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Consolidation of (Ti,Mo)(C,N)–Ni cermets by glass encapsulated hot isostatic pressing
Materials Science and Engineering A 528 (2011) 4453–4461
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Consolidation of (Ti,Mo)(C,N)–Ni cermets by glass encapsulated hot isostatic pressing N. Rodriguez ∗ , J.M. Sanchez, M. Aristizabal CEIT and TECNUN, Paseo de Manuel Lardizabal 15, 20018 San Sebastian, Spain
a r t i c l e
i n f o
Article history: Received 4 October 2010 Received in revised form 2 February 2011 Accepted 2 February 2011
Keywords: TiMoCN cermets Thermal analysis Hot isostatic pressing Uncombined carbon Grain size refinement
a b s t r a c t Microstructural evolution of (Ti,Mo)(C,N)–Ni cermets consolidated by hot isostatic pressing (HIP) has been analysed. HIP processing allows full densification of these materials at lower temperatures than those normally employed in vacuum sintering cycles (VS). Solution and reprecipitation phenomena are limited and a nanometric fraction of (Ti,Mo)(C,N) grains is retained in the microstructure leading to a significant increase in hardness with respect to vacuum sintered materials (from 11 to 14 GPa). HIP-ed cermets show more tendency to uncombined carbon precipitation than those obtained by VS. In studied systems, carbon precipitation can be related either to an excess of carbon in the initial mixture of powders or to destabilization of carbonitride phase during HIP. Control of the C/N ratio has been carried out by the adequate selection of powder mixtures and the design of the thermal treatments. It has been proved that free carbon in these cermets can be avoided including presintering cycles under hydrogen before encapsulation. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Titanium carbonitride based cermets are materials with high potential for tribological use. These low density composite materials combine high hardness with relatively high toughness and excellent chemical stability, properties that ensure tool life and tool enhanced performance [1–4]. Among the strategies used for improving the wear resistance of titanium carbonitride cermets, microstructural refinement is one of the most thoroughly investigated. Different technologies have been applied for this purpose starting from the use of ultrafine powders in the powder mixtures (even in the nanometric range), to the application of assisted sintering methods to reduce the densification temperatures (i.e. hot isostatic pressing (HIP), spark plasma sintering (SPS), etc.) [5–8]. Apart from powder processing issues (i.e. agglomeration, etc.), the main difficulties reported during consolidation of different ultrafine grained materials by HIP or SPS are related to the incomplete carbothermal reduction of the powder oxides and to the poor wettability of such oxides in contact with the liquid phase [9,10]. In Ti(C,N) cermets case, a high tendency towards graphite precipitation has been identified as a critical issue in HIPed samples. For its detrimental effect on mechanical properties, the presence of free carbon in the microstructure should be avoided and, in this sense, a closer control on the composition during HIP
∗ Corresponding author. Tel.: +34 943 21 28 00; fax: +34 943 21 30 76. E-mail address: [email protected] (N. Rodriguez). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.02.008
is essential. On the other hand, microstructural development during vacuum sintering at high temperature of these materials is well established [1]. The dissolution and reprecipitation of the carbides in binder phase lead to the formation of a typical core-rim structure, where N-rich nuclei appear surrounded by C-rich outer shells with a well defined interface between them [1,11,12]. However, effect of HIP processing on microstructural development of these materials is not widely described yet. Under these considerations, this work is aimed at analysing densification processes and microstructural evolution during consolidation of titanium carbonitride cermets by hot isostatic pressing having vacuum sintered materials as a reference. 2. Materials and methods According to Tables 1 and 2, three different alloys were selected, C1 based on a mixture of carbides, TiC and Mo2 C, and a carbonitride, Ti(C0.7 ,N0.3 ), and C2 and C3 based on two different (Ti,Mo)(C,N) prealloyed powders. In the case of cermet 2, a (Ti,Mo)(C,N) prealloyed powder provided by Treibacher and a nickel powder provided by Inco were mixed. On the other hand, cermet 3 is a commercial mixture of powders provided by Ugicarb consisting of a different (Ti,Mo)(C,N) prealloyed powder, nickel and molybdenum. The terms “grade A” for the (Ti,Mo)(C,N) powder in cermet 2 and “grade B” for the (Ti,Mo)(C,N) powder in cermet 3 are referred to the powder grades used in each case. In all cases Ni was employed as metallic binder. The characterization of the pre-reacted powders by XRD in a previous investigation showed significant differences
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N. Rodriguez et al. / Materials Science and Engineering A 528 (2011) 4453–4461 Table 3 Thermal treatments during consolidation of the cermets.
Table 1 Composition of powder mixtures (wt%). Ref.
Ti (C0.7 ,N0.3 )
(Ti,Mo)(C,N)
TiC
Mo2 C
Mo
C
Ni
C1 C2 C3
39.9 – –
– 70(A) 70(B)
8.1 – –
20.6 – –
– – 3.9
1.3 – –
30 30 26.1
Table 2 Composition of powder mixtures (at%). Ref.
Ti
Mo
C
N
Ni
C1 C2 C3
31.6 32.6 32.2
8.0 7.7 9.7
32.1 31.8 34.2
7.9 7.6 6.0
20.3 20.4 17.9
in their composition. Thus, the carbonitride powder used in cermet C2 (grade A) is actually a combination of (Ti,Mo)(C0.7 ,N0.3 ) and (Ti,Mo)(C0.93 ,N0.07 ) phases, whereas the carbonitride powder grade B (in cermet C3) is quasi-monophasic with a (Ti,Mo)(C0.85 ,N0.15 ) stoichiometry [13]. In relation to atomic composition of these mixtures, cermets C1 and C2 are quite similar while, for commercial “ready to press” C3 powder, lower contents of nickel and nitrogen are remarkable (see Table 2). The oxygen content of the C1 powder was the highest at 1.77 wt%, being the oxygen content of C2 and C3, 0.8 and 0.9 wt% respectively. Milling-mixing step of initial powders was carried out in a planetary mill for 7 h in hexane with 3.5 wt% of paraffin as organic binder. Afterwards, powder slurries were dried at 90 ◦ C in an isothermal bath for 1 h. The melting points of these systems were determined by differential scanning calorimetry. Tests were carried out at 1450 ◦ C for 1 h under vacuum (≈10−1 mbar) with a heating and cooling rate of 10 ◦ C/min. Green compacts were obtained by double action pressing at 160 MPa using a cylindrical steel die and punches. Sintering to full density was carried out by glass encapsulated hot isostatic pressing (GEHIP). During the encapsulation method pressed samples were coated with h-BN and subsequently introduced in Pyrex glass tubes [14]. GEHIP was carried out at two different temperatures, 1250 ◦ C and 1350 ◦ C, with 200 MPa of applied pressure for 45 min (using Ar as the pressing fluid) [14,15]. In some cases, specimens were previously presintered in H2 either at 850 or 1200 ◦ C. Table 3 summarises the conditions of thermal treatments carried out to densify the samples. In order to compare HIP with a more conventional processing route of these materials, results obtained in a previous work for samples densified by presintering in H2 and subsequent vacuum sintering cycles are reported [16]. Standard ISO 3369 was used for density measurements using ethylic alcohol instead of distilled water. Since complex chemical reactions take place during sintering of the cermets, theoretical
Ref.
GH(1) GH(2) GH(3) GH(4)
Thermal treatments Presintering in H2
Encapsulation
HIP
– – 850 ◦ C, 30 min 1200 ◦ C, 30 min
Yes Yes Yes Yes
1250 ◦ C, 200 MPa 1350 ◦ C, 200 MPa 1350 ◦ C, 200 MPa 1350 ◦ C, 200 MPa
density cannot be estimated considering a straightforward relationship between the chemical composition and the theoretical density of initial compounds. For practical reasons in this work, theoretical density was established in 6.5 g/cm3 taking into account Archimedes density values corresponding to fully dense samples obtained by vacuum sintering (<0.02 vol% of porosity according to ISO 4505) [16]. Phase identification was carried out by X-ray diffraction (XRD) (with Ni-filtered CuK␣ radiation). Additionally, the mean crystallite size of (Ti,Mo)(C,N) crystals in the nanometric range was characterized by using the Williamson–Hall approach on XRD peak broadening data [17]. This method assumes that intrinsic broadening (strain broadening and effects of small crystallite size) can be approximated considering the profiles to be Cauchy–Cauchy. Instrumental broadening correction was calculated assuming Gaussian–Gaussian (GG profiles) [18]. To determine the total carbon content of the samples a LECO CS-200 equipment was employed. The measurement technique is based on the combustion of the sample in an oxygen atmosphere to convert the total carbon to CO2 . Then the sample gas flows into a non-dispersive infrared (NDIR) detection cell. The N and O contents were measured with a LECO-TC-400 equipment employing the inert gas fusion method. In this case, the principle of operation is based on the fusion of the sample in a high-purity graphite crucible. The oxygen in the sample reacts with the carbon from the crucible to form carbon monoxide (CO); nitrogen is released as molecular nitrogen (N2 ). Oxygen is detected as either carbon dioxide (CO2 ), CO or both, using infrared (IR) detection, while nitrogen is determined using a thermal-conductivity (TC) cell. In each case a minimum of three measurements per sample were carried out. The sintered specimens were ground and polished down to 1 m diamond paste for microstructural analyses, which were carried out by optical microscopy (OM) and field-emission scanning electron microscopy (FE-SEM). The presence and distribution of porosity were evaluated considering the criteria specified in standard ISO 4505. Compositional analyses were performed by energy dispersive spectroscopy (EDS) in the FE-SEM. Vickers hardness was also measured by applying a load of 10 kgf (ISO 3878).
Fig. 1. DSC curves of the alloys studied: (a) heating and (b) cooling curves.
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3. Results and discussion 3.1. Liquid phase formation Results of calorimetric experiments are shown in Fig. 1. During heating (see Fig. 1(a)), cermets C1, based on a mixture of carbides, and C2, based on a (Ti,Mo)(C,N) prealloyed powder, present an endothermic peak at temperatures near to 1290 ◦ C. At higher temperatures, a second peak associated with liquid phase formation is clearly observed in alloy C1, whereas an exothermic peak is detected in the case of C2. Finally, in cermet C3, with a monophasic prealloyed carbonitride, only one peak is observed showing that melting occurs at a temperature near to 1350 ◦ C (see Table 4). The results confirm that liquid formation depends on the chemical composition of starting powders. Thus, the lowest melting points correspond to compositions containing TiC (C1) or C-rich carbonitride phases (alloy C2) where the eutectic found is close to that occurring in the pseudo-ternary TiC–“MoC”–Ni (at 1280 ◦ C) [19] being noteworthy that these data are consistent with XRD analyses of raw prealloyed powders and with the presence of (Ti,Mo)(C0.93 ,N0.07 ) phase in carbonitride grade A used in composition C2. On the other hand, melting phenomena at higher Table 4 Onset and peak temperatures determined by DSC. Cermet
C1 C2 C3 a
DSC, heating
DSC, cooling
Tonset (◦ C)
Tpeak (◦ C)
Tonset (◦ C)
Tpeak (◦ C)
1270 1323 1288 1364a 1339
1283 1343 1298 1397a 1353
1338 1359 1335 1382 1344 1313
1331 1355 1323 1371 1332 1286
Exothermic peak.
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temperatures observed in cermets C1 and C3 are related to the presence of nitrogen in the lattice of the carbonitride phase. According to data published for similar systems, as the nitrogen content of the carbonitride phase increases, melting occurs at higher temperatures [3,19,20]. Denitridation of high nitrogen-containing alloys during heating in vacuum could explain small differences between the melting point of composition based on Ti(C0.7 ,N0.3 ) (C1), established at 1343 ◦ C, and that of composition based on prealloyed (Ti,Mo)(C0.85 ,N0.15 ) (1353 ◦ C) (see Table 4) [21]. As composition C2 is concerned, the absence of the last peak around 1350 ◦ C is explained bearing in mind the exothermic phenomenon previously described, probably related to the formation of an intermetallic compound (exothermic peak point at 1397 ◦ C according to Table 4). The presence of the exothermic peak might be hiding the melting point related to carbonitride phase. In the carbonitride-based systems, intermetallic phases might be a consequence of a nitrogen partial pressure below the equilibrium pressure of the carbonitride phase for a long time during sintering. In these conditions, it can be assumed that considerable amounts of Ti are dissolved in the binder alloy, so that the precipitation of compounds like Ni3 Ti could be possible [21,22]. In detail, the thermodynamic description of Ti–C–Ni system establishes the equilibrium reaction: liquid ↔ Ni + TiC + Ni3 Ti at 1295 ◦ C [23,24]. Therefore, the first Table 5 Chemical composition of ␣ -phase and binder phase of cermets C1, C2 and C3 densified by consecutive presintering under H2 (850 ◦ C, 30 min) and vacuum sintering (1450 ◦ C, 1 h) (expressed as weight percentage). EDS analyses were carried out in transmission electron microscope. Cermet
C1 C2 C3
Carbonitride phase (rim)
Binder phase
Ti
Mo
Ti
Mo
73 ± 2 70 ± 5 68 ± 5
27 ± 2 30 ± 5 32 ± 5
4.5 ± 0.4 3.7 ± 0.1 3.4 ± 0.3
10.3 ± 0.6 10.1 ± 0.9 11.2 ± 0.8
Fig. 2. Microstructures of cermets based on prealloyed powders, C2 and C3, consolidated by treatments GH(1) and GH(2) (see Table 3).
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melting point of cermet C2 (1298 ◦ C) further supports the idea of subsequent Ni3 Ti formation. Although, on heating, clear differences are found in the DSC curves of the different compositions, the behaviour observed on cooling is similar. Strongest exothermic peak is observed at 1330 ◦ C in all cases (see Fig. 1(b)). The described trend can be related to the homogenization of the chemical composition of the cermets during sintering. This is consistent with the TEM analysis carried out on these samples (see Table 5). Composition analyses show no significant differences in the composition of the binder and the ␣ -phases on sintered materials. It can be assumed that the dissolution of metal atoms from hard phases into the Ni during sintering evolves to a state close to the thermodynamic equilibrium as evidenced by the similarity in the final composition of three cermets despite the different starting compounds used [11]. Taking into account these results and with the objective of analysing the effect of the liquid phase on the processing of the cermets, two different temperatures were selected for HIP cycles: 1250 and 1350 ◦ C respectively. 3.2. Densification In the microstructure of the studied cermets, uncombined carbon related porosity is observed in first hipping experiments at 1250 and 1350 ◦ C respectively. Optical micrographs show rounded
Fig. 3. EDS analysis of uncombined carbon presence in cermet C3 densified in solid state by GH(1). The carbon rich zones where different EDS analyses were carried out are pointed by red arrows. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
Fig. 4. Relative change in the carbon content of specimens after thermal treatments (see Table 3).
porosity in most of the cases making the identification of free carbon as described in ISO 4505 difficult (see Fig. 2(a)–(c)). However, as can be observed in Fig. 3, the characterization of the samples by FE-SEM at a higher magnification confirms the presence of carbon rich areas even when the densification takes place in solid state (EDS analyses pointed by red arrows). In addition, changes in the morphology of the precipitation depending on the HIP temperature can be identified. Thus, in the presence of a liquid phase, the C-rich regions appear in clusters (Fig. 2(d)), whereas, in solid state, free carbon zones are finer and homogeneously distributed in the microstructure (Fig. 2(a)–(c)). These data suggest that, apart from being a nucleation and growth process, uncombined carbon presence in (Ti,Mo)(C,N)–Ni cermets is enhanced by the presence of the liquid phase. This is likely related to the higher solubility of carbon bearing phases in the liquid, which increases with temperature [3]. In order to control the total carbon content of the specimens, a presintering step under hydrogen has been introduced in the processing route of the cermets before encapsulation [13]. Fig. 4 shows the reduction in carbon content of the samples subjected to this thermal treatment (ref. GH(3) and GH(4) according to Table 3). As can be observed, presintering at a higher temperature gives rise to increments in the level of decarburization experienced by the specimens, i.e. presintering parameters can be adjusted to control the carbon losses induced in this step. In cermets C1 and C2 carbon losses due to presintering are more pronounced, while cermet C3 shows a more stable behaviour being less affected by this treatment. Hence, cermets containing free carbon and carbonitride powders with high C/N ratios lead to a stronger interaction with H2 inducing higher carbon losses [13]. According to the observations reported by several authors, this behaviour can be explained considering the high solubility of these phases in Ni-rich binder phase, thus resulting in a greater reactivity of available carbon with the hydrogen in the atmosphere [20]. On the other hand, the decarburization observed in commercial grade C3, even without presintering, can be related to a high level of initial oxidation of the powder mixture due to a long storage time. The carbothermal reduction of NiO at low temperature during encapsulation cycle, that takes place under vacuum (10−1 mbar), explains the diminution in carbon levels determined in treatments GH(1) and GH(2). In any case, fully prealloyed carbonitride based cermet, C3, shows a more stable behaviour with no significant changes in carbon content with the treatment atmosphere. Furthermore, the Archimedes density values can be related to the final carbon contents previously described (see Fig. 5). Thus, in cermets C1 and C2, carbon losses result in a density increase, whereas, for composition C3, with the slightest modification in carbon content, density remains unchanged. The metallographic inspection of the samples by optical microscopy also confirms this
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Table 6 Carbon, nitrogen and oxygen final contents of cermets obtained by different methods (expressed as weight percentage) (Cf , Nf and Of ). Ci and Ni are initial contents of the powder mixtures (C1, C2 and C3, GEHIP(4): see Table 3; C2, PR1200 + VS: cermet C2 obtained by presintering under H2 (1200 ◦ C, 30 min) and subsequent vacuum sintering (1450 ◦ C, 1 h)).
Fig. 5. Density of the samples obtained by different methods (see Table 3).
tendency. As can be observed in Fig. 6, micrographs show the total absence of uncombined carbon in cermets based on prealloyed powders densified by GH(4) (see Table 3), while a reduction in Ctype porosity is encountered in cermet C1 densified in the same cycle. Moreover, the level of porosity of the samples obtained by GH(4) is established in 0.6 vol% in all cases (according to ISO 4505). Therefore, the presintering step is effective controlling the carbon content and the phases present in the microstructure; this leads to an increase in the density and a reduction in the porosity measured in the specimens. Presintering-HIP processing allows full densifi-
Ref.
Ci
Ni
Cf
Nf
Of
C1, GEHIP(4) C2, GEHIP(4) C3, GEHIP(4) C2, GEHIP(2) C2, PR1200 + VS
9.7 9.6 10.2 9.6 9.6
3.0 2.7 2.2 2.7 2.7
8.2 ± 0.03 8.15 ± 0.03 8.84 ± 0.1 9.56 ± 0.07 7.46
3.57 ± 0.25 3.11 ± 0.14 3.0 3.20 ± 0.09 2.82 ± 0.09
0.82 ± 0.11 0.93 ± 0.04 1.2 0.50 ± 0.07 0.58 ± 0.07
cation of these cermets at a lower temperature than that employed in vacuum sintering cycles (i.e. 1450 ◦ C) [16]. On the other hand, carbon rich areas still present in cermet C1, even with a presintering cycle in H2 at 1200 ◦ C, are related to the higher proportion of carbides and free carbon in this composition. At the hipping temperature (in this case 1350 ◦ C), the solubility of TiC and Mo2 C phases in the liquid phase is very high compared to nitrogen-containing phases [3,19]. On cooling, a certain amount of Ti and Mo remains alloyed in the Ni-rich binder phase, whereas the rest precipitates as the ␣ phase onto the undissolved carbonitride grains. With respect to the final content of interstitial elements, HIP-ed cermets show higher contents of carbon, oxygen and nitrogen than those obtained by vacuum sintering (see Table 6). Higher levels of carbon and oxygen are due to the modification of the gas–solid
Fig. 6. Evolution of the porosity of the alloys as a function of the thermal treatments during consolidation (see Table 3).
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Fig. 7. XRD patterns obtained in the surface of C2 densified by GH(2) and GH(4) compared with that obtained in the interior of alloy C2 densified by GH(2) (solid circles correspond to TiN (ICDD card: 06-0642) and open triangles to Ti(C0.7 ,N0.3 ) (ICDD card: 42-1489)).
reactions induced by the glass encapsulation process, since, once the capsule is closed, the carbothermal reduction of powder oxides stops [8]. On the other hand, the increase observed in the nitrogen content could be related to the interaction with the h-BN layer used as a barrier to avoid the infiltration of molten glass in the compacts. XRD analyses of the surface of HIP-ed samples evidence the presence of TiN regardless of the thermal treatment applied (GEHIP or presintering-GEHIP) or the initial composition of the cermet. To illustrate the nitridation of the specimens during HIP, diffraction patterns obtained at the surface of different samples are compared with the pattern obtained at the interior of cermet C2 (see Fig. 7). Furthermore, EDS analyses of alloy C2 support this idea (see Fig. 8). An increase in the N peak intensity is observed in the spectra obtained at the surface when compared with the results observed in the interior (compare Spectrum 1 (surface) with Spectrum 2 (interior) in Fig. 8). As described in previous works, GEHIP processing increases the uncombined carbon presence in Ti(C,N) based cermets [8]. With prealloyed powders of the type (Ti,Mo)(C,N) a similar behaviour has been observed which is likely related to an excess of carbon
Fig. 9. Micrographs of composition C2 by optical microscopy: (a) GEHIP-ed sample and (b) vacuum sintered material. Total carbon content of the GEHIP-ed material is 8.5 wt% and total carbon content of the vacuum sintered material is 9.3 wt%.
content in the powder mixtures and to a high nitrogen activity inside the glass capsules during HIP. The absence of uncombined carbon in samples processed by vacuum sintering with a high total carbon content, even higher than that of the HIP-ed specimens, is in agreement with the hypothesis that the carbon content is not
Fig. 8. EDS analyses at the surface of cermet C2 densified by GH(2) (see Table 3). The Spectrum 1 evidences the presence of nitrogen at the superficial layer.
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Fig. 10. Nitrogen partial pressure vs. atomic fraction of nitrogen in the carbonitride phase at 1350 ◦ C [25].
the only variable that determines precipitation in these systems. In Fig. 9 optical micrographs of alloy C2 show uncombined carbon in GEHIP-ed sample, whereas in the vacuum sintered, only A-type porosity is encountered (according to ISO 4505). According to the domain of stability of Ti(C,N) phase suggested by Pastor [25], free carbon presence is possible by destabilization of the carbonitride phase under the combined effect of high carbon potential and high isostatic pressure. Inside the glass capsule (typically closed under vacuum, aprox. 10−1 mbar) nitrogen evolves from the compacts as the carbonitride phase is heated. The decomposition proceeds increasing the N2 pressure inside the capsule. As isostatic external pressure increases before reaching the eutectic point of the alloys, it is possible that, under such conditions, the nitrogen activity reaches the equilibrium condition for free carbon precipitation (see Fig. 10). The high nitrogen content of HIP-ed samples is consistent with this assumption. The highest nitrogen concentration found in cermet C1 is also in agreement with the higher tendency to free carbon precipitation in this alloy (see Table 6 and Fig. 6). As the nitrogen content on the (Ti,Mo)(C,N) cubic lattice increases, there are fewer sites for carbon atoms, hindering the incorporation of this species to the carbonitride phase. 3.3. Microstructural evolution As can be observed in FE-SEM micrographs, in samples densified in solid state (HIP cycles carried out at 1250 ◦ C), the most relevant microstructural feature is the presence of nanometric carbonitride grains generated during milling step and retained in the microstructure (Fig. 11(a), see blue arrows). Solution and reprecipitation phenomena are limited due to the low temperature employed; consequently, carbonitride phase grain growth is inhibited. The incipient formation of the ␣ phase can also be detected at the surface of larger ceramic grains (see Fig. 11(b), blue arrows). In addition, a higher degree of coalescence of ceramic grains, as well as, a heterogeneous distribution of binder phase where the combination of both characteristics results in the presence of aggregates in the microstructure (see Fig. 11(c)) is observed. On the other hand, nanometric carbonitride grains have almost disappeared from the microstructure of cermets processed in liquid phase (HIP at 1350 ◦ C), which also present a more homogeneous distribution of the binder phase around (Ti,Mo)(C,N) grains (see Fig. 12). A higher fraction of ␣ phase and a more homogeneous composition in the ceramic cores are also observed in these images (as an example of these effects Figs. 11(b) and 12(b) can be compared). This microstructural evolution is explained considering an Ostwald ripening process in which the finest carbonitride grains are preferentially dissolved in the liquid leading to both,
Fig. 11. Microstructural characteristics of cermet C2 densified by sintering in solid state (GH(1) according to Table 3): (a) nanometric carbonitride grains (pointed by blue arrows), (b) ␣ formation (pointed by blue arrows) and (c) coalescence of ceramic grains and binder pooling (red and blue arrows respectively). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
the rim formation and the elimination of nanometric grains. Furthermore, the more homogeneous distribution of the binder phase is due to the good wetting behaviour of the liquid phase in the system. Diffraction peak line broadening observed only for the case of samples processed in solid state (GH(1) according to Table 3) also confirms the elimination of nanometric carbonitride grains when
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Fig. 12. BSE-FEG SEM images of cermets C1, C2 and C3 densified by GH(4) according to Table 3.
Fig. 13. (a) Comparison of peak widths corresponding to cermet C3 densified by vacuum sintering at 1450 ◦ C (VS) and by GE-HIP in solid and liquid phases: GH(1) and GH(4) respectively (see Table 3), and (b) plot of the intrinsic broadening for peaks corresponding to the carbonitride phase in cermets C2 and C3 processed by GH(1) (see Table 3) (Williamson–Hall representation).
the liquid appears; no difference is observed between the diffraction peak of the vacuum sintered sample (at 1450 ◦ C) and that corresponding to HIP-ed sample in the presence of a liquid (GH(4) according to Table 3) (see Fig. 13(a)). The broadening observed in solid state has been used to calculate the mean crystallite size of carbonitride crystals in cermets densified in solid state being 92 and 82 nm respectively for compositions C2 and C3 (see linear fittings in Fig. 13(b)).
The abovementioned microstructural changes affect hardness. In Fig. 14 it can be observed that the highest hardness values correspond to samples processed in solid state, i.e. GH(1). On the other hand, the hardness loss of around 100 kg/mm2 observed when the samples are densified in the presence of a liquid, GH(2), is due to both, free carbon precipitation and microstructural coarsening; meanwhile, the hardness increase for samples processed with an intermediate presintering step is related to the
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materials is due to the good wetting behaviour of the liquid phase which also activates dissolution and reprecipitation phenomena. As expected, these phenomena lead to the growth of the finest carbonitride grains and to the formation of ␣ shells surrounding the undissolved carbonitride cores. Acknowledgments FMD-CARBIDE is gratefully acknowledged for their financial and technical support and for providing the raw materials. N. Rodriguez would also like to thank the Ministerio de Educación y Ciencia within the framework of Programa Torres Quevedo for the financial support given to this work. References Fig. 14. Vickers hardness (10 kg) of samples consolidated by different methods (see Table 3) (PR1200 + VS corresponds to samples densified by presintering at 1200 ◦ C under H2 and subsequent vacuum sintering at 1450 ◦ C, 1 h [16]).
reduction of uncombined carbon. Cermets obtained by presintering at 1200 ◦ C and subsequent GEHIP at 1350 ◦ C, i.e. GH(4), show hardness values around HV10 1320 in the case of compositions C1 and C2 and HV10 1390 for the C3, similar to those reported for GH(1) (HV10 1300 and 1430 for cermets C1 and C3 respectively). The higher hardness of cermet C3 is due to its slightly lower binder content (26.1 wt% Ni vs. 30 wt% Ni in cermets C1 and C2), since the C/N ratio obtained after processing is very similar for the three compositions. Moreover, samples processed by GH(4) show hardness values more than 100 kg/mm2 higher than those obtained by conventional vacuum sintering (see data circled in Fig. 14); the differences being related to the microstructural coarsening at 1450 ◦ C during vacuum sintering (100 ◦ C higher than the temperature of GEHIP). Hardness values of cermets obtained by GH(4) are similar to those reported by Gille et al. for WC–Co compositions with a similar metallic binder content (21 vol%) and a WC grain size between 0.5 and 0.8 m [26]. 4. Conclusions Fully dense pseudo-binary (Ti,Mo)(C,N)–Ni cermets have been obtained by combining presintering under hydrogen and glass encapsulated hot isostatic pressing (GEHIP). In these materials GEHIP processing induces a higher tendency to free carbon precipitation than vacuum sintering, this phenomenon being controlled by introducing a presintering step under hydrogen before encapsulation. Cermets consolidated in solid state present microstructures characterized by heterogeneous distribution of the binder phase and the presence of nanometric carbonitride grains. The microstructural homogenization observed in liquid phase sintered
[1] H. Doi, in: E.A. Almond, C.A. Brookes, R. Warren (Eds.), Proceedings of the International Conference of Science of Hard Materials. Institute of Physics Conference Series 75, Adam Hilger Ltd., Bristol, England, 1986, p. 489. [2] P. Ettmayer, Annual Review of Materials Science 19 (1989) 145–164. [3] P. Ettmayer, H. Kolaska, W. Lengauer, K. Dreyer, International Journal of Refractory Metals and Hard Materials 13 (1995) 343–351. [4] J.L. Ellis, C.G. Goetzel, ASM Handbook. Volume 2: Properties and Selection: Non Ferrous Alloys and Special Purpose Materials, ASM International, 1993, pp. 978–1007. [5] M.A. Alvarez, Densificación de cermets de Ti(C,N) basados en mezclas de polvos ultrafinos, Universidad de Navarra, Escuela Superior de Ingenieros, 2007. [6] J.R. Groza, A. Zavaliangos, Materials Science and Engineering A 287 (2000) 171. [7] J. Jung, S. Kang, Acta Materialia 52 (2004) 1379. [8] J.M. Sánchez, M. Alvarez, N. Rodriguez, M. Aristizabal, Materials Science and Engineering A 500 (2009) 225. [9] P. Feng, W. Xiong, L. Yu, Y. Zheng, Y. Xia, International Journal of Refractory Metals and Hard Materials 22 (2004) 133. [10] Y. Zheng, S. Wang, M. You, H. Tan, W. Xiong, Materials Chemistry and Physics 92 (2005) 64. [11] H.O. Andren, Materials and Design 22 (2001) 491. [12] M.G. Gee, M.J. Reece, B. Roebuck, Journal of Hard Materials 3 (1992) 119–142. [13] N. Rodriguez, J.M. Sanchez, International Journal of Refractory Metals and Hard Materials 27 (2009) 164. [14] F. Castro, I. Iturriza, European Patent, 9202931(3). [15] F. Castro, I. Iturriza, Journal of Materials Science Letters 9 (1990) 600. [16] N. Rodriguez, Sinterización de cermets basados en carbonitruros mixtos de Titanio y Molibdeno, Universidad de Navarra, 2009. [17] G.K. Williamson, W.H. Hall, Acta Metallurgica 1 (1953) 22. [18] H.P. Klug, L.E. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, John Wiley and Sons, New York, 1974. [19] L. Chen, W. Lengauer, P. Ettmayer, K. Dreyer, H.W. Daub, D. Kassel, International Journal of Refractory Metals and Hard Materials 18 (2000) 307. [20] P. Wally, S. Binder, P. Ettmayer, W. Lengauer, Journal of Alloys and Compounds 230 (1995) 53–57. [21] P. Ettmayer, H. Kolaska, K. Dreyer, Powder Metallurgy International 23 (1991) 224. [22] T.B. Massalski, Binary Alloy Phase Diagrams, American Society for Metals, 1986. [23] L. Backerud, B. Carlsson, R. Oskarsson, M. Mikus, Scandinavian Journal of Metallurgy 3 (1974) 225. [24] E.R. Stover, The Nickel Titanium Carbon System, Massachusetts Institute of Technology, 1956. [25] H. Pastor, Materials Science and Engineering A 106 (1988) 401–409. [26] G. Gille, B. Szesny, K. Dreyer, H. van den Berg, J. Schmidt, T. Gestrich, G. Leitner, International Journal of Refractory Metals and Hard Materials 20 (2002) 3.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Consolidation of (Ti,Mo)(C,N)–Ni cermets by glass encapsulated hot isostatic pressing N. Rodriguez ∗ , J.M. Sanchez, M. Aristizabal CEIT and TECNUN, Paseo de Manuel Lardizabal 15, 20018 San Sebastian, Spain
a r t i c l e
i n f o
Article history: Received 4 October 2010 Received in revised form 2 February 2011 Accepted 2 February 2011
Keywords: TiMoCN cermets Thermal analysis Hot isostatic pressing Uncombined carbon Grain size refinement
a b s t r a c t Microstructural evolution of (Ti,Mo)(C,N)–Ni cermets consolidated by hot isostatic pressing (HIP) has been analysed. HIP processing allows full densification of these materials at lower temperatures than those normally employed in vacuum sintering cycles (VS). Solution and reprecipitation phenomena are limited and a nanometric fraction of (Ti,Mo)(C,N) grains is retained in the microstructure leading to a significant increase in hardness with respect to vacuum sintered materials (from 11 to 14 GPa). HIP-ed cermets show more tendency to uncombined carbon precipitation than those obtained by VS. In studied systems, carbon precipitation can be related either to an excess of carbon in the initial mixture of powders or to destabilization of carbonitride phase during HIP. Control of the C/N ratio has been carried out by the adequate selection of powder mixtures and the design of the thermal treatments. It has been proved that free carbon in these cermets can be avoided including presintering cycles under hydrogen before encapsulation. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Titanium carbonitride based cermets are materials with high potential for tribological use. These low density composite materials combine high hardness with relatively high toughness and excellent chemical stability, properties that ensure tool life and tool enhanced performance [1–4]. Among the strategies used for improving the wear resistance of titanium carbonitride cermets, microstructural refinement is one of the most thoroughly investigated. Different technologies have been applied for this purpose starting from the use of ultrafine powders in the powder mixtures (even in the nanometric range), to the application of assisted sintering methods to reduce the densification temperatures (i.e. hot isostatic pressing (HIP), spark plasma sintering (SPS), etc.) [5–8]. Apart from powder processing issues (i.e. agglomeration, etc.), the main difficulties reported during consolidation of different ultrafine grained materials by HIP or SPS are related to the incomplete carbothermal reduction of the powder oxides and to the poor wettability of such oxides in contact with the liquid phase [9,10]. In Ti(C,N) cermets case, a high tendency towards graphite precipitation has been identified as a critical issue in HIPed samples. For its detrimental effect on mechanical properties, the presence of free carbon in the microstructure should be avoided and, in this sense, a closer control on the composition during HIP
∗ Corresponding author. Tel.: +34 943 21 28 00; fax: +34 943 21 30 76. E-mail address: [email protected] (N. Rodriguez). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.02.008
is essential. On the other hand, microstructural development during vacuum sintering at high temperature of these materials is well established [1]. The dissolution and reprecipitation of the carbides in binder phase lead to the formation of a typical core-rim structure, where N-rich nuclei appear surrounded by C-rich outer shells with a well defined interface between them [1,11,12]. However, effect of HIP processing on microstructural development of these materials is not widely described yet. Under these considerations, this work is aimed at analysing densification processes and microstructural evolution during consolidation of titanium carbonitride cermets by hot isostatic pressing having vacuum sintered materials as a reference. 2. Materials and methods According to Tables 1 and 2, three different alloys were selected, C1 based on a mixture of carbides, TiC and Mo2 C, and a carbonitride, Ti(C0.7 ,N0.3 ), and C2 and C3 based on two different (Ti,Mo)(C,N) prealloyed powders. In the case of cermet 2, a (Ti,Mo)(C,N) prealloyed powder provided by Treibacher and a nickel powder provided by Inco were mixed. On the other hand, cermet 3 is a commercial mixture of powders provided by Ugicarb consisting of a different (Ti,Mo)(C,N) prealloyed powder, nickel and molybdenum. The terms “grade A” for the (Ti,Mo)(C,N) powder in cermet 2 and “grade B” for the (Ti,Mo)(C,N) powder in cermet 3 are referred to the powder grades used in each case. In all cases Ni was employed as metallic binder. The characterization of the pre-reacted powders by XRD in a previous investigation showed significant differences
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N. Rodriguez et al. / Materials Science and Engineering A 528 (2011) 4453–4461 Table 3 Thermal treatments during consolidation of the cermets.
Table 1 Composition of powder mixtures (wt%). Ref.
Ti (C0.7 ,N0.3 )
(Ti,Mo)(C,N)
TiC
Mo2 C
Mo
C
Ni
C1 C2 C3
39.9 – –
– 70(A) 70(B)
8.1 – –
20.6 – –
– – 3.9
1.3 – –
30 30 26.1
Table 2 Composition of powder mixtures (at%). Ref.
Ti
Mo
C
N
Ni
C1 C2 C3
31.6 32.6 32.2
8.0 7.7 9.7
32.1 31.8 34.2
7.9 7.6 6.0
20.3 20.4 17.9
in their composition. Thus, the carbonitride powder used in cermet C2 (grade A) is actually a combination of (Ti,Mo)(C0.7 ,N0.3 ) and (Ti,Mo)(C0.93 ,N0.07 ) phases, whereas the carbonitride powder grade B (in cermet C3) is quasi-monophasic with a (Ti,Mo)(C0.85 ,N0.15 ) stoichiometry [13]. In relation to atomic composition of these mixtures, cermets C1 and C2 are quite similar while, for commercial “ready to press” C3 powder, lower contents of nickel and nitrogen are remarkable (see Table 2). The oxygen content of the C1 powder was the highest at 1.77 wt%, being the oxygen content of C2 and C3, 0.8 and 0.9 wt% respectively. Milling-mixing step of initial powders was carried out in a planetary mill for 7 h in hexane with 3.5 wt% of paraffin as organic binder. Afterwards, powder slurries were dried at 90 ◦ C in an isothermal bath for 1 h. The melting points of these systems were determined by differential scanning calorimetry. Tests were carried out at 1450 ◦ C for 1 h under vacuum (≈10−1 mbar) with a heating and cooling rate of 10 ◦ C/min. Green compacts were obtained by double action pressing at 160 MPa using a cylindrical steel die and punches. Sintering to full density was carried out by glass encapsulated hot isostatic pressing (GEHIP). During the encapsulation method pressed samples were coated with h-BN and subsequently introduced in Pyrex glass tubes [14]. GEHIP was carried out at two different temperatures, 1250 ◦ C and 1350 ◦ C, with 200 MPa of applied pressure for 45 min (using Ar as the pressing fluid) [14,15]. In some cases, specimens were previously presintered in H2 either at 850 or 1200 ◦ C. Table 3 summarises the conditions of thermal treatments carried out to densify the samples. In order to compare HIP with a more conventional processing route of these materials, results obtained in a previous work for samples densified by presintering in H2 and subsequent vacuum sintering cycles are reported [16]. Standard ISO 3369 was used for density measurements using ethylic alcohol instead of distilled water. Since complex chemical reactions take place during sintering of the cermets, theoretical
Ref.
GH(1) GH(2) GH(3) GH(4)
Thermal treatments Presintering in H2
Encapsulation
HIP
– – 850 ◦ C, 30 min 1200 ◦ C, 30 min
Yes Yes Yes Yes
1250 ◦ C, 200 MPa 1350 ◦ C, 200 MPa 1350 ◦ C, 200 MPa 1350 ◦ C, 200 MPa
density cannot be estimated considering a straightforward relationship between the chemical composition and the theoretical density of initial compounds. For practical reasons in this work, theoretical density was established in 6.5 g/cm3 taking into account Archimedes density values corresponding to fully dense samples obtained by vacuum sintering (<0.02 vol% of porosity according to ISO 4505) [16]. Phase identification was carried out by X-ray diffraction (XRD) (with Ni-filtered CuK␣ radiation). Additionally, the mean crystallite size of (Ti,Mo)(C,N) crystals in the nanometric range was characterized by using the Williamson–Hall approach on XRD peak broadening data [17]. This method assumes that intrinsic broadening (strain broadening and effects of small crystallite size) can be approximated considering the profiles to be Cauchy–Cauchy. Instrumental broadening correction was calculated assuming Gaussian–Gaussian (GG profiles) [18]. To determine the total carbon content of the samples a LECO CS-200 equipment was employed. The measurement technique is based on the combustion of the sample in an oxygen atmosphere to convert the total carbon to CO2 . Then the sample gas flows into a non-dispersive infrared (NDIR) detection cell. The N and O contents were measured with a LECO-TC-400 equipment employing the inert gas fusion method. In this case, the principle of operation is based on the fusion of the sample in a high-purity graphite crucible. The oxygen in the sample reacts with the carbon from the crucible to form carbon monoxide (CO); nitrogen is released as molecular nitrogen (N2 ). Oxygen is detected as either carbon dioxide (CO2 ), CO or both, using infrared (IR) detection, while nitrogen is determined using a thermal-conductivity (TC) cell. In each case a minimum of three measurements per sample were carried out. The sintered specimens were ground and polished down to 1 m diamond paste for microstructural analyses, which were carried out by optical microscopy (OM) and field-emission scanning electron microscopy (FE-SEM). The presence and distribution of porosity were evaluated considering the criteria specified in standard ISO 4505. Compositional analyses were performed by energy dispersive spectroscopy (EDS) in the FE-SEM. Vickers hardness was also measured by applying a load of 10 kgf (ISO 3878).
Fig. 1. DSC curves of the alloys studied: (a) heating and (b) cooling curves.
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3. Results and discussion 3.1. Liquid phase formation Results of calorimetric experiments are shown in Fig. 1. During heating (see Fig. 1(a)), cermets C1, based on a mixture of carbides, and C2, based on a (Ti,Mo)(C,N) prealloyed powder, present an endothermic peak at temperatures near to 1290 ◦ C. At higher temperatures, a second peak associated with liquid phase formation is clearly observed in alloy C1, whereas an exothermic peak is detected in the case of C2. Finally, in cermet C3, with a monophasic prealloyed carbonitride, only one peak is observed showing that melting occurs at a temperature near to 1350 ◦ C (see Table 4). The results confirm that liquid formation depends on the chemical composition of starting powders. Thus, the lowest melting points correspond to compositions containing TiC (C1) or C-rich carbonitride phases (alloy C2) where the eutectic found is close to that occurring in the pseudo-ternary TiC–“MoC”–Ni (at 1280 ◦ C) [19] being noteworthy that these data are consistent with XRD analyses of raw prealloyed powders and with the presence of (Ti,Mo)(C0.93 ,N0.07 ) phase in carbonitride grade A used in composition C2. On the other hand, melting phenomena at higher Table 4 Onset and peak temperatures determined by DSC. Cermet
C1 C2 C3 a
DSC, heating
DSC, cooling
Tonset (◦ C)
Tpeak (◦ C)
Tonset (◦ C)
Tpeak (◦ C)
1270 1323 1288 1364a 1339
1283 1343 1298 1397a 1353
1338 1359 1335 1382 1344 1313
1331 1355 1323 1371 1332 1286
Exothermic peak.
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temperatures observed in cermets C1 and C3 are related to the presence of nitrogen in the lattice of the carbonitride phase. According to data published for similar systems, as the nitrogen content of the carbonitride phase increases, melting occurs at higher temperatures [3,19,20]. Denitridation of high nitrogen-containing alloys during heating in vacuum could explain small differences between the melting point of composition based on Ti(C0.7 ,N0.3 ) (C1), established at 1343 ◦ C, and that of composition based on prealloyed (Ti,Mo)(C0.85 ,N0.15 ) (1353 ◦ C) (see Table 4) [21]. As composition C2 is concerned, the absence of the last peak around 1350 ◦ C is explained bearing in mind the exothermic phenomenon previously described, probably related to the formation of an intermetallic compound (exothermic peak point at 1397 ◦ C according to Table 4). The presence of the exothermic peak might be hiding the melting point related to carbonitride phase. In the carbonitride-based systems, intermetallic phases might be a consequence of a nitrogen partial pressure below the equilibrium pressure of the carbonitride phase for a long time during sintering. In these conditions, it can be assumed that considerable amounts of Ti are dissolved in the binder alloy, so that the precipitation of compounds like Ni3 Ti could be possible [21,22]. In detail, the thermodynamic description of Ti–C–Ni system establishes the equilibrium reaction: liquid ↔ Ni + TiC + Ni3 Ti at 1295 ◦ C [23,24]. Therefore, the first Table 5 Chemical composition of ␣ -phase and binder phase of cermets C1, C2 and C3 densified by consecutive presintering under H2 (850 ◦ C, 30 min) and vacuum sintering (1450 ◦ C, 1 h) (expressed as weight percentage). EDS analyses were carried out in transmission electron microscope. Cermet
C1 C2 C3
Carbonitride phase (rim)
Binder phase
Ti
Mo
Ti
Mo
73 ± 2 70 ± 5 68 ± 5
27 ± 2 30 ± 5 32 ± 5
4.5 ± 0.4 3.7 ± 0.1 3.4 ± 0.3
10.3 ± 0.6 10.1 ± 0.9 11.2 ± 0.8
Fig. 2. Microstructures of cermets based on prealloyed powders, C2 and C3, consolidated by treatments GH(1) and GH(2) (see Table 3).
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melting point of cermet C2 (1298 ◦ C) further supports the idea of subsequent Ni3 Ti formation. Although, on heating, clear differences are found in the DSC curves of the different compositions, the behaviour observed on cooling is similar. Strongest exothermic peak is observed at 1330 ◦ C in all cases (see Fig. 1(b)). The described trend can be related to the homogenization of the chemical composition of the cermets during sintering. This is consistent with the TEM analysis carried out on these samples (see Table 5). Composition analyses show no significant differences in the composition of the binder and the ␣ -phases on sintered materials. It can be assumed that the dissolution of metal atoms from hard phases into the Ni during sintering evolves to a state close to the thermodynamic equilibrium as evidenced by the similarity in the final composition of three cermets despite the different starting compounds used [11]. Taking into account these results and with the objective of analysing the effect of the liquid phase on the processing of the cermets, two different temperatures were selected for HIP cycles: 1250 and 1350 ◦ C respectively. 3.2. Densification In the microstructure of the studied cermets, uncombined carbon related porosity is observed in first hipping experiments at 1250 and 1350 ◦ C respectively. Optical micrographs show rounded
Fig. 3. EDS analysis of uncombined carbon presence in cermet C3 densified in solid state by GH(1). The carbon rich zones where different EDS analyses were carried out are pointed by red arrows. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
Fig. 4. Relative change in the carbon content of specimens after thermal treatments (see Table 3).
porosity in most of the cases making the identification of free carbon as described in ISO 4505 difficult (see Fig. 2(a)–(c)). However, as can be observed in Fig. 3, the characterization of the samples by FE-SEM at a higher magnification confirms the presence of carbon rich areas even when the densification takes place in solid state (EDS analyses pointed by red arrows). In addition, changes in the morphology of the precipitation depending on the HIP temperature can be identified. Thus, in the presence of a liquid phase, the C-rich regions appear in clusters (Fig. 2(d)), whereas, in solid state, free carbon zones are finer and homogeneously distributed in the microstructure (Fig. 2(a)–(c)). These data suggest that, apart from being a nucleation and growth process, uncombined carbon presence in (Ti,Mo)(C,N)–Ni cermets is enhanced by the presence of the liquid phase. This is likely related to the higher solubility of carbon bearing phases in the liquid, which increases with temperature [3]. In order to control the total carbon content of the specimens, a presintering step under hydrogen has been introduced in the processing route of the cermets before encapsulation [13]. Fig. 4 shows the reduction in carbon content of the samples subjected to this thermal treatment (ref. GH(3) and GH(4) according to Table 3). As can be observed, presintering at a higher temperature gives rise to increments in the level of decarburization experienced by the specimens, i.e. presintering parameters can be adjusted to control the carbon losses induced in this step. In cermets C1 and C2 carbon losses due to presintering are more pronounced, while cermet C3 shows a more stable behaviour being less affected by this treatment. Hence, cermets containing free carbon and carbonitride powders with high C/N ratios lead to a stronger interaction with H2 inducing higher carbon losses [13]. According to the observations reported by several authors, this behaviour can be explained considering the high solubility of these phases in Ni-rich binder phase, thus resulting in a greater reactivity of available carbon with the hydrogen in the atmosphere [20]. On the other hand, the decarburization observed in commercial grade C3, even without presintering, can be related to a high level of initial oxidation of the powder mixture due to a long storage time. The carbothermal reduction of NiO at low temperature during encapsulation cycle, that takes place under vacuum (10−1 mbar), explains the diminution in carbon levels determined in treatments GH(1) and GH(2). In any case, fully prealloyed carbonitride based cermet, C3, shows a more stable behaviour with no significant changes in carbon content with the treatment atmosphere. Furthermore, the Archimedes density values can be related to the final carbon contents previously described (see Fig. 5). Thus, in cermets C1 and C2, carbon losses result in a density increase, whereas, for composition C3, with the slightest modification in carbon content, density remains unchanged. The metallographic inspection of the samples by optical microscopy also confirms this
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Table 6 Carbon, nitrogen and oxygen final contents of cermets obtained by different methods (expressed as weight percentage) (Cf , Nf and Of ). Ci and Ni are initial contents of the powder mixtures (C1, C2 and C3, GEHIP(4): see Table 3; C2, PR1200 + VS: cermet C2 obtained by presintering under H2 (1200 ◦ C, 30 min) and subsequent vacuum sintering (1450 ◦ C, 1 h)).
Fig. 5. Density of the samples obtained by different methods (see Table 3).
tendency. As can be observed in Fig. 6, micrographs show the total absence of uncombined carbon in cermets based on prealloyed powders densified by GH(4) (see Table 3), while a reduction in Ctype porosity is encountered in cermet C1 densified in the same cycle. Moreover, the level of porosity of the samples obtained by GH(4) is established in 0.6 vol% in all cases (according to ISO 4505). Therefore, the presintering step is effective controlling the carbon content and the phases present in the microstructure; this leads to an increase in the density and a reduction in the porosity measured in the specimens. Presintering-HIP processing allows full densifi-
Ref.
Ci
Ni
Cf
Nf
Of
C1, GEHIP(4) C2, GEHIP(4) C3, GEHIP(4) C2, GEHIP(2) C2, PR1200 + VS
9.7 9.6 10.2 9.6 9.6
3.0 2.7 2.2 2.7 2.7
8.2 ± 0.03 8.15 ± 0.03 8.84 ± 0.1 9.56 ± 0.07 7.46
3.57 ± 0.25 3.11 ± 0.14 3.0 3.20 ± 0.09 2.82 ± 0.09
0.82 ± 0.11 0.93 ± 0.04 1.2 0.50 ± 0.07 0.58 ± 0.07
cation of these cermets at a lower temperature than that employed in vacuum sintering cycles (i.e. 1450 ◦ C) [16]. On the other hand, carbon rich areas still present in cermet C1, even with a presintering cycle in H2 at 1200 ◦ C, are related to the higher proportion of carbides and free carbon in this composition. At the hipping temperature (in this case 1350 ◦ C), the solubility of TiC and Mo2 C phases in the liquid phase is very high compared to nitrogen-containing phases [3,19]. On cooling, a certain amount of Ti and Mo remains alloyed in the Ni-rich binder phase, whereas the rest precipitates as the ␣ phase onto the undissolved carbonitride grains. With respect to the final content of interstitial elements, HIP-ed cermets show higher contents of carbon, oxygen and nitrogen than those obtained by vacuum sintering (see Table 6). Higher levels of carbon and oxygen are due to the modification of the gas–solid
Fig. 6. Evolution of the porosity of the alloys as a function of the thermal treatments during consolidation (see Table 3).
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Fig. 7. XRD patterns obtained in the surface of C2 densified by GH(2) and GH(4) compared with that obtained in the interior of alloy C2 densified by GH(2) (solid circles correspond to TiN (ICDD card: 06-0642) and open triangles to Ti(C0.7 ,N0.3 ) (ICDD card: 42-1489)).
reactions induced by the glass encapsulation process, since, once the capsule is closed, the carbothermal reduction of powder oxides stops [8]. On the other hand, the increase observed in the nitrogen content could be related to the interaction with the h-BN layer used as a barrier to avoid the infiltration of molten glass in the compacts. XRD analyses of the surface of HIP-ed samples evidence the presence of TiN regardless of the thermal treatment applied (GEHIP or presintering-GEHIP) or the initial composition of the cermet. To illustrate the nitridation of the specimens during HIP, diffraction patterns obtained at the surface of different samples are compared with the pattern obtained at the interior of cermet C2 (see Fig. 7). Furthermore, EDS analyses of alloy C2 support this idea (see Fig. 8). An increase in the N peak intensity is observed in the spectra obtained at the surface when compared with the results observed in the interior (compare Spectrum 1 (surface) with Spectrum 2 (interior) in Fig. 8). As described in previous works, GEHIP processing increases the uncombined carbon presence in Ti(C,N) based cermets [8]. With prealloyed powders of the type (Ti,Mo)(C,N) a similar behaviour has been observed which is likely related to an excess of carbon
Fig. 9. Micrographs of composition C2 by optical microscopy: (a) GEHIP-ed sample and (b) vacuum sintered material. Total carbon content of the GEHIP-ed material is 8.5 wt% and total carbon content of the vacuum sintered material is 9.3 wt%.
content in the powder mixtures and to a high nitrogen activity inside the glass capsules during HIP. The absence of uncombined carbon in samples processed by vacuum sintering with a high total carbon content, even higher than that of the HIP-ed specimens, is in agreement with the hypothesis that the carbon content is not
Fig. 8. EDS analyses at the surface of cermet C2 densified by GH(2) (see Table 3). The Spectrum 1 evidences the presence of nitrogen at the superficial layer.
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Fig. 10. Nitrogen partial pressure vs. atomic fraction of nitrogen in the carbonitride phase at 1350 ◦ C [25].
the only variable that determines precipitation in these systems. In Fig. 9 optical micrographs of alloy C2 show uncombined carbon in GEHIP-ed sample, whereas in the vacuum sintered, only A-type porosity is encountered (according to ISO 4505). According to the domain of stability of Ti(C,N) phase suggested by Pastor [25], free carbon presence is possible by destabilization of the carbonitride phase under the combined effect of high carbon potential and high isostatic pressure. Inside the glass capsule (typically closed under vacuum, aprox. 10−1 mbar) nitrogen evolves from the compacts as the carbonitride phase is heated. The decomposition proceeds increasing the N2 pressure inside the capsule. As isostatic external pressure increases before reaching the eutectic point of the alloys, it is possible that, under such conditions, the nitrogen activity reaches the equilibrium condition for free carbon precipitation (see Fig. 10). The high nitrogen content of HIP-ed samples is consistent with this assumption. The highest nitrogen concentration found in cermet C1 is also in agreement with the higher tendency to free carbon precipitation in this alloy (see Table 6 and Fig. 6). As the nitrogen content on the (Ti,Mo)(C,N) cubic lattice increases, there are fewer sites for carbon atoms, hindering the incorporation of this species to the carbonitride phase. 3.3. Microstructural evolution As can be observed in FE-SEM micrographs, in samples densified in solid state (HIP cycles carried out at 1250 ◦ C), the most relevant microstructural feature is the presence of nanometric carbonitride grains generated during milling step and retained in the microstructure (Fig. 11(a), see blue arrows). Solution and reprecipitation phenomena are limited due to the low temperature employed; consequently, carbonitride phase grain growth is inhibited. The incipient formation of the ␣ phase can also be detected at the surface of larger ceramic grains (see Fig. 11(b), blue arrows). In addition, a higher degree of coalescence of ceramic grains, as well as, a heterogeneous distribution of binder phase where the combination of both characteristics results in the presence of aggregates in the microstructure (see Fig. 11(c)) is observed. On the other hand, nanometric carbonitride grains have almost disappeared from the microstructure of cermets processed in liquid phase (HIP at 1350 ◦ C), which also present a more homogeneous distribution of the binder phase around (Ti,Mo)(C,N) grains (see Fig. 12). A higher fraction of ␣ phase and a more homogeneous composition in the ceramic cores are also observed in these images (as an example of these effects Figs. 11(b) and 12(b) can be compared). This microstructural evolution is explained considering an Ostwald ripening process in which the finest carbonitride grains are preferentially dissolved in the liquid leading to both,
Fig. 11. Microstructural characteristics of cermet C2 densified by sintering in solid state (GH(1) according to Table 3): (a) nanometric carbonitride grains (pointed by blue arrows), (b) ␣ formation (pointed by blue arrows) and (c) coalescence of ceramic grains and binder pooling (red and blue arrows respectively). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)
the rim formation and the elimination of nanometric grains. Furthermore, the more homogeneous distribution of the binder phase is due to the good wetting behaviour of the liquid phase in the system. Diffraction peak line broadening observed only for the case of samples processed in solid state (GH(1) according to Table 3) also confirms the elimination of nanometric carbonitride grains when
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Fig. 12. BSE-FEG SEM images of cermets C1, C2 and C3 densified by GH(4) according to Table 3.
Fig. 13. (a) Comparison of peak widths corresponding to cermet C3 densified by vacuum sintering at 1450 ◦ C (VS) and by GE-HIP in solid and liquid phases: GH(1) and GH(4) respectively (see Table 3), and (b) plot of the intrinsic broadening for peaks corresponding to the carbonitride phase in cermets C2 and C3 processed by GH(1) (see Table 3) (Williamson–Hall representation).
the liquid appears; no difference is observed between the diffraction peak of the vacuum sintered sample (at 1450 ◦ C) and that corresponding to HIP-ed sample in the presence of a liquid (GH(4) according to Table 3) (see Fig. 13(a)). The broadening observed in solid state has been used to calculate the mean crystallite size of carbonitride crystals in cermets densified in solid state being 92 and 82 nm respectively for compositions C2 and C3 (see linear fittings in Fig. 13(b)).
The abovementioned microstructural changes affect hardness. In Fig. 14 it can be observed that the highest hardness values correspond to samples processed in solid state, i.e. GH(1). On the other hand, the hardness loss of around 100 kg/mm2 observed when the samples are densified in the presence of a liquid, GH(2), is due to both, free carbon precipitation and microstructural coarsening; meanwhile, the hardness increase for samples processed with an intermediate presintering step is related to the
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materials is due to the good wetting behaviour of the liquid phase which also activates dissolution and reprecipitation phenomena. As expected, these phenomena lead to the growth of the finest carbonitride grains and to the formation of ␣ shells surrounding the undissolved carbonitride cores. Acknowledgments FMD-CARBIDE is gratefully acknowledged for their financial and technical support and for providing the raw materials. N. Rodriguez would also like to thank the Ministerio de Educación y Ciencia within the framework of Programa Torres Quevedo for the financial support given to this work. References Fig. 14. Vickers hardness (10 kg) of samples consolidated by different methods (see Table 3) (PR1200 + VS corresponds to samples densified by presintering at 1200 ◦ C under H2 and subsequent vacuum sintering at 1450 ◦ C, 1 h [16]).
reduction of uncombined carbon. Cermets obtained by presintering at 1200 ◦ C and subsequent GEHIP at 1350 ◦ C, i.e. GH(4), show hardness values around HV10 1320 in the case of compositions C1 and C2 and HV10 1390 for the C3, similar to those reported for GH(1) (HV10 1300 and 1430 for cermets C1 and C3 respectively). The higher hardness of cermet C3 is due to its slightly lower binder content (26.1 wt% Ni vs. 30 wt% Ni in cermets C1 and C2), since the C/N ratio obtained after processing is very similar for the three compositions. Moreover, samples processed by GH(4) show hardness values more than 100 kg/mm2 higher than those obtained by conventional vacuum sintering (see data circled in Fig. 14); the differences being related to the microstructural coarsening at 1450 ◦ C during vacuum sintering (100 ◦ C higher than the temperature of GEHIP). Hardness values of cermets obtained by GH(4) are similar to those reported by Gille et al. for WC–Co compositions with a similar metallic binder content (21 vol%) and a WC grain size between 0.5 and 0.8 m [26]. 4. Conclusions Fully dense pseudo-binary (Ti,Mo)(C,N)–Ni cermets have been obtained by combining presintering under hydrogen and glass encapsulated hot isostatic pressing (GEHIP). In these materials GEHIP processing induces a higher tendency to free carbon precipitation than vacuum sintering, this phenomenon being controlled by introducing a presintering step under hydrogen before encapsulation. Cermets consolidated in solid state present microstructures characterized by heterogeneous distribution of the binder phase and the presence of nanometric carbonitride grains. The microstructural homogenization observed in liquid phase sintered
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