Deep Cryogenic Treatment of Cold Work Tool Steel

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, 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, 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 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 resistance 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 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 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+T1 +T2 B Q+T1 +C+T2 C Q+C+T1 +T2 D Q+C+T1 +D ∗ E Q+C +T1 +D Table 3. Heat treatment parameters for the studied steels Treatment X155CrMoV12 1 X110CrMoV8 2 Q 980 ◦C×35min 1040 ◦C×1.5h ◦ ◦ T1 500 C×3h 500 C×3h ◦ ◦ T2 510 C×3h 540 C×3h D 240 ◦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 treatmentC∗ 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. 660 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 ◦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. Sam- ples 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 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 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 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 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 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 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 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 ◦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 663 Figure 3.

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