/C coated

specimens made of EN353 steel, failed at different number of cycles, are shown in Figure

4.44. As observed for 20MnCr5 steel, on a macroscopic scale, the fracture surface generally

comprises of three regions: (i) Outermost region exhibiting grainy appearance and having a

number of ratchet marks, indicating multiple sites of crack origin, (ii) Intermediate, donut

shaped region, exhibiting river-patterns, implying secondary stage transgranular mode of

crack propagation [240,292] and (iii) Innermost region, featuring final – ductile fracture.

The region of final fracture is more or less concentric with the cross-section of specimen

and its extent decreases with the increase in magnitude of alternating stress applied during

fatigue test. As observed earlier for other steels, the fatigue crack in this steel too was

always found to initiate at the outer surface.

Fractographic studies conducted at higher magnification revealed that the cracks

originated at the outer surface through intergranular decohesion along prior austenitic grain

boundaries [Refer arrow mark in Figure 4.45], followed by stable – transgranular

propagation for some time. As discussed earlier in section 4.2.4.2, this type of crack

initiation and propagation are commonly observed fracture modes in the carbon rich case of

the carburized steels.

In the micrograph shown in Figure 4.45, the width (c) and depth (a) of stable

transgranular crack propagation region (marked with dashed line) are found to be 363µm

and 165µm respectively. Using these values in Equation 2.21, the fracture toughness, KIC of

the case carburized EN353 steel works out to be 17.38 mMPa . Therefore, the fracture

toughness of EN353 is slightly better than 20MnCr5 steel, but is considerably lower than

that observed for SAE8620 steel in similar state.

163

Figure 4.44: Optical fractographs showing macroscopic features in uncoated and WC/C coated specimens made of EN353 steel, fatigued at various loads: (a): Coated specimen tested at 843 MPa, (b): Uncoated specimen tested at 925 MPa, (c): Coated specimen tested at 726 MPa (d): Uncoated specimen tested at 860 MPa (e): Coated specimen tested at 655 MPa, and(f): Uncoated specimen tested at 765 MPa.

164

It is observed that the region of unstable crack propagation contains a substantial

percentage of intergranular decohesion along prior austenitic grain boundaries [Refer

Figure 4.45 and Figure 4.46]. Upon reaching the region with relatively low carbon content,

the crack is observed to advance almost entirely by cleavage-like transgranular fracture,

where the surface morphology is characterized by river patterns [Figure 4.47]. The

comparison of Figure 4.46 and Figure 4.47 with Figure 4.29 reveals that the fraction of

intergranular mode of crack propagation observed in EN353 steel is somewhat less than

that observed in 20MnCr5 steel specimens.

Figure 4.45: Micrograph of specimen shown in Figure 4.44 (d), indicating crack initiation

by intergranular cracking (marked with arrow).

In line with the observations made on other steels investigated in the present work, no

significant differences in crack initiation and propagation mechanisms could be observed

among the case carburized – uncoated as well as WC/C coated specimens made of EN353

steel, which had failed at comparable number of fatigue cycles. Only the stress level for

their occurrence gets modified in the presence of WC/C coating.

165

Figure 4.46: Scanning electron micrograph of specimen shown in Figure 4.44 (a),

indicating a mix of transgranular and intergranular cracking.

Figure 4.47: Optical fractograph of specimen shown in Figure 4.44 (c), indicating various

regions of crack propagation. The fractograph is constructed as a mosaic by stitching together four individual fractographs.

166

4.4.5 Effect of WC/C coating on fatigue behaviour of EN353 steel

It is evident from the results obtained [Table 4.25 and Figure 4.42] that the fatigue

behaviours of case carburized (uncoated) and case carburized – WC/C coated specimens

made of EN353 steel are quite identical in the high cycle regime. However, the coated

specimens exhibit slightly inferior fatigue performance in comparison to their uncoated

counterparts. Unlike 20MnCr5 steel specimens, the endurance limits are observed to be

quite comparable for the uncoated and coated specimens made of EN353 steel.

The coating process had led to a slight compromise in hardness [Figure 4.2], but

significant gain in terms of residual compressive stress [Refer section 4.4.2]. However, as

the cracks propagate through intergranular decohesion, the residual stresses induced in the

thin coating seem to have little influence on them. The presence of internal oxidation along

prior austenitic grain boundaries offers innumerable crack initiation sites on the surface,

which are not effectively plugged by the thin, hard and brittle coating. Rather, any crack

developed in the coating as a result of poor cohesion among the underlying grains would

tend to aggravate the problem by providing site for viable crack growth on the surface,

whose hardness has already been compromised somewhat by the coating process.

For the sake of comparison with other steels, a more aggravated situation of this kind

was encountered for 20MnCr5 steel, where the extent of grain boundary oxidation was

more severe, while the crack propagation was almost entirely intergranular in nature. Both

these factors seemed to severely impair the fatigue performance of 20MnCr5 steel

specimens in coated conditions. On the other hand, in specimens made of SAE8620 steel,

the benefits of enhanced residual compressive stresses upon coating with WC/C layer could

be harnessed because of largely transgranular crack propagation and insignificant amount

of internal oxidation at the surface. In the backdrop of these observations, the effect of

WC/C coating on fatigue behaviour of EN353 steel appears to lie in-between the effects

observed in SAE8620 and 20MnCr5 steels.

167

4.5 Testing and analysis of SCM420 steel specimens

The influence of case carburization and physical vapour deposition of WC/C layer on

the fatigue performance of SCM420 steel is presented in the following sub-sections. The

observed fatigue behaviour of specimens in various states is correlated with the influence of

material processing on the resulting microstructure, hardness, residual stresses and crack

propagation mechanisms.


The variation in microstructure with depth in case carburized specimens made of

SCM420 steel is captured in the mosaic of metallographs shown in Figure 4.48. The

specimen was first etched with 3% nital for revealing the microstructure, followed by

immersion in potassium metabisulfite solution [Refer Table 3.2] for deposition of anodic

film to enhance the colour contrast among various constituent phases.

In the micrograph, one can observe fine plate martensitic structure in the form of

bluish, needle-shaped morphology, along with retained austenite, which provides light

purple background. Since the martensitic start (Ms) temperature increases with depth due to

negative carbon gradient, the highest volume fraction of martensite is generally found to

occur at some depth below the surface during quenching. In Figure 4.48, the highest

volume fraction of plate martensite can be observed, first upto a depth of about 60µm,

followed by another band at a depth of around 160µm to 230µm. The formation of outer

martensitic layer is attributed to relatively higher quenching rate at the surface, along with

facilitation of martensitic transformation due to introduction of tensile stresses in the

surface layer during the quenching of steel, when the sub-surface layer undergoes

martensitic transformation. Fine carbide precipitates within the carbon-rich case were

revealed upon colour-etching with Klemm’s – I solution, as shown in Figure 4.49.

168

Figure 4.48: Mosaic of metallographs, showing variation of microstructure with depth in case carburized and tempered specimens made of SCM420 steel, etched in 3% Nital for 4 seconds, followed by immersion in potassium metabisulfite solution for 9 seconds.

169

The observed band of retained austenite at depths ranging from 60µm to 160µm is

similar to that reported for specimens made of EN353 steel [Refer Figure 4.38]. However,

the fraction of retained austenite observed in SCM420 steel specimens is considerably less

than that observed in EN353 specimens, which is also reflected in the corresponding

hardness measurements made on these steels [Refer Figure 4.2].

The martensitic start temperature (Ms) corresponding to chemical composition in the

core of SCM420 steel [Refer Table 3.1] is estimated to be 380°C [Refer Equation 4.1],

which corresponds to around 95% volume fraction of martensite for quenching carried at

120°C [272]. The metallographs recorded near the core of SCM420 steel specimens also

confirm such a high volume fraction of martensite. Further, a higher value of Ms

temperature for core is helpful in bringing-about a more complete martensitic

transformation near the surface due to introduction of tensile stresses at an early stage of

quenching. This is in line with the observations reported in the preceding paragraphs.

Figure 4.49: Metallograph of SCM420 specimen, colour etched with Klemm’s – I reagent for 2 minutes to reveal the presence of carbides (marked with arrows). Green tint employed for contrast enhancement.

170


The results of X-ray diffraction residual stress measurements, recorded by following

the procedure outlined in sections 2.5.2 and 3.7.2, are summarized in Table 4.27 and shown

graphically in Figure 4.50. The residual stress is found to be -261.7±3.6 MPa for the

uncoated and -531.8±3.8 MPa for the coated specimens. Among various steels analyzed in

the present study, the magnitude of residual stress observed for case carburized specimens

made of SCM420 steel is the second highest after SAE8620 steel. This is attributed to the

presence of high volume fraction of martensite in the outermost layers of these specimens.

The deposition of WC/C coating on the surface of case carburized specimens led to an

enhancement of 103% in the magnitude of residual compressive stresses. This gain in

compressive stresses is third highest after SAE8620 and 20MnCr5 steels, among the four

steels investigated in the present work.

Table 4.27: Determination of lattice strain in SCM420 steel specimens through X-ray

diffraction Uncoated Specimen Coated Specimen S. No.

Tilt, ψ sin2ψ d-spacing Tilt, ψ sin2ψ d-spacing 1 -40.00 0.4132 1.17099 -40.00 0.4132 1.17110 2 -35.09 0.3305 1.17113 -35.09 0.3305 1.17140 3 -29.86 0.2479 1.17127 -29.86 0.2479 1.17172 4 -22.99 0.1525 1.17139 -22.99 0.1525 1.17203 5 -16.71 0.0827 1.17149 -16.71 0.0827 1.17231 6 0.00 0 1.17149 0.00 0 1.17254 7 16.71 0.0827 1.17118 16.71 0.0827 1.17213 8 22.99 0.1525 1.17096 22.99 0.1525 1.17179 9 29.86 0.2479 1.17076 29.86 0.2479 1.17143 10 35.09 0.3305 1.17058 35.09 0.3305 1.17111 11 40.00 0.4132 1.17043 40.00 0.4132 1.17078

171

Figure 4.50: d vs sin2ψ plot for uncoated and coated specimens made of SCM420 steel



fatigue test specimens made of SCM420 steel in various states, viz. green, case carburized

(uncoated) and case carburized – WC/C coated, are reported in Table 4.28, Table 4.29 and


Table 4.28: Results of fatigue tests conducted on specimens made of SCM420 steel in

green state Stress (MPa) Cycles to failure Average Std. Dev.

N1 N2 Nf Green σ

Green 445 8267 15672 11970 5236 410 23560 39863 31712 11528 350 82389 176732 129561 66711 310 147364 336448 241906 133703 280 269671 694063 481867 300090

172

Table 4.29: Results of fatigue tests conducted on case-carburized (uncoated) specimens made of SCM420 steel



980 9278 17209 13244 5608 945 14790 27367 21079 8893 900 36273 98375 67324 43913 850 102703 280382 191543 125638 815 238450 616892 427671 267599 780 567392 1002832 785112 307903

Table 4.30: Results of fatigue tests conducted on case-carburized, tempered and WC/C coated specimens made of SCM420 steel


WC/C σ WC/C 940 6748 14673 10711 5604 915 17783 32453 25118 10373 880 58749 118637 88693 42347 850 98834 252016 175425 108316 820 240679 536284 388482 209024

The change in fatigue life as a function of maximum alternating stress is represented

in terms of linear model Y = A + BX [121], obtained by fitting individually the set of

fatigue data recorded for SCM420 specimens in various states. The estimated values of

parameters A and B appearing in the model [Refer section 4.2.3] are summarized in Table

4.31. The fatigue data obtained through experiments and fitted linear models are

represented in Log (Nf) vs Stress (S) graph, along with hyperbolic bands of 95% confidence

limits in Figure 4.51. The S-N curves for the steel specimens in various states are found to

exhibit linear trend on semi-log scale.

Table 4.31: Values of parameters in the linear model for fatigue life of SCM420 steel Material & State Parameter A Parameter B

SCM420 (Green) 8.2849 -0.0094 SCM420 (Uncoated) 13.0251 -0.0092 SCM420 (WC/C coated) 16.1665 -0.0129

173

The endurance limit of the specimens in various states was ascertained by conducting

staircase tests for 2×106 cycles, as described earlier for other steel grades employed in this

study. The results of these tests are summarized in Table 4.32. The results obtained by

performing statistical analysis of staircase test data through Dixon-Mood method [Refer

section 2.3.2.4] are given in Table 4.33.

The endurance limit for the SCM420 steel in case carburized (uncoated as well as

WC/C coated) conditions is found to be nearly three times higher than the endurance limit

observed for SCM420 steel specimens in green state. The comparison of case carburized

(uncoated) and case carburized - WC/C coated specimens reveals around 2.6%

improvement in endurance limit upon coating.

The S-N curves shown in Figure 4.51 indicate that the coated specimens exhibit a

relatively lesser decline in the fatigue strength with increasing number of cycles (Nf) than

their uncoated counterparts. The S-N curves of the uncoated and coated specimens cross-

over each other at a stress cycle of around 2.4×105. Therefore, the deposition of coating on

case carburized specimens improves their fatigue performance in the high cycle fatigue

regime beyond Nf = 2.4×105, but for Nf < 2.4×105, the uncoated fatigue specimens exhibit

better fatigue performance than their coated counterparts. The observed behaviour is quite

similar to that reported for SAE8620 steel in section 4.2.3.

Table 4.32: Results of staircase tests conducted on specimens made of SCM420 steel

Green Case Carburized WC/C Coated S. No. Stress (MPa)

Cycles Stress (MPa)

Cycles Stress (MPa)

Cycles

1 280 269671 815 238450 820 536284 2 260 2000000* 780 1002832 795 2000000* 3 280 694063 745 2000000* 820 240679 4 260 564378 780 2000000* 795 974782 5 240 780672 815 616892 770 784560 6 220 2000000* 780 567392 745 2000000* 7 240 2000000* 745 836471 770 2000000* 8 260 2000000* 710 2000000* --- --- 9 --- --- 745 2000000* --- ---


174

Table 4.33: Results of Dixon-Mood statistical analysis of staircase test data obtained for SCM420 steel specimens

i Green Case Carburized WC/C Coated






( )( )∑= iDMDM niA , = 9

( )( )∑= iDMDM niB ,2 = 23

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.69 (≥0.3)

Mean stress, (µs) = 255 MPa



( )( )∑= iDMDM niA , = 4

( )( )∑= iDMDM niB ,2 = 6

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.5 (≥0.3)




( )( )∑= iDMDM niA , = 3

( )( )∑= iDMDM niB ,2 = 5

( )2,

2,

∑∑ −

iDM

DMiDMDM

n

AnB

= 0.67 (≥0.3)


Std. dev. (σs)= 28.17 MPa

175

Figure 4.51: S-N graphs for SCM420 steel specimens in green, case carburized (uncoated) and case carburized - WC/C coated states.

The contribution of case carburization as well as WC/C coating towards percentage

change in the fatigue strength of specimens subjected to a given number of stress cycles

(Nf), as estimated from Equation 4.3, is given in Table 4.34 and depicted graphically in

Figure 4.52.

176

Table 4.34: Effect of case carburization and subsequent WC/C coating on the fatigue strength of SCM420 steel.

SfGreen Sf

Uncoated SfWC/C %Change in Fatigue Strength

A=8.2849 A=13.0251 A=16.1665


B=-0.00937 B=-0.00916 B=-0.01289

Case Carburized

vs Green

WC/C Coated

vs Green 1000 564 1094 1021 94 81

10000 457 985 944 115 106

100000 351 876 866 150 147

1000000 244 767 789 214 223

Figure 4.52: Percentage change in the fatigue strength of uncoated and coated SCM420 steel specimens with respect to specimens in green state.

It is observed from Figure 4.52 that both the uncoated and coated specimens show

substantial improvement in fatigue performance as compared to specimens in green state.

This enhancement in fatigue performance is attributed to the presence of high volume

177

fraction of martensite in the outermost layers and the associated high magnitude of residual

compressive stresses induced in the uncoated as well as WC/C coated specimens.

The percentage gain in fatigue strength of both types of specimens increases with

increasing number of cycles, with a relatively steeper increase for the coated specimens. At

stresses higher than 835 MPa, which corresponds to a fatigue life of 2.4×105 cycles, the

uncoated specimens exhibit better performance than the coated specimens. However, the

coated specimens exhibit better fatigue performance at stress levels below 835 MPa (Nf >

2.4×105 cycles). The observed gain in fatigue strength increases with the increase in

number of cycles until the endurance limit of the coated specimens is reached.


Fractographic observations were made on case carburized, uncoated as well as WC/C

coated specimens made of SCM420 steel, which had failed at different stress cycles. The

low magnification fractomacrographs are shown in Figure 4.53 (a) to (d). High

magnification micrographs of the crack-initiation regions (identified with arrow-marks),

recorded in the same orientation, are placed adjascent to the respective fractomacrographs.

It can be observed from the fractographs that on a macroscopic scale, the fracture

surface is comprised of an outermost, donut-shaped region, dominated by transgranular

mode of crack advancement, which is surrounding the inner region of rapid fracture. The

fracture surface topography corresponding to major portion of fatigue life of the specimens

is thus dominated by transgranular mode of crack advancement, which is quite similar to

that observed for specimens made of SAE8620 steel [Refer section 4.2.4.2], but is contrary

to that observed for 20MnCr5 [Refer section 4.3.4] and EN353 [Refer section 4.4.4] steel

specimens.

From the magnified fractographs shown in the right column of Figure 4.53, it is

revealed that the crack propagates along multiple plateaus, which have concave or convex

morphology and are at different elevations with respect to each other, joined by tear ridges.

Though such features are not apparent for the specimen shown in Figure 4.53 (d), but a

178

further magnified view of the crack initiation site shown in Figure 4.54 is indicative of

crack propagation along two different planes, with slightly different orientations.

On the basis of measurements taken on the stable crack propagation region marked

with dashed curve in Figure 4.53 (a), the value of fracture toughness, KIC [Refer Equation

2.21] for the case carburized specimens made of SCM420 steel is estimated as

21.64 mMPa . Among the four different steels investigated in this study, the fracture

toughness of SCM420 is the second highest after SAE8620.

A composite, magnified fractograph covering a complete diametric swath of the

specimen shown in Figure 4.53 (d) is provided in Figure 4.55, to reveal fracture surface

topography over the entire cross-section. The specimen survived for more than 5×105

cycles in fatigue and reveals a substantial extent of stage-I crack propagation region, shown

further magnified in Figure 4.56, which is typical of specimens tested under high cycle

fatigue [295]. When viewed at still higher magnification [Refer Figure 4.57], the region was

found to be dominated by factory-roof morphology, implying crystallographic orientation

of crack-propagation [298], accompanied by a change in direction at the grain boundaries.

The fracture surface of the specimen tested under low cycle fatigue [Figure 4.53 (b)]

exhibited crack-initiation on multiple planes, which can be seen in both the views of

composite fractomicrograph in Figure 4.58. The cracks are identified by arrow-marks in the

top and side views. Crack-initiation on multiple planes for the specimens tested under low

cycle fatigue is also confirmed by formation of a “double edge”, as shown in Figure 4.59.

These observations imply that multiple cracks are formed when the coated specimens are

subjected to low cycle fatigue. This observation is quite similar to that reported in section

4.2.4.2 for SAE8620 steel specimens, and can be attributed to substantial straining of the

substrate under high applied stresses, which leads to formation of multiple cracks within the

coating, some of which grow to macroscopic scale. This phenomenon seems to be

responsible for relatively poor fatigue performance of the coated specimens in comparison

to the uncoated ones under low cycle fatigue.

179

Figure 4.53: Optical fractographs of SCM420 steel specimens. (a): Uncoated specimen

tested at 980 MPa, (b): Coated specimen tested at 940 MPa, (c): Uncoated specimen tested at 780 MPa, and (d): Coated specimen tested at 820 MPa. Fractomicrographs given in the right column provide magnified views of crack initiation sites, captured by holding the specimen in same orientation as in the left column.

180

Figure 4.54: Magnified optical fractograph of specimen shown in Figure 4.53 (d), depicting crack initiation site

Figure 4.55: Optical fractograph covering entire cross-section of fractured specimen shown in Figure 4.53 (d). The fractograph is constructed as mosaic by joining together six individual fractographs.

181

Figure 4.56: Optical fractograph of a stage-I crack propagation site in the specimen

shown in Figure 4.53 (d).

Figure 4.57: Optical fractograph showing cleavage-like crystallographically oriented

stage-I fatigue fracture exhibiting “factory-roof” morphology, recorded at the region marked with rectangle in Figure 4.56.

182

Figure 4.58: Composite optical fractograph showing formation of multiple cracks in specimen shown in Figure 4.53 (b), tested under low-cycle fatigue.

Figure 4.59: Optical fractograph showing multiple-plane cracking for specimens tested under low-cycle fatigue.

183

4.5.5 Effect of WC/C coating on fatigue behaviour of SCM420 steel

The volume fraction of martensite in the outermost layers, as well as the magnitude of

residual compressive stresses, observed in SCM420 steel specimens are quite similar to

those observed for specimens made of SAE8620 steel. Similar to the observations made on

all other steels examined in the present study, the coating led to substantial enhancement in

the residual stress field in case carburized SCM420 steel. However, unlike 20MnCr5 and

EN353 steel specimens, hardly any internal oxidation could be observed in the specimens

made of SCM420 steel. Further, the nature of crack propagation was found to be largely

transgranular, similar to that observed in SAE8620 steel specimens, though the region of

stable crack propagation was considerably smaller than that observed for SAE8620

specimens processed in a similar manner.

An examination of the S-N graph given in Figure 4.51 reveals that in comparison to

their case carburized (uncoated) counterparts, the case carburized – WC/C coated

specimens made of SCM420 steel had poor fatigue performance under low cycle fatigue,

but exhibited improved performance under high cycle fatigue. The calculations based on

fitted linear fatigue models indicate that the two graphs intersect at a value of Nf = 2.4×105.

The observed fatigue performance, along with metallographic and fractographic

observations, clearly indicates a strong similarity in the fatigue behaviour of SCM420 and

SAE8620 steels [Refer section 4.2]. Therefore, it is apparent that for case carburized and

tempered substrates made of SCM420 steel, the deposition of WC/C coating will lead to

considerable enhancement in fatigue performance under infinite life regime.

Observations concerning the influence of WC/C coating on fatigue behaviour of

SAE8620 and SCM420 steels in this study are quite similar to the results reported by Su

and co-workers [10,11] in their studies on fatigue behaviour of normalized AISI 1045 steel

substrates, coated with different types of PVD coatings, viz. Titanium Nitride (TiN),

Titanium Carbonitride (Ti(C, N)), Chromium Nitride (CrN) and Chromium Carbonitride

(Cr(C, N)). The authors observed that when the substrate is subjected to high stresses,

cracks develop within the coating, which act as stress raisers and aid in early failure of the

184

specimen. However, for specimens surviving more than 5×105 cycles under fatigue, the

coating provided effective protection to the surface. Ferreira et al. [193], in their study

involving fatigue behaviour of 42CrMo4 steel coated with various types of deposits, also

reported similar observations. They attributed the deteriorated performance under low cycle

fatigue to the formation of cracks within the coating.

In the present study, the cracks were found to originate at multiple sites in specimens

tested under low cycle fatigue. This has been attributed to cracking of the hard and brittle

coating, when the substrate is subjected to large strains under low-cycle fatigue. Since the

thickness of coating (2µm) is more than ten times the average surface roughness, Ra (<

0.2µm) [Refer section 3.2] of the specimens, the cracks developed within the coating are

likely to act as stress raisers, thereby promoting pre-mature failure of the specimen under

fatigue. On the other hand, substantial regions of slow, stage-I fatigue crack propagation

have been observed in the specimens tested under high cycle fatigue. It is under such

conditions that the large residual compressive stresses introduced by the coating are able to

play a constructive role in enhancing the fatigue performance of the specimens. Besides,

when the applied stresses are low, the coating provides an effective “seal” against crack

initiation at the surface.

en 353

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