DEEP CRYOGENIC TREATMENT OF COLD WORK
TOOL STEEL

M. Pellizzari and A. Molinari

Department of Materials Engineering

University of Trento

Via Mesiano 77

38050 Trento

Italy

Abstract

Deep Cryogenic Treatment (DCT) was applied to two different cold work
tool steels, X155CrMoV121 and X110CrMoV82, to improve their wear re-
sistance. Several heat treatment cycles were investigated, by carrying out
DCT both after quenching, after tempering and between quenching and tem-
pering. Deep cryogenic treatment always reduces wear rate of the two steels,
even if it does not influence significantly hardness. The effect is more pro-
nounced when cryogenic cycle is carried out immediately after quenching.
The influence of DCT on tempering curves was also investigated and the
effect on retained autenite transformation was highlighten. By means of
DSC and dilatometry, tempering transformations were investigated, confirm-
ing that DCT mainly enhances destabilization of martensite, by activating
carbon clustering and transition carbide precipitation.

Keywords:

Deep Cryogenic Treatment (DCT), wear, tempering, secondary hardening

INTRODUCTION

Several papers have been published about the benefits arising from deep

cryogenic treatment at –196

◦

C (DCT) on the properties of tool steels. In

particular, depending on the material, wear resistance may increase up to
100%. [1, 2, 3, 4, 5].

The first hypothesis, based on the transformation of retained austenite

on DCT, was denied by Carlsson which demonstrated that the benefits in-

657

658

6TH INTERNATIONAL TOOLING CONFERENCE

troduced by this treatment are much higher than those provided by a cold
treatment (CT) carried out at higher temperature (–84

°F) [6]. For most steels,

in fact, M

f

lies above –84

°F, so that the almost complete elimination of re-

tained austenite is induced still by a cold treatment. Furthermore, a certain
amount of retained austenite was found to be positive for wear resistance,
because of its influence on toughness of steel [7].

A new mechanism, involving the second microstructural constituent in

the quenched steel, i.e. martensite, was proposed by Meng et. al. [8] to
explain the improvement of properties of DC treated steels. During soaking
in LN, carbon clustering in virgin martensite is activated, which promotes
the formation of a more homogeneous and fine carbide precipitation on sub-
sequent tempering. This effect was called "low temperature conditioning"
of virgin martensite.

Papandopulo et. al. ascribed this conditioning to the change in martensite

lattice parameter and to an increase in the lattice defects density, which con-
stitute preferential nucleation sites for carbide precipitation during reheating
to room temperature and tempering [9]. The authors of the present paper
also find a high density of dislocations after deep cryogenic of AISI H13 hot
work tool steel [10]. The same effect was observed by other authors, which
found a finer and more homogeneous carbide precipitation in different steels
after DCT [1, 2, 3, 4, 5, 6, 6, 7, 8, 9, 10, 11, 12].

In this work DCT was applied to two cold work tool steels, namely

X155CrMoV12 and X110CrMoV8, to investigate its effect on wear resis-
tance and tempering transformations. Wear tests were carried out on a lab-
oratory apparatus, on specimens treated in different conditions, by combin-
ing quenching, tempering, stress relieving and DCT in several modes, to
optimize the treatment schedule. To study the effect of DCT on temper-
ing transformations, differential scanning calorimetry and dilatometry were
used [13, 14, 15].

EXPERIMENTAL PROCEDURES

MATERIALS AND HEAT TREATMENTS

The nominal composition of the investigated materials is reported in Ta-

ble 1.

Specimens were heat treated according to the cycles indicated in Tables 2

and 3.

Deep Cryogenic Treatment of Cold Work Tool Steel

659

Table 1.

Nominal composition of the steel (wt%)

DIN

C

Si

Mn

Cr

Mo

V

other

X155CrMoV12 1

1.55

0.30

0.30

11.5

0.7

1.0

X110CrMoV8 2

1.10

0.90

0.40

8.30

2.10

0.50

+Al, Nb

Table 2.

Heat treatment variants investigated (Q=quenching, T=tempering, C

∗

=controlled

DCT for 35h, C=direct immersion in liquid nitrogen for 14h)

Code

Treatment

A

Q + T

1

+ T

2

B

Q + T

1

+ C + T

2

C

Q + C + T

1

+ T

2

D

Q + C + T

1

+ D

E

Q + C

∗

+ T

1

+ D

Table 3.

Heat treatment parameters for the studied steels

Treatment

X155CrMoV12 1

X110CrMoV8 2

Q

980

◦

C ×35min

1040

◦

C ×1.5h

T

1

500

◦

C ×3h

500

◦

C ×3h

T

2

510

◦

C ×3h

540

◦

C ×3h

D

240

◦

C ×3h

300

◦

C ×3h

After quenching (Q), part of the samples were twice tempered (T

1

+T

2

)

using the parameters reported in Table 3 (treatment A). The remaining were
cryogenically treated (C), prior to double tempering (T

1

+T

2

) (treatment

C) or tempering plus stress relieving (T

1

+D) (treatment D). In one case

DCT was carried out between two tempering stages (treatment B). During
treatment C samples were directly immersed in liquid nitrogen for a soaking
time of 14 hours. Finally a controlled subzero treatment, i.e. C

∗

, was

realized in treatment D. During the deep cryogenic treatment C

∗

the samples

were cooled down to liquid nitrogen temperature in an industrial plant, with
controlled cooling rate (30

◦

C /min). Soaking time was 35 hours.

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6TH INTERNATIONAL TOOLING CONFERENCE

TEMPERING CURVES

Tempering curves were obtained by treating as quenched and quenched

and DC treated 12 × 8 × 5 mm specimens, 1 hour at different temperatures
(250

◦

C , 350

◦

C , 450

◦

C , 520

◦

C and 580

◦

C ), and measuring HRc hardness.

DIFFERENTIAL SCANNING CALORIMETRY

Differential scanning calorimetry was carried out using a Perkin Elmer

DSC7 apparatus in Argon atmosphere in order to prevent oxidation. The
mass of the specimen between 30 and 60 mg. Samples were heated from
room temperature to 725

◦

C with a heating rate of 5

◦

C /min and subsequently

cooled down to 25

◦

C . The baseline was obtained by a second scanning in

the same experimental conditions. The subtracted signal was considered to
be representative of the irreversible transformation occurring on tempering.

DILATOMETRY

Dilatometry was carried out using a Linseis horizontal dilatometer. Sam-

ples of 4 × 4 × 10 mm

3

were heated from 25

◦

C to 850

◦

C at a constant rate

of 10

◦

C /min. Argon was supplied in order to avoid oxidation. The dilato-

metric results provide a measurement of length change of the specimen
during heating. In order to determine the transformation temperature of
each phase transformation, data were arithmetically averaged over time and
subsequently differentiated with respect to temperature [13].

TRIBOLOGICAL TEST

Block on disc dry sliding wear tests were carried out using an Amsler tri-

bometer. As counterpart material a block of X210Cr12 hardened to 61.3 HRc
was selected. A load of 150 N and a sliding speed of 0.21 m/s for a total slid-
ing distance of 5000 metres were set up. The sample disc (40 mm external
diameter, 10 mm track width) was periodically weighted using a precision
balance (10

−4

g). Wear rate was determined as the slope of the line inter-

polating the experimental points of the wear curve (cumulative mass loss of
the disc vs sliding distance).

Deep Cryogenic Treatment of Cold Work Tool Steel

661

RESULTS

EFFECT OF DCT ON TEMPERING

TRANSFORMATIONS

Tempering curves.

The tempering curves of as quenched and quenched

and DC treated X155CrMoV121 are reported in Fig. 1. Hardness of DC

Figure 1.

Effect of DCT on the tempering curve of X155CrMoV12 1.

treated material is 1 to 3 HRc points higher than that of as quenched one
and DCT suppresses secondary hardening peak. Both effects can be at-
tributed to the transformation of retained austenite on DCT. In fact, retained
austenite decreases hardness of quenched microstructure and is responsi-
ble for secondary hardening transformations in this steel: at about 500

◦

C ,

submicroscopic chromium and molybdenum carbides precipitate from re-
tained austenite, decreasing its content of carbon and alloying elements and
favouring its transformation in martensite on subsequent cooling down to
room temperature (V step of tempering).

The two tempering curves of X110CrMoV82, in as quenched and quenched

and DC treated conditions, are reported in Fig. 2. Curves are very similar,
the one effect of DCT is to slightly increase hardness in the whole temper-
ature range, up to 500

◦

C . Secondary hardening peak does not disappear in

DC tretaed steel. Again, retained austenite transformation on DCT may be

662

6TH INTERNATIONAL TOOLING CONFERENCE

Figure 2.

Effect of DCT on the tempering curve of X110CrMoV8 2.

responsible for the increase in hardness (the very low amount of retained
austenite, due to the low austenitization temperature, takes account for the
poor increase in hardness after DCT). Secondary hardening peak does not
disappear after DCT since in this steel this phenomenon is due to alloy car-
bide precipitation from martensite at 490

◦

C and not to retained austenite, as

in the previous steel.

The effect of DCT on hardness of tempered steel is then correlated to re-

tained austenite transformation and to its possible role on secondary harden-
ing. Depending on the retained austenite amount after quenching, hardness
of DC treated steel increases with respect to as quenched steel. At the same
time, secondary hardening due to retained austenite transformation is sup-
pressed. DCT does not influence secondary hardening due to precipitation
of carbide form martensite.

Differential scanning calorimetry and dilatometry.

Figure 3 shows

DSC diagram of as quenched X110CrMo82 in the range 25

◦

C to 700

◦

C . It

displays three peaks, partially overlapped. Two of them (I and II) are well
pronounced, whilst the third one (III) is quite small. The peak at the highest
temperature can be attributed to the alloy carbide precipitation from marten-
site, which is responsible for secondary hardening. The other two peaks may
belong to cementite precipitation and to transformation of retained austen-

Deep Cryogenic Treatment of Cold Work Tool Steel

663

Figure 3.

DSC curves of X110CrMoV8 2 after quenching (Q).

ite. Figure 4 shows the dilatometry diagram of the same material. The first

Figure 4.

Relative length change and its derivative of X110CrMoV8 2 after quenching

(Q).

derivative of the relative length change shows again the three effects. They

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6TH INTERNATIONAL TOOLING CONFERENCE

appear at higher temperatures than in DSC because of the higher heating
rate used: 10

◦

C /min in dilatometric measurements and 5

◦

C /min in DSC.

In addition to the alloy carbide precipitation at high temperature, which
occurs with a dimensional contraction (III), the precipitation of cementite
from martensite (I) and the decomposition of retained austenite (II) can be
resolved thanks to their opposite effect on dimensional change, since cemen-
tite precipitation occurs with a volume decrease and the retained austenite
transformation occurs with a volume increase. The combination of DSC and
Dilatometry allows tempering transformations to be highligthen.

Figures 5 and 6 compare DSC and dilatometry of X110CrMo82 in the as

quenched (same as Fig. 3) and quenched and DC treated condition.

The

main differences are:

a broad effect appears at low temperature after DCT, not observed
in the as-quenched steel. This can be attributed to precipitation of
transition carbides from martensite. The freshly formed martensite,
in fact, is a metastable constituent at room temperature, which results
supersaturated in carbon. The atoms in the tetragonal interstices, if
activated, are thus subjected to redistribution by segregation and clus-
tering, giving rise to the precipitation of transition carbides. DCT
is known to enhance tempering transformations of martensite, acti-
vating carbon clustering at very low temperature and the subsequent
precipitation of transition carbides (I step of tempering) [8, 11, 16].
The preliminary tempering stage, which involves carbon redistribu-
tion through segregation and clustering at 30

◦

C –150

◦

C in as quenched

martensite, is then activated by DCT

a weaker effect due to cementite precipitation in DCT treated steel than
in as quenched material. The peak for precipitation of cementite from
transition carbides starts at about 350

◦

C in the DSC and at 450

◦

C in

the dilatometric curve. The lower heat flow measured in DCT material
than in as quenched one might be a direct consequence of the prelim-
inary transition carbide precipitation. It may be reasonably assumed
that the sum of the two effects in DCT material has an enthalpy content
comparable to that of the only cementite precipitation in as quenched
one. The energy release during the transformation of virgin martensite
into ferrite and cementite, depends on the metastability of martensite
which, in turn, depends on carbon content. Energy release may occur

Deep Cryogenic Treatment of Cold Work Tool Steel

665

Figure 5.

DSC curves of X110CrMoV8 2 after quenching (Q) and quenching plus DCT

(Q+C).

Figure 6.

Relative length change and its derivative of X110CrMoV8 2 after quenching

(Q) and quenching plus DCT(Q+C).

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6TH INTERNATIONAL TOOLING CONFERENCE

either in one step only (as quenched material) or may be shared into
two steps in DCT material, when transition carbides precipitation is
activated;

the peak due to retained austenite transformation is more evident in
DC treated material than in as quenched one, despite the lower content
of this constituent. This may be an apparent contradiction. We have
to consider that DCT does not fully transform retained austenite into
martensite, some mechanically stabilized austenite remaining untrans-
formed in the interlath regions. Therefore some retained austenite is
still present after DCT, which transforms on tempering. The decrease
of intensity of peak due to cementite precipitation, which is partially
overlapped to that of retained austenite decomposition, may allow the
effect of this last transformation to be better detected on DSC diagram.
This interpretation is supplied by dilatometry diagram of DCT steel,
which shows a smaller contribution due to cementite precipitation
(some volume contraction has already occurred during transition car-
bide precipitation) and a smaller contribution due to retained austenite
at about 610

◦

C , Fig. 6.

the peak of secondary carbides precipitation is higher in DC treated
steel. In both tratement variants considered in Fig. 5 it overlaps that of
cementite precipitation and therefore is not possible to discern a dis-
tinct peak for this tranformation. The higher heat evolution observed
in the cryo treated steel above 600

◦

C / corresponds only apparently to

a lower linear contraction at about 750

◦

C /. The higher contraction

observed in the untreated steel, in fact, mainly arises from the precipi-
tation of cementite rather than of secondary carbides. Thus the results
of differential scanning calorimetry and dilatometry are congruent,
i.e., the high temperature peak given by alloy carbide precipitation is
almost unchengad by DCT.

On the basis of these experiments and of their interpretation, which needs

for further investigation, it may be confirmed that DCT partially reduces
the amount of retained austenite and enhances martensite decomposition on
tempering, producing a more homogeneous precipitation of smaller carbides
[11, 17].

Deep Cryogenic Treatment of Cold Work Tool Steel

667

EFFECT OF DCT ON WEAR RESISTANCE

In Table 4 the values of the wear rate calculated for the two steels in the

different heat treatment states are reported with the relative increment with
respect to the conventional heat treatment Q+T

1

+T

2

.

Table 4.

Values of the wear rate (WR) and the relative increment

(

∆WR %=



(WR−WR

0

)

WR

0



× 100) (WR

0

is referred to the treatment variant Q+T

1

+T

2

)

Treatment

X155 CrMo 12 1

X110 CrMoV 8 2

WR (g/m ×10

−6

)

∆WR%

WR (g/m×10

−6

)

∆WR%

Q + T

1

+ T

2

10.3

—

9.10

—

Q + T

1

+ C + T

2

9.46

−8.2

8.18

−19.2

Q + C + T

1

+ T

2

9.77

−5.2

7.56

−25.3

Q + C + T

1

+ D

7.83

−24.0

5.88

−41.9

Q + C

∗

+ T

1

+ D

7.69

−25.3

5.84

−42.4

Results show that no correlation exists between hardness and wear rate,

Table 5. Data shows that DCT exerts just a little influence on hardness after

Table 5.

hardness values (HRc) for the studied steels in the different heat treatment condi-

tions

Treatment

X155 CrMoV 12 1

X110 CrMoV 8 2

Q + T

1

+ T

2

60,8

61,4

Q + T

1

+ C + T

2

61,4

61,2

Q + C + T

1

+ T

2

61,3

61,5

Q + C + T

1

+ D

61,8

61,4

Q + C

∗

+ T

1

+ D

61,0

61,3

high temperature tempering. In any case the differences measured are in the
range of the experimental error. Therefore, the increase in wear resistance
has to be attributed to the improved microstructure, in terms of distribution
and size of secondary carbides, as reported by other authors [11, 16, 17].

DCT reduces the wear rate of X155CrMo121 and X110CrMoV82. The

best results are obtained for the latter steel whose secondary hardening is
due to alloy carbide precipitation from martensite. Among the investigated

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6TH INTERNATIONAL TOOLING CONFERENCE

treatment variants, that with DCT prior to tempering is the most effective.
Data in Table 4 evidence that the improvement in wear resistance is further
enhanced if the second tempering (T

2

) is replaced by a stress relieving (D)

carried out at lower temperature. Very limited wear resistance improvement
occurs when DCT is carried out after a preliminary stress relieving. This
procedure is frequently applied on treating large parts, in order to release
quenching stresses before subzero cooling. This is aimed at preventing
distorsion or even cracking during subzero cooling. This problem is of
major importance in case of direct immersion of dies with complex shapes
into liquid nitrogen (treatment C), whereas it has been proved to be very
limited by controlled DCT (treatment C

∗

). Hence, despite of the good results

concerning the measured wear rate achieved with treatment C, specific field
test should be made in order to assess the real entity of distorsion or cracking.
The negative effect of a stress relieving prior to DCT can be attributed to
the transformations occurring on stress relieving, which activates the early
stages of martensite decomposition. After stress relieving, DCT works on a
partially destabilized martensite, and therefore its effect is poor. Obviously
DCT works even less when carried out after tempering.

The highest reduction in wear rate obtained by replacing the second tem-

pering with a stress relieving has to be ascribed at the better microstruc-
tural characteristics of the cryogenically treated martensite still after the first
tempering stage. A second tempering promotes carbide coalescence thus
reducing the benefits introduced by DCT on carbide size and distribution.

CONCLUSIONS

Deep Cryogenic Treatment increases wear resistance of cold work tool

steel, in particular when carried out just after quenching, prior to temper-
ing. Additional benefits are obtained by replacing the second tempering
step by a stress relieving at lower temperature. The maximum effect has
been observed in X110CrMoV82 steel, in which secondary hardening oc-
curs through precipitation of carbides rather than through the decomposition
of retained austenite, like in X155CrMoV121.

DCT does not affect hardness of high temperature tempered steels. Con-

trarily, it slightly increases hardness of stress relieved steels and tempering
resistance up to secondary hardening temperature. This arises from the
transformation of retained austenite on DCT and from the enhancement of
carbide precipitation in martensite. These transformations were investigated

Deep Cryogenic Treatment of Cold Work Tool Steel

669

by DSC and dilatometry. The combination of the two experimental tech-
niques allows mechanisms of martensite destabilization to be individuated.
DCT favours carbon clustering in martensite and transition carbide precip-
itation, and this results in a more homogeneous and fine precipitation of
carbides, after tempering. The microstructural improvement of tempered
martensite is responsible for the increased wear resistance.

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