H-O-T - Proven Heat Treatment and Surface Engineering Technologies

HEAT TREATMENT
PROVEN TECHNOLOGIES AND INDUSTRIAL DEVELOPMENTS IN HEAT TREATMENT
AND SURFACE ENGINEERING OF TOOLS AND DIES
J. Bach, F. Dambacher . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1
PROCESS OPTIMISATION FOR DEEP COLD TREATMENT OF TOOL STEELS
P. F. Stratton . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .11
WEAR BEHAVIOUR OF DEEP CRYOGENIC TREATED HIGH SPEED STEELS
M. Kalin, V. Leskovšek, J. Vižintin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .21
IMPROVING TOOL PRODUCT PERFORMANCE THOUGH THE USE
OF INTENSIVE QUENCHING PROCESSES
M.A. Aronov, J.A. Powell, N.I. Kobasko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .27
VACUUM FURNACE - INTEGRATED "SUB-ZERO" TREATMENT
B. Zieger, R. Stein . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .33
THE MAIN PRINCIPLES OF INTENSIVE QUENCHING OF TOOLS AND DIES
N.I. Kobasko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .39
OPTIMIZING THE VACUUM-HEAT-TREATMENT OF HOT-WORK TOOL STEELS BY
LINEAR ELASTIC FRACTURE MECHANICS
V. Leskovšek, B. Šuštaršiè, G. Jutriša, D. Baksa, J. Kopaè . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .45
MICROSTRUCTURAL EVOLUTION OF A RAPIDLY SOLIDIFIED HIGH-SPEED
STEEL DURING HIGH TEMPERATURE ANNEALING
C. Stotter, H. Leitner, R. Ebner, P. Ramminger, E. Brandstätter . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .53
HIGH TEMPERATURE BRAZING OF HIGH TEC COMPONENTS - A THERMAL
PROCESSING WITH SPECIFIC REQUIREMENTS
I. Reinkensmeier, H. Buursen . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .63
IMPROVED PRODUCTION OF STEEL PARTS BY INTENSIVE QUENCHING
N.I. Kobasko, L.C.F. Canale, G.E. Totten . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .69
THE USE OF A NEW TYPE OF LARGE VACUUM BATCH FURNACE FOR THE
HEAT TREATMENT OF MOULDS AND DIES
Th. Wingens, B. Edenhofer, O. Irretier . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .77
SINGLE AND MULTI-CHAMBER VACUUM FURNACES FOR TOOLS AND DIES HEAT TREATMENT
D. Siniarski, M. Korecki . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .83
INFLUENCE OF THE PRETREATMENT THERMAL TREATMENT ON THE MICROSTRUCTURE
AND PROPERTIES OF X37CRMOV5-1 STEEL
J. Jasiñski, R. Torbus, E. Kasprzycka, B. Bogdañski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .91
INFLUENCE OF QUENCHING OILS COMPOSITION ON THE COOLING RATE
L. Pedišiæ, B. Matijeviæ, B. Periæ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .97
THE INFLUENCE OF MULTIPLE SOLUTION ANNEALING ON KINETICS
OF STRUCTURAL TRANSFORMATION OF MARAGING STEELS
I. Kladariæ, D. Krumes, R. Markoviæ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .107
CONTAINS

PROVEN TECHNOLOGIES AND INDUSTRIAL DEVELOPMENTS
IN HEAT TREATMENT AND SURFACE ENGINEERING
OF TOOLS AND DIES
J. Bach, F. Dambacher
1. INTRODUCTION
Increasing requirements in terms of the quality of the products as well as the
processing of new materials require even more extreme processing conditions. Tools
and materials often do not any longer fulfil the rising requirements satisfactorily. The
consequence is in particular the increase of wear of the tool surfaces. Beside the
optimal selection of the materials as well as their heat treatment technologies of the
surface treatment gain strongly significance. The presented contribution concerns
itself with the demand-fair heat treatment of selected steels, the barrier layer
treatment by plasma nitriding, as well as a special surface treatment, the PVD-
coating.
2. BASICS
2.1 Heat treatment technology
An iron carbon alloy with a carbon content < 2.06 % is called steel. Beside the main
alloying element carbon the characteristics of the steel can be affected by different
alloying elements considerably. Heat treatment of steels is based on a lattice
transformation at a temperature of 911°C (for pure iron). At this temperature the
cubically body-centered α-iron (ferrite) converts into the cubically face-centered γ-iron
(austenite). Apart from the different type of lattice these two types differ in their
carbon solubility. In the ferrite the solubility limit is about 0.02 weight- %, higher
contents lead to the formation of cementite (iron carbide Fe3C).
H-O-T
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Kleinreuther Weg 118
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ÄRTE- UND BERFLÄCHEN ECHNIKH O T
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IFHTSE 2005

In the austenite lattice however up to 2 weight-% C can be solved. With the
exceeding of the transformation temperature the cementite dissolves and carbon is
distributed evenly in the austenite. During rapid cooling solved carbon can’t seperate
itself as iron carbide, but remains in a not stable obligation solution. With falling
below an intended temperature martensite formation begins with turning the
austenite lattice into a tetragonally distorted lattice, in which obligation-solved carbon
leads to high spanning and thus high hardness.
For reaching the martensite transformation the following points must be fulfilled:
warming up the steel to temperatures higher than transformation
temperature
cooling velocity higher than critical cooling velocity
falling below a critical temperature (martensite starting-temperature)
These terms essentially depend thereby on the composition of the steel. Usually the
critical cooling speed becomes smaller with rising alloy content. For tools with large
cross sections this even makes the martensite-transformation possible, also in the
core. However usually also a degradation of the critical temperature takes place with
rising alloy content. This leads usually to the fact that the martensite formation is not
completely finished at ambient temperature, and in the structure a not negligible
share of not converted austenite (retained austenite) is present. This is unwanted, on
the one hand because of the small hardness. On the other hand transformation is
possible during employment (change of measure). The share of retained austenite
can be reduced to a negligible share by deep-freezing or repeated tempering.
nadeliger Martensit
Restaustenit
nadeliger Martensit
Restaustenit
nadeliger Martensit
Restaustenit
Figure 1: Retained austenite and acicular martensite, caused by incorrect heat treatment
At best after quenching the structure consists of pure, very brittle martensite. By a
tempering following the hardening is it possible to adjust a demanded relationship
between hardness and toughness. During the tempering process the martensite
converts into ferrite under formation of finest carbide eliminations. This is connected
with a lost in hardness and an increase in toughness. With highly alloyed materials
an increase in hardness (secondary hardness) after tempering is possible. This is
caused by the formation of special carbides, which are reached by special alloying
elements.
acicular martensite
retained austenite
2

The most modern heat treatment process is the vacuum process, which is
accomplished in special vacuum furnaces with pressures of ~0,1mbar. As quenching
media inert gas (e.g. N2, He) is used, flowing with overpressure into the vacuum
chamber. The large advantage of the vacuum heat treatment is an oxidation-free
surface after the treatment. Due to the small quenching effect this procedure is
suitable only for high-alloyed steel.
Hardness[HV30]
Tempering temperature [°C]
Hardness[HV30]
Tempering temperature [°C]
Figure 2: Temper curves of unalloyed and alloyed steels. With highly alloyed materials an
increase in hardness is possible.
2.2 Plasma nitriding technology
Nitriding belongs to the procedures, with which the hardening of the tool surface is
possible. Contrary to other procedures, like for example case-hardening, laser-
hardening or the induction-hardening, an increase in hardening is reached without
conversion of the steel structure. That fact makes treatment of finaloutlined tools
possible without the necessity for remachining. Nitriding designates a
thermochemical diffusion process, in which nitrogen intrudes into the surface of the
workpiece. This leads to an increase of hardness in the outer zone up to several 100
m. Depending on material and nitriding process the formation of a compound layer
consisting of nitrides (up to a thickness of 20 m) at the surface is possible. After the
compound layer a diffusion layer follows with a characteristic hardness profile
depending on the nitrogen concentration behaviour.
0
100
200
300
400
500
600
700
0 50 100 150 200 250 300 350 400 450 500
Nhd [ m]
hardness[HV0,2]
Nhd
0
100
200
300
400
500
600
700
0 50 100 150 200 250 300 350 400 450 500
Nhd [ m]
hardness[HV0,2]
Nhd
Figure 3: Hardness profile after a nitriding process
3

In figure 4 the nitriding process is schematically represented:
cathode -
+
+ ION
Fe N
Fe2N
Fe4N
glowing seam
workpiece
furnace wall
FeN
anode +
N
cathode -
+
+ ION
Fe N
Fe2N
Fe4N
glowing seam
workpiece
furnace wall
FeN
anode +
N
Figure 4: Schematic structure of a plasma nitriding plant
During the plasma nitriding process the nitrogen gas is ionized in a gas-discharge in
a vacuum chamber and the gas ions are accelerated toward the tool switched as
cathode. At the tool-surface the ions penetrate into the workpiece outer zone. The
large advantage of the plasma nitriding is, that the results are well adjustable by
variation of the processing parameters. A demand-fairly nitrated outer zone supplies
an increase of the wear-resisting quality (due to the hard outer zone), an
improvement of the endurance limit (by induced compression stresses) as well as an
increase of the corrosion resistance (due to the session layer; does not apply to
stainless steel).
2.3 PVD-coating technology
PVD (physical vapour deposition) designates all procedures of the physical
separation of thin layers over the vapor phase. The raw material for the layers in the
vacuum is transferred into the gaseous phase. The gas particles condense at the
surface which can be coated. By controlled addition of nitrogen different types of
coatings can be formed.
Table 1: Comparison of different types of PVD- coatings
coating TiN TiCN TiAlN CrN WCC®
DLC
colour golden blue-grey
black-
violett
silver-
grey
black-
grey
black-
grey
approx.
thickness
m
1-5 1-5 1-5 1-10 1-3 1-3
micro-
hardness
HV0,5
2400 3000 3300 1900 1500 2800
4

The advantages of PVD coatings are in detail:
very high surface hardness, with receipt of the toughness of the base
material
small friction losses and good sliding qualities
small layer thicknesses→ Accuracy to size remains
A common process is the so-called ARC- process. In a high-vacuum-chamber an arc
between evaporator (cathode) and the surrounding chamber wall (anode) is ignited.
Due to the high current density in the toe of the arc it comes to an emission like an
explosion of atoms, ions and clusters from the evaporation material (e.g. Ti). Up to
90% of the emitted particles thereby are ionized. The tool which will be coated is put
during the coating on negative potential. Due to ion bombardment it comes to a
compression of the surface and to an improvement of the layer adhesion. The
production of layers with good layer adhesion requires process temperatures of ~
450 °C. This presupposes good tempering properties of for instance 500°C, in order
to avoid inadvertent material transformations.
2.4 Characteristic comparison of the different treatment procedures
The procedures presented in chapter 2 differ in their mechanical characteristics, like
the maximally attainable hardness or the maximally attainable layer thickness. With
the classical heat treatment it is possible to achieve a high hardness over a large
cross section (depending upon steel quality several 100 mm can be achieved). With
the plasma nitriding a hardness increase is possible however is limited to layer
thicknesses of some tenths mm. The highest hardness can be achieved by a PVD-
coating. However these layers have a thickness of only about 10 m.
hardness, HV
thickness,m
0 500 1000 1500 2000 2500 3000
10000
1000
100
10
1
0,5
PVD-coatings
Plasma nitriding
Heat treated steels
hardness, HV
thickness,mthickness,m
0 500 1000 1500 2000 2500 3000
10000
1000
100
10
1
0,5
PVD-coatings
Plasma nitriding
Heat treated steels
Figure 5: Comparison of the individual treatment procedures of tools
5

3. Technology of heat and surface treatment
3.1 From the semi-finished material to the tool
The mechanical tool manufacturing has large influence on the quality of the heat
treatment. For an optimal result the best way is to keep the presented production run
(Figure 6).
0
200
400
600
800
1000
1200
z.B. 1.2379
z.B 1.2767
1. Rough machine cutting 4. Heat treatment
2. Low-stress annealing 5. End measure
3. Finishing (with admeasurement) 6. Surface treatment
0
200
400
600
800
1000
1200
z.B. 1.2379
z.B 1.2767
1. Rough machine cutting 4. Heat treatment
2. Low-stress annealing 5. End measure
3. Finishing (with admeasurement) 6. Surface treatment
Figure 6: Production run during the tool production
By the generally first accomplished rough machine cutting, tool tensions are induced
which should be removed to eliminate unexpected warping after heat treatment. After
the rough machine cutting a low-stress annealing should take place. During the
following finishing it is to be made certain that sufficient admeasure is left around to
be able to adjust a warping or a change of measure arising after the heat treatment.
During the rework in the recompensed condition attention should be paid to the fact
that the working temperature does not exceed the temperature of tempering.
Otherwise the forming of a soft starter zone or in the worst case a formation of a new
hardness zone is possible.
3.2 Heat treatment-fair tool engineering
Heat treatment of tools is always connected with a growing volume, since the lattice
of the martensite is larger than that of the ferrite. Since the cooling speed is not alike
over the tool cross section everywhere, but reduces itself from the surface to the
core, tensions in the tool are being produced themselves, which can entail a warping
of the construction unit. With critical tool engineering, e.g. sharp edges or large cross
section transitions these tensions can lead to tearing of the product (Figure 7). During
the tool engineering attention should always be given a heat treatment-fair
construction (Figure 8).
6

Figure 7: Tear danger due to large cross-sections or sharp edges
Figure 8: Heat treatment-fair construction of critical ranges
With ledeburitic chromium steels, as in practice within the range of the cold work
tools predominantly used 1.2379 also the linedness of the structure is absolutely to
be considered for tool engineering.
Figure 9: Microstructure of a ledeburitic chromium steel.
Inclination to cracking
critical uncritical
7

During the heat treatment and the associated changes of measure, tensions in the
workpiece are called, which are in equilibrium. During the following rework it can
occur now that a carbide line becomes split. Thus the tension household of the
workpiece is changed, the consequence is warping. In practice it satisfactorily works
to make discharge cuts into the semi-finished material, before the heat treatment.
Thus the arising hardness tensions can be reduced, and so the danger of warping
during the rework is smaller.
3.3 Heat and surface treatment of die casting and plastic moulds
Depending upon demand of the moulds heat treatment will be arranged optimal. In
principle two routes are differentiated. With the achievement-optimized route as rapid
a cooling as possible is selected. Thus a smaller intermediate stage portion in the
structure is reached. However it is to be counted on a higher warping. With the
distortion-optimized route by the slower cooling speed a smaller warping is realized.
However it is to be counted on a higher portion of intermediate stage structures,
which affects the service lives of the forms negatively. Therefore the achievement-
optimized route is to be preferred. In practice a hardness of 42-46 HRC works the
best way. For the better evaluation of the thermal treatment in practice satisfactorily
worked, to equip the form with two sample bodies, from which the evaluation of the
annealing structure and the hardening structure can be taken (Figure 10).
Figure 10: Sample bodies at the back of a die casting mould
8

Figure 11: Annealing and hardening microstructures
On the basis of these sample bodies statements can be done, whether a bad
structure is caused by a bad heat treatment or a bad steel. In the represented
example the hardening structure shows a strong allocation of the grain boundaries
with carbides exhibits. This is to be recognized also with the annealing structure. In
this case, therefore, is not to be seen in the heat treatment. In the area of the die
casting moulds the existence of heat cracks is the main failure reason. By a heat
treatment following nitriding the formation of heat cracks can be shifted to higher tool
lives. However, then arising cracks lead due to the higher depth to the total failure of
the tool. A too deep nitriding leads due to more bring in internal voltages however
likewise to an earlier failure. For this reason nitriding depth of ~0,1mm worked in the
best way in practice. In practice die casting moulds often are treated with the so-
called UniTwin® procedure. The recompensed material is first plasma-nitrided and
coated afterwards. With the base material alloy and heat treatment-dependent
hardness, toughness and wear-resisting quality are reached. The following plasma
nitriding process promotes a higher stability due to the higher surface hardness in
relation to abrasive wear and serves for the supporting effect of the layer. The
pressure internal voltages brought in by nitrating promote the resistance against
fatigue of the base material. The PVD- coating is characterised by very high surface
hardness and small friction resistance. This offers an additional protection against
adhesion. As a result one receives a material group with a high adhesive strength of
the hard material layer on a load- carrying substrate with a high resistance in relation
to surfaces and fatigue wear. Such treated forms are additionally characterised
besides by small cold welding bar as well as very good removability.
9

Figure 12: UniTwin® - treated die casting mould.
References
1. Qualitätssicherung bei PVD und CVD- Hartstoffbeschichtungen: Anforderungen
an beschichtete Werkzeuge und Bauteile; VDI Richtlinie 3824, Blatt 2: (Februar
1997)
2. R. Chatterjee-Fischer; „Wärmebehandlung von Eisenwerkstoffen- Nitrieren und
Nitrocarbouerieren“; Expert Verlag, Renningen-Malmsheim (1995)
3. H. Hougardy: Die Umwandlung der Stähle (Teil 1+2); Verlag Stahl Eisen M.B.H.
Düsseldorf (1975)
4. H. Kunst: „Verschleiß metallischer Werkstoffe und seine Verminderung durch
Oberflächenschichten“; Expert Verlag, Grafenau (1982)
5. Michael Mack; Oberflächentechnik Verschleißschutz; Die Bibliothek der Technik,
Band 38; Verlag moderne Industrie AG & Co., Landsberg (1990)
10

PROCESS OPTIMISATION FOR DEEP COLD TREATMENT
OF TOOL STEELS
P F Stratton
BOC, c/o 42 Park Lea, Bradley, UK
Paul.Stratton@boc.com
ABSTRACT
Deep cold treatment has been applied to many materials, but is best
understood in the treatment of tools to improve their wear resistance.
Even here, inconsistencies in processing techniques have led to different
conclusions as to the exact mechanisms involved. This paper reviews
current thinking and attempts to explain these differences in terms of the
processing parameters and to elucidate the implications for the practical
application of the technique for optimum properties. An optimised
processing route is recommended and a possible further improvement in
the process is suggested.
Keyword: deep cold treatment, tool steels
1. INTRODUCTION
Despite much research over many years, deep cold treatment remains
something of a mystery. It has been reported to improve the properties of
everything from tool steels to golf balls and from nuns’ habits to copper
spot welding electrodes. The greatest improvement reported is in wear
properties. However, many of these miraculous transformations can be
ascribed to the phenomenon of structure stabilisation. It has been
suggested that the deep cold orders the structure, eliminating voids and
dislocations so that slip is less likely [1, 2]. This is the reason for deep cold
treating precision engineering parts that are exposed to rapid temperature
changes but must not deform in service. Typical examples are gun barrels
and parts of racing cars and bikes [3-8].
It might be supposed that the same argument could be applied to
improvements in the properties of polymers and natural fibres. One
example of using polymeric fibres in an extreme situation is in the sails of
racing dinghies (Figure 1). They are subjected to high oscillating loads and
usually fail by being stretched beyond the point where the sail is efficient.
11
st
IFHTSE 2005

BOC tested cold treatment of polymer fibres by running two sets of sails
on identical boats for an entire season, alternating sail sets for each race
to eliminate any effects of seamanship. One set were deep cold treated in
liquid nitrogen for 24 hours and the others were untreated. If indeed the
cold treatment did “regularise” the structure, then less stretching should
have occurred in the treated sails. Both sets were measured at regular
intervals. The deep cold treated sails fared slightly better, but the
improvement was within the statistical error of the experiment. Therefore,
there is no proven case as yet for cold treating polymeric fibres.
Figure 1. One of the test sails
2. TOOL TREATMENTS
The situation and its explanation in tool steels would, at first sight, appear
to be far more clear-cut. The treatment of steels at temperatures in the
range −80 o
C to −120 o
C is sufficient to fully transform any austenite
retained in the quenched microstructure and has been extensively used
for this purpose for many years [9]. Examples include case hardening
steels such as EN36 (832 H13, 655 H13) and EN 39 (835 H15); hardening
tool steels such as M-2 and D-2; and stabilisation of components,
particularly for the aerospace and roll making industries. The stabilisation
effect can be attributed entirely to removal of the retained austenite as it
eliminates the 4% volume change when environmental conditions
transform austenite to martensite.
Although the transformation of small volumes of well dispersed retained
austenite increases hardness, it may be counter-productive with regard to
wear as it is offset by a decrease in toughness and the ability to stop
micro-cracking. It has been reported that deep cold treatment at −196 ºC
12

(77 °K) combined with different austenitising temperatures, and hence
varying the volume of initial retained austenite, can optimise the fracture
toughness and hardness for a particular application [10-13]. However,
these studies used only short duration deep cold treatment that is known
not to optimise wear resistance [2].
3. DEEP COLD FOR WEAR RESISTANCE
The benefits of deep cold treatment (−196 °C) on the wear properties of
tools have been known for some years [14, 15], but inconsistencies
experienced by its users have limited its acceptance by European industry.
Its application has, however, been growing rapidly in the USA where there
are many specialist treatment companies. In recent years some new
theories have been developed to explain how deep cold treatment
improves wear resistance. These theories can be applied to the processing
route to help to eliminate the inconsistencies.
Most of the treatment routes cool the components in cold nitrogen gas
before finally immersing them in liquid nitrogen at −196°C, although some
processors consider it better if the components never touch the liquid,
only the cold gas. Since 1965, when commercial deep cryogenic
treatments first became available, a number of reports have been
published showing the improved performance of some tools steels treated
in this way. The most frequently quoted is that from Barron shown in
Table I [16]. However, wear is not a simple thing to measure in the
laboratory, where testing parameters significantly affect the result, as
Figure 2 shows. The actual wear experienced by a tool may be quite
different in practice [17].
Table 1. The improvement in wear for various tool steels after deep cold
treatment
Steel (AISI
No.)
Improvement in wear
rate (%)
D-2 817
S-7 503
52100 420
O-1 418
A-10 264
M-1 225
H-13 209
M-2 203
T-1 176
CPM-10V 131
P-20 130
440 121
13

Figure 2. The effect of sliding speed on the wear rate of tool steel.
It has been reported that some machine elements, such as slitting blades
used in paper cutting, have lasted six times longer after deep cryogenic
treatment. Results from field trials (Table II) on punches, stamping dies,
and milling cutters, such as those shown in Figure 3, have also shown
significant improvement, supporting the experimental data [18, 19]. Deep
cryogenic treatment is said to have one important advantage over surface
treatments aimed at increasing wear resistance, such as chromium
plating, titanium nitride coating or nitriding. This is its ability to change the
entire structure of the tool, not just its surface, so that the benefits cannot
be negated by subsequent finishing operations or regrinds [7].
From the start the cold treatment process was dogged by inconsistency.
It would work on one component, but not on another similar one in the
same material. Today those problems have been solved to some extent
through an understanding of the mechanisms involved in the wear
improvement.
TableII. Field trial of wear improvements in deep cold treated tools
Tool Type Tool Material
(AISI No)
Improvement in wear
rate (%)
Stamping die D-2 1000
Punch M-7 600
End mill M-42 450
Drills M-42, M-7, C-2 300
Milling cutters M-7 250
Drill M-42 200
Punch M-2 100
14

Figure 3. A selection of deep cold treated tools
4. THE MECHANISM OF WEAR RESISTANCE IMPROVEMENT IN
DEEP COLD TREATMENT
There have been several studies of the effect of deep cold on wear and the
mechanism that may be giving rise to the improvement [2, 15, 20-25].
Unfortunately some of the papers do not give all the processing
parameters and every processing route specified is different. The effects
of deep cryogenic processing are seen as occurring in several stages. In
the first stage, down to −130 o
C, retained austenite is transformed in
exactly the same way as in conventional sub-zero treatment, increasing
the hardness. However, Meng et al suggest that the lattice parameters of
the martensite formed are different from that formed in conventional
treatments, which may well account for its subsequent lack of response to
extended exposure to deep cold temperatures [25]. As in the conventional
treatment, the transformation is not time-dependent [2, 24].
In the second stage, which occurs at deep cryogenic temperatures
(typically −196 o
C, the temperature of liquid nitrogen), there is a time-
dependent decomposition of the primary martensite. This decomposition
causes some initial softening but nucleates numerous coherent nano-
carbides [23]. During the subsequent tempering operation the fine
ε-carbides formed and precipitated at these sites are the reason for the
increased wear resistance of the treated tools. It has been shown that the
longer the exposure to cryogenic temperatures, the more nano-carbides
are formed [24].
It is suggested, however, that it is only the primary martensite that
decomposes, and not that with the higher c/a ratio produced by the
transformation in the first stage. This mechanism goes some way to
explaining the inconsistency of the results. If a component initially has a
high retained austenite level then the transformation in the first stage will
dramatically increase the hardness, but not necessarily the wear
resistance, compared with the original state. If a component has only a
low retained austenite level then the carbide formation engendered by the
15

second stage would dramatically increase the wear resistance compared to
its original state, but without altering the hardness. It is also possible to
deep cold treat after high temperature tempering with less, but still
significant, fine ε-carbides precipitation and it has been reported that M2
that has been deep cryogenically treated twice shows further improvement
after the second treatment [25]. The “popular” literature contains many
references into the need for low cooling rates and accurate cooling curve
control down to liquid nitrogen temperatures [26]. This need appears to
be driven by the desire of the equipment manufactures to sell expensive
equipment capable of such control rather than by any evidence in the
“technical” literature. It is however, important to cool and reheat slowly
enough to avoid cracking through differential contraction/expansion and
the effects of the 4% volume change on transformation of the retained
austenite to martensite, but it is not critical to ε-carbide formation. It is
often reported that deep cold treatment should follow immediately after
quenching, but it has also been suggested that a short warm (60 ºC)
ageing after quenching can reduce the tendency to cracking [23]. It has
been shown more ε-carbides are produced by a longer deep cold
treatment at cryogenic temperatures and, as might be expected, the
subsequent wear rate also falls with increasing exposure time [24, 25].
The recommended minimum time at low temperature is 24 hours,
however, extended or multiple treatments are known to be beneficial. Tool
steels that normally exhibit secondary hardening do not do so after deep
cold treatment so they can be tempered at a lower temperature to
maintain hardness [24,27]. However, the wear rates are lower when the
steel is tempered at its normal tempering temperature because the
morphology of the larger carbides is improved.
5. RECOMMENDED PROCESS STEPS
It is not possible to recommend a single process for every tool steel, nor
even a single cycle for all tools manufactured from the same steel. Each
tool needs to be separately assessed and an individual process route
devised for it that will depend on the combination of hardness, toughness
and wear resistance required in service. The cryogenic treatment is just
one step in that process and must be integrated into the processing route
[28]. BOC recommends that tools be treated in specially designed
equipment using liquid nitrogen as the refrigerant. The liquid nitrogen is
supplied by BOC either via small portable storage containers or static,
externally sited, vacuum insulated vessels for larger volumes. The whole
process cycle can be automatically controlled for greater consistency and
reproducibility. In addition to deep cold capability, some of the units on
the market also incorporate a temper/stress relief facility. This releases
the heat treatment furnaces for further treatments.
16

Figure 4. A schematic of the recommended processing route
.
The steps in the processing route for maximum wear resistance in Figure 4
are:
1. Heat to an austenitising temperature that will minimise retained
austenite in the tool steel being treated
2. Hold for the recommended time for the steel
3. Quench at a rate sufficient to give a fully martensitic structure
4. Condition at 60ºC for a maximum of one hour and immediately go to
step 5
5. Cool to liquid nitrogen temperature (−196ºC) at a rate slow enough to
prevent cracking, preferably in a nitrogen atmosphere to avoid
condensation
6. Hold at liquid nitrogen temperature for a minimum of 24 hours,
preferably in a nitrogen atmosphere to avoid condensation
7. Reheat to room temperature at a rate slow enough to prevent cracking,
preferably in a nitrogen atmosphere to avoid condensation
8. Temper at the temperature recommended for the steel being treated.
6. FURTHER IMPROVEMENTS
There is insufficient driving force at liquid nitrogen temperature (−196ºC)
to form the nuclei of the ε-carbides in the martensite formed at low
temperature [2, 22]. This limits the useful application of the deep cold
process to tool steels that have been austenitised at lower temperature to
minimise retained austenite formation, but which, however, also
minimises hardness. If the driving force were increased by cooling to a
lower temperature, then it might be possible to form carbides in the
martensite formed at low temperatures, thus maximising performance
with a combination of high hardness and fine ε-carbide dispersion. The
obvious choice of refrigerant would be liquid helium at −269 ºC (4 ºK).
However, it is inevitable that treatment times would be very long as there
is little atomic movement at such a low temperature.
17

7. CONCLUSIONS
As part of an optimised heat treatment cycle, deep cold treatment can
dramatically improve measured wear by the precipitation of fine ε-carbides
in the primary martensite. The transformation of retained austenite to
martensite is a minor additional benefit. In many practical uses of tools
this increase in measured wear translates into longer tool life.
REFERENCES
1. P Paulin “Mechanism and Applicability of Heat Treating at Cryogenic
Temperatures”, Industrial Heating, August 1992, pp22-27
2. Collins, D.N., “Deep cryogenic treatment of tool steels: a review”,
Heat Treatment of Metals, 1996, Vol.23, No.2.
3. “Freezing barrels for better accuracy”, Guns & Ammo, April 1996, pp
66-67
4. R T Dohacz, “Rocket Science: adding power and reliability with
cryogenics”, Drag Racing USA, August 2003, pp 28-30
5. P Paulin, “Cryo-Rifles: Deep Cryogenic Stress Relief”, Precision
Shooting, March 1995, pp 74-76
6. “Deep Freeze, Deep Secret?”, Race & Rally, January, 1996, pp 26-56
7. Paulin, P., “Cold cuts”, Cutting Tool Engineering, Vol.44, No.5, August
1992.
8. R Schiradelly and FJ Diekman, “Cryogenics; the racers edge”, Heat
Treating Progress, November 2001, pp 43-49.
9. Moore, C, “Development of the BOC Ellenite process (cold treatment
of metals with liquid nitrogen), Heat Treatment ’73, The Metals
Society,1975, Book No 163 pp 157-161
10. R Mahmudi, HM Ghasemi and HR Faraji, “Effects of cryogenic
treatments on the mechanical properties and wear behaviour of high-
speed steel M2”, Heat Treatment of Metals, 2000, Vol.27, No.3,
pp.69-72
11. V Leskovsek, B Liščić and B Ule, “Some aspects of sub-zero
tempering at vacuum heat treatment of HSS”, Proceedings of the
21st Conference 5-8 November 2001 CD ROM, Heat Treating Society
12. V.Leskovsek and B. Ule, “Influence of deep cryogen treatment on
microstructure, mechanical properties and dimensional changes in
vacuum heat-treated high-speed steel”, Heat Treatment of Metals,
2002, Vol.29,No.3, pp72-76
13. AI Wojcieszynski, “Cryogenic treatment: a mystery or misery of heat
treatment”, Proceedings of the 19th
Heat Treating Society conference,
pp 237-243, 1999
14. Keen, AR, “Cryogenic treatment to improve wear resistance of steel
by the “Cryotough” process”, Metals Australasia, August 1982, Vol.14,
No.7. 12-12,21
18

15. Reasbeck, RB, “Improved tool life by the Cryotough treatment”,
Metallurgia, April 1989, Vol.56, No.4, pp 178-179
16. RF Barron, “Cryogenic treatment of metals to improve wear
resistance", Cryogenics, Vol.22, No.8. pp 409-413, 1982
17. RC Lasky, “The effects of cryotempering on tool steels”,
http://www.nitrofreeze.com/toolsteels.html
18. http://www.cryoeng.com/images/Field%20Test%20Results.pdf
19. http://www.diversifiedcryogenics.com/testresults.htm
20. L Alexandru, G Coman and V Bulancea, “The change of the
substructure elements and the redistribution of the alloying elements
by means of cryotreatments in alloy tool steels”, Proceedings of the
5th
International Congress on Heat Treatment of Materials, Vol.2,
pp.901-908, 1986
21. L Alexandru, C Baciu and G Ailincai, “Contributions on the study of the
increase of durability of the high-alloyed tool steels by thermal
treatments at cryogenic temperatures”, Memoires et Etudes Sci. Rev.
Metall. 1990, Vol.87, No.6, pp.383-339
22. Dormer, J., “The cryogenic treatment of tool steels”, Thesis, National
University of Ireland, University College Dublin, August 1994.
23. F Meng, K Tagashira, R Azuma amd H Sohma, “Role of eta-carbide
precipitation’s in the wear resistance improvements of Fe-12-Cr-Mo-
V-1.4C tool steel by cryogenic treatment”, ISIJ International, Vol.34,
No.2, pp 205-210, 1994
24. DN Collins and J Dormer, “Deep cryogenic treatment of a D2 cold
work tool steel”. Heat Treatment of Metals, Vol.24, No.3, pp 71-74,
1997
25. D Yun, L Xiaoping and X Hongshen, “Deep cryogenic treatment of
high speed steel and its mechanism”, Heat Treatment of Metals,
Vol.25, No.3, pp 55-59, 1998
26. DL Hallum, “Cryogenic tempering delivers better cutting tool
durability”, American Machinist, May, 1996, pp140-141
27. DN Collins and J Dormer, “Deep cryogenic treatment of an ASP 23
high speed steel”, Proceedings of the 18th
Heat Treating Society
conference, pp.255-258, 1998
28. DN Collins and G O'Rourke, “The response of tool steels to deep
cryogenic treatment: effect of alloying elements”, Proceedings of the
18th
Heat Treating Society conference, pp.229-247, 1998
29. http://www.cryogenic.co.nz/index.cfm?fuseaction=dsp_content&page
_id=23
19

WEAR BEHAVIOUR OF DEEP CRYOGENIC TREATED
HIGH SPEED STEELS
M. Kalin 1, *
, V. Leskovšek 2
, J. Vižintin 1
1
Center for Tribology and Technical Diagnistics, University of Ljubljana,
Bogišičeva 8, 1000 Ljubljana, Slovenia (* mitjan.kalin@ctd.uni-lj.si )
2
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
ABSTRACT
Tools for the cold-working applications are typically made from the high-speed steels.
However, due to wear and plastic deformation their performance in several applications is not
adequate and should be further improved. By using appropriate combination of vacuum heat-
treatment in conjunction with a deep-cryogenic treatment (duplex treatment) the
microstructure of high-speed steel matrix can be substantially changed and the hardness and
fracture toughness can be modified and optimised. In the present work we have investigated
the effect of four different tempering temperatures of vacuum and cryogenic treated ESR AISI
M2 high speed steel on the resulting combinations of microstructure, hardness and toughness
and their effect on the wear mechanisms at different loads. The results showed that at
relatively high loads the different treatments resulted in an order of magnitude difference of
wear resistance, while at low loads the selected treatments were efficient enough to keep the
wear within the mild wear regime and small variations between the samples. However, the
overall wear transition did not occur at any load used or any sample treatment, although some
small differences in wear mechanisms can be seen, primarily depending on the fracture
toughness of the samples.
Key words: cryogenic and duplex treatment, high speed steel, wear, friction, microstructure
1. INTRODUCTION
The wear resistance of high-speed steels is largely influenced by the microstructure, which
consequently affects the hardness and fracture toughness of the material. In hardened and
tempered condition, the important constituents of high-speed steel are the matrix, i.e.
tempered martensite and retained austenite, and un-dissolved eutectic carbide particles. These
phases importantly affect the wear-resistance of the steel. However, when applying the deep-
cryogenic treatment followed by single tempering, the matrix of high-speed steel can be
additionally tailored [1-4]. By using appropriate combination of vacuum heat-treatment in
conjunction with a deep-cryogenic treatment the microstructure of high-speed steel can be
substantially modified, which consequently affect the ratio between the hardness and fracture
toughness. Moreover, a semi-empirical equation (Eq.1) was recently derived for the HSS [5],
providing the relation between the fracture toughness KIc, Rockwell-C hardness HRC and
21
st
IFHTSE 2005

quantified microstructural parameters. Several microstructural parameters, like mean distance
between undissolved eutectic carbide particles dp, volume fractions of undissolved eutectic
carbides fcarb, volume fractions of retained austenite faust and some other material properties
like modulus of elasticity and hardness were considered. The resulting calculated fracture
toughness agreed very well with the experimental results [6], which confirmed a strong
dependence between microstructure, hardness and fracture toughness [7].
( ) ( )



+⋅⋅⋅⋅





−
⋅= 





−
aust
6
1
carbpIc f1fdE
53HRc
HRc
1.363K (Eq. 1)
In our previous study, we have investigated the effect of deep cryogenic treatment on the wear
and friction behaviour of high-speed steel by comparing the conventionally treated samples
and the additionally deep cryogenic treated samples. The deep cryogenic treated samples with
an optimal ratio of the above mentioned parameters showed significantly better results [8]. In
the present work we present the effect of deep cryogenic treatment of vacuum hardened ESR
M2 high speed steel (tempered at four different temperatures, and initially austentized) at
three different loads, to cover the load range important from the application point of view.
The effect of resulting hardness, toughness and microstructure on the acting wear mechanisms
in reciprocating sliding tests were investigated.
2. EXPERIMENTAL
2.1 Material characterisation and duplex treatment
Investigated high-speed steel M2 was delivered in the shape of rolled, soft annealed bar of φ
20 mm x 4000 mm. The bar was cut into metallographic samples of φ 20x9 mm. All the
samples were vacuum heat treated in a horizontal vacuum furnace with uniform high-pressure
gas quenching using N2 at a pressure of 5 bars. After the last preheat the specimens were
heated to the austenitizing temperature of 1230 °C, soaked for 2 minutes, and gas quenched to
25 °C with the cooling parameter λ800-500 = 0,55. The specimens were then removed from the
furnace for a subsequent deep-cryogenic treatment, followed by a single tempering cycle for
one hour at four different and carefully selected temperatures, i.e. 500, 540, 550, and 600o
C to
obtain specific hardness and toughness properties. The deep cryogenic treatment of
metallographic specimens was performed by a controlled immersion of individual test
specimens in liquid nitrogen (-196o
C). After equalization of the temperature (when the liquid
N2 ceased boiling) the specimens were soaked in liquid nitrogen for one hour.
After the duplex treatment of the samples, they were ground and polished for further analyses.
The resulting Ra roughness of all samples was better than 0.01 m. The Rockwell-C hardness
was measured on the metallographic specimens using a Wilson 4JR hardness tester. The
microstructures of the investigated metallographic specimens were assessed as described in
detail elsewhere [9] and the fracture toughness was calculated using Eq. (1) [5]. The results
are presented in Table I.
Table I. Results of Rockwell hardness and fracture toughness for investigated samples
Samples at different
tempering temperatures
Rockwell-C hardness
( HRc )
Fracture toughness KIC
( MPa m1/2
)
A-500 65.2 8.5
B-540 65.8 8.0
C-550 65.5 8.8
D-600 64.4 9.5
22

The micrographs in Figure 1. show the microstructure of the samples that were later used in
wear experiments. It can be seen that deep cryogenic treatment following vacuum hardening
results in a continuation of the austenite-martensite transformation. The retained austenite is
not visible in the matrix of tempered martensite. From the micrographs in Figure 1. it can be
also seen that a rod-like carbide precipitation occurs during the tempering that follows the
deep cryogenic treatment. These precipitates are approximately 20 to 40 nm long and 5 to 10
nm thick. The rod-like carbide precipitation in the investigated high-speed steel occurs for
given vacuum heat-treatment conditions in larger quantities during the tempering following
the deep cryogenic treatment. The size of these particles precipitated in the matrix of freshly
formed martensite depends on the tempering temperature. They are coarser after tempering at
500 °C and finer after tempering at 600 °C.
As can be seen from the Table 1, in the whole range of the tempering temperatures used, there
is no significant variation in fracture toughness. This could be attributed to the fact that the
deep-cryogenic treatment following the vacuum hardening results in an almost complete
transformation of retained austenite to martensite. It seems that the subsequent single
tempering does not have any significant impact on the fracture toughness. The secondary
hardness maximum and fracture toughness minimum is attained by single tempering at
540°C. Namely, the net effect of tempering is attributed to a combination of stress relief and a
reduction in the ductility due to the secondary hardening peak.
Figure 1. Microstructure of metallographic specimens, which were vacuum hardened, deep cryogenic treated at –
196 o
C, and finally single tempered at temperatures as follows:
(a) A-500 at 500 o
C, (b) B-540 at 540 o
C, (c) C-550 at 550 o
C, and (d) D-600 at 600 o
C.
2.2 Wear experiments
Steel samples for the wear experiments were the discs (φ 20x9 mm), same as used in
metallographic investigation, having four different types of vacuum treatment. To obtain a
reasonably high and measurable amount of wear, almost two times harder (≈ 16,7 GPa)
material was selected for the counter samples, i.e. silicon nitride ceramic. The ceramic balls
(Cerbec Corp., East Granby, CT) were standard bearing balls (grade 5) with a diameter of
12.7 mm and had fine polished surfaces with Ra roughness value better than 0.03 m.
Wear experiments were performed with a reciprocating sliding test machine at a constant
frequency of 2 Hz and stroke length 6.8 mm. The steel discs were stationary in the test
machine, while the counter balls were sliding in a reciprocating motion. Relative contact
velocity was 27,2 mm/s. The total sliding distance in each test was approximately 200 m,
corresponding to 28,800 loading cycles. Three loads, 5, 20 and 50 N were used, which
resulted in contact pressures of 1,0, 1,5 and 2,0 GPa. All experiments were performed under
room ambient conditions (≈ 20 o
C and ≈ 50 % relative humidity). Prior to wear experiments
the samples were ultrasonically cleaned in acetone and ethanol. After the wear experiments,
23

further cleaning and profilometric analyses, the discs were sputter coated with gold and
examined by Scanning Electron Microscopy (JEOL JSM-T330A, Tokyo, Japan).
3. RESULTS AND DISCUSION
3.1. Wear
Figure 2a shows the wear volume of the four different steel samples for the three loads used.
It can be seen that at 20 and 50 N the wear is reduced as the tempering temperature increased
from 500 to 600 o
C (samples A-500 to D-600). This change is more pronounced at the highest
load used, where the difference in wear between the sample A-500 and D-600 reached about
an order of magnitude. On the other hand, at the smallest load of 5 N, a slight increase in wear
volume can be observed as the tempering temperature is increased (samples A-500 to D-600),
however, the change is very low and varies within the scatter of the results. The steady and
relatively small changes in wear with the tempering temperature increase suggest that there
has no wear transition occurred under the testing conditions at any load or type of sample
treatment. Moreover, corresponding dimensionless wear coefficients were all in the range of
10-6
or just slightly above it, which indicate the mild wear regime in the selected tests. Figure
2b shows the coefficient of friction for all three loads used as a function of tempering
temperature. It can be seen that coefficient of friction decreases as the tempering temperature
increases. The coefficient of friction decreases with increased contact load, which suggests
the faster and more “thorough” deformation of surface asperities, in terms of running-in
process. An exemption from these two clear trends is the behaviour of sample D-600 at the
lowest load of 5 N, which is again consistent with the wear data (Figure 2a), where also a
small increase in wear was observed, as discussed above.
0
0,002
0,004
0,006
0,008
0,01
0,012
0,014
0,016
A-500 B-540 C-550 D-600
Wearvolume(mm
3
)
5 N
20 N
50 N
0,4
0,5
0,6
0,7
0,8
0,9
1
A-500 B-540 C-550 D-600
Coefficientoffriction
5 N
20 N
50 N
Figure 2. (a) Wear volume and (b) coefficient of friction in experiments at three selected loads: 5, 20 and 50 N
3.2. SEM analyses
Figure 3 shows the SEM micrographs of the tested samples at the lowest load used, 5 N. We
can see that the surfaces are very little damaged and that there is no significant difference
between them. Worn surface of the sample B-540 (Figure 3b) seem to be slightly more
damaged, but in the areas where it is not damaged is also quite smooth. This could indicate
the relatively low ability of plastic deformation and immediate damage (fracture) of the
surface, when the stresses locally exceeded the limit values. Such observation is in agreement
with the lowest fracture toughness and the highest hardness of this sample among the tested
samples, Table 1. To a smaller extent, but similar type of damage can also be observed in
sample A-500. In addition, the surfaces of samples A-500 and B-540 have less structured
topography than those of samples C and D, indicating distinction in microstructure, and lower
deformation and/or smearing at the surface, again in agreement with the hardness and fracture
toughness data. On the other hand, sample D appears to be more plastically deformed, and
24

also several scratches can be observed on the surface in direction of sliding, which is, again,
consistent with it's the highest fracture toughness and the lowest hardness among the tested
samples, respectively. Moreover, in some locations (see Figure 3d) clear formation of voids
and cracks can be observed in the matrix in the vicinity of carbides due to the mismatch of
their elastic properties. This might be one of the reasons for slightly higher wear observed
with this sample. Nevertheless, the wear of all samples tested at 5 N was rather low (Figure
2a) and except few scratches and slight plastic deformation/smearing, no other types of
damage could be observed.
Figure 3. SEM Micrographs of samples tested 5 N:
(a) A-500 at 500 o
C, (b) B-540 at 540 o
C, (c) C-550 at 550 o
C, and (d) D-600 at 600 o
C.
More wear and damage can, however, be observed at higher loads. Steel samples tested at 20
and 50 N have very similar appearance, only the extent of damage was higher at 50 N. Figure
4 shows the worn surfaces of the four steel samples after the tests at 50 N. Sample A-500
seems to be the most damaged, with several deformation ridges that are extended in the
direction of sliding, suggesting ductile behaviour. On the contrary, fractures of the ridges with
sharp edges, which subsequently delaminate, suggest also a degree of brittle behaviour of this
sample and causing the highest wear in our investigation (Figure 2a). The worn surface of
sample B-540 is clearly less damaged than A-500. However, there are also many ridges
observed on the surface, but in contrast to sample A-500, they are shorter and thinner and the
distances between them are smaller. In addition, the ridges appear to be fractured at their ends
and tent to extend in direction perpendicular to sliding. Nevertheless, the wear resistance of
this surface is obviously higher than that of A-500. On the sample C-550 even more ridges
can be observed (Figure 4c), although the wear is lower than in previous two cases, Figure 2a.
However, the ridges are smoother and smeared, thus more plastically deformed indicating
more ductile behaviour than those on samples A-500 and B-540. Therefore, less wear debris
are formed and consequently the wear was lower, Figure 2a. In distinction, the worn surface
of the sample D-600 is much less damaged than all other surfaces. Figure 4d shows one of the
most damaged areas, while most of the wear scar was even less damaged. However, the few
ridges observed deform plastically to a great extent and due to their prolongation in direction
of sliding they seem to smear easy. The ductile behaviour of samples C-550 and D-600,
which agrees with their high fracture toughness, protect the surfaces from high wear, as can
be deduced also from the low wear. In distinction, at the sample B-540 that has lower fracture
toughness, the deformation ridges appear more of brittle nature, they form wear debris easier
and the wear was higher, Figure 2a.
From our results it appears that even a small difference in fracture toughness plays the very
important role, especially at high loads, however, it is also clear that fracture toughness alone
is not the only relevant parameter. Namely, as the tempering temperature increases from
samples B-540 to D-600, hardness clearly decreased, while the fracture toughness increased,
and the wear decreased. On the other hand, sample A-500 has higher fracture toughness than
25

sample B-540, but the wear was higher (although not significantly). However, the hardness of
A-500 was relatively much lower. Moreover, since a slight change from brittle to ductile wear
behaviour can be observed from SEM analyses between the samples B-540 and C-550, this
could suggest that a certain threshold value for toughness is needed for the toughness to
become predominant; in our case at 8,7 MPa √m. In this case, even much lower hardness
(sample D-600) did not make the wear resistance worse. The effect of microstructure was
primarily related to the resulting combination (ratio) of fracture toughness and hardness and
particularly to increased fracture toughness with increased tempering temperature.
Figure 4. SEM Micrographs of samples tested 50 N:
(a) A-500 at 500 o
C, (b) B-540 at 540 o
C, (c) C-550 at 550 o
C, and (d) D-600 at 600 o
C.
4. Conclusions
1. Microstructure, fracture toughness and hardness interplay the decisive role for the wear
resistance of the HSS. In our experiments higher fracture toughness appear to be more
important for better wear resistance than hardness. The microstructure affect the results
primarily indirectly through improved fracture toughness.
2. At low load (contact pressure 1.0 GPa) the wear resistance of all samples was practically
the same and the wear corresponded to mild wear regime.
3. At high loads, the wear resistance of the samples was significantly different, resulting in
an order of magnitude difference. Wear decreases as the tempering temperature increases.
4. Plastic deformation with formation of ridges that behave more brittle or ductile (in
accordance with their hardness and fracture toughness) was the predominant wear
mechanism. No overall wear transition was found in this work.
References
1. D.J. Kamody, Advanced Materials & Processes 10 (1998) p. 215.
2. P.-L. Yen, and D.J. Kamody, Industrial Heating 1 (1997) p. 40.
3. F. Meng, K. Tagashira, R. Azuma and H. Sohma, ISIJ International 34 (1994) p. 205.
4. M. Pellizzari, and A. Molinari, in The use of steels: Experience and research, edited by J: Bergstrom,
G. Fredriksson, M. Johansson, O. Kotik, and F. Thuvander (Proceedings of the 6th
International tooling
conference, Karlstad University, 10-13 September, 2002) p. 547.
5. V. Leskovšek, B. Ule and B. Liščić, Journal of Materials Processing Technology 127 (2002) p. 298.
6. V. Leskovšek, and B. Ule, Heat Treatment of Metals 3 (2002) p. 72.
7. V. Leskovšek, and B. Ule, Journal of Materials Processing Technology 82 (1998) p. 89.
8. V. Leskovšek, M. Kalin, J. Vižintin, Trans. of Materials and Heat Treatment, vol. 25 (2004) p. 540.
9. G.F. Vander Voort, Metallography; McGraw-Hill Book Company, NY, USA (1984), p. 410.
26

IMPROVING TOOL PRODUCT PERFORMANCE THOUGH THE USE
OF INTENSIVE QUENCHING PROCESSES
M.A. Aronov, J.A. Powell and N.I. Kobasko
IQ Technologies Inc
Akron, Ohio, USA
ABSTRACT
The intensive quenching (IQ) method is an innovative thermal process for hardening steel
parts. In contrast to conventional quenching in oil or polymer, the IQ process is an
environmentally friendly process conducted in highly agitated plain water. One of the major
benefits of the IQ technique is the development of high, beneficial residual compressive
stresses in the part surface layer during quenching. The IQ process is interrupted at the
computer-calculated time when residual surface compressive stresses reach their maximum
value. Residual surface compressive stresses improve part performance characteristics
(strength, fatigue and wear resistance, service life, etc.). The paper describes applications of
the IQ process to a variety of tool products made of shock-resistance cold work AISI S5 steel,
high-speed M2 steel, 52100 steel and others. The paper describes intensive quenching
equipment used for IQ demonstration studies.
Key words: Intensive quenching, tool products, residual surface compressive stresses,
service life, and intensive quenching equipment.
1. INTRODUCTION
In 1983, authors [1, 2] conducted a computer simulation study on the effect of the cooling
rate during quenching on residual surface stresses in the part. The results of calculations
performed from conventional quenching in oil confirmed that the higher the cooling rate
during quenching, the greater the residual surface tensile stresses in the part. However, the
results of calculations showed that when quenching parts intensively (with a much higher
heat extraction rate than in oil) the residual compressive stresses develop in the part surface
layer. Table I presents the data on residual stress conditions obtained for different Bi numbers
[1, 2]. The Bi number is equal to h⋅R/λ, where h is a heat transfer coefficient on the part
surface, R is a part dimension characteristic (for example, a radius for a cylindrical part), and
λ is the part material thermal conductivity. Thus, for a given part, the Bi number characterizes
the heat extraction rate. As seen from the table, for Bi numbers below 10 (that is a range for
conventional quenching), the residual surface stresses are tensile.
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IFHTSE 2005
27

For Bi numbers exceeding 20 (for intensive quenching conditions), there are residual surface
compressive stresses and they are more compressive with the increase of the Bi number. The
computer simulations showed also that this fact is correct even for through hardened parts.
However, at a time of conducting this study, the heat-treating community did not accept these
results. This is mainly because the findings presented in [1, 2] contradicted accepted heat-
treating knowledge and practice and because of the lack of experimental data supporting the
results.
Table I. Hoop residual stresses at the surface of steel parts vs. Biot number Bi
Biot Number Bi 0 2 5 10 20 40 80
Hoop residual stresses, MPa 0 + 300 + 400 +200 - 30 - 280 - 630
Over the last several years, IQ Technologies Inc. has conducted a numerous experimental IQ
studies of part residual stress conditions after intensive quenching. We considered a variety of
steel parts made of different steels. This paper summarizes these results.
2. INTENSIVE QUENCHING EQUIPMENT USED FOR IQ DEMONSTRATION
STUDIES
For IQ trials, we used two types of IQ equipment: a) a pre-production high-velocity single-
part quenching IQ system where we quench parts one-by-one out of a neutral salt bath
furnace, and b) production batch type IQ units where we quenched parts in batches in an IQ
water tank out of an atmosphere furnace. The pre-production high-velocity IQ system was
specifically designed for the implementation of a so-called IQ-3 quenching process [5]. When
applying the IQ-3 quenching technique, the water flow velocity along the surface of the part
being quenched is so high and the convection heat transfer is so great that any boiling
processes (both film boiling and nucleate boiling) are fully eliminated. The convection heat
transfer prevails from the very beginning of the quench. Therefore, the IQ-3 process is also
called “direct convection cooling.” The system is able to provide optimum IQ-3 quenching
conditions to a variety of steel products.
Figure 1 presents a schematic of the high-velocity IQ system. The IQ system works as
follows. Initially, the IQ system is at an idle condition: the pump, 2, is “ON” pushing the
water from the tank, 1, through the 3-way valve, 3, and a bypass pipe, 10, back to the tank.
The loading table, 7, with an attached fixture is in the lower position. A hot part to be
quenched is put into the lower section of the fixture that is attached to the loading table. The
air cylinders, 8, move the loading table, 7, with the part upward into stationary upper section
of the fixture. The upper section of the fixture (not shown) is a pipe that is attached to the
tube, 11. The lower end of the upper fixture has a flange with an attached rubber ring. When
the loading table is at the upper position, the rubber ring is held against it providing sealing of
the system. As soon as the part is in position within the upper section of the fixture, the three-
way valve, 3, switches the water flow from the idle position into the piping, 4, for intensive
quenching of the part. After the quench is completed, the 3-way valve, 3, switches the water
flow back to the bypass pipe, 10, and the air cylinders lower the loading table with the part.
Note that when the system is in quenching mode, the water flow may be split in two flows or
streams after passing the 3-way valve. A shut-off valve, 5, and a flow meter, 6 control each
water flow path. The reason for this is that when quenching bearing rings it is necessary to
control two water flows: one along the ID surface and along the ring OD surface. The high-
28

velocity IQ system is capable of intensively quenching steel parts up to 15 cm in diameter
and up to 40 cm in length. Figure 2 presents a schematic of one of IQ Technologies full-scale
production IQ systems installed at the Akron Steel Treating Co. of Akron, Ohio. The system
includes a Surface
Figure 1. High-Velocity IQ System Schematic
Figure 2. Production 6,000-gallon IQ System Installed at Akron Steel Treating Co.
Combustion atmosphere furnace having a work-zone of 91cm×91cm×122cm (36”×36”×48”)
and the IQ quench tank of 22.7 m3
(6,000 gallons) across the aisle. The mild steel IQ tank is
equipped with four 46 cm (18”) propellers that are rotated by four motors. The tank uses
plain water with 8% to 10% sodium nitrite solution as the quenchant. The quenchant flow
1 2
3
4
5
6
7
8
9
1
0
1
1
1
2
29

velocity in the tank is about 1.5 m/sec (5 ft/sec) as it passes over the parts. A chiller maintains
the quenchant temperature within the required limit. The production IQ system is designed
for quenching loads of up to 1,135 kg (2,500 lb). Our second production IQ unit is built by
AFC-Holcroft Co. of Wixom, Michigan. It is an integral quench furnace of
91cm×91cm×182cm (36”×36”×72”) equipped with a 41.6 m3
(11,000 gallons) IQ water tank.
The unit is installed at the Euclid Heat Treating Co. of Cleveland, Ohio.
When quenching steel parts in batches, we implement a so-called IQ-2 intensive quenching
process. During the IQ-2 process, an initial film boiling stage of heat transfer is fully
eliminated due to a high water agitation rate and due to the presence of the salt in water.
Very intensive nucleate boiling mode of heat transfer starts practically immediately after the
load is immersed into the IQ water tank.
3. INTENSIVE QUENCHING OF SHOCK RESISTING PUNCHES
To evaluate an effect of intensive quenching on the punch service life improvement, we
quenched twelve square 17.5x17.5 mm (11/16”x11/16”) S5 steel punches (Figure 3). The
punches were austenitized in a neutral salt bath furnace and quenched in batches (six punches
per batch) in the IQ system. Twelve identical punches made from the same steel heat were oil
quenched out of a vacuum furnace in accordance with the current technology. Figure 4
presents the results of measurements of residual surface stresses for the oil-quenched punches
and for the intensively quenched punches conducted by the Lambda Laboratory of
Cincinnati, Ohio. As seen from the figure, there are residual surface tensile stresses about 200
MPa in the oil-quenched punch, while there are very high residual compressive stresses in the
range of about –1,000 to –500 MPa in the intensively quenched punch surface layer.
Figure 3. Square S5 Steel Punch Figure 4. Punch Circumferential Stresses
Both the intensively quenched punches and the oil-quenched punches were put into
the field by the end-user. In the field-testing, the punches punched 17.5 mm (11/16”) holes
through 15.9 mm (5/8”) and 19.1 mm (¾”) thick 1085 steel material using a 19.1 mm (¾”)
square female die in a single station 250-ton mechanical press. “Service life” (as defined by
the punch user) is when chipping or wear is “excessive” and the punched holes are no longer
acceptable. The press cycled every 15 seconds. Oil quenched punches lasted approximately 1
hour and made an average of approximately 450 holes in the 1085 material. While intensively
quenched punches lasted approximately 2 hours and made on average approximately 900 holes. Thus,
-1000
-800
-600
-400
-200
0
200
400
0 0.1 0.2 0.3 0.4 0.5 0.6
Depth, mm
ResidualStress,MPa
Oil Quenching
Intensive
Quenching
30

the IQ process improved the service life of the S5 punches by about two times. Table II below
summarizes the punch properties improvements due to the intensive quenching process.
Table II. Improvement of S5 Steel Punch Sample Properties *
Property Oil Quench Intensive Quench
Hardness, As-quenched 62-63 63-64
HRC As-tempered 60-61 60-61
Impact strength, @72o
F 1.36 4.08
N⋅m @100o
C 3.4 6.12
Residual stresses, MPa 200 -900
*As measured by Case Western Reserve University of Cleveland, Ohio, USA.
4. INTENSIVE QUENCHING OF OTHER STEEL PARTS
Table III presents the results of mechanical property improvements obtained for other parts
that were intensively quenched. Table IV presents the data on residual surface compressive
stresses. Note that most of specimens were through hardened.
Table III. Part Property Improvements
Steel Part Property/Performance
Characteristic
Improvement
Surface hardness 5-10%
Core hardness 20-50%
Springs, shafts, bearing
rollers, bearing rings,
fasteners, sprockets Hardened depth 50-600%
Forklift forks,
fasteners, springs
Strength (core) 20-30%
Punches, dies, fasteners Toughness 30-300%
Punches, coil springs,
leaf springs, forklift
forks
Service life/fatigue resistance 50-200%
Table IV. Experimental residual surface compressive stresses for different steel parts
Part Residual Surface Compressive Stresses,
MPa
52100 bearing ring -136
52100 roller, D = 76 mm - 348
4140 Kingpin, D = 46 mm - 563
S5 Punch, D = 38 mm - 750
5160H Torsion Bar, D = 36 mm - 311
1547 Cylinder, D = 51 mm - 515
31

CONCLUSIONS
1. Both the computer simulations and experimental data show that when applying the IQ
process there are always high compressive residual stresses in the part surface layer
regardless of whether the part is quenched through or has a mixed structure in the core.
2. The experimental results showed that the value of residual compressive stress when the
Biot number is in the range of Bi = 20 to 80 depends on the content of carbon in the steel,
i.e. on the start temperature of the martensitic transformation. The higher the martensitic
start temperature the higher the residual compressive stresses on a surface of the
quenched products.
3. With increased intensity of cooling within the martensite temperature range, the
mechanical properties of a material are improved.
4. Both high compressive residual stresses at the surface of components and improved
mechanical properties from intensively quenched steel parts increased the part service
life.
5. It is very important to continue experimental and computational studies of the established
findings.
REFERENCES
1. N.I.Kobasko, W.S.Morhuniuk, Study of thermal and stress-strain state at heat
treatment of power plant products, Kyiv, Znanie, 1983, 16 p
2. Kobasko, N.I., Intensive Steel Quenching Methods. – In a Handbook "Theory and
Technology of Quenching", B.Liscic, H.M.Tensi, W.Luty (Eds.), Berlin, Springer-
Verlag, 1992. – pp.367-389.
3. N.I.Kobasko, Steel quenching in liquid media under pressure, Kyiv, Naukova Dumka,
1980, 206 p
4. N.I.Kobasko, W.S.Morhuniuk, V.V.Dobrivecher, Software “Tandem-Hart Analysis”,
commercially available from Intensive Technologies Ltd. Kyiv, Ukraine (e-mail:
managers@itl.kiev.ua, www.itl.kiev.ua )
5. M. A. Aronov, N. I. Kobasko, J. A. Powell, “Practical Application of Intensive
Quenching Process for Steel Parts”, Proceeding of The 2000 Heat Treating
Conference, St. Louse, (2000).
32

VACUUM FURNACE – INTEGRATED “SUB-ZERO” TREATMENT
B. Zieger, R. Stein
SCHMETZ GmbH, Holzener Strasse 39, 58708 Menden, Germany
ABSTRACT
The vacuum heat treatment with overpressure gas quenching is more and more
accepted due to considerable advantages compared to the traditional oil and salt
bath processes. Continuous further developments and new concepts like multi-
directional cooling systems, a separate quenching chamber and „sub-zero“ systems
lead towards an oxidation free and low distortion vacuum heat treatment for a broad
range of parts and materials. Short and energy saving processes guarantee a high
economic efficiency and environmental compatibility.
The „sub-zero“ system which is integrated into the standard vacuum furnace
achieves a heat treatment result with a high conversion of retained austenite in fully
automatic hardening and tempering processes.
Key words: vacuum heat treatment, cryogenic treatment, tool steel, corrosion resistant
steel
1. GENERAL INFORMATION
Since the introduction of the heat treatment into vacuum furnaces with overpressure gas
quenching more than 25 years ago continuous new developments and new concepts
have lead to a technology that has many advantages compared to the conventional salt
bath process:
• no decarburization
• no oxidation of parts – bright surfaces
• defined temperature guidance with load thermocouples – reproducible results
• complete documentation of the load’s time/temperature process actual values
• fully automatic, man less heat treatment process
• high temperature uniformity – low distortion level
Today a broad range of materials is heat treated in different processes in the vacuum
furnace (fig. 1). Due to its high flexibility and the above mentioned advantages the
vacuum furnace is in operation with big success at numerous sub-contracting services
and the tool manufacturers as well as in the automotive industry, the aircraft technology,
the medical technology and so on.
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IFHTSE 2005
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Figure 1. SCHMETZ vacuum furnace with overpressure gas quenching
The vacuum furnace is heated up through convection and radiation. The convectional
heating up in the lower temperature range serves for the faster and constant heating up
of the load in the hot zone. In the upper temperature range radiation heat transfer can
only be used for the heating up.
The exact control of the actual temperature in the hot zone through the heating
thermocouples and in the parts through the load thermocouples is one advantage of the
vacuum heat treatment. The load thermocouples enable the measuring of the part’s
temperature in its core and guarantee in this way the exact determination of the holding
time. The fully automatic process and the documentation of the heat treatment by means
of recorder (printer) guarantees reproducible results.
As to the quenching process of the heat treatment the following is demanded:
• hardenability of steel
• quenching as fast as necessary but also as slow as possible
• constant cooling down of the load
• temperature difference in the part as low as possible
The fulfilment of these demands makes possible the main target: a fully martensitic
hardening with the lowest possible distortion. The engineering of the furnace design has
a considerable influence on the constant cooling down process.
2. VACUUM FURNACE WITH INTEGRATED “SUB-ZERO” SYSTEM – SCHMETZ
SYSTEM *COOL PLUS*
The „sub-zero“treatment of steels at the hardening and tempering of tools is a well-
known and established process. In this process the tools‘ characteristics are improved
due to the specific optimizing of metallurgical structures like for example the reduction of
the retained austenite by means of a „sub-zero“ treatment. In practice so called cooling
machines, realizing a „sub-zero“ treatment in liquid nitrogen, are used. This conventional
technology is connected with a manual handling out of the hardening and tempering
furnaces into the cooling machines.
The integration of a „sub-zero“ system into the standard vacuum hardening furnace
realizes a man less, fully automatical hardening and tempering process. Due to the fact
that the parts do not have contact with the surrounding atmosphere during the „sub-
34

zero“ treatment and the following tempering process no surface corrosion can occur. In
this process the absolute metallic bright surface typical for the vacuum heat treatment is
guaranteed. A smoothly controlled „sub-zero“ treatment is also guaranteed by means of
the eligible „sub-zero“ gradient by using load thermocouples. Thus the risk of cracks due
to a „rough“ „sub-zero“ treatment like in the usual processes is restricted considerably.
The principle of the SCHMETZ system *COOL PLUS* is the inlet and the gasification of
liquid nitrogen through a nozzle system into the graphite insulated hot zone of the
vacuum furnace (fig 2). During the gasification the volume of nitrogen is increased by
700 times. In the load space the cold gaseous nitrogen is distributed constantly due to a
“circulator” (convection fan). During this process the load loses the heat energy. By
means of a gas outlet pipe the heat energy is drained off together with the „used“
nitrogen.
Fig 2 scheme: “sub-zero” system for vacuum furnace
The absolute reproducible, fully automatic process does not require any handling
between the single processes. Low staff costs, no risk of accidents (combustions caused
by splashes and vapour of liquid nitrogen) and no space need for additional cooling
machines are additional advantages besides the continuous documentation with a
complete load thermocouple recording (fig. 3).
Figure 3. Process documentation: hardening and “sub-zero” treatment down to –100°C
of a 420 kg load, material: 1.2379 = X 153CrVMo12= D2
35

Extensive examinations on stainless steels and tool steels, which were “sub-zero”
treated in a vacuum furnace were carried out.
3. „SUB-ZERO“ TREATMENT OF STAINLESS STEELS
In the hardening and tempering process of stainless steels the possibility of the increase
of the final hardness by means of the „sub-zero“ treatment between the hardening and
tempering process was examined. With an increase of the final hardness the chip
forming can be optimised and thus the production costs of the mechanical further
treatment, which follows the heat treatment, can be reduced.
In magazine loaded „combs“ of the material 1.4021 = X 20Cr 13 = 420 a load gross
weight of 100 kg was hardened in a vacuum furnace at 1035 °C (Fig 4). The „sub-zero“
treatment down to –100 °C is followed by a tempering process at 150 °C (fig. 5). In this
process the hardening and tempering process corresponds to the standard process.
Figure 4. Load: hardening and tempering process with “sub-zero” treatment of
stainless steels
Figure 5. Process cycle: hardening and tempering process with “sub-zero” treatment of
stainless steels
The final hardness of 640 HV after a standard hardening and tempering process could
be increased by an additional „sub-zero“ treatment to 690 HV. Similar results were
produced with “razor heads” of 1.4122 = X35CrMo17.
36

4. „SUB-ZERO“ TREATMENT OF TOOL STEELS
The aim of the „sub-zero“ treatment is the optimization of the dimensional stability due to
the reduction of the retained austenite content at the hardening and tempering process
of tool steels like 1.2379 = X 153 CrVMo 12= D2.
Test parts (diameter 180 mm, height 80 mm) with a load gross weight of 2 kg were
secondary hardened at 1060°C in a vacuum furnace with the dimension of the useful
space 600 x 900 x 600 mm (w x l x h) and „sub-zero“ treated at different temperatures
(0°C, -50 °C, -100 °C) (fig. 6). A part of the „sub-zero“ treated test parts with different
temperatures were afterwards tempered at 510 °C. With a „sub-zero“ treatment of –
100°C the retained austenite content can even be reduced to clearly below 10 % without
tempering (fig. 7).
Figure 6. Load: hardening and “sub-zero” treatment of 1.2379 = X 153CrVMo12= D2
Figure 7. Retained austenite at the hardening and tempering process with “sub-zero”
treatment of 1.2379 = X 153CrVMo12= D2
By means of the modular integration of the „sub-zero“ system into the standard vacuum
furnace the process time and costs of the hardening and tempering process of steels
37

can be minimized considerably, through reduction of the tempering process and a higher
dimensional stability of the parts if necessary.
5. CONCLUSION
Continuous further developments in the vacuum technology open in an ever broader
sector of materials and parts the advantages of a modern, environmental friendly, fully
automatic heat treatment cycle with a continuous documentation of the actual
temperature sequence of the part.
38

THE MAIN PRINCIPLES OF INTENSIVE QUENCHING OF TOOLS
AND DIES
N.I. Kobasko
Intensive Technologies Ltd, Kyiv, Ukraine
IQ Technologies Inc., Akron, USA
ABSTRACT
The paper discusses in details the main principles of intensive quenching process as applied to
steel tools, dies and other steel parts. The paper considers one- and two-step intensive
quenching process for tools and dies. When applying a one- step intensive quenching method,
the cooling is interrupted at the moment of time when the part surface compressive stresses
are at their maximum value and the core has not reached the marten site start temperature.
When applying a two-step intensive quenching technique, the duration of the first step of
cooling depends on the duration of the “self- regulated thermal process”. At the second step of
steel parts within the marten site range are cooled very rapidly. The paper presents also a new
method of calculation and optimisation of the process of quenching.
Key words: Intensive quenching, one-and two-step quenching, low hardenability steel,
optimisation, service life, cheap materials
1. INTRODUCTION
At present time the three main principles are used when developing intensive quenching of
steel parts. The first principle means that the reason of additional strengthening
(superstrengthening) of a material is high cooling rate within the martensite range. Detailed
information is published in [1] The second principle allows choosing conditions of cooling for
creation of the maximal compressive stresses at the surface of the quenched steel parts. It
means that very intensive cooling should be stopped at the moment of achievement of the
maximum compressive stresses at the surface [2] The third principle means, that the chemical
composition of steel should be such that after intensive cooling optimum depth of the
quenched layer could be formed [3]. Below are three examples, which were realized in the
practice.
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IFHTSE 2005
39

2. DESIGN OF INDUSTRIAL QUENCH PROCESSES
It is required to determine the speed of movement of the conveyor, which would provide
temperature 650 ºF (343 ºC) at the core of the part when it should be delivered from the
quenchant. To make these calculations we are using equation presented in [4, 5], i.e.
K
TT
TT
b
KnLaL
w
m
m






−
−
+Ω
==
0
ln
τ
; (1)
where: a is average thermal diffusivity of the material for the range of temperatures mTT −0 ;
Kn is Kondratjev number (dimensionless value), Ω =0.48 for cylinder-shaped bodies, b=1 if
the core’s temperature is determined, T0 is austenitizing temperature or temperature at the
time of immersing the part into the quenchant, Tm is temperature of the medium, if convection
prevails, or temperature of boiling if nucleate boiling prevails.
Table IKondratjev number Kn for 10% aqueous solutions of UCON A and UCON E at
temperature of 90 ºF (~32 ºC) and speed of the stream of 80 fpm (~0.4 m/s).
Temperature of the core of probes is 1300 ºF (704 ºС) [5]
Probe
diameter in
inches (mm)
UCON A UCON E Kn
0.5
(12.7)
0.424
0.417
0.412
0.408 0.415
1
(25.4)
0.546
0.526
0.488
0.488 0.512
1.5
(38.1)
0.578
0.556
0.523
0.514 0.543
The average thermal diffusivity a of the overcooled austenite within the temperature range of
1550 ºF – 200 ºF (840ºC – 100 ºC) is equal to 5.36·10-6
m2
/s. Having all above-stated facts we
calculate what speed of the conveyor should be to provide the core’s temperature 650 ºF (343
ºC) for a cylindrical part made of AISI 4140 steel and having diameter of 25 mm and height
of 50 mm. At the time of immersion the part has the same temperature of 1550 ºF through all
cross-sections. To answer this question, it is necessary just to determine Kondratjev number
Kn. For the determination of Kn there must be available an experimental database.
40

Figure 1. Industrial installation for the implementation of IQ-2 technology [4]
І– loading steel parts to conveyor for their heating in heater 1; II – chute with intensive
cooling devices; ІІІ – loading of quenchant to quenching tank with two conveyors; ІV–
unloading of steel parts from heater 2; TR1, ТR2, ТR3, ТR4, ТR5– speed control units for
conveyors 1, 2, 3, 4 and 5 operated by the control device; HT1, HT2– heaters 1 and 2; WQ1–
washing and quenching device; PM1, PM2– pumps 1 and 2; CL1, CL2– coolers 1 and 2; F1–
filter; BX1– container for quenched parts
3. CRITERION DETERMINING THE ABSENCE OF NON- STATIONARY NUCLEATE
BOILING
We can draw the criterion determining the absence of nucleate boiling at the surface of a part
to be quenched on the basis of the generalized dependence for the determination of the
duration of non-stationary nucleate boiling, i.e., self-regulated thermal process. As is already
known, the specified dependence has the following form:
a
K
b
II
I






+Ω=
ϑ
ϑ
τ ln (2)
In this formula the value of Ω determines the duration of irregular thermal process and is
quite a small value. The duration of the established non-stationary nucleate boiling is
determined basically by the second term of dependence (2), i.e
II
I
b
ϑ
ϑ
ln . To avoid nucleate
boiling, it is necessary that the second part of formula (2) is equal to zero, i.e., 0ln =
II
I
b
ϑ
ϑ
.
We have obtained equations for Iϑ and IIϑ , which can be presented as:
( ) 3.0
021





 −
=
R
I
I
ϑϑλ
β
ϑ (3)
and ( )[ ] 3.01
uhIIconvII ϑϑα
β
ϑ += . (4)
41

Equating Iϑ and IIϑ , we are obtaining the criterion for determining the absence of non-
stationary nucleate boiling:
( )
( )[ ] 3.0
3.0
02
uhIIconv
I
R
ϑϑα
ϑϑλ
+≡




 −
or
( )
uhI
I
Bi
ϑϑ
ϑϑ
+
−
= 02
, (5)
because in formula (2) III ϑϑ ≡ .
Equation (5) is the basic criterion that determines the absence of non-stationary nucleate
boiling (self-regulated thermal process) at steel quenching [6,7].
999999999999999999888888888888888888877777777777777777777777776666
999999999999999999999999888888888888888777777777777788888877777666
AAAA99999999999999999999999988888888888888877777778888888877777666
AAAAAAAAAAAA999999999999999999998888888888888888788888888877777666
AAAAAAAAAAAAAAAAAAAAA999999999999999999888888888888888888887777766
AAAAAAAAAAAAAAAAAAAAAAAAAAAAAAAAAAA9999999988888888888888887777666
BBBBBBBBABBBBBBBBBBAAAAAAAAAAAAAAAAAAA9999999999999999888887777666
BBBBBBBBBBBBBBBBBBBBBBBBBBBBAAAAAAAAAAAAA9999999999999988888777666
BBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBAAAAAAAAA99999999999998888777666
CCCCCCBBCCCCCCCCCCCCBBBBBBBBBBBBBBBBBAAAAAAAAA99999999998888777666
CCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBBAAAAAAAAA999999999888877766
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBBAAAAAAAAAAAAAA999888777766
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBAAAAAAAAAAAAAA999888877776
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBBAAAAAAAAAAAAA999888777777
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBAAAAAAAABBBAAAA9998877777777777
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBBBBBBBAABBBBBBAA99988877777777777
CCCCCCCCCCCCCCCCCCCCCCCCCCCCCBBBBBBBBBBBBBBBBBBBBBBBBBBBAA99988877777777777
CCBBBBBBBBBCCBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBCCCBBBBAA99888877778888888
BBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBABBBBBBCCCCBBBA9988888888888888888
BBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBBAAAAAAAAAABBCCCDCBBA998888888888888888888
AAAAAAAAABBBBBBBBBBBBBBBBBBBBBBBAAAAAAAABBBCCCCCCDB987777777788888888888888
AAAAAAAAAAAAAAABBBBBBBBBBBBBBBBAAAAABBBBBCCDCA87788777777778888888888888888
AAAAAAAAAAAAAAAAAAAAAAAAAAABBBBBBBBBBBCCBBA98777777777888888888888888888888
EDCCBBBBBBBBBBBAAAAAAAAAAABBBBAAAAAAA99998766667777778888888888888888888888
776AA9999999999998888888888888777777677777777777 8888888888888888
566666666666666666666677666666667777777777777777 8888888888888888
666677777777777777777777777777777777777777778888 8888888888888888
888888888888888
8888888888888
LEGEND
5 –1000 - -800
6 -800 - -600
7 -600 - -400
8 -400 - -200
9 -200 - 0
A 0 - 200
B 200 - 400
C 400 - 600
Figure 2. Hoop residual stress distribution on the cross section of the stamp after its partial self-
tempering and final cooling to room temperature (α=20000 W/m2
K)
4. QUENCH PROCESS OPTIMIZATION
The optimal residual stress distribution in the quenched steel part occurs in case of optimal depth
of the hardened layer. In this case high compressive stresses at the surface and less tensile
stresses in the core are observed. It is fair for any size of a part if the condition (6) is met:
.const
D
DI
opt
= (6)
Where: DI is the ideal critical diameter or specific size, Dopt is size of the steel part with the
optimal stress distribution. Ideal critical diameter can be calculated using equation (7):
5.0
ln 







+Ω
=
θ
τMba
DI , (7)
a is average thermal diffusivity (m2
/s);
42

mτ is limit time of the core cooling from the austenitizing temperature to martensite start temperature,
providing the formation of 99% or 50% martensite, 48.0=Ω for a bar (or cylinder),
mM
m
TT
TT
−
−
= 0
θ , T0 is austenitizing temperature, b is parameter depending only on form of steel
part; Tm is temperature of quenchant; TM is martensite start temperature at limit time of cooling.
Table II. Commercial and industrially tested technologies using intensive quenching [8]
Steel parts Steel for IQ
process
Steels and technologies which were
replaced by IQ
Cylindrical and conical gears of trucks
and tractors
58(55PP) 30KhGT (and other carburizing
steels, carburizing for 10 hr.
Cylindrical gears of electric driven train
transmissions and locomotives (m=10
mm)
ShKh4 20KhN3A, 20Kh2N4А, carburizing
for 30 hr.
Small modulated gears (m=4-6 mm) with
splined openings (solar, satellite ones)
58(55PP) 18KhGT and others, carburizing for
15 hr.
Rear wheel truck half-axles 47GT 40KhGRT and others, through
hardening in oil
Dies for punching the bearing bolls ShKh4 Enhanced alloyed steels, through
hardening in oil
Table III. Steels of low hardenability for Intensive Quenching (IQ) [8]
Steel,
GOST C Mn Si Ni Cr Mo Ti Al Cu
58(55PP)
GOST
1050
0.55-
-0.63
0.2 0.1-
-0.3
0.25 0.15 - - - 0.2
47GT 0.44-
-0.51
0.95-
-1.25
0.10-
-0.25
0.25 0.25 - 0.06-
-0.12
- 0.30
ShKh2 1.15-
-1.25
0.15-
-0.30
0.15-
-0.30
0.10 0.15 <0.03 0.06-
0.12
0.015-
0.03
0.12
ShKh4
GOST 801
0.95-
-1.05
0.15-
-0.30
0.15-
-0.30
0.30 0.25 - - - 0.25
45S 0.42-
0.48
0.17-0.32 0.40-0.65 0.20 0.25 - - - 0.15
70PP 0.66-
-0.73
0.15-
-0.30
0.15-
-0.30
0.25 0.25 - - - 0.25
115PP 1.10-
-1.20
0.40-
-0.60
0.15-
-0.30
0.20 0.25 - 0.06-
-0.12
- 0.20
43

CONCLUSIONS
1. Delayed transformation austenite into martensite at the first step of cooling and very rapid cooling
within the martensite range at the second step results in decreasing of distortion of steel parts and
increasing mechanical properties of a material.
2. Intensive cooling from austenite temperature till the time of the formation of the optimal
quenched layer and the maximal compressive stresses at the surface with the subsequent
tempering of the quenched layer also reduces distortion and increases mechanical properties
of the materials.
3. Low-hardenability steels, which provide optimal depth of the quenched layer in conditions of
intensive cooling, reduce distortion and increase service life of steel parts similarly to items 1
and 2.
4. The software and original technique of calculations of optimal conditions of quenching
depending on the shape and the sizes of parts, conditions of cooling and chemical
composition of steel has been developed.
5. More detailed calculations are carried out on the basis of the software TANDEM developed in
Ukraine.
6. At calculations of the optimal depth of quenched layer it is possible to use chemical
compositions of steels published in “Worldwide Guide to Equivalent Irons and Steels, 4th
Edition, (William C. Mack, Coordinating Editor), ASM International, 2002 [9].
References
1. N.I.Kobasko, Steel Superstrengthening Phenomenon, Journal of ASTM, February 2005,
Vol. 2, No 2, Paper ID JAI 12824, Available on line at www.astm.org
2. N.I.Kobasko, Intensive Steel Quenching Methods, In a Handbook: Theory and Technology of
Quenching, B. Liscic, H.M. Tensi, and W. Luty, Ed., Springer – Verlag, Berlin, 1992, p 367. – 389.
3. N.I.Kobasko, Quench Process Optimization, Proc.of the 6th
International Conf. on Heat
Treating of Materials (OTTOM-6), 16-20 May,2005, Kharkov, Ukraine
4. Ukrainian Patent No 27059
5. N.I.Kobasko, G.E.Totten, Design of Industrial Quenching Processes, Proc. of the 14th
IFHTSE Congress, Beijing, China, 2004.
6. Ukrainian Patent No 56189
7. US Patent #6,364,974B1
8. B.K.Ushakov, K.Z.Shepeliakovskii, V.M.Fedin, A.A.Kuznetsov, N.Yu.Kuznetsova,
Development of Through Surface Hardening Method for Heavy-Load Products and
Machine Parts, Steel No. 11, 2001, рр. 64.-68.
9. Worldwide Guide to Equivalent Irons and Steels, 4th Edition, (William C. Mack,
Coordinating Editor), ASM International, 2002.
44

OPTIMIZING THE VACUUM-HEAT-TREATMENT OF HOT-WORK
TOOL STEELS BY LINEAR ELASTIC FRACTURE MECHANICS
Vojteh Leskovšek1
, Borivoj Šuštaršič1
, Gorazd Jutriša1
, Dani Baksa2
, Janez Kopač3
1
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2
Unior Kovaška industrija d.d., Zreče, Slovenia
3
Faculty of Mechanical Engineering, Aškerčeva 6, 1000 Ljubljana, Slovenia
ABSTRACT
Linear elastic fracture mechanics was used to optimise the vacuum-heat-treatment procedures
for conventional hot-work AISI H11 tool steel. The fracture toughness was determined with
non-standard, circumferentially notched and fatigue-precracked tensile-test specimens. The
fracture-testing method is sensitive to changes caused by variations in the microstructure
resulting from the austenitizing and tempering temperatures as well as the homogeneity of the
material itself. The combined tempering diagram– Rockwell-C hardness, Fracture toughness
KIc, Tempering temperature– was used for the choice of the vacuum-heat-treatment
parameters necessary to obtain the best properties for a given application with respect to the
investigated steel.
Key words: conventional hot work tool steel, vacuum heat-treatment, fracture toughness,
hardness, microstructure
1. INTRODUCTION
The process parameters, the work material and the tool material determine the dominant
damage mechanism. For this reason, improving the tool’s performance requires a detailed
knowledge of the relevant damage mechanisms. It is also clear that the tool material itself
plays a very important role, and that the properties’ profile of the tool material greatly
influences its lifetime. Despite the enormous variety of tooling operations there are some
basic properties of tool materials that are common to almost all applications. These properties
are the toughness, which prevents instantaneous fracture of the tool or tool edges due to local
overload, and the hardness, which must be sufficiently high to avoid local plastic deformation.
st
IFHTSE 2005
45

Hardness and toughness are more or less mutually exclusive properties, which means the
prevention of instantaneous tool failures is often connected with a critical hardness level that
must not be exceeded for a specific application. The hardness and the toughness of hot-work
tool steel depend a lot on the vacuum-heat-treatment procedure. Hardness is closely related to
ductility and toughness, in particular the latter. In this paper the influence of the austenitizing
and tempering temperatures on the hardness and fracture toughness of conventional hot-work
AISI H11 tool steel is investigated and discussed.
2. THEORY
According to ref. [1] toughness and ductility are the most relevant properties in terms of
resistance to total failure as a result of overloading. Toughness and ductility are two different
material properties, even though both– unfortunately– are sometime denominated as
toughness. The opposite of both properties is, however, the same, i.e., brittleness. No
standardised tests for the determination of toughness or ductility are in common use; often,
data determined with different test methods are available, which makes them difficult to
compare, and this can lead to confusion. Toughness and ductility are different characteristics,
and for this reason it is necessary to distinguish between them [1]. Their importance for tool-
steel performance depends a lot on the geometry of the tool [1]. In the case of un-notched
specimens or specimens with smooth notches, the ductility and fracture stress are the relevant
material properties; however, if sharp notches or cracks are present, fracture toughness is the
most relevant property. The conclusion, therefore, is that the tool steel should be optimised in
terms of ductility and fracture stress for un-notched regions and in terms of fracture toughness
for notched regions. The toughness depends a lot on the hardness, and the hardening
mechanism is different in as-quenched and fully-heat-treated tool steels. In the as-quenched
tool work-hardening and solid-solution hardening, mostly due to carbon in the solid solution,
mainly affect the steel’s hardness. Tempering leads to the precipitation of carbide particles
and significantly decreases the carbon content in the solid solution and the dislocation density.
The hardness of fully-heat-treated tool steels is therefore mainly affected by precipitates that
cause precipitation hardening and, to small extent, solid-solution hardening. The work-
hardening and grain refinement seem to play only a minor role [2].
The most reliable measure of toughness is the plain-strain fracture toughness. The minimum
size of the specimens depends on the yield stress and the fracture toughness of the tested
material, both of which are necessary for a plane-strain deformation. A fatigue crack of a
46

defined length is propagated from a mechanical notch in the specimens ensuring that the
notch effect is a maximum and equal for all tests. The same value of fracture toughness
should be found for tests on specimens of the same material with different geometries and
with a critical combination of crack size and shape and fracture stress. Within certain limits,
this is indeed the case, and information about the fracture toughness obtained under standard
conditions can be used to predict failure for different combinations of stress and crack size
and for different geometries [3].
3. EXPERIMENTAL
3.1 Material and vacuum heat treatment
Conventional hot-work AISI H11 tool steel delivered in the shape of plates with dimensions
263 mm x 220 mm x 25 mm, cut from forged-and-soft-annealed master blocks with
dimensions 263 mm x 220 mm x 4000 mm and the following chemical composition (mass
content in %): 0.39 % C; 1.06 % Si; 0.32 % Mn; 0.019 % P; 0.004 % S; 4.91 % Cr; 0.11 %
Ni; 1.17 % Mo; 0.37 % V; and 0.011 % Ti was used. The KIc-test specimens,
circumferentially notched and fatigue-precracked tensile-test specimens, were cut from these
plates in the short transverse direction. A round notch with a fatigue crack at the notch root
was at the same distance (60 mm) from the surface of the master block in all the KIc-test
specimens. The specimens were heat treated in a horizontal vacuum furnace with uniform
high-pressure gas-quenching using nitrogen (N2) at a pressure of 1.05 bar. After the last
preheat (850 °C) the specimens were heated (10 °C/min) to the austenitizing temperatures
1000 °C, 1020 °C and 1050 °C, soaked for 20 minutes, gas quenched to a temperature of 100
°C. First temper was performed at 540 °C and second at different temperatures between 540
°C and 620 °C as shown in Fig. 2, each time for 2 hours, respectively. For each group of
vacuum-heat-treatment conditions from A to C, five KIc-test specimens were tested for each
second tempering temperature.
3.2 Hardness and fracture-toughness tests
The Rockwell-C hardness (HRc) was measured on the individual groups of the KIc-test
specimens using a Wilson 4JR hardness machine. Circumferentially notched and fatigue-
precracked tensile-test specimens with the dimensions indicated in Fig. 1 were used for this
investigation [4].
47

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