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 resistance. Several heat treatment cycles were investigated, by carrying out DCT both after quenching, after tempering and between quenching and tempering. Deep cryogenic treatment always reduces wear rate of the two steels, even if it does not influence significantly hardness. The effect is more pronounced 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, confirming 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-

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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, Mf lies above –84°F, so that the almost complete elimination of retained 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 subsequent 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 constitute 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 resistance and tempering transformations. Wear tests were carried out on a laboratory apparatus, on specimens treated in different conditions, by combining quenching, tempering, stress relieving and DCT in several modes, to optimize the treatment schedule. To study the effect of DCT on tempering 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 Table 1. Specimens were heat treated according to the cycles indicated in Tables 2 and 3.

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Deep Cryogenic Treatment of Cold Work Tool Steel Table 1. Nominal composition of the steel (wt%)

DIN X155CrMoV12 1 X110CrMoV8 2

C

Si

Mn

Cr

Mo

V

other

1.55 1.10

0.30 0.90

0.30 0.40

11.5 8.30

0.7 2.10

1.0 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 B C D E

Q + T1 + T2 Q + T1 + C + T2 Q + C + T1 + T2 Q + C + T1 + D Q + C∗ + T1 + D

Table 3. Heat treatment parameters for the studied steels

Treatment Q T1 T2 D

X155CrMoV12 1 980 ◦C×35min 500 ◦C×3h 510 ◦C×3h 240 ◦C×3h

X110CrMoV8 2 1040 ◦C×1.5h 500 ◦C×3h 540 ◦C×3h 300 ◦C×3h

After quenching (Q), part of the samples were twice tempered (T1 +T2 ) using the parameters reported in Table 3 (treatment A). The remaining were cryogenically treated (C), prior to double tempering (T 1 +T2 ) (treatment C) or tempering plus stress relieving (T1 +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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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 ◦Cand 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 ◦Cwith 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. Samples of 4 × 4 × 10 mm3 were heated from 25 ◦Cto 850 ◦Cat a constant rate of 10 ◦C/min. Argon was supplied in order to avoid oxidation. The dilatometric 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 tribometer. 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 sliding 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 interpolating the experimental points of the wear curve (cumulative mass loss of the disc vs sliding distance).

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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 attributed to the transformation of retained austenite on DCT. In fact, retained austenite decreases hardness of quenched microstructure and is responsible for secondary hardening transformations in this steel: at about 500 ◦C, submicroscopic chromium and molybdenum carbides precipitate from retained 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 temperature range, up to 500 ◦C. Secondary hardening peak does not disappear in DC tretaed steel. Again, retained austenite transformation on DCT may be

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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 carbide precipitation from martensite at 490 ◦Cand not to retained austenite, as in the previous steel. The effect of DCT on hardness of tempered steel is then correlated to retained austenite transformation and to its possible role on secondary hardening. 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 suppressed. 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 ◦Cto 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 martensite, 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

Figure 3.

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DSC curves of X110CrMoV8 2 after quenching (Q).

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

Figure 4. (Q).

Relative length change and its derivative of X110CrMoV8 2 after quenching

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

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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 cementite 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 clustering, giving rise to the precipitation of transition carbides. DCT is known to enhance tempering transformations of martensite, activating 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 redistribution through segregation and clustering at 30 ◦C–150 ◦Cin 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 ◦Cin the DSC and at 450 ◦Cin the dilatometric curve. The lower heat flow measured in DCT material than in as quenched one might be a direct consequence of the preliminary 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

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Figure 5. (Q+C).

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

Figure 6. Relative length change and its derivative of X110CrMoV8 2 after quenching (Q) and quenching plus DCT(Q+C).

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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 untransformed 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 carbide 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 distinct 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 precipitation 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].

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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+T1 +T2 . Table 4. Values of the wear   rate (WR) and the relative increment 0) (∆WR %= (WR−WR × 100) (WR0 is referred to the treatment variant Q+T1 +T2 ) WR0 X155 CrMo 12 1 WR (g/m ×10−6 ) ∆WR%

Treatment Q + T1 + T2 Q + T1 + C + T2 Q + C + T1 + T2 Q + C + T1 + D Q + C∗ + T 1 + D

10.3 9.46 9.77 7.83 7.69

— −8.2 −5.2 −24.0 −25.3

X110 CrMoV 8 2 WR (g/m×10−6 ) ∆WR% 9.10 8.18 7.56 5.88 5.84

— −19.2 −25.3 −41.9 −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. tions

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

Treatment Q + T1 + T2 Q + T1 + C + T2 Q + C + T1 + T2 Q + C + T1 + D Q + C∗ + T1 + D

X155 CrMoV 12 1

X110 CrMoV 8 2

60,8 61,4 61,3 61,8 61,0

61,4 61,2 61,5 61,4 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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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 (T2 ) 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 tempering with a stress relieving has to be ascribed at the better microstructural 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 tempering. 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 occurs 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. Contrarily, 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

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by DSC and dilatometry. The combination of the two experimental techniques allows mechanisms of martensite destabilization to be individuated. DCT favours carbon clustering in martensite and transition carbide precipitation, 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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