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Investigation into the potential of a composite combining toughness and plastic deformation resistance
International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Investigation into the potential of a composite combining toughness and plastic deformation resistance T. Viatte b
a,*
, S. Bolognini b, T. Cutard c, G. Feusier b, D. Mari b, W. Benoit
b
a Teledyne Metalworking Products, 1 Teledyne Place,LaVergne, TN 37086, USA Swiss Federal Institute of Technology, Institut de G enie Atomique, PHB Ecublens, CH-1015Lausanne, Switzerland c Present address: Ecole des Mines d'Albi, Campus Jarlard, route de Teillet, F-81000Albi, France
Received 9 March 1998; accepted 16 August 1998
Abstract The hard materials considered in this work cover the range from the Ti(C,N)±Mo2 C±Ni model cermet to the WC±Co model hardmetal. Special emphasis is given to the intermediate systems Ti(C,N)±Mo2 C±Co and the hybrid Ti(C,N)±WC±Mo2 C±Co with dierent Ti(C,N)/WC ratios. This paper illustrates how the sintered microstructure and the mechanical behavior of these grades are related to their initial composition. Image analysis based on SEM observations is used to measure morphologic and structural characteristics. Some of the hybrid materials exhibit a very ®ne structure and a high contiguity, giving rise to high hardness and excellent high temperature mechanical behavior. Ó 1999 Published by Elsevier Science Ltd. All rights reserved. Keywords: Hardmetal; Cermet; Microstructure; Mechanical properties; Image analysis
1. Introduction This paper presents data from an ongoing project, which is a collaboration between Stellram S.A. (a unit of Allegheny-Teledyne Company) and the Swiss Federal Institute of Technology of Lausanne (EPFL, Switzerland). The aim of this research is to correlate the sintered microstructure and the high temperature mechanical behavior with the initial composition of several cemented carbides or carbonitrides intended for cutting tools applications. A complete description of the experimental strategy is presented by Mari et al. in this issue [1]. The WC±Co material is known to be tough, but it deforms plastically at high temperature. This constitutes a limitation to the use of this material as a cutting tool. In high speed machining, the temperature of the edge gets high enough to induce plastic deformation of the tool nose, and the cutting quality is rapidly lost. The Ti(C,N)±Mo2 C±Ni withstands higher temperatures with little plastic deformation, but it is much more brittle than WC±Co [2]. The substitution of nickel by cobalt
*
Corresponding author. Tel.: +1-615-641-4491; fax: +1-615-6414268.
gives a material, the Ti(C,N)±Mo2 C±Co cermet, with mechanical properties between those of WC±Co and Ti(C,N)±Mo2 C±Ni. The brittleness of Ti(C,N)±Mo2 C± Co, however, is still much higher than that of WC±Co [3,4]. It is pparent that the Ti(C,N) based hard skeleton is more brittle than that based on WC. At high temperature however, the Ti(C,N) based skeleton has better resistance to cobalt penetration at the interfaces. The idea of using a hybrid Ti(C,N)±WC skeleton appears naturally to combine low temperature toughness and high temperature creep resistance. For this reason, the combination of the two materials has been further advanced, with the choice of dierent hybrid refractory skeletons Ti(C,N)±WC bound with a pure cobalt metallic phase. The results on including WC grains in a carbonitride based cermet have been studied by dierent authors [5±7], who have found indications of an increase in toughness due to the WC grains, but little is known on the high temperature mechanical behavior of these materials. The change in the mechanical properties when going from a Ti(C,N) based to a WC based material was never studied across the whole range of compositions. Indeed, for compositions where the cubic carbonitrides are in the minority, sintering is known to generate nonequilibrium states for the cubic phases. In this case, the
0263-4368/99/$ ± see front matter Ó 1999 Published by Elsevier Science Ltd. All rights reserved. PII: S 0 2 6 3 - 4 3 6 8 ( 9 8 ) 0 0 0 4 4 - 4
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spontaneous decomposition of the cubic carbonitride grains makes it more dicult to sinter without porosity. Furthermore, this eect generally results in the formation of a layer free of cubic carbide close to the surface of the material, where the composition turns out to be a cobalt rich WC±Co alloy [8,9]. For high speed machining applications, most of the bene®t of adding cubic grains in WC±Co grades is lost due to the poor resistance to plastic deformation of this cobalt rich layer.
binder. In the last four grades, the WC has been introduced with a volume ratio WC/(TiCN+Mo2 C) 25/75 or 75/25 for two dierent cobalt levels. For the four grades containing WC, the volume ratio Ti(C,N)/Mo2 C has been maintained so as to be identical to that in the medium molybdenum cermet grades. In the adopted designation of the grades T refers to Ti(C,N), Mo to Mo2 C, W to WC. The numbers refer to the volumic percentage of the considered phase. 2.2. Image analysis of the sintered microstructures
2. Experimental
The titanium carbonitride based cermets have a more complex microstructure than WC±Co grades. As previously described by Yoshimura [10], during heating a solid state dissolution of the molybdenum carbide occurs in the metallic phase. This dissolution is followed by the formation of the (Ti,Mo)(C,N) rims surrounding the undissolved Ti(C,N) grains of the structure. This core±rim structure has been extensively studied by authors such as Roebuck [11] or more recently Andren [12] and Lindhal [7]. The use of scanning electron microscopy allows one to distinguish up to four dierent phases, which are listed here in order of increasing brightness in the SEM image: 1. The core of the carbonitride grains. In this work, this phase is designated as c0 . The composition is Ti(C1ÿx ,Nx ); it is a face centered cubic structure (NaCl type). Because this phase has the lowest electronic density, it appears as the darkest areas in the structure revealed by secondary electrons microscopy [11]. 2. The rim of the carbonitride grains designated as c00 phase in this work. This phase also is a fcc structure that nucleates and grows during sintering in coher-
2.1. Materials The powders were prepared using standard powder metallurgy methods and commercial powders from HCST Starck. Two qualities of Ti(C,N) powders were used, one with a nominal C/N ratio of 70/30, and the other with a nominal C/N ratio of 50/50. This second powder was used in order to raise the level of nitrogen in the material when WC was added to the composition. The molybdenum was introduced as Mo2 C powder. The powders were attritor milled for 21 h in acetone with hardmetal balls. Wax was added to help the compacting of the samples by uniaxial pressing. The samples were sintered for 2 h at 1450°C. The densi®cation diculties for the compositions containing high WC and low Ti(C,N) were overcome by sintering under a 5 MPa argon pressure (sinterhip). In Table 1 the initial compositions of the 12 grades selected for this paper are presented. The selection was made in order to study the in¯uence of the type of binder (Co or Ni), of the amount of cobalt binder, and of the amount of molybdenum with both Co or Ni Table 1 Initial composition of the cermet grades considered Designation
1 2 3 4 5 6 5 7 8 9 10 11 12
Generic
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18
Low Mo, low Ni Medium Mo, low Ni High Mo, low Ni Low Mo, low Co Medium Mo, low Co High Mo, low Co Medium mo, low Co Medium Mo, medium Co Medium Mo, high Co Low W, low Co Low W, high C High W, low Co High W, high Co
TiC1ÿx Nx
0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.5 0.5 0.5 0.5
% Volume TiCN
Mo2 C
Co
Ni
WC
90.9 87.2 80.5 90.4 87.3 80.6 87.3 80.6 75.6 61.6 53.2 26.4 22.8
3.0 6.4 13.0 3.2 6.4 13.0 6.4 6.4 6.4 4.2 4.2 1.8 1.8
± ± ± 6.4 6.4 6.4 6.4 13.0 18.0 6.0 18.0 6.0 18.0
6.0 6.4 6.5 ± ± ± ± ± ± ± ± ± ±
± ± ± ± ± ± ± ± ± 28.20 24.60 65.80 57.40
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
ence with the c0 core. The composition is (Ti1ÿy ,Moy )(C1ÿx , Nx ) or (Ti1ÿyÿz ,Moy ,Wz )(C1ÿx ,Nx ). Because this phase contains the heavy Mo or W atoms, the c00 rim appears brighter than the c0 core of the grains on the SEM images. In this paper, the symbol c is used to designate the cubic phases c0 or c00 , when no distinction between the two phases is needed. 3. The metallic binder phase (Co or Ni, with Mo, W, Ti atoms in solid solution), or b phase. It generally appears brighter than the rim. However, when heavy W atoms are present in both the rim and the binder, the contrast between these two phases vanishes. 4. The hexagonal WC grains or a phase. These grains, having the highest electronic density, appear as the brightest grains on the SEM image. The image analysis cannot be fully automatic for the following two reasons: · The low contrast between the dierent phases, and the irregular shape of the grains. · In many cases, the interface between two contiguous hard grains is dicult to detect. For these reasons, the pictures had to be visually interpreted, the border of the grains redrawn, and the type of phase identi®ed before being able to use the image analysis calculation. A gray level is assigned to each identi®ed phase. This ®rst visual interpretation of the picture is the most time-consuming, and the most critical in regard to possible errors in the ®nal calculations. The greatest uncertainty comes from the distinction between the metallic binder and the predominant c00 phase.
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Based on the calculations, the following data were evaluated: · The area fraction of each of the four phases on the picture. From this ratio is calculated the corresponding volume ratio of the phase in the sintered material. · The contiguity C of the hard grains. It is calculated here as the ratio between the length of interface (a± a; a±c,c±c) between the hard grains and the total perimeter of these grains. · The mean grain size distribution of each type of grain. The cross-sectional area of each grain is measured and the size is calculated as the diameter of the disc having the same surface. Because it is based on a surface calculation, this value is referred to as ds . The average value ds is the center of the gaussian function ®tting the ds distribution on a logarithmic scale. ds is an apparent diameter that under-estimates the real mean diameter of the grains. The true grain size d that is presented in Table is derived from ds p2 by the geometric relation: d 3=2ds deduced from Underwood [13]. The value of the mean linear inter cept diameter d1 can also be derived p from ds by using the geometric relation: d1 2=3ds also deduced from [13]. · The mean free path in the b phase, calculated from k Vb = 1 ÿ Vb 1 ÿ Cdl as used by Lee and Gurland in [14], where Vb is the volume fraction of the binder, and dl is the mean particle size (mean linear incercept) over the carbide and carbonitride particles. This value of k is dependent upon the estimation of the binder volume fraction Vb .
Table 2 Data calculated from analysis of images presented in Fig. 2 Designation
1 2 3 4 5 6 5 7 8 9 10 11 12
% Volume calculated from image analysis
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18
Cores c0
Rims c00
30.8 26.1 32.1 17.2 22.2 19.6 22.2 16.7 12.7 28.0 15.1 9.1 7.0
59.7 65.1 61.3 76.4 70.7 71.2 70.7 67.7 63.1 51.1 55.9 27.9 23.0
Volume % ratio
Mean diameter (lm)
Grains a
Binder b
Rim/ Rim+core (core+rim) c
13.80 14.21 57.42 56.27
9.5 8.8 6.7 6.4 7.1 9.2 7.1 15.6 24.2 7.0 14.9 5.6 13.7
66.0 71.4 65.7 81.6 76.1 78.4 76.1 80.2 83.2 64.6 78.8 75.4 76.7
0.71 0.58 0.46 0.94 0.85 0.55 0.85 0.81 0.76 0.52 0.41 0.58 0.57
k (lm)
Contiguity
Grains a
Binder b
Grains a,c
0.47 0.55 0.67 0.63
0.177 0.138 0.129 0.163 0.161 0.137 0.161 0.234 0.246 0.149 0.153 0.153 0.165
0.717 0.729 0.829 0.739 0.731 0.727 0.731 0.573 0.342 0.825 0.669 0.836 0.608
The phase symbols a, b, c stand for WC grains, metallic binder phase and cubic carbonitride grains respectively. k designates the mean free path calculated for the binder phase.
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2.3. Room temperature measurements The hardness (Hv ) and the critical-stress intensity factor K1c were measured from 30 kg load Vickers indentation tests. The K1c values were calculated from the measurements of the crack length induced by indentation following the Warren equation [15]: r Hv P 1 K1c 0:087 9:806 4l where Hv is the Vickers hardness, P the load in kgf, and l the average length of the four indentation cracks (in lm). The room temperature elastic properties were investigated by classical ultrasonic pulse±echo measurements and by indentation tests. The Young's modulus E, the shear modulus G and the Poisson ratio m were thus determined from the measurements of the velocity of longitudinal (VL ) and transversal (VT ) ultrasonic waves (with 10 MHz transducers) from the following relations: s E 1 ÿ m 2 VL q 1 m 1 ÿ 2m and
s G VT q
3
where q is the density of the material. The magnetic properties of the sintered samples were measured using conventional sigmameter and coercimeter. In the case of the grades with a nickel binder, a very poor reproducibility was obtained on the measurement of the coercive ®eld, therefore no result are reported in this paper. 2.4. High temperature measurements Three point bend tests were performed using 3:5 7 35 mm3 samples. A modi®ed Instron equipment was used, where the sample is inductively heated under vacuum. The temperature is measured using a pyrometer pointing on the sample itself. The speed of the column is 1 l sÿ1 , corresponding to a strain rate of 1:5 10ÿ5 sÿ1 . 3. Results 3.1. Results of image analysis The original SEM pictures used for the image analysis are presented in Fig. 1. It is apparent that the higher the W content in the composition, the lower the contrast
between the c00 rim and the binder phase. The images on Fig. 2 are the same pictures after image enhancement for phase and grain identi®cation. The four assigned gray levels correspond to the four phases listed previously. The hard skeleton is predominantly composed of the cubic c grains in the case of the two pictures in the fourth row of Fig. 2 (T62±W28±Mo4±Co6 and T53± W25±Mo4±Co18); the hexagonal a grains are isolated in this structure. On the contrary, the hexagonal a grains are predominant in the hard skeleton of the two structures in the ®fth row of Fig. 2 (T26±W66±Mo2±Co6 and T23±W57±Mo2±Co18); this time the cubic grains c are isolated. The main results of the calculations from the image analysis are presented in the Table 2. These values reveal or con®rm several important facts about the eect of composition on the microstructure of the TiCN based cermet. The ®rst fact is that in most of the cases, the c00 rim constitute the predominant phase in the structure, occupying a volume up to ®ve times bigger than the c0 cores. The grain growth inhibitor eect of the molybdenum is clearly con®rmed when comparing the mean diameter of the c grains for dierent molybdenum contents. This eect is observed for both the cobalt and nickel based cermet. It must be noted that the volume ratio between the core and the rim is not aected by the molybdenum content in the grade. Indeed, for a high molybdenum cermet the (Ti,Mo)(C,N) rims are richer in molybdenum but smaller in size compared to a low molybdenum cermet. By using EDS analysis, the Mo content in these rims was found to be 5 and 19 at% for T±Mo3±Co6 and T±Mo13±Co6 respectively, 6 and 24 at% for T±Mo3±Ni6 and T±Mo13±Ni6 respectively. As described by Yoshimura et al. [10], during the ®rst stages of solid phase sintering the c0 cores are surrounded by a Mo-rich c00 rim. This rim limitates any further dissolution of the small grains, and thus any growth of the structure. This description is in agreement with the mechanisms described by Lindau and Stjernberg [16]. In the case of the low molybdenum cermets, the dissolution of the smallest c grains continues during sintering, with a reprecipitation forming the c00 rim around the biggest grains. It must be noted that for the nickel based cermets, the inhibition of the grain growth is accompanied by an increase of the contiguity. In the case of the cobalt based cermets, no such eect on contiguity is observed. This indicates that the building of the hard skeleton during sintering and the resultant morphology of the structure are dependent on the type of metallic binder present. This information is helpful when comparing the mechanical properties of the hard skeleton of TiCN±Mo±Ni and TiCN±Mo±Co cermets.
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Fig. 1. SEM images of the dierent cermet compositions in this study.
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T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Fig. 2. Images reworked from SEM pictures in Fig. 1 as prepared for image analysis.
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
The volume ratio between the rim and the core is much higher with a cobalt than with a nickel binder. The grain size is also larger. This is evidence that coarsening of the structure takes place through dissolution of the smallest c grains in the cobalt and reprecipitation of larger c00 rims. This result is somewhat surprising since the solubility of the TiC in the cobalt is much lower than in the nickel. On the other hand, the solubility of Mo2 C and WC is higher in the cobalt than in nickel [17]. More expected is the decrease in the contiguity of the hard grains observed when the volume of cobalt is higher. The introduction of WC grains and the raise of the nitrogen content in cobalt based cermets also has important eects on their structure. The most spectacular is the increase of the contiguity of the hard grains. The values of Table 2 clearly indicate that the contiguity is enhanced when a part of the c grains is replaced by WC grains. This is observed with 6 vol% as well as with 18 vol% of binder phase. Compared to TiCN, the WC grains exhibit a much higher wettability by the cobalt. The more regular shape of the WC grains also allows a better rearrangement of the structure. In addition to the eect on contiguity, a severe reduction of the diameter of the c grains, and a decrease of the volume ratio between the rim and the core of these grains is observed. This is probably due to a slower coarsening of the rim around the Ti(C,N) grains rich in nitrogen during sintering [10], and to the growth inhibition eect of the tungsten atoms in the binder phase [16]. It is apparent from Table 2 that increasing the WC content induces an increase of the WC average grain size. This can be understood assuming that the dissolved WC reprecipitates on the predominant phase of the structure. The reprecipitation can either take place on the c0 rims, raising the tungsten content in this phase, or give rise to the growth of the WC grains themselves. This is con®rmed in Table 2 when comparing the ®nal volume of WC grains in the structure with the initial volume in the composition. In the case of a high initial WC content, most of the WC remains present as hexagonal a grains. In the case of a low WC grade, the amount of a grains in the structure is only half of the initial WC volume. At the same time, the c0 rims appear to be brighter, indicating a higher tungsten content. Similar observations were obtained by Suzuki and Matsubara [6] and by Nishimura et al. [5] in the case of nickel based cermets containing dierent amounts of WC grains. The enrichment of the inner part of the c0 rim with heavy atoms has already been observed in nickel and cobalt base cermets [7]. As proposed in the model by Cutard and Viatte [2], this is expected to enhance the resistance of the hard skeleton to plastic deformation at high temperature.
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3.2. Mechanical properties The role of molybdenum for the control of the cermet mechanical properties has been presented in a previous paper [2] for the case of Ti(C,N)±Mo±Ni, and in a second paper [4] for the Ti(C,N)±Mo±Co. The present paper focuses on the evolution of the mechanical properties when progressing from pure TiCN±Mo±Co to pure WC±Co. Two conventional WC±Co grades with 6 and 18 vol% Co respectively were used as comparisons. The room temperature mechanical properties and the magnetic characteristics of these materials are presented in Table 3. Important variations are observed and can be related to the morphological and to the chemical characteristics of each phase (core, rim and binder) as a function of the initial composition. As an example, for the grades 5,7 and 8, the increase of the cobalt content is responsible for the increase of the toughness, the decrease of the hardness and the decrease of Young's modulus. However, this eect is small, compared to the one observed when increasing the binder content in the grades containing WC: going from grade 9 to grade 10, from 11 to 12 and from 13 to 14. Figure 3 indicates that the materials can be classi®ed into groups, that show dierent proportionality between the hardness and the Palmqvist crack length. Clearly, the higher the position of the line, the more brittle the material is. This ®gure illustrates that the presence of cubic crystals in the composition makes the material more brittle, and that using a nickel binder makes it generally more brittle than with a cobalt binder. When comparing grades 9, 11, 13, or comparing 10, 12, 14, the higher the WC content, the higher the toughness of the material. This means that changing the nature and morphology of the hard phases can also improve the toughness of the material. In Fig. 4, the relation between toughness and hardness is presented when varying the WC content. Clearly, the nickel based cermets are the most brittle. Replacing nickel binder with cobalt gives the ®rst improvement in the hardness/toughness relation. Comparing grade 5 to grade 9, and grade 8 to grade 10, the addition of the ®rst 25 vol% WC improved the hardness but proportionally reduced the toughness. The hardening eect is a logical result of decreasing the grain size, but more WC interfaces seem to be needed to further improve the toughness of the material together with its hardness. In the WC±Co, the coercitivity is known to be inversely proportional to the mean free path in the cobalt phase. The Fig. 5 indicates that the same is found in the case of the Ti(C,N)±WC±Mo2 C±Co. Only grade 6 is somewhat out of the trend, probably due to the low magnetic saturation that aects the reading of its coercivity.
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Table 3 Room temperature magnetic and mechanical properties of the samples Designation 1 2 3 4 5 6 5 7 8 9 10 11 12 13 14
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18 W94±Co6 W82±Co18
Hc (Oe)
Magnetic saturation (%)
Hardness (Hv 30)
K1c (MPa mÿ1=2 )
na na na 176 171 154 171 135 124 310 251 264 162 380 159
78.5 51.4 0.0 95.3 76.6 40.2 76.6 89.7 85.0 91.2 87.8 99.5 96.8 97.0 96.0
1575 1640 1725 1711 1698 1777 1698 1589 1415 1907 1538 1805 1445 2000 1430
7.70 7.38 7.00 6.76 8.41 7.23 8.41 8.94 9.56 7.70 9.39 8.71 11.23 9.20 12.50
Palmqvist cracks (lm) 146 165 192 204 132 186 132 106 85 177 96 130 63 128 50
E (GPa)
G (GPa)
m
460 450 434 461 455 449 455 436 416 na na na na na na
192 187 179 192 189 185 189 180 171 na na na na na na
0.2 0.21 0.22 0.20 0.21 0.21 0.21 0.21 0.22 na na na na na na
Hc the coercive ®eld. The magnetic saturation is given as a percentage of the theoretical value for each Co content. K1c is the value calculated from the length of the indentation Palmqvist cracks. na not analyzed.
Fig. 3. Relation between the hardness and the Palmqvist crack length for the dierent groups of material. In index is the grade designation from Table 1. The thin lines are ``iso-K1c '' curves calculated from Eq. (2).
Figure 6 establishes that for each cobalt content, the improvement in hardness is correlated with the increase of the coercitivity of the grade. This means that even when drastically changing the composition of the hard phase of the material ± from Ti(C,N)±Mo2 C±Co to WC±Co ± the mean free path in the binder phase remains an important factor to control the hardness of the material. However, the relation between the hardness and the mean free path is far from being as regular as the one measured by Matsubara in the case of pure Ti(C,N)±Mo±Ni [18].
3.3. High temperature mechanical properties In the Fig. 7, the high resistance to plastic deformation of T±Mo13±Ni6 at elevated temperature is indicated by the value ot the yield strength, that is about four times higher for this grade than for the WC±Co6. Comparing with T±Mo13±Co6, it appears that the replacement of the nickel by a cobalt binder lowered the resistance to plastic deformation of the material. In the T±W28±Mo4±Co6, the addition of WC grains in the structure and the raise of the nitrogen content appear to
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
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Fig. 4. Evolution of the Vickers hardness and the fracture toughness of the cermets when adding WC in substitution of Ti(C,N)±Mo2 C in Ti(C,N)± Mo2 C±Co. The K1c value is calculated from the Palmqvist crack length using Eq. (2). In index is the grade designation from Table 1. The values for T±Mo6±Ni6 are given as comparison.
Fig. 5. Relation between coercivity and hardness for dierent TiCN±WC±Mo±Co and WC±Co. In index is the grade designation from Table 1.
have balanced this eect. At 1200°C the yield strength for this grade is as high as that of T±Mo13±Ni6. 4. Discussion It has been noted that due to the grain growth inhibitor eect of the molybdenum, the structure of the T± Mo13±Ni6 is very ®ne and exhibits a high contiguity. Among the tested grades, this one is the most brittle at room temperature. It also exhibits the highest resistance to plastic deformation at high temperature as previously presented by Cutard and Viatte [2]. In that paper, it was proposed that the resistance to plastic deformation is due to the extreme ®neness of the grains of the hard skeleton, and to the particular core±rim structure of the
grains in the Ti(C,N)±Mo2 C±Ni cermets. In the case of the cermet T±Mo13±Co6 the grains are larger and the contiguity is lower. The high temperature bending test results in [6] indicate that this grade exhibits lower resistance to plastic deformation than T±Mo13±Ni6, although still better than the one of the W±Co6. The composition T±W28±Mo4±Co6 is very promising, since the contiguity and grain size are the same as in the best Ni cermet Ti±Mo13±Ni6. The hardness and toughness are higher than in Ti±Mo13±Ni6 since the binder is changed to cobalt. The high temperature measurements presented in Fig. 6 con®rm that the resistance to plastic deformation is as good as in Ti± Mo13±Ni6. Surprisingly, the addition of WC and Co in this material did not extend its plasticity range. Even at temperatures as high as 1200°C this material cannot
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Fig. 6. Relation between the mean free path k and the coercive ®eld Hc for WC±Co and dierent TiCN±WC±Mo±Co. In index is the grade designation from Table 1.
Fig. 7. Stress±strain curve from the three-point-bend tests performed at high temperature. The deformation rate is 1 lm sÿ1 corresponding the strain rate e_ 1:5 sÿ5ÿ1 .
withstand a plastic deformation higher than 0.6% without breaking. In the same conditions, the T±Mo13± Ni6 could be deformed at up to 1% before breaking, and the T±Mo13±Co6 up to 4%. A description of this superplastic-like behavior in the case of cobalt-cermet has been given in [19], in which the penetration of the cobalt into the carbide grain boundaries accelerates the decohesion of the Ti(C,N) skeleton. At these high temperatures, the major part of the deformation is attributed to grain boundary sliding, in addition to a residual contribution due to the deformation of the carbonitride grains. The results presented here indicate
that adding WC grains in the structure could hinder this sliding mechanism to occur. Such a bene®cial eect of WC grains on the high temperature resistance of a cobalt-cermet is unexpected. The WC/WC interfaces are known to be severely dissolved by cobalt at temperatures as low as 900°C [20,21], this is apparently not the case with the more complex WC/(Ti,Mo,W)(C,N) interfaces. The higher nitrogen/carbon ratio in the composition also contributes to raise the creep resistance of T±W28± Mo4±Co6. Indeed, the model presented in [2] explains that the deformation of the c grains is more dicult
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
when a high amount of molybdenum is present in the c00 rim, or when a high amount of nitrogen is present in the c0 core of the carbonitride grains. Indeed, the simultaneous shearing of the core and the rim would impose the formation of numerous Mo±N bonds (or W±N in the present case). The poor anity of these atoms makes this mechanism energetically unfavorable. The diusion of the nitrogen atoms back to the inner core of the c0 grain, and the crossed diusion of carbon atoms toward the c0 ±c00 interface, are needed to make possible the deformation of the grain. The energy needed for this diffusion to occur is the factor limiting the deformation. The particular core±rim structure of the Ti(C,N)±WC± Mo2 C±Co cermets, and the weak bonding in the system nitrogen±molybdenum or nitrogen±tungsten would thus be one of the reasons for the exceptional resistance of this material to plastic deformation at high temperature. 5. Conclusion The range of composition between titanium carbonitride based cermet and tungsten carbide based hardmetal has been explored. The structure of several materials in this range was studied in relation to their initial composition. Image analysis was used for this purpose, revealing that the contiguity of the structures increases, and the grain size decreases, when high amount of tungsten carbide is added to the structure of titanium based cermets. The enhancement of the hardness is a natural consequence of the increase in contiguity. More unexpected is the enhancement of the resistance to plastic deformation at temperature as high as 1200°C. An explanation is proposed that takes into consideration the role of nitrogen, molybdenum and tungsten in the control of the deformation of the carbonitride structure. The resulting material oers a new and very promising combination of mechanical properties, with an increase of the toughness provided by the cobalt binder, and the same resistance to plastic deformation as the best nickel bound cermet tested in this program. References [1] Mari D, Bolognini S, Feusier G, Viatte, Benoit W. Experimental strategy to study the mechanical behaviour of cermets for cutting tools. Int J Refract Met Hard Mater 1998;16(4).
89
[2] Cutard T, Viatte T, Feusier G, Benoit W. Microstructure and high temperature mechanical properties of TiC0.7N0.3±Mo2C± Ni cermets. Mater Sci Engng A 1996;A209:218±27. [3] Viatte T, Cutard T, Bolognini S, Feusier G, Benoit W. Comportement mecanique a haute temperature du metal dur et de dierents cermets de coupe Mater Techn 1997;9±10:13±20. [4] Bolognini S, Feusier G, Mari D, Viatte T, Benoit W. High temperature mechanical behaviour of Ti(C,N)±Mo±Co. Int J Refract Met Hard Mater 1998;16(4). [5] Nishimura T, Murayama K, Kitada T. Some properties of cermets sintered in nitrogen gas. Nippon Tungsten Review 1984;17:11±7. [6] Suzuki H, Matsubara H. Some properties of Ti(CN)±WC±Ni alloy. J Jap Soc P PM 1986;33(4):199±203. [7] Lindahl P, Mainert T, Jonsson H, Andren H.-O. Microstructure and mechanical properties of a (TiW,Ta,Mo)(C,N)±(Co,Ni)-type cermet. J Hard Mater 1993;4:187±204. [8] Suzuki H, Hayashi K, Tanigushi Y. The beta-free layer formed near the surface of vacuum sintered WC±Beta±Co alloys containing nitrogen. Trans Japan Inst Met. 1981;22(11):758± 764. [9] Yohe WC. The development of cubic-carbide-free surface layers in cemented carbides without nitrogen Int J Refract Met Hard Mater 1994;12:137±44. [10] Yoshimura H, Sugisawa T, Nishigaki K, Doi H. Reaction during sintering of TiC±20TiN±15WC±10TaC±9Mo±5.5Ni±11Co cermet. Int J Refract Met Hard Mater 1983;12:170±174. [11] Roebuck B, Gee MG. TiC and Ti(C,N) cermet microstructures. In: Proceedings of 12th Plansee Seminar, 1987. [12] Andren H.-O, Rolander U, Lindahl P. Phase composition in cemented carbides and cermets. Int J Refract Met Hard Mat 1993;12:107±13. [13] Underwood EE. The mathematical foundations of quantitative stereology. In: Pelissier GE, editor. Stereology and quantitative metallography. ASTM, Special Technical Publication 504, Philadelphia, 1974:3±38. [14] Lee HC, Gurland J. Hardness and deformation of cemented tungsten carbide. Mater Sci Engng. 1978;33:125±83. [15] Warren R, Matzke H. Indentation testing of a broad range of cemented carbides. In: Journal of hard materials. New York: Plenum Press, 1983. [16] Lindau L, Stjernberg KG. Grain growth in TiC±Ni±Mo and TiC± Ni±W cemented carbides. Powder Metall 1976;4:210. [17] Edwards R, Raine T. The solid solubilities of some stable carbides in cobalt, nickel and iron at 1250°C. In: 1st Plansee Seminar, 1952. [18] Matsubara H. Mechanical properties of TiC and Ti(CN) base cermets. J Hard Mater 1991;3(3±4):339±50. [19] Cutard T, Bolognini S, Feusier G, Viatte T, Benoit W. Microstructure and high temperature behaviour of various cermets and hard metals. In: PM96, Stockholm, 1996. [20] Mari D, Ammann JJ, Benoit W, Bonjour C. High temperature deformation of WC±Co hardmetal. In: Proceedings of 9th Ris o International Symposium Metal and Materials Science, Denmark, 1988. [21] Mari D, Krawitz AD, Richardson JW, Benoit W. Residual stress in WC±Co measured by neutron diraction. Mater Sci Engng 1996;A209:197±205.
Investigation into the potential of a composite combining toughness and plastic deformation resistance T. Viatte b
a,*
, S. Bolognini b, T. Cutard c, G. Feusier b, D. Mari b, W. Benoit
b
a Teledyne Metalworking Products, 1 Teledyne Place,LaVergne, TN 37086, USA Swiss Federal Institute of Technology, Institut de G enie Atomique, PHB Ecublens, CH-1015Lausanne, Switzerland c Present address: Ecole des Mines d'Albi, Campus Jarlard, route de Teillet, F-81000Albi, France
Received 9 March 1998; accepted 16 August 1998
Abstract The hard materials considered in this work cover the range from the Ti(C,N)±Mo2 C±Ni model cermet to the WC±Co model hardmetal. Special emphasis is given to the intermediate systems Ti(C,N)±Mo2 C±Co and the hybrid Ti(C,N)±WC±Mo2 C±Co with dierent Ti(C,N)/WC ratios. This paper illustrates how the sintered microstructure and the mechanical behavior of these grades are related to their initial composition. Image analysis based on SEM observations is used to measure morphologic and structural characteristics. Some of the hybrid materials exhibit a very ®ne structure and a high contiguity, giving rise to high hardness and excellent high temperature mechanical behavior. Ó 1999 Published by Elsevier Science Ltd. All rights reserved. Keywords: Hardmetal; Cermet; Microstructure; Mechanical properties; Image analysis
1. Introduction This paper presents data from an ongoing project, which is a collaboration between Stellram S.A. (a unit of Allegheny-Teledyne Company) and the Swiss Federal Institute of Technology of Lausanne (EPFL, Switzerland). The aim of this research is to correlate the sintered microstructure and the high temperature mechanical behavior with the initial composition of several cemented carbides or carbonitrides intended for cutting tools applications. A complete description of the experimental strategy is presented by Mari et al. in this issue [1]. The WC±Co material is known to be tough, but it deforms plastically at high temperature. This constitutes a limitation to the use of this material as a cutting tool. In high speed machining, the temperature of the edge gets high enough to induce plastic deformation of the tool nose, and the cutting quality is rapidly lost. The Ti(C,N)±Mo2 C±Ni withstands higher temperatures with little plastic deformation, but it is much more brittle than WC±Co [2]. The substitution of nickel by cobalt
*
Corresponding author. Tel.: +1-615-641-4491; fax: +1-615-6414268.
gives a material, the Ti(C,N)±Mo2 C±Co cermet, with mechanical properties between those of WC±Co and Ti(C,N)±Mo2 C±Ni. The brittleness of Ti(C,N)±Mo2 C± Co, however, is still much higher than that of WC±Co [3,4]. It is pparent that the Ti(C,N) based hard skeleton is more brittle than that based on WC. At high temperature however, the Ti(C,N) based skeleton has better resistance to cobalt penetration at the interfaces. The idea of using a hybrid Ti(C,N)±WC skeleton appears naturally to combine low temperature toughness and high temperature creep resistance. For this reason, the combination of the two materials has been further advanced, with the choice of dierent hybrid refractory skeletons Ti(C,N)±WC bound with a pure cobalt metallic phase. The results on including WC grains in a carbonitride based cermet have been studied by dierent authors [5±7], who have found indications of an increase in toughness due to the WC grains, but little is known on the high temperature mechanical behavior of these materials. The change in the mechanical properties when going from a Ti(C,N) based to a WC based material was never studied across the whole range of compositions. Indeed, for compositions where the cubic carbonitrides are in the minority, sintering is known to generate nonequilibrium states for the cubic phases. In this case, the
0263-4368/99/$ ± see front matter Ó 1999 Published by Elsevier Science Ltd. All rights reserved. PII: S 0 2 6 3 - 4 3 6 8 ( 9 8 ) 0 0 0 4 4 - 4
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spontaneous decomposition of the cubic carbonitride grains makes it more dicult to sinter without porosity. Furthermore, this eect generally results in the formation of a layer free of cubic carbide close to the surface of the material, where the composition turns out to be a cobalt rich WC±Co alloy [8,9]. For high speed machining applications, most of the bene®t of adding cubic grains in WC±Co grades is lost due to the poor resistance to plastic deformation of this cobalt rich layer.
binder. In the last four grades, the WC has been introduced with a volume ratio WC/(TiCN+Mo2 C) 25/75 or 75/25 for two dierent cobalt levels. For the four grades containing WC, the volume ratio Ti(C,N)/Mo2 C has been maintained so as to be identical to that in the medium molybdenum cermet grades. In the adopted designation of the grades T refers to Ti(C,N), Mo to Mo2 C, W to WC. The numbers refer to the volumic percentage of the considered phase. 2.2. Image analysis of the sintered microstructures
2. Experimental
The titanium carbonitride based cermets have a more complex microstructure than WC±Co grades. As previously described by Yoshimura [10], during heating a solid state dissolution of the molybdenum carbide occurs in the metallic phase. This dissolution is followed by the formation of the (Ti,Mo)(C,N) rims surrounding the undissolved Ti(C,N) grains of the structure. This core±rim structure has been extensively studied by authors such as Roebuck [11] or more recently Andren [12] and Lindhal [7]. The use of scanning electron microscopy allows one to distinguish up to four dierent phases, which are listed here in order of increasing brightness in the SEM image: 1. The core of the carbonitride grains. In this work, this phase is designated as c0 . The composition is Ti(C1ÿx ,Nx ); it is a face centered cubic structure (NaCl type). Because this phase has the lowest electronic density, it appears as the darkest areas in the structure revealed by secondary electrons microscopy [11]. 2. The rim of the carbonitride grains designated as c00 phase in this work. This phase also is a fcc structure that nucleates and grows during sintering in coher-
2.1. Materials The powders were prepared using standard powder metallurgy methods and commercial powders from HCST Starck. Two qualities of Ti(C,N) powders were used, one with a nominal C/N ratio of 70/30, and the other with a nominal C/N ratio of 50/50. This second powder was used in order to raise the level of nitrogen in the material when WC was added to the composition. The molybdenum was introduced as Mo2 C powder. The powders were attritor milled for 21 h in acetone with hardmetal balls. Wax was added to help the compacting of the samples by uniaxial pressing. The samples were sintered for 2 h at 1450°C. The densi®cation diculties for the compositions containing high WC and low Ti(C,N) were overcome by sintering under a 5 MPa argon pressure (sinterhip). In Table 1 the initial compositions of the 12 grades selected for this paper are presented. The selection was made in order to study the in¯uence of the type of binder (Co or Ni), of the amount of cobalt binder, and of the amount of molybdenum with both Co or Ni Table 1 Initial composition of the cermet grades considered Designation
1 2 3 4 5 6 5 7 8 9 10 11 12
Generic
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18
Low Mo, low Ni Medium Mo, low Ni High Mo, low Ni Low Mo, low Co Medium Mo, low Co High Mo, low Co Medium mo, low Co Medium Mo, medium Co Medium Mo, high Co Low W, low Co Low W, high C High W, low Co High W, high Co
TiC1ÿx Nx
0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.3 0.5 0.5 0.5 0.5
% Volume TiCN
Mo2 C
Co
Ni
WC
90.9 87.2 80.5 90.4 87.3 80.6 87.3 80.6 75.6 61.6 53.2 26.4 22.8
3.0 6.4 13.0 3.2 6.4 13.0 6.4 6.4 6.4 4.2 4.2 1.8 1.8
± ± ± 6.4 6.4 6.4 6.4 13.0 18.0 6.0 18.0 6.0 18.0
6.0 6.4 6.5 ± ± ± ± ± ± ± ± ± ±
± ± ± ± ± ± ± ± ± 28.20 24.60 65.80 57.40
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
ence with the c0 core. The composition is (Ti1ÿy ,Moy )(C1ÿx , Nx ) or (Ti1ÿyÿz ,Moy ,Wz )(C1ÿx ,Nx ). Because this phase contains the heavy Mo or W atoms, the c00 rim appears brighter than the c0 core of the grains on the SEM images. In this paper, the symbol c is used to designate the cubic phases c0 or c00 , when no distinction between the two phases is needed. 3. The metallic binder phase (Co or Ni, with Mo, W, Ti atoms in solid solution), or b phase. It generally appears brighter than the rim. However, when heavy W atoms are present in both the rim and the binder, the contrast between these two phases vanishes. 4. The hexagonal WC grains or a phase. These grains, having the highest electronic density, appear as the brightest grains on the SEM image. The image analysis cannot be fully automatic for the following two reasons: · The low contrast between the dierent phases, and the irregular shape of the grains. · In many cases, the interface between two contiguous hard grains is dicult to detect. For these reasons, the pictures had to be visually interpreted, the border of the grains redrawn, and the type of phase identi®ed before being able to use the image analysis calculation. A gray level is assigned to each identi®ed phase. This ®rst visual interpretation of the picture is the most time-consuming, and the most critical in regard to possible errors in the ®nal calculations. The greatest uncertainty comes from the distinction between the metallic binder and the predominant c00 phase.
81
Based on the calculations, the following data were evaluated: · The area fraction of each of the four phases on the picture. From this ratio is calculated the corresponding volume ratio of the phase in the sintered material. · The contiguity C of the hard grains. It is calculated here as the ratio between the length of interface (a± a; a±c,c±c) between the hard grains and the total perimeter of these grains. · The mean grain size distribution of each type of grain. The cross-sectional area of each grain is measured and the size is calculated as the diameter of the disc having the same surface. Because it is based on a surface calculation, this value is referred to as ds . The average value ds is the center of the gaussian function ®tting the ds distribution on a logarithmic scale. ds is an apparent diameter that under-estimates the real mean diameter of the grains. The true grain size d that is presented in Table is derived from ds p2 by the geometric relation: d 3=2ds deduced from Underwood [13]. The value of the mean linear inter cept diameter d1 can also be derived p from ds by using the geometric relation: d1 2=3ds also deduced from [13]. · The mean free path in the b phase, calculated from k Vb = 1 ÿ Vb 1 ÿ Cdl as used by Lee and Gurland in [14], where Vb is the volume fraction of the binder, and dl is the mean particle size (mean linear incercept) over the carbide and carbonitride particles. This value of k is dependent upon the estimation of the binder volume fraction Vb .
Table 2 Data calculated from analysis of images presented in Fig. 2 Designation
1 2 3 4 5 6 5 7 8 9 10 11 12
% Volume calculated from image analysis
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18
Cores c0
Rims c00
30.8 26.1 32.1 17.2 22.2 19.6 22.2 16.7 12.7 28.0 15.1 9.1 7.0
59.7 65.1 61.3 76.4 70.7 71.2 70.7 67.7 63.1 51.1 55.9 27.9 23.0
Volume % ratio
Mean diameter (lm)
Grains a
Binder b
Rim/ Rim+core (core+rim) c
13.80 14.21 57.42 56.27
9.5 8.8 6.7 6.4 7.1 9.2 7.1 15.6 24.2 7.0 14.9 5.6 13.7
66.0 71.4 65.7 81.6 76.1 78.4 76.1 80.2 83.2 64.6 78.8 75.4 76.7
0.71 0.58 0.46 0.94 0.85 0.55 0.85 0.81 0.76 0.52 0.41 0.58 0.57
k (lm)
Contiguity
Grains a
Binder b
Grains a,c
0.47 0.55 0.67 0.63
0.177 0.138 0.129 0.163 0.161 0.137 0.161 0.234 0.246 0.149 0.153 0.153 0.165
0.717 0.729 0.829 0.739 0.731 0.727 0.731 0.573 0.342 0.825 0.669 0.836 0.608
The phase symbols a, b, c stand for WC grains, metallic binder phase and cubic carbonitride grains respectively. k designates the mean free path calculated for the binder phase.
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2.3. Room temperature measurements The hardness (Hv ) and the critical-stress intensity factor K1c were measured from 30 kg load Vickers indentation tests. The K1c values were calculated from the measurements of the crack length induced by indentation following the Warren equation [15]: r Hv P 1 K1c 0:087 9:806 4l where Hv is the Vickers hardness, P the load in kgf, and l the average length of the four indentation cracks (in lm). The room temperature elastic properties were investigated by classical ultrasonic pulse±echo measurements and by indentation tests. The Young's modulus E, the shear modulus G and the Poisson ratio m were thus determined from the measurements of the velocity of longitudinal (VL ) and transversal (VT ) ultrasonic waves (with 10 MHz transducers) from the following relations: s E 1 ÿ m 2 VL q 1 m 1 ÿ 2m and
s G VT q
3
where q is the density of the material. The magnetic properties of the sintered samples were measured using conventional sigmameter and coercimeter. In the case of the grades with a nickel binder, a very poor reproducibility was obtained on the measurement of the coercive ®eld, therefore no result are reported in this paper. 2.4. High temperature measurements Three point bend tests were performed using 3:5 7 35 mm3 samples. A modi®ed Instron equipment was used, where the sample is inductively heated under vacuum. The temperature is measured using a pyrometer pointing on the sample itself. The speed of the column is 1 l sÿ1 , corresponding to a strain rate of 1:5 10ÿ5 sÿ1 . 3. Results 3.1. Results of image analysis The original SEM pictures used for the image analysis are presented in Fig. 1. It is apparent that the higher the W content in the composition, the lower the contrast
between the c00 rim and the binder phase. The images on Fig. 2 are the same pictures after image enhancement for phase and grain identi®cation. The four assigned gray levels correspond to the four phases listed previously. The hard skeleton is predominantly composed of the cubic c grains in the case of the two pictures in the fourth row of Fig. 2 (T62±W28±Mo4±Co6 and T53± W25±Mo4±Co18); the hexagonal a grains are isolated in this structure. On the contrary, the hexagonal a grains are predominant in the hard skeleton of the two structures in the ®fth row of Fig. 2 (T26±W66±Mo2±Co6 and T23±W57±Mo2±Co18); this time the cubic grains c are isolated. The main results of the calculations from the image analysis are presented in the Table 2. These values reveal or con®rm several important facts about the eect of composition on the microstructure of the TiCN based cermet. The ®rst fact is that in most of the cases, the c00 rim constitute the predominant phase in the structure, occupying a volume up to ®ve times bigger than the c0 cores. The grain growth inhibitor eect of the molybdenum is clearly con®rmed when comparing the mean diameter of the c grains for dierent molybdenum contents. This eect is observed for both the cobalt and nickel based cermet. It must be noted that the volume ratio between the core and the rim is not aected by the molybdenum content in the grade. Indeed, for a high molybdenum cermet the (Ti,Mo)(C,N) rims are richer in molybdenum but smaller in size compared to a low molybdenum cermet. By using EDS analysis, the Mo content in these rims was found to be 5 and 19 at% for T±Mo3±Co6 and T±Mo13±Co6 respectively, 6 and 24 at% for T±Mo3±Ni6 and T±Mo13±Ni6 respectively. As described by Yoshimura et al. [10], during the ®rst stages of solid phase sintering the c0 cores are surrounded by a Mo-rich c00 rim. This rim limitates any further dissolution of the small grains, and thus any growth of the structure. This description is in agreement with the mechanisms described by Lindau and Stjernberg [16]. In the case of the low molybdenum cermets, the dissolution of the smallest c grains continues during sintering, with a reprecipitation forming the c00 rim around the biggest grains. It must be noted that for the nickel based cermets, the inhibition of the grain growth is accompanied by an increase of the contiguity. In the case of the cobalt based cermets, no such eect on contiguity is observed. This indicates that the building of the hard skeleton during sintering and the resultant morphology of the structure are dependent on the type of metallic binder present. This information is helpful when comparing the mechanical properties of the hard skeleton of TiCN±Mo±Ni and TiCN±Mo±Co cermets.
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Fig. 1. SEM images of the dierent cermet compositions in this study.
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T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Fig. 2. Images reworked from SEM pictures in Fig. 1 as prepared for image analysis.
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
The volume ratio between the rim and the core is much higher with a cobalt than with a nickel binder. The grain size is also larger. This is evidence that coarsening of the structure takes place through dissolution of the smallest c grains in the cobalt and reprecipitation of larger c00 rims. This result is somewhat surprising since the solubility of the TiC in the cobalt is much lower than in the nickel. On the other hand, the solubility of Mo2 C and WC is higher in the cobalt than in nickel [17]. More expected is the decrease in the contiguity of the hard grains observed when the volume of cobalt is higher. The introduction of WC grains and the raise of the nitrogen content in cobalt based cermets also has important eects on their structure. The most spectacular is the increase of the contiguity of the hard grains. The values of Table 2 clearly indicate that the contiguity is enhanced when a part of the c grains is replaced by WC grains. This is observed with 6 vol% as well as with 18 vol% of binder phase. Compared to TiCN, the WC grains exhibit a much higher wettability by the cobalt. The more regular shape of the WC grains also allows a better rearrangement of the structure. In addition to the eect on contiguity, a severe reduction of the diameter of the c grains, and a decrease of the volume ratio between the rim and the core of these grains is observed. This is probably due to a slower coarsening of the rim around the Ti(C,N) grains rich in nitrogen during sintering [10], and to the growth inhibition eect of the tungsten atoms in the binder phase [16]. It is apparent from Table 2 that increasing the WC content induces an increase of the WC average grain size. This can be understood assuming that the dissolved WC reprecipitates on the predominant phase of the structure. The reprecipitation can either take place on the c0 rims, raising the tungsten content in this phase, or give rise to the growth of the WC grains themselves. This is con®rmed in Table 2 when comparing the ®nal volume of WC grains in the structure with the initial volume in the composition. In the case of a high initial WC content, most of the WC remains present as hexagonal a grains. In the case of a low WC grade, the amount of a grains in the structure is only half of the initial WC volume. At the same time, the c0 rims appear to be brighter, indicating a higher tungsten content. Similar observations were obtained by Suzuki and Matsubara [6] and by Nishimura et al. [5] in the case of nickel based cermets containing dierent amounts of WC grains. The enrichment of the inner part of the c0 rim with heavy atoms has already been observed in nickel and cobalt base cermets [7]. As proposed in the model by Cutard and Viatte [2], this is expected to enhance the resistance of the hard skeleton to plastic deformation at high temperature.
85
3.2. Mechanical properties The role of molybdenum for the control of the cermet mechanical properties has been presented in a previous paper [2] for the case of Ti(C,N)±Mo±Ni, and in a second paper [4] for the Ti(C,N)±Mo±Co. The present paper focuses on the evolution of the mechanical properties when progressing from pure TiCN±Mo±Co to pure WC±Co. Two conventional WC±Co grades with 6 and 18 vol% Co respectively were used as comparisons. The room temperature mechanical properties and the magnetic characteristics of these materials are presented in Table 3. Important variations are observed and can be related to the morphological and to the chemical characteristics of each phase (core, rim and binder) as a function of the initial composition. As an example, for the grades 5,7 and 8, the increase of the cobalt content is responsible for the increase of the toughness, the decrease of the hardness and the decrease of Young's modulus. However, this eect is small, compared to the one observed when increasing the binder content in the grades containing WC: going from grade 9 to grade 10, from 11 to 12 and from 13 to 14. Figure 3 indicates that the materials can be classi®ed into groups, that show dierent proportionality between the hardness and the Palmqvist crack length. Clearly, the higher the position of the line, the more brittle the material is. This ®gure illustrates that the presence of cubic crystals in the composition makes the material more brittle, and that using a nickel binder makes it generally more brittle than with a cobalt binder. When comparing grades 9, 11, 13, or comparing 10, 12, 14, the higher the WC content, the higher the toughness of the material. This means that changing the nature and morphology of the hard phases can also improve the toughness of the material. In Fig. 4, the relation between toughness and hardness is presented when varying the WC content. Clearly, the nickel based cermets are the most brittle. Replacing nickel binder with cobalt gives the ®rst improvement in the hardness/toughness relation. Comparing grade 5 to grade 9, and grade 8 to grade 10, the addition of the ®rst 25 vol% WC improved the hardness but proportionally reduced the toughness. The hardening eect is a logical result of decreasing the grain size, but more WC interfaces seem to be needed to further improve the toughness of the material together with its hardness. In the WC±Co, the coercitivity is known to be inversely proportional to the mean free path in the cobalt phase. The Fig. 5 indicates that the same is found in the case of the Ti(C,N)±WC±Mo2 C±Co. Only grade 6 is somewhat out of the trend, probably due to the low magnetic saturation that aects the reading of its coercivity.
86
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Table 3 Room temperature magnetic and mechanical properties of the samples Designation 1 2 3 4 5 6 5 7 8 9 10 11 12 13 14
T±Mo3±Ni6 T±Mo6±Ni6 T±Mo13±Ni6 T±Mo3±Co6 T±Mo6±Co6 T±Mo13±Co6 T±Mo6±Co6 T±Mo6±Co13 T±Mo6±Co18 T±W28±Mo4±Co6 T±W25-Mo4±Co18 T±W66±Mo2±Co6 T±W57±Mo2±Co18 W94±Co6 W82±Co18
Hc (Oe)
Magnetic saturation (%)
Hardness (Hv 30)
K1c (MPa mÿ1=2 )
na na na 176 171 154 171 135 124 310 251 264 162 380 159
78.5 51.4 0.0 95.3 76.6 40.2 76.6 89.7 85.0 91.2 87.8 99.5 96.8 97.0 96.0
1575 1640 1725 1711 1698 1777 1698 1589 1415 1907 1538 1805 1445 2000 1430
7.70 7.38 7.00 6.76 8.41 7.23 8.41 8.94 9.56 7.70 9.39 8.71 11.23 9.20 12.50
Palmqvist cracks (lm) 146 165 192 204 132 186 132 106 85 177 96 130 63 128 50
E (GPa)
G (GPa)
m
460 450 434 461 455 449 455 436 416 na na na na na na
192 187 179 192 189 185 189 180 171 na na na na na na
0.2 0.21 0.22 0.20 0.21 0.21 0.21 0.21 0.22 na na na na na na
Hc the coercive ®eld. The magnetic saturation is given as a percentage of the theoretical value for each Co content. K1c is the value calculated from the length of the indentation Palmqvist cracks. na not analyzed.
Fig. 3. Relation between the hardness and the Palmqvist crack length for the dierent groups of material. In index is the grade designation from Table 1. The thin lines are ``iso-K1c '' curves calculated from Eq. (2).
Figure 6 establishes that for each cobalt content, the improvement in hardness is correlated with the increase of the coercitivity of the grade. This means that even when drastically changing the composition of the hard phase of the material ± from Ti(C,N)±Mo2 C±Co to WC±Co ± the mean free path in the binder phase remains an important factor to control the hardness of the material. However, the relation between the hardness and the mean free path is far from being as regular as the one measured by Matsubara in the case of pure Ti(C,N)±Mo±Ni [18].
3.3. High temperature mechanical properties In the Fig. 7, the high resistance to plastic deformation of T±Mo13±Ni6 at elevated temperature is indicated by the value ot the yield strength, that is about four times higher for this grade than for the WC±Co6. Comparing with T±Mo13±Co6, it appears that the replacement of the nickel by a cobalt binder lowered the resistance to plastic deformation of the material. In the T±W28±Mo4±Co6, the addition of WC grains in the structure and the raise of the nitrogen content appear to
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
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Fig. 4. Evolution of the Vickers hardness and the fracture toughness of the cermets when adding WC in substitution of Ti(C,N)±Mo2 C in Ti(C,N)± Mo2 C±Co. The K1c value is calculated from the Palmqvist crack length using Eq. (2). In index is the grade designation from Table 1. The values for T±Mo6±Ni6 are given as comparison.
Fig. 5. Relation between coercivity and hardness for dierent TiCN±WC±Mo±Co and WC±Co. In index is the grade designation from Table 1.
have balanced this eect. At 1200°C the yield strength for this grade is as high as that of T±Mo13±Ni6. 4. Discussion It has been noted that due to the grain growth inhibitor eect of the molybdenum, the structure of the T± Mo13±Ni6 is very ®ne and exhibits a high contiguity. Among the tested grades, this one is the most brittle at room temperature. It also exhibits the highest resistance to plastic deformation at high temperature as previously presented by Cutard and Viatte [2]. In that paper, it was proposed that the resistance to plastic deformation is due to the extreme ®neness of the grains of the hard skeleton, and to the particular core±rim structure of the
grains in the Ti(C,N)±Mo2 C±Ni cermets. In the case of the cermet T±Mo13±Co6 the grains are larger and the contiguity is lower. The high temperature bending test results in [6] indicate that this grade exhibits lower resistance to plastic deformation than T±Mo13±Ni6, although still better than the one of the W±Co6. The composition T±W28±Mo4±Co6 is very promising, since the contiguity and grain size are the same as in the best Ni cermet Ti±Mo13±Ni6. The hardness and toughness are higher than in Ti±Mo13±Ni6 since the binder is changed to cobalt. The high temperature measurements presented in Fig. 6 con®rm that the resistance to plastic deformation is as good as in Ti± Mo13±Ni6. Surprisingly, the addition of WC and Co in this material did not extend its plasticity range. Even at temperatures as high as 1200°C this material cannot
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T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
Fig. 6. Relation between the mean free path k and the coercive ®eld Hc for WC±Co and dierent TiCN±WC±Mo±Co. In index is the grade designation from Table 1.
Fig. 7. Stress±strain curve from the three-point-bend tests performed at high temperature. The deformation rate is 1 lm sÿ1 corresponding the strain rate e_ 1:5 sÿ5ÿ1 .
withstand a plastic deformation higher than 0.6% without breaking. In the same conditions, the T±Mo13± Ni6 could be deformed at up to 1% before breaking, and the T±Mo13±Co6 up to 4%. A description of this superplastic-like behavior in the case of cobalt-cermet has been given in [19], in which the penetration of the cobalt into the carbide grain boundaries accelerates the decohesion of the Ti(C,N) skeleton. At these high temperatures, the major part of the deformation is attributed to grain boundary sliding, in addition to a residual contribution due to the deformation of the carbonitride grains. The results presented here indicate
that adding WC grains in the structure could hinder this sliding mechanism to occur. Such a bene®cial eect of WC grains on the high temperature resistance of a cobalt-cermet is unexpected. The WC/WC interfaces are known to be severely dissolved by cobalt at temperatures as low as 900°C [20,21], this is apparently not the case with the more complex WC/(Ti,Mo,W)(C,N) interfaces. The higher nitrogen/carbon ratio in the composition also contributes to raise the creep resistance of T±W28± Mo4±Co6. Indeed, the model presented in [2] explains that the deformation of the c grains is more dicult
T. Viatte et al. / International Journal of Refractory Metals & Hard Materials 17 (1999) 79±89
when a high amount of molybdenum is present in the c00 rim, or when a high amount of nitrogen is present in the c0 core of the carbonitride grains. Indeed, the simultaneous shearing of the core and the rim would impose the formation of numerous Mo±N bonds (or W±N in the present case). The poor anity of these atoms makes this mechanism energetically unfavorable. The diusion of the nitrogen atoms back to the inner core of the c0 grain, and the crossed diusion of carbon atoms toward the c0 ±c00 interface, are needed to make possible the deformation of the grain. The energy needed for this diffusion to occur is the factor limiting the deformation. The particular core±rim structure of the Ti(C,N)±WC± Mo2 C±Co cermets, and the weak bonding in the system nitrogen±molybdenum or nitrogen±tungsten would thus be one of the reasons for the exceptional resistance of this material to plastic deformation at high temperature. 5. Conclusion The range of composition between titanium carbonitride based cermet and tungsten carbide based hardmetal has been explored. The structure of several materials in this range was studied in relation to their initial composition. Image analysis was used for this purpose, revealing that the contiguity of the structures increases, and the grain size decreases, when high amount of tungsten carbide is added to the structure of titanium based cermets. The enhancement of the hardness is a natural consequence of the increase in contiguity. More unexpected is the enhancement of the resistance to plastic deformation at temperature as high as 1200°C. An explanation is proposed that takes into consideration the role of nitrogen, molybdenum and tungsten in the control of the deformation of the carbonitride structure. The resulting material oers a new and very promising combination of mechanical properties, with an increase of the toughness provided by the cobalt binder, and the same resistance to plastic deformation as the best nickel bound cermet tested in this program. References [1] Mari D, Bolognini S, Feusier G, Viatte, Benoit W. Experimental strategy to study the mechanical behaviour of cermets for cutting tools. Int J Refract Met Hard Mater 1998;16(4).
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[2] Cutard T, Viatte T, Feusier G, Benoit W. Microstructure and high temperature mechanical properties of TiC0.7N0.3±Mo2C± Ni cermets. Mater Sci Engng A 1996;A209:218±27. [3] Viatte T, Cutard T, Bolognini S, Feusier G, Benoit W. Comportement mecanique a haute temperature du metal dur et de dierents cermets de coupe Mater Techn 1997;9±10:13±20. [4] Bolognini S, Feusier G, Mari D, Viatte T, Benoit W. High temperature mechanical behaviour of Ti(C,N)±Mo±Co. Int J Refract Met Hard Mater 1998;16(4). [5] Nishimura T, Murayama K, Kitada T. Some properties of cermets sintered in nitrogen gas. Nippon Tungsten Review 1984;17:11±7. [6] Suzuki H, Matsubara H. Some properties of Ti(CN)±WC±Ni alloy. J Jap Soc P PM 1986;33(4):199±203. [7] Lindahl P, Mainert T, Jonsson H, Andren H.-O. Microstructure and mechanical properties of a (TiW,Ta,Mo)(C,N)±(Co,Ni)-type cermet. J Hard Mater 1993;4:187±204. [8] Suzuki H, Hayashi K, Tanigushi Y. The beta-free layer formed near the surface of vacuum sintered WC±Beta±Co alloys containing nitrogen. Trans Japan Inst Met. 1981;22(11):758± 764. [9] Yohe WC. The development of cubic-carbide-free surface layers in cemented carbides without nitrogen Int J Refract Met Hard Mater 1994;12:137±44. [10] Yoshimura H, Sugisawa T, Nishigaki K, Doi H. Reaction during sintering of TiC±20TiN±15WC±10TaC±9Mo±5.5Ni±11Co cermet. Int J Refract Met Hard Mater 1983;12:170±174. [11] Roebuck B, Gee MG. TiC and Ti(C,N) cermet microstructures. In: Proceedings of 12th Plansee Seminar, 1987. [12] Andren H.-O, Rolander U, Lindahl P. Phase composition in cemented carbides and cermets. Int J Refract Met Hard Mat 1993;12:107±13. [13] Underwood EE. The mathematical foundations of quantitative stereology. In: Pelissier GE, editor. Stereology and quantitative metallography. ASTM, Special Technical Publication 504, Philadelphia, 1974:3±38. [14] Lee HC, Gurland J. Hardness and deformation of cemented tungsten carbide. Mater Sci Engng. 1978;33:125±83. [15] Warren R, Matzke H. Indentation testing of a broad range of cemented carbides. In: Journal of hard materials. New York: Plenum Press, 1983. [16] Lindau L, Stjernberg KG. Grain growth in TiC±Ni±Mo and TiC± Ni±W cemented carbides. Powder Metall 1976;4:210. [17] Edwards R, Raine T. The solid solubilities of some stable carbides in cobalt, nickel and iron at 1250°C. In: 1st Plansee Seminar, 1952. [18] Matsubara H. Mechanical properties of TiC and Ti(CN) base cermets. J Hard Mater 1991;3(3±4):339±50. [19] Cutard T, Bolognini S, Feusier G, Viatte T, Benoit W. Microstructure and high temperature behaviour of various cermets and hard metals. In: PM96, Stockholm, 1996. [20] Mari D, Ammann JJ, Benoit W, Bonjour C. High temperature deformation of WC±Co hardmetal. In: Proceedings of 9th Ris o International Symposium Metal and Materials Science, Denmark, 1988. [21] Mari D, Krawitz AD, Richardson JW, Benoit W. Residual stress in WC±Co measured by neutron diraction. Mater Sci Engng 1996;A209:197±205.