M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
565–574
INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE
TOUGHNESS ON THE DYNAMIC WEAR PROPERTIES OF
COATED TOOL STEELS
VPLIV TRDOTE IN LOMNE @ILAVOSTI PODLAGE NA
DINAMI^NE OBRABNE LASTNOSTI OPLA[^ENEGA
ORODNEGA JEKLA
Marko Sedla~ek
1*
, Barbara [etina Bati~
1
, Damir ^esnik
2
, Bojan Podgornik
1
1
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2
Hidria Rotomatika d.o.o., Spodnja Kanomlja 23, 5281 Spodnja Idrija, Slovenia
Prejem rokopisa – received: 2018-11-13; sprejem za objavo – accepted for publication: 2019-03-11
doi:10.17222/mit.2018.242
The aim of this work was to determine the impact of various heat-treatment processes and parameters on the substrate hardness
and fracture toughness combination and subsequently on the wear properties and subsurface deformation of coated tool steel
under dynamic impact loading.
Powder-metallurgy steels used as a substrate material were vacuum heat treated under six different conditions and combined
with deep cryogenic treatment and nitriding with the aim to modify the hardness and fracture toughness. A TiAlN monolayer
with 2-μm thickness was used for all the substrates. In order to evaluate the impact failure of the coating a ball-on-plate impact
fatigue test was designed and used. It was shown that the use of a cryogenic treatment can increase the fracture toughness from
10 % up to 67 % while maintaining the same hardness for some cold-work tool steels, when on others it can have a negative
effect. Regarding the static load-carrying capacity, the most important feature of the substrate is its hardness, which at working
hardness of 63–64 HRc already provides good static load-carrying capacity of the coated substrate. The increased fracture
toughness achieved by the deep cryogenic treatment can at very high and/or very low hardness of the substrate have a negative
impact on the dynamic wear properties of the coated surfaces. On the other hand, in the case of vacuum heat treatment that
ensures an adequate working hardness of 63–64 HRc, the deep cryogenic treatment improves the impact wear resistance of
coated surfaces.
Keywords: heat treatment, fracture toughness, hardness, wear, hard-coating
Cilj raziskave je bil dolo~iti vpliv razli~nih toplotnih obdelav in parametrov na trdoto in lomno `ilavost ter posledi~no na
obrabno odpornost in podpovr{insko deformacijo opla{~enih orodnih jekel pri dinami~nem udarnem obremenjevanju.
Jekla, izdelana na osnovi pra{ne metalurgije, so bila toplotno obdelana s {estimi razli~nimi parametri toplotne obdelave v
kombinaciji s kriogenskim podhlajevanjem in nitriranjem z namenom spreminjati njihovo trdoto in lomno `ilavost. Na jekla je
bila nane{ena trda prevleka TiAlN z debelino 2 μm. Da bi ocenili odpornost prevle~enega jekla na udarno obremenjevanje so
bili na~rtovani in izvedeni udarni testi s pomo~jo kroglice. Rezultati so pokazali, da lahko pri nekaterih jeklih kriogensko
podhlajevanje pove~a lomno `ilavost od 10 % do 67 % ob nespremenjeni trdoti. Najpomembnej{a lastnost podlage je njena
trdota, ki pri trdoti 63–64 HRc `e zagotavlja dobro stati~no nosilnost opla{~ene podlage. Pove~ana lomna `ilavost, dose`ena z
globoko kriogensko obdelavo, lahko pri zelo visoki in/ali zelo nizki trdoti podlage negativno vpliva na dinami~ne obrabne
lastnosti opla{~enih povr{in. Po drugi strani pa, v primeru toplotne obdelave, ki zagotavlja zadostno trdoto (63–64 HRc),
globoka kriogenska obdelava izbolj{a udarno obrabno odpornost opla{~enih povr{in.
Klju~ne besede: toplotna obdelava, lomna `ilavost, trdota, obraba, trda prevleka
1 INTRODUCTION
The metal-processing industry is constantly faced
with demands for greater productivity and lower costs,
which requires better sustainability of tools and the use
of different material-modification techniques. With the
use of new materials (HSS and AHSS), tools are even
more exposed to very demanding contact conditions,
including high impact loads, high contact pressures,
elevated temperatures and high wear. The processing of
harder materials puts a lot of stress on the contact
surfaces, exposing them to a combination of cyclic,
mechanical, chemical and tribological loads, which can
result in fatigue, chipping and wear of the tool.
1–2
In
order to face all those problems a variety of approaches,
from changing tool material and heat treatment, different
geometry, design, manufacturing parameters and work
material, can be applied.
3–4
Although each of those
parameters has its role, the largest impact comes from
the tool material and its microstructure. The most used
material in precision punching tools is nowadays based
on powder-metallurgy (P/M) steels. P/M steels, which
have fine and uniform microstructure, in some cases
could have a higher toughness but lower abrasion resist-
ance than conventionally produced steel.
5
Tools are
normally heat treated to obtain a microstructure of tem-
pered martensite in a uniform distribution of carbides,
which enable sufficient fracture toughness at the working
Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574 565
UDK 67.017:620.17:620.178.16:66.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 53(4)565(2019)
*Corresponding author's e-mail:
marko.sedlacek@imt.si

hardness and acceptable wear resistance.
6
A problem
with lower abrasion resistance is usually solved by the
use of techniques of surface modification, and to a
limited extent also with the application of hard coat-
ings.
7–9
It is illusory to expect that merely hard coating
deposition will solve the abrasion resistance. Since
applied hard coatings are normally very thin and brittle,
the majority of the contact load is carried by the sub-
strate material. In order to achieve the right load-carrying
capacity of the substrate material, it must reflect high
hardness and at the same time, sufficient fracture tough-
ness. This can be achieved by using different heat-treat-
ment processes and parameters, which have an influence
on the microstructure of a tool steel and therefore on its
mechanical and tribological properties.
10
Especially
vacuum heat treatment, deep cryogenic treatment and
pulse plasma nitriding have shown very positive effects
on the performance of high-speed steels.
11–15
If the
substrate is not hard enough to carry the load, elastic and
plastic deformation will occur in the substrate under the
contact, leading to the failure of the coating. Therefore,
on a harder substrate, a higher contact loading can be
applied without the coating failure due to fracture,
spalling or delamination.
16
However, the fracture tough-
ness and resistance to crack initiation and propagation of
the substrate are equally if not more important than the
hardness and wear resistance.
17–20
Combining those pro-
perties is therefore essential to ensure the load-carrying
capacity and wear resistance of the substrate.
The aim of our work was to investigate the influence
of different vacuum heat-treatment parameters, a combi-
nation of deep cryogenic treatment, nitriding and
HRc/K
Ic
ratio on the dynamic wear properties of two
coated cold-work tool steels and one high-speed steel.
2 EXPERIMENTAL PART
2.1 Material, heat treatment and coating
Three different materials produced by powder-metal-
lurgy methods were used for this investigation. Two
materials denoted as A and B were from cold-work tool
steel, and one from high-speed steel, denoted as C. The
chemical composition is shown in Table 1.
Table 1: Chemical composition (in w/%) of used powder-metallurgy
steels
Material C Si Mn Cr Mo V W Co
A 0.85 0.55 0.40 4.35 2.80 2.10 2.55 4.50
B 2.47 0.55 0.40 4.20 3.80 9.00 1.00 2.00
C 1.64 0.60 0.30 4.80 2.00 4.80 10.40 8.00
In order to give three different relationships between
the fracture toughness (K
I
c) and the hardness (HRC), the
vacuum heat-treatment parameters were chosen to
achieve the maximum hardness of the substrate (denoted
as treatment 1), the maximum toughness hardness (de-
noted as treatment 2) and a maximum ratio of fracture
toughness at the working range hardness (denoted as
treatment 3). The heat-treatment parameters were chosen
according to the existing tempering charts.
Additionally, some heat treatments of samples A, B
and C were combined with a deep cryogenic treatment in
such a way that the samples were first heated to the
austenitizing temperature, and hardened with a stream of
gaseous nitrogen at a pressure of 5 bar followed by
immediate controlled sub-cooling of the material in
liquid nitrogen at a temperature of –196 °C for 25 h. On
the other hand, for samples C a combination of vacuum
heat treatment and plasma nitriding was also chosen and
denoted with the letter N. Nitriding in plasma was
carried out for2hat520°Cin5%N
2
–95%H
2
. The
temperatures, time of austenitization and tempering and
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
566 Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574
Table 2: Temperatures, time of austenitization and tempering for material A
Heat treatment
Treatment Austenitization Hardening Undercooling Tempering
A1 1130 °C/6 min N2
, 5 bar – 2×2 h/520 °C + 1×2 h/500 °C
A2 1070 °C/20 min N2, 5 bar – 2×2 h/585 °C + 1×2 h/555 °C
A3 1100 °C/20 min N2
, 5 bar – 2×2 h/500 °C + 1×2 h/470 °C
A1-U 1130 °C/6 min N2
, 5 bar –196 °C/25 h 1×2 h/520 °C
A2-U 1070 °C/20 min N2
, 5 bar –196 °C/25 h 1×2 h/585 °C
A3-U 1100 °C/20 min N2, 5 bar –196 °C/25 h 1×2 h/500 °C
Table 3: Temperatures, time of austenitization and tempering for material B
Heat treatment
Treatment Austenitization Hardening Undercooling Tempering
B1 1180 °C/6 min N2
, 5 bar – 2×2 h/540 °C + 1×2 h/510 °C
B2 1180 °C/6 min N2
, 5 bar – 2×2 h/500 °C + 1×2 h/470 °C
B3 1180 °C/6 min N2
, 5 bar – 2×2 h/560 °C + 1×2 h/530 °C
B1-U 1180 °C/6 min N2, 5 bar –196 °C/25 h 1×2 h/540 °C
B2-U 1180 °C/6 min N2
, 5 bar –196 °C/25 h 1×2 h/500 °C
B3-U 1180 °C/6 min N2
, 5 bar –196 °C/25 h 1×2 h/560 °C

combinations of treatments for material A are presented
in Table 2, for material B in Table 3 and for material C
in Table 4.
As a hard protective coating a TiAlN monolayer with
a hardness of 3300 HV0.05 (measured by nano-indenta-
tion – Fischerscope H100C) was deposited using a PVD
procedure. The coating thickness was approximately
2 μm and was deposited on the substrate at a temperature
of  450 °C on polished heat-treated samples (R
a
= 0.05
– 0.10 μm) with dimensions of (20 × 20 × 8) mm. Details
of the coating deposition process are given in the paper
of M. Panjan et al.
21
2.2 Fracture toughness, hardness, coating adhesion
and impact wear
The fracture toughness was determined with the use
of round, circumferentially notched tensile-test speci-
mens and a fatigue pre-crack in the root of the notch
made in the rotating-bending mode before the heat
treatment.
22–23
The fracture toughness is calculated based
on the measured rapture force and diameter of the in-
stantly fractured part of the specimen. The measurements
of fracture toughness were conducted on at least 12 pa-
rallel K
Ic
-specimens. Rockwell C hardness measurements
using an Instron B2000 machine were made on each K
Ic
test specimen. For each sample up to six measurements
were made circumferentially on the widest part of the K
Ic
test specimen.
Adhesion and static load-carrying capacity of coat-
ings applied on the different vacuum heat-treated sub-
strates were determined using the Rockwell C indenta-
tion test. The test is considered a very quick and easy test
of coatings adhesion, which is defined with the standard
VDI 3198.
24
In order to evaluate the influence of different vacuum
heat treatments and the HRc/K
Ic
ratio of the substrate on
dynamic impact wear resistance of the coating ball-on-
plate impact wear test was designed. For this purpose an
Instron 8802 dynamic testing machine was used. In the
dynamic impact tests carried out at room temperature, a
tungsten carbide ball with a diameter of 32 mm was used
as a counterpart, ensuring that most of the wear occurred
on the coated sample. Due to the configuration of the
machine the ball was fixed during testing, while the
coated plate (20 mm × 20 mm × 8 mm) was impacting
against the stationary ball at a frequency of 30 Hz, as
shown in Figure 1. During the testing the impacting
force was changing in a sinusoidal wave and reached a
peak compressive value of 5.5 kN, which corresponds to
the contact pressure of 3.5 GPa. In the lowest position
the ball and the plate were completely separated with a
gap of  0.5 mm. The duration of the test was limited to
300,000 cycles. For each substrate at least three repeated
tests were performed. To avoid wear and transfer of the
tungsten carbide to the coating surface a new ball and
lithium grease were used for each test. After the test, the
wear of the coating was measured using optical profilo-
metry with an Alicona Infinitefocus G4.
3 RESULTS AND DISCUSSION
3.1 Hardness and fracture toughness
The results of the hardness and fracture toughness for
two different cold-work tool steel (denoted as A and B)
and one from high-speed steel (denoted as C) are pre-
sented in Figures 2 to 4. Figure 2 presents the influence
of different heat treatments on the hardness and the
fracture toughness of the cold-work tool steel with a
lower content of C and a higher content of W. It can be
seen that in the case of the highest austenitizing tempe-
rature (A1; T
A
= 1130 °C,) a hardness of 65.8 HRc and a
fracture toughness K
Ic
= 6.1 MPa·m
1/2
were achieved. By
reducing the austenitization temperature (A2; T
A
=
1070 °C), where the maximum fracture toughness was
expected, a fracture toughness K
Ic
= 12.7 MPa·m
1/2
and a
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574 567
Table 4: Temperatures, time of austenitization and tempering for material C
Heat treatment
Treatment Austenitization Hardening Undercooling Tempering Nitriding
C1 1180 °C/2 min N2
, 5 bar – 2×2 h/530 °C + 1×2 h/500 °C –
C2 1030 °C/20min N
2
, 5 bar – 2×2 h/570 °C + 1×2 h/540 °C –
C3 1100 °C/20 min N2
, 5 bar – 2×2 h/560 °C + 1×2 h/530 °C –
C1-U 1180 °C/2 min N2
, 5 bar –196 °C/25 h 1×2 h/530 –
C2-U 1030 °C/20 min N2
, 5 bar –196 °C/25 h 1×2 h/570 °C –
C3-U 1100 °C/20 min N2, 5 bar –196 °C/25 h 1× 2 h/560 °C –
C1-N 1180 °C/2 min N2, 5 bar – 2× 2 h/530 °C 2 h/520 °C 5 % N
2-95%H
2
C2-N 1030 °C/20 min N2, 5 bar – 2× 2 h/570 °C 2 h/520 °C 5 % N
2-95%H
2
C3-N 1100 °C/20 min N2, 5 bar – 2×2 h/560 °C 2 h/520 °C 5 % N
2-95%H
2
Figure 1: Schematic representation and actual impact dynamic testing

hardness of 59.3 HRc were obtained. In the case of the
heat treatment where the optimal combination of work-
ing hardness and fracture toughness was expected (A3;
T
A
= 1100 °C), a hardness of 63.9 HRc and a fracture
toughness K
Ic
= 10.1–10.2 MPa·m
1/2
were achieved.
Combining a vacuum heat treatment with a deep
cryogenic treatment in liquid nitrogen for 25 h led to an
increase in the fracture toughness at practically the same
hardness. In the case of the first series of samples
hardened at the highest austenitizing temperature (A1U),
at practically the same hardness of 65.0 HRc, a 25-hour
deep cryogenic treatment in liquid nitrogen increased the
fracture toughness from 6.1 MPa to 10.4 MPa. For the
second series of samples with the lowest austenitizing
temperature (A2U), a 25-hour deep cryogenic treatment
in liquid nitrogen led to an approximately 12 % increase
in fracture toughness, from 12.7 MPa·m
1/2
to 14.2
MPa·m
1/2
, at an even higher hardness of 59.5 HRc. For
the optimal combination of fracture toughness at the
working hardness (treatment A3U), the deep cryogenic
treatment resulted in a 16 % improvement in the fracture
toughness, increasing it from 11.0 MPa·m
1/2
to 12.8
MPa·m
1/2
and in a slight increase in hardness. The effect
of the deep cryogenic treatment on the change in
hardness and fracture toughness, including scatter, is
graphically displayed in a combined HRc/K
Ic
diagram in
Figure 2.
Despite the intensive research on cold-work and
high-speed steels that has shown that a deep cryogenic
treatment can lead to an improved material performance,
including fracture toughness and wear resistance, the
cause is still not fully understood, since the authors
reported contradictory results. P. Baldissera et al. suggest
that the improvement obtained by deep cryogenic treat-
ment is mainly contributed to the complete elimination
of retained austenite and the formation of very small
carbides dispersed in the tempered martensitic struc-
ture.
25
On the other hand, G. Gavriljuk et al. reported that
the full austenite-to-martensite transformation does not
occur in high-carbon steels, with the martensitic transfor-
mation at low temperature being accompanied by plastic
deformation of virgin martensite.
26
According to the
B. Podgornik et al. the improved fracture toughness is a
result of the plastic deformation of the primary marten-
site combined with a reduced amount of dissolved car-
bon in the martensite and finer and more homogeneous
carbides precipitation.
27
The same as for the cold-work tool steel A, was also
for the cold work tool steel B with a higher content of C
and V, heat treatment determined so it resulted in the
highest hardness (B1), highest fracture toughness (B2)
and optimal combination of fracture toughness at the
working hardness (treatment B3). With the highest aus-
tenitizing temperature (B1) a hardness of 66.0 HRc and a
fracture toughness of 11.2 MPa·m
1/2
were reached. With
the same austenitizing temperature but lower tempering
temperature (B2) a maximum fracture toughness of 15.0
MPa·m
1/2
and a hardness of 64.6 HRc were achieved. A
hardness of 64.6 HRc and a toughness of 6.10 MPa·m
1/2
were achieved with treatment B3.
For material B, combining vacuum heat treatment
with a deep cryogenic treatment in liquid nitrogen for
25 h did not lead to an increase in the fracture toughness.
For treatment B1U the hardness decreased by 2.5 % on
64.3 HRc and the fracture toughness up to 10 %, to
10 MPa·m
1/2
. In the case of treatment B2U, where the
initial fracture toughness was the highest, a 40 % de-
crease was recorded (8.9 MPa·m
1/2
) when combining a
heat treatment with a cryogenic treatment. Treatment B3U,
on the other hand, resulted in a slight increase in the
fracture toughness from 10.6 MPa·m
1/2
to 11.5 MPa·m
1/2
,
but it decreases the hardness to only 63 HRc. As reported
in the paper of B. Podgornik et al., a high volume fraction
of undissolved eutectic carbides (18–20 %) was observed
and related to the neglected effect of plastic deformation
of the primary martensite during the deep cryogenic
treatment.
27
The undissolved eutectic carbides are mainly
of more stable MC type, probably leading to a change in
the precipitation type and kinetics, which might lead to
impairment of the properties when including a deep
cryogenic treatment.
27
The results are shown in Figure 3.
For high-speed steel it can be seen (Figure 4) that in
the case of the highest austenitizing temperature (C1;
T
A
= 1180 °C,) a hardness of 68.2 HRc and a fracture
toughness K
Ic
= 7.6 MPa·m
1/2
were achieved. By reduc-
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
568 Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574
Figure 2: Effect of deep cryogenic treatment on the change in hard-
ness and fracture toughness of cold-work tool steel A
Figure 3: Effect of deep cryogenic treatment on the change in hard-
ness and fracture toughness of cold-work tool steel B

ing the austenitization temperature (C2; T
A
= 1030 °C),
where the maximum fracture toughness was expected, a
fracture toughness K
Ic
= 10.0 MPa·m
1/2
and a hardness of
60.8 HRc were obtained. In the case of the heat
treatment where the optimal combination of working
hardness and fracture toughness was expected (C3; T
A
=
1100 °C), a hardness of 64.1 HRc and a fracture tough-
ness K
Ic
= 8.7 MPa·m
1/2
were achieved.
Combining a vacuum heat treatment with a deep
cryogenic treatment in liquid nitrogen for 25 h did not
lead to a significant increase in the fracture toughness or
hardness. In the case of the first series of samples
hardened at the highest austenitizing temperature (C1U),
at slightly lower hardness 67.7HRc, a 25-hour deep cryo-
genic treatment in liquid nitrogen increased the fracture
toughness from 7.6 MPa·m
1/2
to 8.5 MPa·m
1/2
, which
represents a 10 % improvement in the fracture
toughness. An even lower change was recorded for the
second series of samples with the lowest austenitizing
temperature (C2U), where a 25-hour deep cryogenic
treatment in liquid nitrogen led to an approximately 4 %
increase in the fracture toughness, from 10.0 MPa·m
1/2
to
10.4 MPa·m
1/2
, at a slightly higher hardness of 61.4 HRc.
For the optimal combination of fracture toughness at the
working hardness (treatment C3U), a deep cryogenic
treatment resulted in a less than 5 % improvement in the
fracture toughness, increasing it from 8.7 MPa·m
1/2
to
9.1 MPa·m
1/2
and in a slight increase in the hardness
(from 64.1 HRc to 64.7 HRc). The effect of the deep
cryogenic treatment on the change in the hardness and
the fracture toughness is graphically displayed in a com-
bined HRc/K
Ic
diagram in Figure 4. As explained in the
paper of Duh et al. the small effect of the deep-cryogenic
treatment on the hardness and fracture toughness of steel
C can be attributed to the very small volume fraction of
retained austenite in the steel after quenching, but mainly
due to the high volume fraction of undissolved eutectic
carbides, consequently reducing the amount of marten-
site matrix and diminishing the effect of the martensite
plastic deformation and capturing of the immobile
carbon atoms by gliding dislocations during the deep
cryogenic treatment.
28
When combining all three different vacuum heat
treatments with plasma nitriding we observed a surface
hardness increase. The nitriding diffusion zone has a
thickness  65 μm. With the samples with the highest
austenitizing temperature (C1N) additional nitriding
resulted in an increase of the hardness to 1450 HV
0.05
,
wherein the hardness of the base material decreased to
920 HV
0.05
, 67.5 HRc. For the samples with the lowest
austenitizing temperature (C2N) the hardness increased
to 1280 HV
0.05
, while the hardness of the base stayed the
same (722 HV
0.05
,  61 HRc). The same was true for the
optimal combination of fracture toughness at the work-
ing hardness (treatment C3N), where for the same
hardness of the base (804 HV
0.05
, 64 HRc) the hardness
of the surface increased to over 1350 HV
0.05
. The in-
fluence on the nitriding of the fracture toughness could
not be determined, because of the inadequacy of the
method with a cylindrical specimen having a notch.
From the literature it is known that nitriding deteriorates
the ductility of the surface layer.
28
3.2 Coating adhesion
The comparison of the Rockwell C adhesion test
results for a monolayer TiAlN coating deposited on
differently vacuum heat-treated substrates is shown in
Figures 5 and 6. In general, all the substrates, regardless
of the type of vacuum heat treatment used, either com-
bined with a cryogenic treatment or nitriding exhibit
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574 569
Figure 4: Effect of deep cryogenic treatment on the change in
hardness and fracture toughness of high-speed steel C
Figure 5: Influence of heat treatment (1, 2, 3), cryogenic treatment
(U), nitriding (N) and used material (A, B, C) on: a) coating adhesion,
b) hardness (*) – for nitrided samples values are for information only

good adhesion (HF1 to HF3). None of the samples
reached unacceptable damage level, but there are visible
differences between the different vacuum heat treat-
ments, deep cryogenic treatment and nitriding combina-
tions used.
In general, it can be seen that hardness of the sub-
strate material is the main impact factor. When hardness
is higher than or equal to 64 HRc (Figure 5b), the
adhesion is good with only a network of radial cracks
observed around the indentation (Figure 6a). When, on
the other hand, the hardness is too low, for cold-work
tool steel A and high-speed steel C, also circular cracks
and a small coating delamination can be recorded,
indicating a too low load-carrying capacity of the steel
substrate to be able to support the deposited coating
(Figure 6b).
Combining a vacuum heat treatment with a deep
cryogenic treatment in liquid nitrogen has no influence
on the coating adhesion for cold-work tool steel A. On
the other hand, for cold-work tool steel B, with the use of
a deep cryogenic treatment, the adhesion slightly dete-
riorates when the hardness of the substrate is the highest
(B1U) and the lowest (B2U). This can be explained with
a decrease of the hardness with the use of a deep cryo-
genic treatment.
The best adhesion of the coating was recorded on
cold-work tool steel B and high-speed steel C with heat
treatment 1 and for the high-speed steel C when the
vacuum heat treatment was combined with the nitriding
(Figure 6c), indicating that the substrate hardness has
the most important role on the adhesion as it provides
sufficient load-carrying capacity of the steel substrate to
support the top coating. This correlates well with the
findings of C. Zhang et al. that the adhesion between the
coating and the substrate increases linearly with the
increase in the substrate hardness.
29
In Figure 5b the hardness values for the nitrided
samples are given in Rockwell. Those values were con-
verted from the Vickers measurement and are for infor-
mation only, because the values are out of the conversion
range.
3.3 Impact wear
In Figure 7a the impact wear properties of the inves-
tigated materials are presented in the form of coating
wear depth as a function of the different substrate heat
treatment. To facilitate the presentation hardness (Figure
7a) and fracture toughness (Figure 7c) as a function of
different heat treatments are presented. It can be seen
that for all the used materials the hardness of the sub-
strate plays the most important role. It can be seen that
the worst wear was recorded when the heat treatment
resulted in the lowest hardness (treatment 2), regardless
of the used material. The lowest wear was recorded when
the heat treatment gave the highest hardness (treatment
1), with the only exception being high-speed steel C,
where the lowest impact wear was recorded with the
vacuum heat treatment that provided the highest fracture
toughness at the working hardness (treatment C3; 64.1
HRc and K
Ic
= 8.7 MPa·m
1/2
). Treatment C2, which has
the lowest hardness overall, resulted in the highest wear
rate among all the investigated samples.
The deep cryogenic treatment, although providing
increased fracture toughness, had a negative impact on
the coating impact wear resistance when combined with
the vacuum heat treatment that provided the maximum
hardness (treatment 1) or the maximum fracture tough-
ness (treatment 2), for all the investigated materials. On
the other hand, for the cold-work tool steel A the vacuum
heat treatment A3, which at working hardness of 63-64
HRc ensures maximum fracture toughness, the deep
cryogenic treatment increases both the hardness and the
fracture toughness and thus has a positive impact on the
dynamic impact wear resistance of the coatings. By
simultaneously increasing the hardness and the fracture
toughness (64.2 HRc and K
Ic
= 12.4 MPa·m
1/2
), the wear
of the coating decreased by  30 % and became com-
parable to the case of the vacuum heat treatment A1.
For high-speed steel C, combining the vacuum heat
treatment with the nitriding has a positive effect on the
impact wear properties. Introducing nitriding drastically
decreased the impact wear compared to the heat treated
with or without the deep cryogenic treatment (C and
C_U), which can be attributed to the influence of the
increased hardness. When comparing the influence of the
hardness on the impact wear of the nitride samples, the
influence of hardness is still visible. The vacuum heat
treatment that resulted in the highest hardness (C1N),
resulted in the lowest impact wear, and the vacuum heat
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
570 Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574
Figure 6: Rockwell C indentation for: a) A1 b) C2 and c) C1N

treatment that resulted in lowest hardness (C2N) resulted
in the highest impact wear.
A material’s hardness has long been erroneously used
as a criterion for wear resistance. Two materials having
different microstructures can have identical hardness
values can have significantly different wear rates. This is
clearly seen if we compare samples A1 and B1, which
have similar hardness ( 66 HRc) but have significantly
different wear rates. As it was reported by Moore et al.,
the material microstructure has a greater influence on the
wear resistance than the bulk hardness.
30
Assuming that
the harder the material, the greater the wear resistance is
technically correct, but cannot be used generally. With
some degree of caution, hardness may still be used for
the same chemical compositions of the steel. The results
clearly show that instead of only hardness, the chemical
compositions and the resulting metallurgical structures
should also be taken into account.
With the use of SEM and EDS, all the wear craters
were analysed, which confirmed that for all cases the
main wear mechanism was abrasion. In Figure 8a it can
be seen that the wear is gradual and without any evident
coating cracking or delamination observed inside the
impacted wear scar. If the substrate was too soft to be
able to support the coating, the impact wear could ex-
ceed the coating thickens (Figure 8b). When this
happened, the hardness was again playing a critical role
since the substrates with lower hardness resulted in
higher wear (B2 > B1).
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574 571
Figure 7: a) Impact wear in dependence of different substrate heat treatment and used steel, b) hardness depending on the heat treatment, c)
fracture toughness depending on the heat treatment
Figure 8: SEM micrograph of the wear crater for: a) sample A1, b) C2

To better understand the subsurface deformation,
EBSD studies were performed on selected specimens.
Figure 9 shows the results of EBSD mappings of the
subsurface region of specimen B3. The map size is
123.8 μm × 100.6 μm, acquired at step size of 0.2 μm.
The image presents a band contrast map, a measure of
the contrast of the Kikuchi pattern. The carbides have
sharp diffractions and seem largely unaffected by the
impact, while the martensitic matrix has much poorer
diffractions, hinting at large deformations of the crys-
talline lattice. The coating shows no diffraction, meaning
that it is most likely amorphous.
The diffraction patterns from the carbides and the
martensite matrix taken at different locations below the
surface are shown in Figure 10. As J. Perret et al. have
shown, the diffractions directly below the impact are too
poor to be recognized, due to distorted metal crystalline
lattice.
31
To better understand the situation near the impact,
EBSD maps of higher magnification were recorded.
Figure 11 shows two such maps: the map in Figures 11
a and 11c represents a region directly below the surface,
and the map in Figure 11b and 11d has been recorded at
approximately 2 mm above the impact point. The upper
images are band contrast maps, and the lower images are
phase maps, with yellow representing the carbide phase
and green representing the martensitic phase. No post-
processing noise filtering has been applied. The image
sizes are 24.75 μm × 20.1 μm.
The images recorded just below the impact show
signs of wear: the pattern quality is poor for the marten-
sitic matrix, and because of that the indexing is also quite
bad. We can observe little detail of the microstructure of
the matrix. Even the carbides seem affected in the zone
approximately up to 5 μm from the impact. On the other
hand, the map recorded about 2 mm above the impact
shows no signs of damage, and we can clearly see the
martensitic microstructure in the pattern quality map.
All this leads us to believe that the damage is con-
centrated in the sub-surface zone, the affected zone is
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
572 Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574
Figure 11: EBSD maps: a,c) region directly below the surface; b,d)
recorded at approximately 2 mm above the impact point
Figure 9: Band contrast EBSD map. The bright globular features are
carbides, and the martensitic matrix exhibits poorer diffraction.
Figure 10: Kikuchi patterns taken at different distances from the
surface. Left column: martensite diffractions, right column: carbide
diffractions

approximately 5 μm from the surface, and the damage is
mostly concentrated in the martensitic matrix.
4 CONCLUSIONS
With the combination of a vacuum heat treatment and
a deep cryogenic treatment in liquid nitrogen an im-
provement in the fracture toughness from 10 % to 67 %
can be achieved, while maintaining a high hardness for
cold-work tool steel A. For high-speed steel C the im-
provement in the fracture toughness is less pronounced
(4 % to 10 %). When, on the other hand, for cold-work
tool steel B a deep cryogenic treatment has a negative
influence on the fracture toughness, decreasing it by up
to 10 %.
Regarding the static load-carrying capacity, the most
important property of the substrate is its hardness, with
the working hardness of 63–64 HRc already providing
excellent static load-carrying capacity of the coated
substrate. On the other hand, the increase in fracture
toughness achieved by the deep cryogenic treatment
improves the static load-carrying capacity only for the
case where the hardness of the substrates is too low,
below 63 HRc.
In terms of dynamic impact wear resistance of the
coated surface, the hardness of the substrate still has the
greatest impact. The higher the hardness of the sub-
strates, the higher wear resistance of the coated surface
can be expected for the material with lower content of
carbon. For cold-work tool steels A and B an adequate
impact resistance is achieved when hardness of 65 HRc
and fracture toughness of 10 MPa·m
1/2
is reached. For
high-speed steel C the best impact wear resistance is
achieved when the hardness is above 65 HRc and
additional nitriding in plasma was used. The deep
cryogenic treatment and associated increase in fracture
toughness has a negative effect for cases when the
substrate has a very high hardness and a low toughness
or a high fracture toughness but too low hardness. On the
other hand, in the case of vacuum heat treatment that
ensures an adequate working hardness of 63–64 HRC,
the deep cryogenic treatment leads to an improved im-
pact wear resistance of the coated surface. The wear
resistance then becomes equal to the case where the
substrate with a high hardness is used, but at the same
time providing more than 50 % higher fracture toughness
required for proper fatigue resistance of the tool. The
results also clearly show that instead of only hardness,
the chemical compositions and the resulting metallur-
gical structures should also be taken into account.
Acknowledgment
Dr. M. ^ekada from the Jo`ef Stefan Institute, Ljub-
ljana, Slovenia, is greatly acknowledged for the help and
the deposition of the TiAlN coating.
5 REFERENCES
1
J. A. Schey, Tribology in Metalworking: Friction, Lubrication and
Wear, Oxford Publishing, Oxford 1984
2
A. Gåård, Wear in sheet metal forming; PhD thesis, Karlstad
University Studies, Karlstad 2008
3
R. Ebara, K. Takeda, Y. Ishibashi, A. Ogura, Y. Kondo, S. Hamaya,
Microfractography in failure analysis of cold forging dies, Engineer-
ing Failure Analysis, 16 (2009) 6, 1968–1976, doi:10.1016/
j.engfailanal.2008.10.023
4
S. Subramonian, T. Altan, B. Ciocirlan, C. Campbell, Optimum
selection of variable punch-die clearance to improve tool life in
blanking non-symmetric shapes, International Journal of Machine
Tools and Manufacture, 75 (2013) 63–71, doi:10.1016/j.ijmachtools.
2013.09.004
5
Y. Torres, S. Rodriguez, A. Mateo, M. Anglada, L. Llanes, Fatigue
behavior of powder metallurgy high-speed steels: fatigue limit
prediction using a crack growth threshold-based approach, Materials
Science and Engineering A, 387 (2004) 501–504, doi:10.1016/
j.msea.2003.12.072
6
V. Leskov{ek, B. [u{tar{i~, G. Jutri{a, The influence of austenitizing
and tempering temperature on the hardness and fracture toughness of
hot-worked H11 tool steel, Journal of Materials Processing
Technology, 178 (2006) 1–3, 328–334, doi:10.1016/j.jmatprotec.
2006.04.016.
7
V. Leskov{ek, B. Podgornik, M. Jenko, A PACVD duplex coating for
hot-forging applications, Wear, 266 (2009) 3–4, 453–460
doi:10.1016/j.wear.2008.04.016
8
L. Lu, Q. Wang, B. Chen, Y. Ao, D. Yu, C. Wang, Microstructure and
cutting performance of CrTiAlN coating for high-speed dry milling,
Transactions of Nonferrous Metals Society of China, 24 (2014)6 ,
1800–1806, doi:10.1016/S1003-6326(14)63256-8
9
F. Sergejev, P. Peetsalu, A. Sivitski, M. Saarna, E. Adoberg, Surface
fatigue and wear of PVD coated punches during fine blanking
operation, Engineering Failure Analysis, 18 (2011) 71689–1697,
doi:10.1016/j.engfailanal.2011.02.011
10
B. Podgornik, S. Hogmark, O. Sandberg, Proper coating selection for
improved galling performance of forming tool steel, Wear, 261
(2006) 15–21, doi:10.1016/j.wear.2005.09.005
11
V. Leskov{ek, B. Podgornik, Vacuum heat treatment, deep cryogenic
treatment and simultaneous pulse plasma nitriding and tempering of
P/M S390MC steel, Materials Science and Engineering: A, 531
(2012) 119–129, doi:10.1016/j.msea.2011.10.044
12
N. Popandopulo, L. T. Zhukova, Transformation in high speed steels
during cold treatment, Met. Sci. Heat Treat., 22 (1980) 10, 708–710,
doi:10.1007/BF00700561
13
B. Podgornik, F. Majdi~, V. Leskov{ek, J. Vi`intin, Improving
tribological properties of tool steels through combination of
deep-cryogenic treatment and plasma nitriding, Wear, 288 (2012)
88–93, doi:10.1016/j.wear.2011.04.001
14
B. Podgornik, V. Leskov{ek, Microstructure and origin of hot-work
tool steel fracture toughness deviation, Metall Mater Trans A, 44
(2013) 694–702, doi:10.1007/s11661-013-1921-6
15
G. Prieto, W. R. Tuckart, J. E. Perez Ipiña, Influence of a cryogenic
treatment on the fracture toughness of an AISI 420 martensitic
stainless steel, MTAEC9, 51 (2017), 591, doi:10.17222/mit.2016.126
16
P. Hedenqvist, M. Olsson, S. Jacobson, S. Söderberg, Failure mode
analysis of TiN-coated high speed steel, In situ scratch adhesion
testing in the scanning electron microscope, Surface and Coatings
Technology, 41 (1990) 31–49, doi:10.1016/0257-8972(90)90128-Y
17
M. Bayramoglu, H. Polat, N. Geren, Cost and performance evalu-
ation of different surface treated dies for hot forging process, J.
Mater. Process. Technol., 205 (2008) 394–403, doi:10.1016/
j.jmatprotec.2007.11.256
18
I. Campos, M. Farah, N. López, G. Bermúdez, G. Rodríguez, C. Villa
Velázquez, Evaluation of the tool life and fracture toughness of
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574 573

cutting tools boronized by the paste boriding process, Appl. Surf.
Sci., 254 (2008) 2967–2974, doi:10.1016/j.apsusc.2007.10.038
19
T. Atkins, Toughness and processes of material removal, Wear, 267
(2009) 1764–1771, doi:10.1016/j.wear.2009.04.010
20
B. Podgornik, V. Leskov{ek, B. Arh, The effect of heat treatment on
the mechanical, tribological and load-carrying properties of
PACVD-coated tool steel, Surface and Coatings Technology, 232
(2013) 528–534, doi:10.1016/j.surfcoat.2013.06.019
21
M. Panjan, M. ^ekada, P. Panjan, F. Zupani~, W. Kölker, Depen-
dence of microstructure and hardness of TiAlN/VN hard coatings on
the typ of substrate rotation, Vacuum, 86 (2012) 6, 699–702,
doi:10.1016/j.vacuum.2011.07.052
22
B. Ule, V. Leskov{ek, B. Tuma, Estimation of plain strain fracture
toughness of AISI M2 steel from precracked round-bar specimens,
Engineering Fracture Mechanics, 65 (2000) 559–572, doi:10.1016/
S0013-7944(99)00105-8
23
V. Leskov{ek, B. Ule, B. Li{~i}, Relations between fracture
toughness, hardness and microstructure of vacuum heat-treated
high-speed steel, Journal of Materials Processing Technology, 127
(2002) 298–308, doi:10.1016/S0924-0136(02)00280-7
24
Verein Deutscher Ingenieure Normen, VDI 3198, VDI-Verlag,
Dusseldorf, 1991
25
P. Baldissera, D. Delprete, Deep cryogenic treatment: A biblio-
graphic review, Open Mech. Eng. J., 2 (2008) 1–11, doi:10.2174/
1874155X00802010001
26
G. Gavriljuk, W. Theisen, V. V. Sirosh, E. V. Polshin, A. Kortmann,
G. S. Mogilny, Yu. N. Petrov, Ye. V. Tarusin, Low-temperature mar-
tensitic transformation in tool steels in relation to their deep cryo-
genic treatment, Acta Mater., 61 (2013) 1705–1715, doi:10.1016/
j.actamat.2012.11.045
27
B. Podgornik, I. Paulin, B. Zajec, S. Jacobson, V. Leskov{ek, Deep
cryogenic treatment of tool steels, Journal of Materials Processing
Technology, 229 (2016) 398–406, doi:10.1016/j.jmatprotec.2015.
09.045
28
D. Duh, I. Schruff, Optimized heat treatment and nitriding para-
meters for a new hot-work tool steel, Proccedings of 6
th
International
tooling conference, Karlstad, 2002, 479-495
29
C. Zhang, T. Hu, N. Zhang, Influence of substrate hardness on
coating-substrate adhesion, Adv. Mater. Res., 177 (2010) 148–150,
doi:10.4028/www.scientific.net/AMR.177.148
30
M. A. Moore, The relationship between the abrasive wear resistance,
hardness and microstructure of ferritic materials, Wear, 1 (1974) 28,
59–68, doi:10.1016/0043-1648(74)90101-X
31
J. Perret, E. Boehm-Courjault, M. Cantoni, S. Mischler, A. Beau-
douin, W. Chitty, J.-P. Vernot, EBSD, SEM and FIB characterisation
of subsurface deformation during tribocorrosion of stainless steel in
sulphuric acid, Wear, 269 (2010) 383-393, doi:10.1016/j.wear.
2010.04.023
M. SEDLA^EK et al.: INFLUENCE OF THE SUBSTRATE HARDNESS AND FRACTURE TOUGHNESS ...
574 Materiali in tehnologije / Materials and technology 53 (2019) 4, 565–574