GIGACYCLE FATIGUE FRACTOGRAPHY OF COLD WORK TOOL STEELS PRODUCED BY PM COMPARED TO INGOT METALLURGY

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1 Powder Metallurgy Progress, Vol.8 (2008), No GIGACYCLE FATIGUE FRACTOGRAPHY OF COLD WORK TOOL STEELS PRODUCED BY PM COMPARED TO INGOT METALLURGY Ch. R. Sohar, A. Betzwar-Kotas, Ch. Gierl, B. Weiss, H. Danninger Abstract In this study the fatigue behaviour of cold work tool steels, wrought AISI D2 type (WNr , Böhler grade K110) and of powder metallurgical (PM) variant (Böhler grade K390) with high vanadium content, both acquired from Böhler Edelstahl GmbH, Austria, was studied up to the gigacycle regime employing an ultrasonic fatigue testing system that operates at 20 khz in fully-reversed tension-compression mode. Fracture surfaces were investigated by SEM. It was found that in both cases the S- N curves consistently dropped, i.e. there is no fatigue limit up to cycles. The production route showed to have significant effect on the fatigue behaviour and fracture surfaces. In the case of wrought cold work tool steel, primary alloy carbides and carbide clusters located at or just below the specimen surface were the nucleation sites of fatigue cracks. A few internal failures, due to large carbide clusters, were also obtained. At cycles, fatigue endurance strength of about 400 MPa was observed for wrought cold work steel. In contrast, the PM steel showed significantly higher fatigue endurance strength at loading cycles of about 700 MPa, and in most cases surface failure was observed, beside a few internal failures due to nonmetallic inclusions. The surface crack origins of PM cold work tool steel were mostly difficult to identify; in some cases, nonmetallic inclusions located just at the surface were the crack nucleation sites. Often, holes were observed at the crack origins, most probably deriving from broken-out inclusions, which are assumed to occur during the fracture process, since surface inspection prior to fatigue testing did not reveal any holes. Macroscopically, the fracture surfaces of both steels revealed rather flat morphology with somewhat finer fracture structure in case of the PM steel, and around each crack origin a (semi-) circular fish-eye feature was detected. However, while in wrought tool steel a granular area (GA) was observed around the crackinitiating carbides, such a zone was hardly found in the PM steel. Keywords: tool steels, gigacycle fatigue, carbide clusters, inclusions, ultrasonic frequency fatigue testing INTRODUCTION There are three main reasons for studying tool steels up to the very high cycle fatigue regime: First, it is aimed to produce fatigue data up to the gigacycle regime Christian R. Sohar, Christian Gierl, Herbert Danninger, Vienna University of Technology, Institute of Chemical Technologies and Analytics, Vienna, Austria Agnieszka Betzwar-Kotas, Brigitte Weiss, University of Vienna, Faculty of Physics, Nanostructured Materials, Vienna, Austria

2 Powder Metallurgy Progress, Vol.8 (2008), No contributing to the ongoing discussion about the shape of S-N curves and the question whether a real fatigue limit exists or not for such high strength steels. Furthermore, it turned out recently that fatigue testing at low amplitudes represents a reasonable tool to identify defects within a material, thus, supporting material development. And finally, in some applications of tool steels, such as in engine parts, and plastic moulding processes, but also for some tools, fatigue failure instead of wear represents the lifetime controlling factor. Generally, the fatigue behaviour of steels beyond 10 6 cycles became a focus of investigations after Naito et al. [1] had shown in 1984 that carburized and surface hardened steels do no exhibit a conventional fatigue limit at 10 6 to 10 7 cycles and fail even beyond 10 7 loading cycles at fairly low applied stresses. The present authors [2] gave a review of the existing literature on fatigue behaviour of high strength bearing and spring steels in a recently published article. Summarizing, the most important conclusion is that these high strength steels fail also at stress levels below the conventional fatigue limit, in the long life regime above 10 7 loading cycles. Usually failure of these steels in the very high cycle fatigue regime is associated with internal crack initiation and so-called fish-eye patterns on the fracture surface. The observed crack origins of internal failures were mostly nonmetallic inclusions, such as Al 2 O 3, TiN, SiO 2, MgO, and CaO or sulphides. Tool steels differ from high strength bearing and spring steels in such a way that tool steels contain numerous primary carbides required for high abrasion resistance, often in the same size or larger than the aforementioned crack initiating inclusions. Thus, the defects are much more frequent than for the materials discussed above. Investigations of fatigue behaviour of tool steels are scarce and mostly limited to maximum of 10 6 to 10 7 cycles [2-9]. The observed crack origins were found to be nonmetallic inclusions, but also primary carbides and carbide aggregates (carbide clusters), which were located in the interior or within the surface layer of the specimen. The study presented here aims at comparing the gigacycle fatigue behaviour up to N max = cycles of cold work tool steels produced by PM and ingot metallurgies, evaluating potential crack initiating defects of the materials and discussing probable differences. EXPERIMENTAL Two cold work tool steels have been studied, both acquired from Böhler Edelstahl GmbH, Austria, in cylindrical bars in annealed condition. The wrought AISI D2 type steel (DIN ) was conventional high Cr alloyed tool steel (1.55% C, 0.32% Si, 0.32% Mn, 12.7% Cr, 0.85% Mo, 0.89% V, 0.12% W, 0.12% Co) with a high amount of chromium carbides of type (Cr,Fe) 7 C 3, as presented in Fig.1a. The PM tool steel was a V-rich steel (2.45% C, 0.38% Si, 0.41% Mn, 4.8% Cr, 4.8% Mo, 11% V, 1.4% W, 2.1% Co), containing numerous fine spherical MC type vanadium carbides (Fig.1b), which were significantly smaller than the chromium carbides of steel K110. Hour glass shaped fatigue test specimens (diameter at gage length 4 mm, see also [2]) were machined and the surface longitudinally ground in the gage length prior to the applied heat treatment, for which parameters are presented in Table 1 together with the achieved Rockwell C hardness values. After the heat treatment, the fatigue specimens were longitudinally ground, and subsequently polished using 15 µm diamond suspension down to a depth of µm and 80 to 100 µm, for steel K110 and K390, respectively, in order to remove surface residual stresses [2]. The measured residual stresses after polishing were in the range of -170 to +40 MPa and -95 to -250 MPa for steel K110 and K390, respectively. Final mirror-like finish was achieved by polishing with 6 µm diamond suspension. Specimens were then fatigue tested in an ultrasonic frequency fatigue testing system operating at 20 khz and R = -1, as described in [2]. Due to the high stress amplitudes required for testing of such high strength materials, the samples have to be cooled during testing, which was performed by liquid cooling with noncorrosive coolant

3 Powder Metallurgy Progress, Vol.8 (2008), No (distilled water with corrosion inhibitor). In addition, the acoustic horn and sample-horn coupling site were cooled by compressed air in the case of steel K390. All obtained fracture surfaces were investigated by means of scanning electron microscopy (SEM). Tab.1. Heat treatment conditions and measured hardness and transverse rupture strengths. steel K110 (IM) K390 (PM) initial bar diameter / [mm] austenitizing temp. [ C] / time [min] quenching medium x / hrs tempering at temp. [ C] oil 3x560 T.R.S. [MPa] tempered steel 3600 ± ± 650 Rockwell C hardness tempered steel 58 ± 2 59 ± 2 Fig.1. As-tempered microstructure of steel K110 (cut perpendicular to the rolling direction) (a) and steel K390 (b). FATIGUE RESULTS S-N data Figure 2 presents the S-N curves obtained for the two cold work tool steels. Obviously, the S-N curves drop consistently, i.e. there is no fatigue limit. The fatigue endurance strength of PM steel K390 at Nmax = 1010 is nearly twice that of the conventional steel K110 over the entire cycle number range tested. For steel K110 predominantly primary carbides and carbide clusters (Fig.2Δ) located at/near to the surface gave rise to fatigue cracks. In addition some fractures due to internal carbide clusters (Fig.2{) were also observed. Steel K390 showed nearly exclusively fatigue failure due to surface defects such as nonmetallic inclusions ( ) or holes ( ) originating from decohesion of inclusions located at the specimens surface. However, in many cases it was impossible to identify the crack origins, a problem which was also reported by Marsoner et al. [6]. One specimen failed due to an internal Zr-oxide inclusion ( ). Fatigue endurance strength of steel K390 was about 700 to 800 MPa, and thus, about two times higher than the endurance strength obtained for steel K110, which was 400 MPa.

4 Powder Metallurgy Progress, Vol.8 (2008), No Fig.2. Fatigue data of the two studied cold work tool steels (figures close to the arrows represent the number of runout specimens). Interestingly, the fatigue data of the PM steel (K390) showed considerably more scatter than the data of steel K110, especially at lower amplitudes This can probably be attributed to the fact that for steel K110 the large primary carbides, of which a large number exists, were crack origins; however, fatigue failure of steel K390 was caused by material singularities such as inclusions located at or near the surface. Fractography Figure 3 presents representative images of macroscopic appearance of the fracture surfaces of the two cold work tool steels. Obviously, cracks and ridges point back to the crack origin, however, surface morphology of steel K390 is finer than that of K110. While in the fractograph of steel K110 a larger fatigue area can be detected, in the case of steel K390 this area seems to be smaller and a clear definition of the area was not possible. Fig.3. Comparison of macroscopic appearance of fracture surfaces obtained for (a) conventional cast tool steel K110 and (b) PM steel K390.

5 Powder Metallurgy Progress, Vol.8 (2008), No K110 specimens failed predominantly from carbides or carbide clusters located at or near the surface as shown in Fig.4a, b. The large crack-nucleating primary carbides in steel K110 were surrounded by an area revealing granular surface morphology, referred to as granular area (GA). This area was followed by a rather smooth zone (Fig.4a - area 2a ), probably a result of fatigue crack growth perpendicular to the loading direction. Such a zone was also observed in fractographs of steel K390 (Fig.4c); however a granular area cannot be detected clearly. Figure 4c and d presents representative images of crack initiation by decohesion of an inclusion, which was located at the specimen surface, resulting in a hole. Fractures of steel K390, for which determination of the exact crack origin was impossible, exhibited an appearance rather similar to the one presented in Fig.4c,d, however without significant microstructural features. Fig.4. Crack initiation sites of failed specimens of steel (a), (b) K110 and of PM steel K390 (c), (d). CONCLUSION The presented fatigue study up to the gigacycle fatigue regime has shown that neither for wrought nor for PM tool steels a fatigue limit exists, at least up to 1010 cycles. The manufacturing route has a decisive effect on the fatigue strength of cold work tool steels: Large primary carbides or carbide clusters located at or near the surface or carbide clusters in the interior gave rise to fatigue cracks in the conventional ingot metallurgy tool steels. The PM grade revealed significantly higher fatigue strength over the entire tested

6 Powder Metallurgy Progress, Vol.8 (2008), No cycle number range, which can be attributed to the fact that the primary carbides were too small to initiate cracks. Thus, in contrast to conventional tool steel, in PM tool steel the fatigue cracks nucleated at singularity defects such as nonmetallic inclusions at the surface or holes stemming from decohesion of nonmetallic inclusions, or at other near-surface defects that could not be clearly identified. Fatigue fractography appeared rather similar for the two steels studied. REFERENCES [1] Naito, T., Ueda, H., Kikuchi, M.: Met. Trans. A, vol. 15A, 1984, p [2] Sohar, CR., Betzwar-Kotas, A., Gierl, C., Weiss, B., Danninger, H.: Int. J. Fatigue, vol. 30, 2008, p [3] Berns, H., Trojahn, W.: VDI-Z, vol. 127, 1985, p. 889 [4] Berns, H., Lueg, J., Trojahn, W., Wähling, R., Wisell, H.: Powder Metall. Int., vol. 19, 1987, p. 22 [5] Meurling, F., Melander, A., Tidesten, M., Westin, L.: Int. J. Fatigue, vol. 23, 2001, p. 215 [6] Marsoner, S., Ebner, R., Liebfahrt, W., Jeglitsch, F.: HTM, vol. 57, 2002, p. 283 [7] Marsoner, S., Ebner, R., Liebfahrt, W.: BHM, vol. 148, 2003, p. 176 [8] Fukaura, K., Yokoyama, Y., Yokoi, D., Tsujii, N., Ono, K.: Met.Mat.Trans.A, vol. 35A, 2004, p [9] Sohar, CR., Betzwar-Kotas, A., Gierl, C., Weiss, B., Danninger, H.: Int. J. Fatigue doi: /j.ijfatigue