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High-speed steels: increasing wear resistance by adding ceramic particles
Journal of Materials Processing Technology 92±93 (1999) 15±20
High-speed steels: increasing wear resistance by adding ceramic particles M.M. Oliveiraa,*, J.D. Boltonb b
a INETI-IMP-Department of Materials, Lumiar, 1699 Lisbon, Portugal Department of Mechanical and Manufacturing Engineering, University of Bradford, Bradford, UK
Abstract The main objective of this work was to develop new wear-resistant materials to be used in components for the automotive industry, and study the effects of the ceramic particles on the hardness, bend strength and fracture toughness of the composites, as compared to those of the base material. The composite materials were produced by adding TiC and TiN particles to AISI M3/2 high-speed steel powder containing copper-phosphide and graphite, in order to obtain composite mixtures sinterable in the temperature range of from 1140 to 11708C. It was found that the ceramic particles caused a small increase in the macrohardness but a decrease in the bend strength, since they acted as crack initiators. The fracture toughness was not affected signi®cantly by the presence of the ceramic additions. # 1999 Elsevier Science S.A. All rights reserved. Keywords: High-speed steel; Composites; Titanium carbide; Titanium nitride; Sintering; Wear resistance; Microstructure; Mechanical properties
1. Introduction The results reported in this work were obtained within the framework of the EC BRITE project ``Sintering of highspeed steels containing ceramic or metal coated ceramic powders for wear and fatigue resistance components''. The aim was to use the powder-metallurgy route of the sintering of near-net shapes by developing an economically competitive sintering process to incorporate hard, wear-resistant and cheap ceramic powders into metal matrices. The wearresistant composites produced would be used in standard engineering wear components, especially for the automotive industry: cams, rocker arms, valve seat inserts, and the rotors and stators of high-pressure pumps. In order to maintain the sintering temperatures below 11508C (the maximum attainable in metal link belt furnaces), additions of copper-phosphorus were made, so that activated sintering mechanisms [1] could operate at temperatures below those usually required for full densi®cation of current high-speed steel powders (in the range of 1230± 13208C). The base high-speed steel material chosen was AISI M3/2 and some of the candidate ceramic particles were titanium carbide and titanium nitride. *Corresponding author. Tel.: +351-1-7165181; fax: +351-1-7166568
One purpose of developing current metal matrix composites is the improvement of the wear-resistance of the base material. The ultimate objective is to combine the toughness of the metallic matrix and the hardness of the ceramic particles. Nearly all of the investigations reported on the abrasive wear of metal matrix composites have shown that, although an increase in the volume fraction of the reinforcement generally leads to a decrease in the sliding wear rate, in some cases the reinforcement has been found to produce either no signi®cant improvement or even an increase in the wear rate [2]. This was attributed to a weak bonding at the ceramic/matrix interface, which caused pulling out of the reinforcement particles, creating wear debris that would further aggravate the wear process. The important role of the ceramic/matrix bond, not only on wear-resistance, but also on the bend strength of composite materials, was made evident by previous work [3]. Composites were prepared by adding 10 vol% of Al2O3 particles (20±70 mm) to the same type of base material as that used in the present work. The composites were vacuum sintered and tested in the as-sintered condition. Although the addition of Al2O3 particles resulted in an increase in wearresistance, there was also a drop in bend strength (from 1.6 to 1.2 GPa). When the Al2O3 particles added had a TiN coating (obtained using a CVD technique), there was a signi®cant increase in both the bend strength and the hardness of the
0924-0136/99/$ ± see front matter # 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 4 - 0 1 3 6 ( 9 9 ) 0 0 1 8 0 - 6
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
composite material, their values becoming close to those of the base material, with further increase of wear-resistance. These effects could be explained from the differences observed in the microstructures of the two composite materials: the uncoated Al2O3 particles were not bonded across their interface with the high-speed steel matrix, whilst in the case of the TiN-coated particles, practically no porosity was detected at the interface, and an interfacial layer of vanadium-rich MC type carbides (or carbonitrides) was found [3]. The role of the TiN coating was to introduce extra reactivity at the ceramic/matrix interfaces, creating conditions for better cohesion across those interfaces and improving the sinterability of the composite material. The strength of metal matrix composites is therefore generally lower than that of the base material, due to decreased sinterability of the composites and also to the generally bad quality of the bonding at the ceramic/matrix interface. The important role of the ceramic/matrix interfaces having been established, the ceramic particles used in this work (TiC and TiN) being chosen taking into account the following two conditions [4]: (i) to be stable enough in order to give rise to a dispersion of hard ceramic particles, instead of dissolving into the matrix during the sintering process; (ii) to have some reactivity with the matrix so that a good bonding at the ceramic/matrix interface could be obtained. 2. Experimental work All of the materials studied were processed by a powdermetallurgy route, using the following initial powders: (i) a water atomised and annealed AISI M3/2 high-speed steel powder (1.1C-3.8Cr-5.0W-6.7Mo-2.7V, wt%, bal.Fe); (ii) a copper±phosphorus alloy (Cu3P), which was added (7 wt%) and mixed with the high-speed steel powder in order to lower the sintering temperature by promoting an activated sintering mechanism [1]; (iii) graphite powder, to improve the sinterability of the M3/2Cu3P mixture, the addition found necessary being 0.5 wt% [1]; and (iv) ceramic powders with different nominal mean particle sizes; i.e. three batches of TiC (5 mm; 16 mm; 42 mm) and two batches of TiN (7 and 12 mm). All of the mixtures were prepared by dry blending for 15 min in a Turbula mixer. The ceramic additions (10 vol%) were made by blending into batches of the base material, for 15 min, using the same mixer. Compacts of rectangular shape and approximate dimensions of 26 mm6 mm 4 mm were pressed at pressures in the range of 600± 700 MPa. Sintering was carried out in a vacuum furnace at 11408C and 11708C. Some specimens were hardened and tempered before testing, with austenitisation being carried out at 9508C for 30 min in ¯owing nitrogen, followed by oil quenching. The specimens were subsequently single tempered at temperatures within the range of 450±5508C for 1 h.
The bend strength was determined by three-point bend tests, using specimens with ®nal dimensions of 20 mm 4 mm2 mm, that had been ground longitudinally with diamond wheels on all surfaces. The tensile faces were polished metallographically to 3 mm diamond ®nish. The bend strength (F) was calculated after the ASTM standard B528-76 [5], using the following relationship: F 3Pl= 2bh2 ; where P is the load, l is the span, b is the specimen width and h is the specimen depth. The span used was 16 mm and the cross-head speed 0.l mm/min. The fracture toughness was determined also by a threepoint bend testing method, using dimensions that complied with the standard BS 5447 [6], and being calculated from the following relationship: KIC 3PLY= BW 1:5 ; where P is the load and B(breadth) is 3 mm, W(width) is 6 mm, and 2L(span) is 24 mm, whilst Y is given by: Y 1:93 a=W0:5 ÿ 3:07 a=W1:5 14:53 a=W2:5 25:11 a=W3:5 25:80 a=W4:5 : The test specimens were pre-cracked to give a monotonically-sharp crack of lengtha and an a/W ratio of between 0.35 and 0.6. This was achieved by the arrest of an impact-induced crack by means of a transverse compressive force applied just below a chevron-shaped notch cut by a diamond slitting saw into the test piece [7]. Impact loads were applied by a weight falling onto a cemented carbide wedge placed vertically into the chevron notch. The test pieces were immersed in a nital etching solution after precracking and before testing to stain the pre-crack fracture surface, thus allowing clear identi®cation of the pre-crack and accurate measurement of the pre-crack length (a) to be made by a travelling microscope. Characterisation of the microstructures produced in the various composites and the identi®cation of phases present were carried out by X-ray diffraction and by optical and scanning electron microscopy, including the use of a quantitative image analysis programme. 3. Results Previous experiments with increasing volume fractions of the ceramic particles established that 10 vol% was the maximum content of addition compatible with bend strength levels of above 1 GPa [8]. It was also found that ceramic contents above 10 vol% caused some decrease of the hardness of the composites, due to the presence of porosity, which caused a decrease in the sinterability of the composites. The sintering curves of the composites (Fig. 1) show the effect of the addition of 10 vol% ceramic particles on the
M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
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Fig. 1. Sintering curves of the base material and the composite mixtures. The addition of the ceramic particles decreased the sinterability of the base material, causing an increase of the onset temperature of the sintering plateau.
density of the composites, which maximum was achieved at either 1140 or 11708C. Typical as-sintered microstructures are illustrated in Fig. 2: in the base material, angular M6C and round MC carbides are dispersed in the martensitic matrix, which shows pearlitic areas and isolated copper pools; an eutectic phase of iron-phosphide is also present, both in small areas and along grain boundaries. In the composites, the ceramic particles tended to form agglomerates that involved MC and M6C carbides; pores were found preferentially between agglomerated ceramic particles and more rarely in the ceramic matrix interfaces. The ceramic particles also showed a clearly detectable interaction with the matrix: in all cases, an interfacial layer formed by precipitation of vanadium-rich MC type carbides was observed at the ceramic±matrix interfaces. Whenever TiC and TiN particles showed cracks formed during the sintering process, these were always ®lled with an iron- and phosphorus-rich phase, which had some precipitation of MC carbides, establishing that a liquid phase had been formed which wetted the ceramic particles [1,2]. As shown in Table 1, the addition of the TiC ceramic particles caused an increase in the hardness compared to that of the base material, especially in the case of the composite that contained coarse TiC particles. The heat treatment caused also some hardness increase of the base material and composites, although only limited secondary hardening took place (Fig. 3) because of the low austenitising temperature. The effect of the particle size of the ceramic additions on the bend strength of the composite materials (Table 2) must be compared for the same sintering temperature of 11708C,
Fig. 2. Typical as-sintered microstructures of: (a) the base material; and (b) the TiN composite. M6C and MC carbides are dispersed in a martensitic matrix with dark pearlite areas. Eutectic phases and isolated copper pools are observed. Note the interfacial layer formed by a precipitation of V-rich MC type carbides at the TiN±matrix interfaces.
Table 1 Hardness values (HV30) of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1140 1170 1140 1170 1140 1170 1140 1170 1140 1170
67111 6996 69820 74513 68918 76414 76816 77715 67317 67725 66228 7109
7538 75311 71428 74614 74213 76810 77813 78116 71219 77719 73228 77110
TiC (6 mm) TiC (6 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
corresponding to a fully dense microstructure: it can be seen that coarser ceramic additions decrease the value of the bend strength. The effect of sintering temperature was not the same in all cases: increase in the sintering temperature
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20 Table 3 Fracture toughness KIC of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1170 1170 1140 1170 1170 1170
17.10.1 16.01.4 16.10.8 18.20.9 17.20.8 17.61.0 15.80.9 16.00.1
13.20.3 12.71.0 11.90.9 13.60.6 13.50.4 13.30.6 12.80.8 14.40.9
TiC (6 mm) TiC (16 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
Fig. 3. Tempering curves of the base material and of the composites after hardening from 9508C. Due to the low austenitising temperature, only limited secondary-hardening took place.
Table 2 Three-point bend strength of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1140 1170 1140 1170 1140 1170 1140 1170 1140 1170
1.400.05 1.280.04 1.260.05 1.380.10 1.160.04 1.180.05 1.290.09 1.190.09 1.180.07 1.310.03 1.160.03 1.180.04
1.360.07 1.300.09 1.060.05 1.220.08 1.140.04 1.20.05 1.060.04 1.140.12 1.090.03 1.200.07 1.200.06 1.240.03
TiC (6 mm) TiC (16 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
caused the bend strength to decrease in the base material and in the TiC 42 mm composite, which had attained maximum density at 11408C, but caused an increase in the bend strength of the other composites, which had some porosity left at the lower sintering temperature. The bend tests of heat-treated specimens (Table 2) yielded bend-strength values that tended to be lower than those of the sintered specimens without heat treatment. The composites containing TiC, tested in the as-sintered condition, yielded KIC values not signi®cantly different from those of the base material. (Table 3), whilst slightly lower values were obtained for the composites containing TiN. The results of the fracture-toughness tests of heat-treated specimens, also shown in Table 3, gave KIC values that were signi®cantly lower, both for the base material and for the composites containing TiC and TiN.
In the case of either sintered or heat-treated specimens, the ceramic particles had no signi®cant effect on KIC, and the fracture toughness appeared to be related to the structure of the matrix only. The fracture surfaces of bend-test specimens (Fig. 4) were examined in order to identify failure-initiating sites, which were associated frequently with pores or agglomerates of ceramic particles. Both cleavage and small areas of ductile failure were observed in the matrix of as-sintered samples. Fracture also occurred by cleavage across the M6C carbides and by decohesion around the MC carbides in the base material. In the composites, cleavage was also observed through the ceramic particles. The tensile faces of fractured bend-test specimens were examined in order to search for non-propagating cracks, since static sub-critical cracking has been reported in some high-speed steels [9]. Microcracks at least 200 mm away from the fracture surfaces were detected easily in all specimens, with sizes ranging from 50 to 130 mm and were observed to follow similar paths to those found for fracture surfaces: through M6C carbides and around MC carbides (Fig. 5) or initiated by cracking of the ceramic particles (Fig. 6). 4. Discussion The small increase in hardness caused by the hard ceramic particles additions was attributed to their rather low volume fraction. Since the ceramic particles were well dispersed in the matrix, the hardness of the composites was determined predominantly by that of the base material. Vacuum sintered and heat-treated high-speed steels have bend strengths signi®cantly higher (2 GPa) [10] than those found in this work: it should be noted that the high-speed steel studied here was modi®ed initially by copper±phosphorus additions, yielding microstructures with eutectic areas and grain-boundary carbide and phosphide ®lms [1]. The addition of the ceramic particles caused further degradation of the bend strength, since ceramic particles acted as crack initiators. Crack initiation was seen to occur from TiC and TiN particles or their agglomerates. On the
M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
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Fig. 5. Microcracks formed on carbides on the tensile face of bend-tested specimens (base material). The tensile axis is indicated.
Fig. 6. Microcracks initiated by cracking of the ceramic particles (the tensile face of the bend-test specimen of TiC-16 mm composite material). The tensile axis is indicated.
Fig. 4. The fracture surfaces of bend-tested specimens showing cracks in the carbides and across the ceramic particles: (a) base material; (b) TiC42 mm composite; (c) TiN-12 mm composite.
tensile faces of fractured bend-test specimens, sub-critical microcracks were observed in the ceramic particles and carbides/ceramic agglomerates. Agglomerates of ceramic particles that also contain pores are possible failure initiation sites. The general effect of heat treatment in decreasing the bend strength and increasing the hardness was mainly due to the transformation of retained austenite after tempering and also to a small secondary-hardening effect. The fracture toughness results for all sintered and heattreated samples fell into two scatter bands, one for the as-
sintered and the other for the heat-treated specimens. Neither the presence of the ceramic particles nor the higher levels of porosity in the composite materials appeared to have any effect on KIC, which values met all of the requirements to be valid and agreed with those found by other authors [11]. The reduction caused in KIC by heat treatment is also consistent with the correlation of decreasing fracture toughness with increasing matrix hardness and with the recognition of KIC as a matrix property in high-speed steels. 5. Conclusions 1. Wear-resistant metal matrix composites suitable for automotive applications were produced by adding 10 vol% TiC and TiN ceramic particles to a modi®ed high-speed steel (M3/27%Cu3P0.5%C). These composites were sintered to full density in the temperature range of 1140±11708C. 2. The hardness of the composites reflected mainly the hardness of the matrix, due to the rather low volume
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
fraction of hard ceramic particles. Heat treatment was found to provide a significant hardness increase. 3. The addition of the ceramic particles further reduced the bend strength of the base material, although not significantly for the composites containing the finer additions. 4. Failure initiation in the composites was found to occur from ceramic particles, which acted as crack initiators. 5. The fracture toughness (KIC) of either sintered or heattreated specimens was not affected by porosity or ceramic particles, which supports the assumption that KIC is a matrix property. Quenching and tempered increased the matrix hardness and therefore KIC was lower in the heattreated specimens. Acknowledgements The results reported in this work were obtained in the frame of the research contract EC BRITE ``Sintering of high-speed steels containing ceramic powders for wear and fatigue resistant components''. The authors are grateful to all of the other partners involved: Manganese Bronze (UK),
Metafram (now Sintertech) and Armines (France), and SGM and XYCarb (Holland). References [1] M.M. Oliveira, J.D. Bolton, Powder Met. 38 (1995) 131. [2] C. Jouanny-TreÂsy, Ph.D. Thesis, EÂcole Nationale SupeÂrieure des Mines de Paris (1992) (in French). [3] I.M. Martins, M.M. Oliveira, H. Carvalhinhos, Advances in Powder Metallurgy-1992, MPIF, NJ, 86 (1992) 213. [4] J.D. Bolton, Colloque de la SocieÂte Franc,aise de MeÂtallurgie 1(15) (1990) 1. [5] Standard Test Method for Transverse Rupture Strength of Sintered Metal Powder Specimens, ASTM Standard B 528-76. [6] Methods for Plane Strain Fracture Toughness (KIC) of Metallic Materials', British Standard BS 5447:1977. [7] K. Eriksson, Scand. Met. 4 (1975) 182. [8] M.M. Oliveira, J.D. Bolton, Int. Powder Met. 32 (1996) 37. [9] P.W. Shelton, A.S. Wronski, Metals Sci. 17 (1983) 533. [10] M. Santos, M.M. Oliveira, M.M. Rebbeck, A.S. Wronski, Powder Met. 34 (1991) 93. [11] A.S. Wronski, M.M. Rebbeck, C.S. Wright, W.J.C. Price, T.H. Childs, Advances in Hard Materials Production-HMP'88, Conf. Proc., London 28 (1988) 1.
High-speed steels: increasing wear resistance by adding ceramic particles M.M. Oliveiraa,*, J.D. Boltonb b
a INETI-IMP-Department of Materials, Lumiar, 1699 Lisbon, Portugal Department of Mechanical and Manufacturing Engineering, University of Bradford, Bradford, UK
Abstract The main objective of this work was to develop new wear-resistant materials to be used in components for the automotive industry, and study the effects of the ceramic particles on the hardness, bend strength and fracture toughness of the composites, as compared to those of the base material. The composite materials were produced by adding TiC and TiN particles to AISI M3/2 high-speed steel powder containing copper-phosphide and graphite, in order to obtain composite mixtures sinterable in the temperature range of from 1140 to 11708C. It was found that the ceramic particles caused a small increase in the macrohardness but a decrease in the bend strength, since they acted as crack initiators. The fracture toughness was not affected signi®cantly by the presence of the ceramic additions. # 1999 Elsevier Science S.A. All rights reserved. Keywords: High-speed steel; Composites; Titanium carbide; Titanium nitride; Sintering; Wear resistance; Microstructure; Mechanical properties
1. Introduction The results reported in this work were obtained within the framework of the EC BRITE project ``Sintering of highspeed steels containing ceramic or metal coated ceramic powders for wear and fatigue resistance components''. The aim was to use the powder-metallurgy route of the sintering of near-net shapes by developing an economically competitive sintering process to incorporate hard, wear-resistant and cheap ceramic powders into metal matrices. The wearresistant composites produced would be used in standard engineering wear components, especially for the automotive industry: cams, rocker arms, valve seat inserts, and the rotors and stators of high-pressure pumps. In order to maintain the sintering temperatures below 11508C (the maximum attainable in metal link belt furnaces), additions of copper-phosphorus were made, so that activated sintering mechanisms [1] could operate at temperatures below those usually required for full densi®cation of current high-speed steel powders (in the range of 1230± 13208C). The base high-speed steel material chosen was AISI M3/2 and some of the candidate ceramic particles were titanium carbide and titanium nitride. *Corresponding author. Tel.: +351-1-7165181; fax: +351-1-7166568
One purpose of developing current metal matrix composites is the improvement of the wear-resistance of the base material. The ultimate objective is to combine the toughness of the metallic matrix and the hardness of the ceramic particles. Nearly all of the investigations reported on the abrasive wear of metal matrix composites have shown that, although an increase in the volume fraction of the reinforcement generally leads to a decrease in the sliding wear rate, in some cases the reinforcement has been found to produce either no signi®cant improvement or even an increase in the wear rate [2]. This was attributed to a weak bonding at the ceramic/matrix interface, which caused pulling out of the reinforcement particles, creating wear debris that would further aggravate the wear process. The important role of the ceramic/matrix bond, not only on wear-resistance, but also on the bend strength of composite materials, was made evident by previous work [3]. Composites were prepared by adding 10 vol% of Al2O3 particles (20±70 mm) to the same type of base material as that used in the present work. The composites were vacuum sintered and tested in the as-sintered condition. Although the addition of Al2O3 particles resulted in an increase in wearresistance, there was also a drop in bend strength (from 1.6 to 1.2 GPa). When the Al2O3 particles added had a TiN coating (obtained using a CVD technique), there was a signi®cant increase in both the bend strength and the hardness of the
0924-0136/99/$ ± see front matter # 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 4 - 0 1 3 6 ( 9 9 ) 0 0 1 8 0 - 6
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
composite material, their values becoming close to those of the base material, with further increase of wear-resistance. These effects could be explained from the differences observed in the microstructures of the two composite materials: the uncoated Al2O3 particles were not bonded across their interface with the high-speed steel matrix, whilst in the case of the TiN-coated particles, practically no porosity was detected at the interface, and an interfacial layer of vanadium-rich MC type carbides (or carbonitrides) was found [3]. The role of the TiN coating was to introduce extra reactivity at the ceramic/matrix interfaces, creating conditions for better cohesion across those interfaces and improving the sinterability of the composite material. The strength of metal matrix composites is therefore generally lower than that of the base material, due to decreased sinterability of the composites and also to the generally bad quality of the bonding at the ceramic/matrix interface. The important role of the ceramic/matrix interfaces having been established, the ceramic particles used in this work (TiC and TiN) being chosen taking into account the following two conditions [4]: (i) to be stable enough in order to give rise to a dispersion of hard ceramic particles, instead of dissolving into the matrix during the sintering process; (ii) to have some reactivity with the matrix so that a good bonding at the ceramic/matrix interface could be obtained. 2. Experimental work All of the materials studied were processed by a powdermetallurgy route, using the following initial powders: (i) a water atomised and annealed AISI M3/2 high-speed steel powder (1.1C-3.8Cr-5.0W-6.7Mo-2.7V, wt%, bal.Fe); (ii) a copper±phosphorus alloy (Cu3P), which was added (7 wt%) and mixed with the high-speed steel powder in order to lower the sintering temperature by promoting an activated sintering mechanism [1]; (iii) graphite powder, to improve the sinterability of the M3/2Cu3P mixture, the addition found necessary being 0.5 wt% [1]; and (iv) ceramic powders with different nominal mean particle sizes; i.e. three batches of TiC (5 mm; 16 mm; 42 mm) and two batches of TiN (7 and 12 mm). All of the mixtures were prepared by dry blending for 15 min in a Turbula mixer. The ceramic additions (10 vol%) were made by blending into batches of the base material, for 15 min, using the same mixer. Compacts of rectangular shape and approximate dimensions of 26 mm6 mm 4 mm were pressed at pressures in the range of 600± 700 MPa. Sintering was carried out in a vacuum furnace at 11408C and 11708C. Some specimens were hardened and tempered before testing, with austenitisation being carried out at 9508C for 30 min in ¯owing nitrogen, followed by oil quenching. The specimens were subsequently single tempered at temperatures within the range of 450±5508C for 1 h.
The bend strength was determined by three-point bend tests, using specimens with ®nal dimensions of 20 mm 4 mm2 mm, that had been ground longitudinally with diamond wheels on all surfaces. The tensile faces were polished metallographically to 3 mm diamond ®nish. The bend strength (F) was calculated after the ASTM standard B528-76 [5], using the following relationship: F 3Pl= 2bh2 ; where P is the load, l is the span, b is the specimen width and h is the specimen depth. The span used was 16 mm and the cross-head speed 0.l mm/min. The fracture toughness was determined also by a threepoint bend testing method, using dimensions that complied with the standard BS 5447 [6], and being calculated from the following relationship: KIC 3PLY= BW 1:5 ; where P is the load and B(breadth) is 3 mm, W(width) is 6 mm, and 2L(span) is 24 mm, whilst Y is given by: Y 1:93 a=W0:5 ÿ 3:07 a=W1:5 14:53 a=W2:5 25:11 a=W3:5 25:80 a=W4:5 : The test specimens were pre-cracked to give a monotonically-sharp crack of lengtha and an a/W ratio of between 0.35 and 0.6. This was achieved by the arrest of an impact-induced crack by means of a transverse compressive force applied just below a chevron-shaped notch cut by a diamond slitting saw into the test piece [7]. Impact loads were applied by a weight falling onto a cemented carbide wedge placed vertically into the chevron notch. The test pieces were immersed in a nital etching solution after precracking and before testing to stain the pre-crack fracture surface, thus allowing clear identi®cation of the pre-crack and accurate measurement of the pre-crack length (a) to be made by a travelling microscope. Characterisation of the microstructures produced in the various composites and the identi®cation of phases present were carried out by X-ray diffraction and by optical and scanning electron microscopy, including the use of a quantitative image analysis programme. 3. Results Previous experiments with increasing volume fractions of the ceramic particles established that 10 vol% was the maximum content of addition compatible with bend strength levels of above 1 GPa [8]. It was also found that ceramic contents above 10 vol% caused some decrease of the hardness of the composites, due to the presence of porosity, which caused a decrease in the sinterability of the composites. The sintering curves of the composites (Fig. 1) show the effect of the addition of 10 vol% ceramic particles on the
M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
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Fig. 1. Sintering curves of the base material and the composite mixtures. The addition of the ceramic particles decreased the sinterability of the base material, causing an increase of the onset temperature of the sintering plateau.
density of the composites, which maximum was achieved at either 1140 or 11708C. Typical as-sintered microstructures are illustrated in Fig. 2: in the base material, angular M6C and round MC carbides are dispersed in the martensitic matrix, which shows pearlitic areas and isolated copper pools; an eutectic phase of iron-phosphide is also present, both in small areas and along grain boundaries. In the composites, the ceramic particles tended to form agglomerates that involved MC and M6C carbides; pores were found preferentially between agglomerated ceramic particles and more rarely in the ceramic matrix interfaces. The ceramic particles also showed a clearly detectable interaction with the matrix: in all cases, an interfacial layer formed by precipitation of vanadium-rich MC type carbides was observed at the ceramic±matrix interfaces. Whenever TiC and TiN particles showed cracks formed during the sintering process, these were always ®lled with an iron- and phosphorus-rich phase, which had some precipitation of MC carbides, establishing that a liquid phase had been formed which wetted the ceramic particles [1,2]. As shown in Table 1, the addition of the TiC ceramic particles caused an increase in the hardness compared to that of the base material, especially in the case of the composite that contained coarse TiC particles. The heat treatment caused also some hardness increase of the base material and composites, although only limited secondary hardening took place (Fig. 3) because of the low austenitising temperature. The effect of the particle size of the ceramic additions on the bend strength of the composite materials (Table 2) must be compared for the same sintering temperature of 11708C,
Fig. 2. Typical as-sintered microstructures of: (a) the base material; and (b) the TiN composite. M6C and MC carbides are dispersed in a martensitic matrix with dark pearlite areas. Eutectic phases and isolated copper pools are observed. Note the interfacial layer formed by a precipitation of V-rich MC type carbides at the TiN±matrix interfaces.
Table 1 Hardness values (HV30) of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1140 1170 1140 1170 1140 1170 1140 1170 1140 1170
67111 6996 69820 74513 68918 76414 76816 77715 67317 67725 66228 7109
7538 75311 71428 74614 74213 76810 77813 78116 71219 77719 73228 77110
TiC (6 mm) TiC (6 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
corresponding to a fully dense microstructure: it can be seen that coarser ceramic additions decrease the value of the bend strength. The effect of sintering temperature was not the same in all cases: increase in the sintering temperature
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20 Table 3 Fracture toughness KIC of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1170 1170 1140 1170 1170 1170
17.10.1 16.01.4 16.10.8 18.20.9 17.20.8 17.61.0 15.80.9 16.00.1
13.20.3 12.71.0 11.90.9 13.60.6 13.50.4 13.30.6 12.80.8 14.40.9
TiC (6 mm) TiC (16 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
Fig. 3. Tempering curves of the base material and of the composites after hardening from 9508C. Due to the low austenitising temperature, only limited secondary-hardening took place.
Table 2 Three-point bend strength of the base material and of the TiC and TiN composites, as sintered and as heat-treated Addition
Sintering temperature (8C)
As sintered
Quenched and tempered
±
1140 1170 1140 1170 1140 1170 1140 1170 1140 1170 1140 1170
1.400.05 1.280.04 1.260.05 1.380.10 1.160.04 1.180.05 1.290.09 1.190.09 1.180.07 1.310.03 1.160.03 1.180.04
1.360.07 1.300.09 1.060.05 1.220.08 1.140.04 1.20.05 1.060.04 1.140.12 1.090.03 1.200.07 1.200.06 1.240.03
TiC (6 mm) TiC (16 mm) TiC (42 mm) TiN (7 mm) TiN (12 mm)
caused the bend strength to decrease in the base material and in the TiC 42 mm composite, which had attained maximum density at 11408C, but caused an increase in the bend strength of the other composites, which had some porosity left at the lower sintering temperature. The bend tests of heat-treated specimens (Table 2) yielded bend-strength values that tended to be lower than those of the sintered specimens without heat treatment. The composites containing TiC, tested in the as-sintered condition, yielded KIC values not signi®cantly different from those of the base material. (Table 3), whilst slightly lower values were obtained for the composites containing TiN. The results of the fracture-toughness tests of heat-treated specimens, also shown in Table 3, gave KIC values that were signi®cantly lower, both for the base material and for the composites containing TiC and TiN.
In the case of either sintered or heat-treated specimens, the ceramic particles had no signi®cant effect on KIC, and the fracture toughness appeared to be related to the structure of the matrix only. The fracture surfaces of bend-test specimens (Fig. 4) were examined in order to identify failure-initiating sites, which were associated frequently with pores or agglomerates of ceramic particles. Both cleavage and small areas of ductile failure were observed in the matrix of as-sintered samples. Fracture also occurred by cleavage across the M6C carbides and by decohesion around the MC carbides in the base material. In the composites, cleavage was also observed through the ceramic particles. The tensile faces of fractured bend-test specimens were examined in order to search for non-propagating cracks, since static sub-critical cracking has been reported in some high-speed steels [9]. Microcracks at least 200 mm away from the fracture surfaces were detected easily in all specimens, with sizes ranging from 50 to 130 mm and were observed to follow similar paths to those found for fracture surfaces: through M6C carbides and around MC carbides (Fig. 5) or initiated by cracking of the ceramic particles (Fig. 6). 4. Discussion The small increase in hardness caused by the hard ceramic particles additions was attributed to their rather low volume fraction. Since the ceramic particles were well dispersed in the matrix, the hardness of the composites was determined predominantly by that of the base material. Vacuum sintered and heat-treated high-speed steels have bend strengths signi®cantly higher (2 GPa) [10] than those found in this work: it should be noted that the high-speed steel studied here was modi®ed initially by copper±phosphorus additions, yielding microstructures with eutectic areas and grain-boundary carbide and phosphide ®lms [1]. The addition of the ceramic particles caused further degradation of the bend strength, since ceramic particles acted as crack initiators. Crack initiation was seen to occur from TiC and TiN particles or their agglomerates. On the
M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
19
Fig. 5. Microcracks formed on carbides on the tensile face of bend-tested specimens (base material). The tensile axis is indicated.
Fig. 6. Microcracks initiated by cracking of the ceramic particles (the tensile face of the bend-test specimen of TiC-16 mm composite material). The tensile axis is indicated.
Fig. 4. The fracture surfaces of bend-tested specimens showing cracks in the carbides and across the ceramic particles: (a) base material; (b) TiC42 mm composite; (c) TiN-12 mm composite.
tensile faces of fractured bend-test specimens, sub-critical microcracks were observed in the ceramic particles and carbides/ceramic agglomerates. Agglomerates of ceramic particles that also contain pores are possible failure initiation sites. The general effect of heat treatment in decreasing the bend strength and increasing the hardness was mainly due to the transformation of retained austenite after tempering and also to a small secondary-hardening effect. The fracture toughness results for all sintered and heattreated samples fell into two scatter bands, one for the as-
sintered and the other for the heat-treated specimens. Neither the presence of the ceramic particles nor the higher levels of porosity in the composite materials appeared to have any effect on KIC, which values met all of the requirements to be valid and agreed with those found by other authors [11]. The reduction caused in KIC by heat treatment is also consistent with the correlation of decreasing fracture toughness with increasing matrix hardness and with the recognition of KIC as a matrix property in high-speed steels. 5. Conclusions 1. Wear-resistant metal matrix composites suitable for automotive applications were produced by adding 10 vol% TiC and TiN ceramic particles to a modi®ed high-speed steel (M3/27%Cu3P0.5%C). These composites were sintered to full density in the temperature range of 1140±11708C. 2. The hardness of the composites reflected mainly the hardness of the matrix, due to the rather low volume
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M.M. Oliveira, J.D. Bolton / Journal of Materials Processing Technology 92±93 (1999) 15±20
fraction of hard ceramic particles. Heat treatment was found to provide a significant hardness increase. 3. The addition of the ceramic particles further reduced the bend strength of the base material, although not significantly for the composites containing the finer additions. 4. Failure initiation in the composites was found to occur from ceramic particles, which acted as crack initiators. 5. The fracture toughness (KIC) of either sintered or heattreated specimens was not affected by porosity or ceramic particles, which supports the assumption that KIC is a matrix property. Quenching and tempered increased the matrix hardness and therefore KIC was lower in the heattreated specimens. Acknowledgements The results reported in this work were obtained in the frame of the research contract EC BRITE ``Sintering of high-speed steels containing ceramic powders for wear and fatigue resistant components''. The authors are grateful to all of the other partners involved: Manganese Bronze (UK),
Metafram (now Sintertech) and Armines (France), and SGM and XYCarb (Holland). References [1] M.M. Oliveira, J.D. Bolton, Powder Met. 38 (1995) 131. [2] C. Jouanny-TreÂsy, Ph.D. Thesis, EÂcole Nationale SupeÂrieure des Mines de Paris (1992) (in French). [3] I.M. Martins, M.M. Oliveira, H. Carvalhinhos, Advances in Powder Metallurgy-1992, MPIF, NJ, 86 (1992) 213. [4] J.D. Bolton, Colloque de la SocieÂte Franc,aise de MeÂtallurgie 1(15) (1990) 1. [5] Standard Test Method for Transverse Rupture Strength of Sintered Metal Powder Specimens, ASTM Standard B 528-76. [6] Methods for Plane Strain Fracture Toughness (KIC) of Metallic Materials', British Standard BS 5447:1977. [7] K. Eriksson, Scand. Met. 4 (1975) 182. [8] M.M. Oliveira, J.D. Bolton, Int. Powder Met. 32 (1996) 37. [9] P.W. Shelton, A.S. Wronski, Metals Sci. 17 (1983) 533. [10] M. Santos, M.M. Oliveira, M.M. Rebbeck, A.S. Wronski, Powder Met. 34 (1991) 93. [11] A.S. Wronski, M.M. Rebbeck, C.S. Wright, W.J.C. Price, T.H. Childs, Advances in Hard Materials Production-HMP'88, Conf. Proc., London 28 (1988) 1.