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Improved vacuum heat-treatment for fine-blanking tools from high-speed steel M2
Journal of Materials Processing Technology 82 (1998) 89 – 94
Improved vacuum heat-treatment for fine-blanking tools from high-speed steel M2 V. Leskovs' ek a,*, B. Ule b a
Vacuum Heat Treatment and Surface Engineering Center, Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slo6enia b Uni6ersity of Ljubljana, Faculty of Natural Science and Engineering, Department of Materials and Metallurgy, Aker e6a 12, 1000 Ljubljana, Slo6enia Received 10 March 1997
Abstract The influence of the vacuum heat-treatment process on the hardness, fracture toughness and fine-blanking performances of tools made from high speed steel M2 is investigated. To improve the fine-blanking performances, the fracture toughness and hardness of the AISI M2 high-speed steel was measured after a variety of vacuum heat treatments. © 1998 Published by Elsevier Science S.A. All rights reserved. Keywords: Fine blanking tool; Fracture toughness; Hardness; High speed steel M2; Vacuum heat treatment
1. Introduction Due to their high wear resistance, high speed steels are nowadays used also for fine-blanking tools, especially in case of long series. High-speed steels must withstand compressive stresses and abrasive or adhesive wear, and have a sufficient fracture toughness to resist chipping and cracking. Therefore, vacuum heat-treatment with uniform high-pressure gas quenching is one of the most important operations in the manufacturing of fine blanking tools from high speed steel AISI M2. The importance of optimal vacuum heat-treatment is in the possibility of obtaining an optimum combination of basic characteristic of tools of high speed steels, and of the given working part – fine blanking-tool combination. The fracture toughness and hardness of AISI M2 high-speed steel was measured after a variety of heat treatments. It was found that it is a sensitive function of heat treatment, depending on both the austenitizing and the tempering temperatures. The fracture toughness exhibited an inverse correlation with hardness. The results of this investigation are useful for the optimization of the vacuum heat-treatments of different * Corresponding author. Tel.: +386 61 213780; fax. + 386 61 213780.
tools from high-speed steels, which support the exploitation of tensile impact loads that require the optimal combination of hardness and fracture toughness. 2. Basic characteristics of high-speed tool steel M2 Fine-blanking tool cutting speeds have constantly increased, partly because of the use of new tool materials. However, the lack of toughness has prevented the use of some new materials in certain cutting applications in which enhanced toughness is required. Fractures, as well as macro- and micro-chipping, can destroy the cutting edges. The ability of a steel to resist these phenomena is known as toughness. The toughness that can be achieved by high-speed steel is limited by the microstructural features (carbide size and segregation, strings of inclusions, etc.). When the steel is loaded, stress concentrations appear because of the tool shape and microstructural constituents and cause tool fracture, unless the stress concentrations are relieved by local plastic flow on a micro scale. The ability of the matrix to undergo plastic flow can be altered within wide limits by varying the hardness. Thus, the characteristic microstructural features of steel determine the limits of the toughness levels that can be achieved by vacuum heat treatment.
0924-0136/98/$19.00 © 1998 Published by Elsevier Science S.A. All rights reserved. PII S0924-0136(98)00023-5
V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
90
Fig. 1. Comparative hardness of carbides found in tool steels [2].
To improve the toughness of conventionally-manufactured high-speed steel M2, recent development [1] has focused mainly on vacuum heat treatment procedures. High-speed steels have better resistance to wear in comparison to cold-work tool steels because of the increased matrix and carbide phase hardness. The carbide phase in high-speed steels increases the wear resistance, which latter depends on the total volume of carbides and their hardness. The wear resistance in high-speed steel is mainly determined by vanadium carbides which have a micro-hardness of 2200 to 2400HV [2] (Fig. 1). However, it must not be forgotten that high-speed steels have a greater hot hardness. Even if the workpieces are placed into the tools whilst cold, the working tool surfaces become hot because of the processing in the tool.
3. Experimental procedure The test material was a conventional high-speed steel of AISI M2 type of the same melt. The chemical composition of the steel examined is listed in Table 1. The toughness of high-speed steels is relatively low at the hardness to which they are normally heat treated before use. It is therefore necessary, in optimizing vacuum heat-treatment procedures, that the toughness test methods are capable of responding to changes caused by the variations in microstructure. For this reason, recent investigations [1,3] have focused on the fracture toughness of high-speed steels.
Fig. 2. The geometry of a cylindrical round-notched and pre-cracked tensile specimens.
The geometry of cylindrical round-notched pre-cracked tensile specimens, prepared according recommendations [1,4] is shown in Fig. 2. Accordingly to Grossmann’s concept of a hardenability, the formation of uniform microstructure along the crack front is possible because of the axial symmetry of cylindrical specimens, in comparison with the CT-specimens, where this condition is not fulfilled. For a round-notched pre-cracked specimen, the stress intensity factor is given by Heckel [4] as: KI = P/D 3/2(− 1.27+ 1.72D/d)
(1)
where d is the radius of the uncracked ligament after fatiguing, P is the applied fracture load, and D is the outer diameter of the cylindrical specimen. In order to apply linear-elastic fracture mechanical concepts, the size of the plastic zone at the crack tip must be small compared with the nominal dimensions of the specimen. The size requirement for a valid KIc test is given by Shen Wei et al. [5] as D] 1.5(KIc/sys)
(2)
where sys is the yield stress of the investigated steel. This requirement (Eq. (2)) was fulfilled in all of the present measurements. The fracture surfaces of the cylindrical round-notched and pre-cracked specimens were examined in S.E.M. at low magnification. As is shown in Fig. 3, the fatigue-crack propagation area was sharply separated from the circular central part of the fast-fractured area, so that diameter d of this area was easily measured.
Table 1 Chemical composition of the high-speed steel examined (in wt.%) Material
C
Si
Mn
Cr
Mo
W
V
Co
AISI M2
0.87
0.29
0.30
3.77
4.90
6.24
1.18
0.53
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91
temperatures of 1050, 1100, 1150 and 1230°C, quenched in a flow of gaseous nitrogen at a pressure of 5 bar abs. and double tempered 1 h at temperatures of 510, 540, 570 and 600°C, respectively.
4. Results and discussion
Fig. 3. Fracture surface of a cylindrical round-notched and precracked tensile specimen with the circumferential fatigue-crack propagation area sharply separated from the circular central fast-fractured area.
Cylindrical round-notched tensile specimens with a diameter of 10 mm were machined from soft annealed bars with a Brinell hardness of 255. Specimens were fatigued to produce a sharp circumferential crack at the notch root, then austenitized in a vacuum furnace at
Metallographic examination of the specimens [1] shows that the austenite grain size of all specimens that were gas quenched from the austenitization temperature 1050 and 1230°C was 21 and 8 SG, respectively. The microstructure of the cylindrical round-notched tensile specimens examined by S.E.M. at higher magnification is shown in Fig. 4, consisting of tempered martensite and carbides. The quantity of fine carbide particles decreases with the increase of the austenitizing temperature. In addition, it was also noticed that at higher austenitizing temperatures, particularly at 1230°C (the last column in Fig. 4), primary carbide particles penetrate along the edges of the grains because of partial melting and decrease of surface tension of the matrix–carbides interface. When specimens are tempered above 500°C, the retained austenite is ‘conditioned’ and subsequently decomposes to martensite on quenching [3]. The reaction is completed at 540°C, although small amounts of austenite (less than 3%) remain after tempering at 580°C. If retained austenite affects the fracture tough-
Fig. 4. The microstructure of vacuum hardened and tempered specimens examined by S.E.M.
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Fig. 5. The effect of the austenitizing and tempering temperature on the fracture toughness and hardness of M2 steel (FT-fracture toughness; H-hardness).
ness of M2 steel, the effect is small enough to be obscured by secondary hardening. The fracture toughness would be expected to increase with the increasing content of retained austenite, as observed for underaged microstructures tempered below 500°C. In contrast, specimens austenitized at 1230°C, and thus containing larger amounts of retained austenite than the specimens austenitized at 1050°C, have a lower fracture toughness. Obviously, other factors such as the distribution, morphology and stability of retained austenite may also be changing and have a pronounced effect on the fracture toughness [3]. The fracture toughness properties are summarized in Fig. 5. Specimens austenitized at 1050°C were somewhat tougher than those austenitized at 1230°C. The fracture toughness decreases with increasing tempering temperature up to 540°C and then increases again with further increase in tempering temperature. The fracture toughness is inversely correlated with hardness, as generally observed for steels. Despite the non-standard specimens and pre-cracking technique used, KIc values were reproducible and in good agreement with other published data [3]. Fig. 5 shows that KIc increases with decreasing hardness, as observed earlier [6]. In the temperature range of vacuum heat-treatments examined, the hardness is dependent on the secondary-hardening reaction. It is therefore postulated that carbide particles play a primary role in determining the fracture toughness of AISI M2 high speed steel. The lower fracture toughness of specimens austenitized at 1230°C can be associated with an increased carbon content in the matrix and subsequently with a
more profuse precipitation of secondary carbides. The enhanced volume fraction of carbide particles also reduces the fracture toughness by providing a preferential path for crack propagation. This postulate is supported by the work of Olsson [7], in a high speed steel the final crack growth occurring in the carbide–matrix interfaces. On the basis of the experimental results [1], it was possible to draw the diagram shown in Fig. 6, where the technological parameters of the vacuum heat-treatment, mechanical properties and microstructure of the vacuum heat-treated specimens are combined. From the diagram in Fig. 6, it is evident that the fracture toughness for the tempering temperatures 540, 570 and 600°C, respectively, increases with the decrease of hardening temperature. This is in good agreement with earlier observations [3,6]. On the other hand, for the tempering temperature of 510°C, it was found that the fracture toughness values are very close, although slightly higher than for 600°C, irrespective of the hardening temperatures. On the basis of the curves in Fig. 6, it can be assumed that the high-speed steel M2 hardened from low austenitizing temperatures and tempered at 510°C can achieve an optimal combination of hardness and fracture toughness. The relationship between fracture toughness and austenite grain size, f.i. SG grade 8, shows that at the tempering temperatures 510 and 600°C, the obtained values KIc are 17.8 and 10.6 MNm − 3/2, a difference that is quite important in practice. Different fracture toughness at the same grain size (SG8, 1230/510°C) and virtually nearly constant fracture toughness at the tem-
Fig. 6. Influence of austenite grain size on the fracture toughness of high speed steel M2 (TT-tempering temperature, TA-hardening temperature, HRc-hardness at 510°C).
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pering temperature of 510°C, irrespectively of the austenitizing temperature (1050 to 1230°C, SG21 to SG8), confirms the hypothesis that the austenite grain size is not the only dominant parameter affecting the fracture toughness of high-speed steel M2.
punches and dies were tested on a 250t triple-action hydraulic press and compared with the fine-blanking tools for fine-blanking ratchet wheels conventionally heat-treated in a salt bath [1]. The final hardness of the conventionally heat-treated tools achieved after double tempering at 600°C/1 h, was 58 to 59 HRc, depending on the alloying. The tests [1] showed that higher working hardness (61.5 HRc) and improved fracture toughness of vacuum heat-treated punches and dies—particularly those double tempered at 510°C—had a significant affect on the defects on the cutting edges. The vacuum heat-treated tool life was greater by 15–20%, compared to conventionally heattreated fine blanking tools tempered at 600°C. None of the in-service dimensional instability due to the later-transformed retained austenite was found at tool testing. Namely, X-ray diffraction analysis showed that the content of retained austenite did not exceed 1 vol.% in all of the specimens [1].
5. Tool life tests
6. Conclusions
Long production runs have underlined the importance of improved fine blanking tool life, Fig. 7. The qualities most commonly required from a fineblanking tool are wear resistance and toughness, which are needed to maintain a keen cutting edge, combined with sufficient red hardness. In fine blanking, close dimensional stability is essential. Since gross plastic deformation of a fine blanking tool is not desirable in this respect, the initial compressive yield stress is important. On the basis of the experimental results shown in Fig. 6, it was found that the optimum vacuum heat-treatment of fine-blanking tools from high speed steel M2 needs at least two pre-heating stages (650 and 850°C respectively), a variable hardening temperature (usually 1050 – 1150°C) and double tempering at the same temperature (510°C/1 h). The life of a fine-blanking tool varies considerably depending on the size and design of the blank, the type of blanking steel, and the care and maintenance of the tool. To establish tool life, three tools for a fine-blanking ratchet wheel were selected. The punches and dies were from high-speed steel M2. The blanks, with a material gauge of 4 mm were from AISI C 1045 blanking steel in a spheroidizedannealed condition. The punches and dies were stressrelieved at 650°C/4 h and vacuum heat-treated. Depending on the alloying and the condition of austenitization (1100°C/2 min), a final hardness of 61.5 HRc and a final fracture toughness of 17.3 MNm − 3/2 were obtained after double tempering at 510°C for 1 h (Fig. 6). The vacuum heat-treated
On the basis of a study of the influence of vacuum heat-treatment on mechanical properties and fineblanking performance of AISI M2 high-speed steel, the following conclusions are drawn. (1) The measurements of wear on the cutting edges of punches and dies show that double tempering at 510°C/1 h after vacuum hardening prolongs the life of a fine-blanking tool for ratchet wheels by 15–20%, compared to conventionally heat-treated tools that were hardened at the same austenitizing temperature and tempered twice at 600°C. (2) The different fracture toughness at the same grain size (SG8, 1230–510°C) and the nearly constant fracture toughness at the tempering temperature 510°C, irrespectively of the austenitizing temperatures (1050–1230°C), confirm the hypothesis that the austenite grain size is not the only dominant parameter affecting the fracture toughness of high-speed steel M2. (3) There is a marked improvement in fine-blanking tool life when the tempering temperature is decreased from the currently recommended 600–510°C. It is not the type of heat-treatment process that substantially affects the tool life, but first of all the proper choice of hardening and tempering temperatures. (4) The results presented, confirmed by daily data, justify the use of modern vacuum heat-treatment technology and the use of the newest high-performance high speed steels to achieve great improvements in the lives of fine-blanking tools and overall economy.
Fig. 7. Fine-blanking tool for a ratchet wheel.
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References [1] V. Leskovs' ek, B. Ule, A. Rodc, KZT, 3-4 (1995/39) 397. [2] R. Wilson, Metallurgy and Heat Treatment of Tool Steels, McGraw-Hill Book Company UK, 1975, p. 69. [3] V.T.T. Mihkinen, Mat. Sci. Eng. 78 (1986) 35–45.
.
[4] K. Heckel, Einfu¨hrung in die technische Anwendung der Bruchmechanik, Mu¨nchen, Carl Hanser Verlag, 1970. [5] Shen Wei et al., Eng. Fract. Mech. 16 (1) (1982) 69. [6] G. Hoyle, High Speed Steels, Butterworths, 1988, p. 143. [7] L.R. Olssons, Ph.D Thesis, Chalmers University of Technology, Gothenburg, 1978.
Improved vacuum heat-treatment for fine-blanking tools from high-speed steel M2 V. Leskovs' ek a,*, B. Ule b a
Vacuum Heat Treatment and Surface Engineering Center, Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slo6enia b Uni6ersity of Ljubljana, Faculty of Natural Science and Engineering, Department of Materials and Metallurgy, Aker e6a 12, 1000 Ljubljana, Slo6enia Received 10 March 1997
Abstract The influence of the vacuum heat-treatment process on the hardness, fracture toughness and fine-blanking performances of tools made from high speed steel M2 is investigated. To improve the fine-blanking performances, the fracture toughness and hardness of the AISI M2 high-speed steel was measured after a variety of vacuum heat treatments. © 1998 Published by Elsevier Science S.A. All rights reserved. Keywords: Fine blanking tool; Fracture toughness; Hardness; High speed steel M2; Vacuum heat treatment
1. Introduction Due to their high wear resistance, high speed steels are nowadays used also for fine-blanking tools, especially in case of long series. High-speed steels must withstand compressive stresses and abrasive or adhesive wear, and have a sufficient fracture toughness to resist chipping and cracking. Therefore, vacuum heat-treatment with uniform high-pressure gas quenching is one of the most important operations in the manufacturing of fine blanking tools from high speed steel AISI M2. The importance of optimal vacuum heat-treatment is in the possibility of obtaining an optimum combination of basic characteristic of tools of high speed steels, and of the given working part – fine blanking-tool combination. The fracture toughness and hardness of AISI M2 high-speed steel was measured after a variety of heat treatments. It was found that it is a sensitive function of heat treatment, depending on both the austenitizing and the tempering temperatures. The fracture toughness exhibited an inverse correlation with hardness. The results of this investigation are useful for the optimization of the vacuum heat-treatments of different * Corresponding author. Tel.: +386 61 213780; fax. + 386 61 213780.
tools from high-speed steels, which support the exploitation of tensile impact loads that require the optimal combination of hardness and fracture toughness. 2. Basic characteristics of high-speed tool steel M2 Fine-blanking tool cutting speeds have constantly increased, partly because of the use of new tool materials. However, the lack of toughness has prevented the use of some new materials in certain cutting applications in which enhanced toughness is required. Fractures, as well as macro- and micro-chipping, can destroy the cutting edges. The ability of a steel to resist these phenomena is known as toughness. The toughness that can be achieved by high-speed steel is limited by the microstructural features (carbide size and segregation, strings of inclusions, etc.). When the steel is loaded, stress concentrations appear because of the tool shape and microstructural constituents and cause tool fracture, unless the stress concentrations are relieved by local plastic flow on a micro scale. The ability of the matrix to undergo plastic flow can be altered within wide limits by varying the hardness. Thus, the characteristic microstructural features of steel determine the limits of the toughness levels that can be achieved by vacuum heat treatment.
0924-0136/98/$19.00 © 1998 Published by Elsevier Science S.A. All rights reserved. PII S0924-0136(98)00023-5
V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
90
Fig. 1. Comparative hardness of carbides found in tool steels [2].
To improve the toughness of conventionally-manufactured high-speed steel M2, recent development [1] has focused mainly on vacuum heat treatment procedures. High-speed steels have better resistance to wear in comparison to cold-work tool steels because of the increased matrix and carbide phase hardness. The carbide phase in high-speed steels increases the wear resistance, which latter depends on the total volume of carbides and their hardness. The wear resistance in high-speed steel is mainly determined by vanadium carbides which have a micro-hardness of 2200 to 2400HV [2] (Fig. 1). However, it must not be forgotten that high-speed steels have a greater hot hardness. Even if the workpieces are placed into the tools whilst cold, the working tool surfaces become hot because of the processing in the tool.
3. Experimental procedure The test material was a conventional high-speed steel of AISI M2 type of the same melt. The chemical composition of the steel examined is listed in Table 1. The toughness of high-speed steels is relatively low at the hardness to which they are normally heat treated before use. It is therefore necessary, in optimizing vacuum heat-treatment procedures, that the toughness test methods are capable of responding to changes caused by the variations in microstructure. For this reason, recent investigations [1,3] have focused on the fracture toughness of high-speed steels.
Fig. 2. The geometry of a cylindrical round-notched and pre-cracked tensile specimens.
The geometry of cylindrical round-notched pre-cracked tensile specimens, prepared according recommendations [1,4] is shown in Fig. 2. Accordingly to Grossmann’s concept of a hardenability, the formation of uniform microstructure along the crack front is possible because of the axial symmetry of cylindrical specimens, in comparison with the CT-specimens, where this condition is not fulfilled. For a round-notched pre-cracked specimen, the stress intensity factor is given by Heckel [4] as: KI = P/D 3/2(− 1.27+ 1.72D/d)
(1)
where d is the radius of the uncracked ligament after fatiguing, P is the applied fracture load, and D is the outer diameter of the cylindrical specimen. In order to apply linear-elastic fracture mechanical concepts, the size of the plastic zone at the crack tip must be small compared with the nominal dimensions of the specimen. The size requirement for a valid KIc test is given by Shen Wei et al. [5] as D] 1.5(KIc/sys)
(2)
where sys is the yield stress of the investigated steel. This requirement (Eq. (2)) was fulfilled in all of the present measurements. The fracture surfaces of the cylindrical round-notched and pre-cracked specimens were examined in S.E.M. at low magnification. As is shown in Fig. 3, the fatigue-crack propagation area was sharply separated from the circular central part of the fast-fractured area, so that diameter d of this area was easily measured.
Table 1 Chemical composition of the high-speed steel examined (in wt.%) Material
C
Si
Mn
Cr
Mo
W
V
Co
AISI M2
0.87
0.29
0.30
3.77
4.90
6.24
1.18
0.53
V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
91
temperatures of 1050, 1100, 1150 and 1230°C, quenched in a flow of gaseous nitrogen at a pressure of 5 bar abs. and double tempered 1 h at temperatures of 510, 540, 570 and 600°C, respectively.
4. Results and discussion
Fig. 3. Fracture surface of a cylindrical round-notched and precracked tensile specimen with the circumferential fatigue-crack propagation area sharply separated from the circular central fast-fractured area.
Cylindrical round-notched tensile specimens with a diameter of 10 mm were machined from soft annealed bars with a Brinell hardness of 255. Specimens were fatigued to produce a sharp circumferential crack at the notch root, then austenitized in a vacuum furnace at
Metallographic examination of the specimens [1] shows that the austenite grain size of all specimens that were gas quenched from the austenitization temperature 1050 and 1230°C was 21 and 8 SG, respectively. The microstructure of the cylindrical round-notched tensile specimens examined by S.E.M. at higher magnification is shown in Fig. 4, consisting of tempered martensite and carbides. The quantity of fine carbide particles decreases with the increase of the austenitizing temperature. In addition, it was also noticed that at higher austenitizing temperatures, particularly at 1230°C (the last column in Fig. 4), primary carbide particles penetrate along the edges of the grains because of partial melting and decrease of surface tension of the matrix–carbides interface. When specimens are tempered above 500°C, the retained austenite is ‘conditioned’ and subsequently decomposes to martensite on quenching [3]. The reaction is completed at 540°C, although small amounts of austenite (less than 3%) remain after tempering at 580°C. If retained austenite affects the fracture tough-
Fig. 4. The microstructure of vacuum hardened and tempered specimens examined by S.E.M.
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V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
Fig. 5. The effect of the austenitizing and tempering temperature on the fracture toughness and hardness of M2 steel (FT-fracture toughness; H-hardness).
ness of M2 steel, the effect is small enough to be obscured by secondary hardening. The fracture toughness would be expected to increase with the increasing content of retained austenite, as observed for underaged microstructures tempered below 500°C. In contrast, specimens austenitized at 1230°C, and thus containing larger amounts of retained austenite than the specimens austenitized at 1050°C, have a lower fracture toughness. Obviously, other factors such as the distribution, morphology and stability of retained austenite may also be changing and have a pronounced effect on the fracture toughness [3]. The fracture toughness properties are summarized in Fig. 5. Specimens austenitized at 1050°C were somewhat tougher than those austenitized at 1230°C. The fracture toughness decreases with increasing tempering temperature up to 540°C and then increases again with further increase in tempering temperature. The fracture toughness is inversely correlated with hardness, as generally observed for steels. Despite the non-standard specimens and pre-cracking technique used, KIc values were reproducible and in good agreement with other published data [3]. Fig. 5 shows that KIc increases with decreasing hardness, as observed earlier [6]. In the temperature range of vacuum heat-treatments examined, the hardness is dependent on the secondary-hardening reaction. It is therefore postulated that carbide particles play a primary role in determining the fracture toughness of AISI M2 high speed steel. The lower fracture toughness of specimens austenitized at 1230°C can be associated with an increased carbon content in the matrix and subsequently with a
more profuse precipitation of secondary carbides. The enhanced volume fraction of carbide particles also reduces the fracture toughness by providing a preferential path for crack propagation. This postulate is supported by the work of Olsson [7], in a high speed steel the final crack growth occurring in the carbide–matrix interfaces. On the basis of the experimental results [1], it was possible to draw the diagram shown in Fig. 6, where the technological parameters of the vacuum heat-treatment, mechanical properties and microstructure of the vacuum heat-treated specimens are combined. From the diagram in Fig. 6, it is evident that the fracture toughness for the tempering temperatures 540, 570 and 600°C, respectively, increases with the decrease of hardening temperature. This is in good agreement with earlier observations [3,6]. On the other hand, for the tempering temperature of 510°C, it was found that the fracture toughness values are very close, although slightly higher than for 600°C, irrespective of the hardening temperatures. On the basis of the curves in Fig. 6, it can be assumed that the high-speed steel M2 hardened from low austenitizing temperatures and tempered at 510°C can achieve an optimal combination of hardness and fracture toughness. The relationship between fracture toughness and austenite grain size, f.i. SG grade 8, shows that at the tempering temperatures 510 and 600°C, the obtained values KIc are 17.8 and 10.6 MNm − 3/2, a difference that is quite important in practice. Different fracture toughness at the same grain size (SG8, 1230/510°C) and virtually nearly constant fracture toughness at the tem-
Fig. 6. Influence of austenite grain size on the fracture toughness of high speed steel M2 (TT-tempering temperature, TA-hardening temperature, HRc-hardness at 510°C).
V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
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pering temperature of 510°C, irrespectively of the austenitizing temperature (1050 to 1230°C, SG21 to SG8), confirms the hypothesis that the austenite grain size is not the only dominant parameter affecting the fracture toughness of high-speed steel M2.
punches and dies were tested on a 250t triple-action hydraulic press and compared with the fine-blanking tools for fine-blanking ratchet wheels conventionally heat-treated in a salt bath [1]. The final hardness of the conventionally heat-treated tools achieved after double tempering at 600°C/1 h, was 58 to 59 HRc, depending on the alloying. The tests [1] showed that higher working hardness (61.5 HRc) and improved fracture toughness of vacuum heat-treated punches and dies—particularly those double tempered at 510°C—had a significant affect on the defects on the cutting edges. The vacuum heat-treated tool life was greater by 15–20%, compared to conventionally heattreated fine blanking tools tempered at 600°C. None of the in-service dimensional instability due to the later-transformed retained austenite was found at tool testing. Namely, X-ray diffraction analysis showed that the content of retained austenite did not exceed 1 vol.% in all of the specimens [1].
5. Tool life tests
6. Conclusions
Long production runs have underlined the importance of improved fine blanking tool life, Fig. 7. The qualities most commonly required from a fineblanking tool are wear resistance and toughness, which are needed to maintain a keen cutting edge, combined with sufficient red hardness. In fine blanking, close dimensional stability is essential. Since gross plastic deformation of a fine blanking tool is not desirable in this respect, the initial compressive yield stress is important. On the basis of the experimental results shown in Fig. 6, it was found that the optimum vacuum heat-treatment of fine-blanking tools from high speed steel M2 needs at least two pre-heating stages (650 and 850°C respectively), a variable hardening temperature (usually 1050 – 1150°C) and double tempering at the same temperature (510°C/1 h). The life of a fine-blanking tool varies considerably depending on the size and design of the blank, the type of blanking steel, and the care and maintenance of the tool. To establish tool life, three tools for a fine-blanking ratchet wheel were selected. The punches and dies were from high-speed steel M2. The blanks, with a material gauge of 4 mm were from AISI C 1045 blanking steel in a spheroidizedannealed condition. The punches and dies were stressrelieved at 650°C/4 h and vacuum heat-treated. Depending on the alloying and the condition of austenitization (1100°C/2 min), a final hardness of 61.5 HRc and a final fracture toughness of 17.3 MNm − 3/2 were obtained after double tempering at 510°C for 1 h (Fig. 6). The vacuum heat-treated
On the basis of a study of the influence of vacuum heat-treatment on mechanical properties and fineblanking performance of AISI M2 high-speed steel, the following conclusions are drawn. (1) The measurements of wear on the cutting edges of punches and dies show that double tempering at 510°C/1 h after vacuum hardening prolongs the life of a fine-blanking tool for ratchet wheels by 15–20%, compared to conventionally heat-treated tools that were hardened at the same austenitizing temperature and tempered twice at 600°C. (2) The different fracture toughness at the same grain size (SG8, 1230–510°C) and the nearly constant fracture toughness at the tempering temperature 510°C, irrespectively of the austenitizing temperatures (1050–1230°C), confirm the hypothesis that the austenite grain size is not the only dominant parameter affecting the fracture toughness of high-speed steel M2. (3) There is a marked improvement in fine-blanking tool life when the tempering temperature is decreased from the currently recommended 600–510°C. It is not the type of heat-treatment process that substantially affects the tool life, but first of all the proper choice of hardening and tempering temperatures. (4) The results presented, confirmed by daily data, justify the use of modern vacuum heat-treatment technology and the use of the newest high-performance high speed steels to achieve great improvements in the lives of fine-blanking tools and overall economy.
Fig. 7. Fine-blanking tool for a ratchet wheel.
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V. Lesko6s' ek, B. Ule / Journal of Materials Processing Technology 82 (1998) 89–94
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