- Email: [email protected]
Spark plasma sintering of TiC particle-reinforced molybdenum composites
Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
Contents lists available at SciVerse ScienceDirect
Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Spark plasma sintering of TiC particle-reinforced molybdenum composites R. Ohser-Wiedemann a,⁎, Ch. Weck a, U. Martin a, A. Müller a, H.J. Seifert b,⁎ a b
Institute of Materials Science, TU Bergakademie Freiberg, Gustav-Zeuner-Str. 5, 09599 Freiberg, Germany Institute for Applied Materials, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, Bldg. 681, 76344 Eggenstein-Leopoldshafen, Germany
a r t i c l e
i n f o
Article history: Received 19 July 2011 Accepted 1 December 2011 Keywords: Molybdenum Spark plasma sintering Composites Microstructure Recrystallization
a b s t r a c t In order to improve the recrystallization resistance and the mechanical properties of molybdenum, TiC particle-reinforcement composites were sintered by SPS. Powders with TiC contents between 6 and 25 vol.% were prepared by high energy ball milling. All powders were sintered both at 1600 and 1800 °C, some of sintered composites were annealed in hydrogen for 10 h at 1100 up to 1500 °C. The powders and the composites were investigated by scanning electron microscopy and XRD. The microhardness and the density of composites were measured, and the densification behavior was investigated. It turns out that SPS produces Mo–TiC composites, with relative densities higher than 97%. The densification behavior and the microhardness of all bulk specimens depend on both the ball milling conditions of powder preparation and the TiC content. The highest microhardness was obtained in composites containing 25 vol.% TiC sintered from the strongest milled powders. The TiC particles prevent recrystallization and grain growth of molybdenum during sintering and also during annealing up to 10 h at 1300 °C. Interdiffusion between molybdenum and carbide particles leads to a solid solution transition zone consisting of (Ti1 −x Mox)Cy carbide. This diffusion zone improves the bonding between molybdenum matrix and TiC particles. A new phase, the hexagonal Mo2C carbide, was detected by XRD measurements after sintering. Obviously, this phase precipitates during cooling from sintering temperature, if (Ti1 −x Mox)Cy or molybdenum, are supersaturated with carbon. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Due to their high melting point, pure molybdenum and its alloys are used for high temperature applications in a variety of industries. However, application temperature is limited by the onset of recrystallization, which can be constraint by addition of fine disperse distributed particles, like carbides and oxides in the molybdenum alloys TZM (titanium–zirconium–molybdenum) or MHC (molybdenum–hafnium– carbon). These molybdenum alloys contain only small quantities of particles [1]. Another way to improve the recrystallization resistance and the mechanical properties at higher temperatures of refractory metals is the creation of metal matrix composites. Yoo et al. [2] investigated the recrystallization behavior of molybdenum sheets which were mechanical alloyed with 1.0 wt.% TiC and 0.2 wt.% C. At first, the powders were sintered by hot isostatic pressing (HIP) at 1300 °C, 200 MPa and 5 h. After it, the specimens were forged and rolled. Transmission electron microscopy investigations showed that very fine TiC particles are homogeneously distributed in the molybdenum phase and are pinned on dislocations and grain boundaries. The micro hardness increases up to 500 HV. Moreover,
⁎ Corresponding author. Tel.: + 49 3731 39 2647; fax: + 49 3731 39 3657. E-mail address: [email protected] (R. Ohser-Wiedemann). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.12.001
the formation of the molybdenum carbide Mo2C was detected in these composites. After annealing at temperatures between 1400 and 2000 °C for 10 h in flowing H2 atmosphere, grain and particle growth occur and the microhardness decreases. However, some small particles are still pinned on the dislocations and grain boundaries even at annealing temperature of 1800 ° C. Therefore, the authors assume, the recrystallization was not complete up to this temperature. At temperatures higher than 1800 °C a coarse polycrystalline microstructure with pores was observed and the microhardness decreased considerably. Mo–TiC composites with various TiC fractions are used as material for nuclear applications [3–6]. Mo–30 vol.% TiC has been proven useful as a composite with high mechanical strength at temperatures higher than 1000 K and up to 1900 K under fast neutron flux. Mixed Mo–TiC powders were successful compacted by HIP at 1600 °C, 2 h under a pressure of 160 MPa [3–5], but until now the recrystallization behavior of the composites was not investigated. After sintering a core-rime-structure was formed surrounding the TiC particles. The rim consist of the solid solution (Ti1 −x Mox)Cy which has the same cubic structure as TiC. This suggests that diffusion occurs in the carbide during sintering. The formation of Mo2C carbides was not detected in these investigations. Takida et al. [7] successfully consolidated ZrC particle dispersed molybdenum (1.4 vol.% ZrC) by spark plasma sintering. The particles were dispersed and no particle coarsening was observed after
2
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
annealing at 1797 °C for 1 h. The grain growth during annealing was strongly limited and the material exhibit excellent mechanical properties. The authors in ref. [8–12] investigated W-TiC or W-ZrC composites with particle contend up to 40 vol.% by hot pressing (HP) at 2000 or 2200 °C for 1 h. The results show that mainly fine particles prevent the tungsten grain growth. The mechanical properties at room temperature are improved depending on both, particle fraction and particle size. Similar to Mo–TiC composites, core-rime-structures were formed by diffusion of tungsten into TiC or ZrC. The formation of (Ti1 −xWx)Cy in W-TiC composites improve the strengthening effect of TiC [9]. In W-ZrC composites a hexagonal W2C phase with the same crystal structure like Mo2C (space group of both carbides P63/mmc [15]) was detected by XRD investigations. The phase was observed in only such composites where the carbide particle sizes are less than 1.7 μm. Obviously, this phase precipitates from the (Zr1 −xWx)Cy solid solution during cooling down from sintering temperature, because the solubility of tungsten in (Zr1 −xWx)Cy reduces with decreasing temperature [12]. The motivation of the presented work is to prepare Mo–TiC composites with low TiC volume fractions (6, 12.5 and 25%) using the spark plasma sintering process. During the SPS process, the dies and powders are heated by a pulsed direct current. This is in contrast to hot pressing or hot isostatic pressing and offers the possibility to sinter at higher heating rates, lower sintering temperatures and within shorter sintering times. To get a homogeneous TiC particle distribution in molybdenum, the Mo–TiC powder mixtures were produced by high energy ball milling with different milling parameters. Therefore, a further topic of the work is to investigate how the different milling parameters control the microstructure and hardness of sintered composites. Molybdenum, considered as the matrix material, was systematically investigated according to its SPS sintering behavior in ref. [13]. First results of the compaction of Mo–TiC composites were published in ref. [14]. 2. Experimental Pure molybdenum powders (99.95 wt.%, Plansee Metall GmbH, Austria; Fisher number 3 to 5 μm) were blended either with 6, 12.5 or 25 vol.% TiC powders (TiC STD 120, H.C. Starck GmbH, Germany; Fisher number 1 to 1.5 μm). Afterwards, powder mixtures were milled in a planetary ball mill (Fritsch “Pulverisette 6”) under air. Grinding bowls and balls were made of steel; the ball diameter was 10 mm. To prevent heating-up during milling, the milling was interrupted with a break of 15 min after a milling time of 15 min. The total milling time (without breaks) was 5 h. The ball powder ratio (BPR) and the rotating speed have been varied to change the milling energy. Table 1 gives an overview of all varied milling parameters. X-ray diffraction (XRD) spectra of all Mo–TiC powder mixtures and sintered composites were recorded with a powder diffractometer using Cu-Kα radiation and a graphite monochromator in the diffracted beam. The PDF data base [15] was used for qualitative phase analysis. Using the Rietveld-analysis BGMN (version 3.4.27) [16], the crystallite size (domain size) and micro strain values were derived from the integral breadths of the molybdenum peaks. The broadening due to the instrumental aberrations was separated by a
Table 1 Milling conditions and marking of the prepared Mo–TiC powders (total milling time 5 h). Ball powder ratio
Rotation speed
Powder identification marks
5:1 10 : 1 10 : 1
min− 1 200 200 300
Mo–25% TiC A B C
Mo–12.5% TiC – D –
Mo–6% TiC – E –
Monte-Carlo simulation. The dislocation density was roughly estimated from the mean squared strain, where an approach of Krivoglaz [17, page 357 ff.] was used. The approach is applicable only for cubic metals and their alloys, which exhibits high dislocation densities. Moreover, elastic isotropy of the crystals is supposed. The Mo–TiC powders were filled without any additives in cylindrical graphite dies of 20 mm in diameter. The sintering runs in a FCT-HP D 25 spark plasma sinter equipment (FCT Systeme GmbH, Rauenstein, Germany). The compaction of the powder mixtures started with raising the external pressure up to 67 MPa, followed by heating up under vacuum to temperatures of 1600 or 1800 °C. The heating rate amounted to 300 K/min, and the holding time at maximum temperature was 3 min for all batches. In contrast to the sintering by HP [8–12] or HIP [2–5], the whole process time of SPS needs 15 min only. The pulse pattern of the electric current was adjusted on 2:1 (10 ms pulse time, 5 ms pause time). The compaction pressure has been decreased during the cooling step. In the following, a combination of the powder name and the sintering temperature is used as a notation of composites, e.g. “A-1600” for composite sintered at 1600 °C using powder A. The densities of all consolidated bulk specimens were measured using the Archimedes technique [18,19]. The relative densities were calculated based on the theoretical densities of molybdenum (10.28 g cm − 3) and titanium carbide (4.93 g cm − 3). The microstructure was studied by scanning electron microscopy (SEM) at cross sections (sections parallel to the acting force). The preparation steps were grinding and mechanical polishing up to 1 μm diamonds abrasive. Same cross sections were chemical etched using Murakami's reagent (a solution mixture of 10 g sodium hydroxide, 10 g potassium ferricyanide, and 100 ml of distilled water). The hardness of the Mo–TiC composites was measured at the cross sections by a Vickers microhardness tester (LECO Instrumente GmbH). The indentations were made in a pattern consisting of three parallel lines each with nine measuring points. In order to investigate the tempering resistance of microstructure, some sintered composites were annealed at 1100, 1300 and 1500 °C for 10 h in hydrogen atmosphere. After annealing the microhardness was measured again, and XRD spectra were recorded. 3. Experimental results and discussion 3.1. Characterization of ball milled Mo–TiC powders Scanning electron microscopy investigations showed that the original molybdenum powder particles are of spherical shape, some of them include small pores (see also [13]). The particles agglomerate and tend to form chains. The powders contain a fraction of fine particles with sizes b 1 μm. In contrast, the finer TiC powders consist of some big angular particles and a lot of small fractions before milling. Fig. 1 shows the morphology of powder mixtures after different milling conditions. At low BPR of 5:1 and a rotating speed of 200 rpm (powder A), the original shapes of both powders are still observable (Fig. 1a). In contrast, the morphology of powder particles changes if the BPR was increased up to 10:1 (powders B, D and E). Fig. 1b gives an impression of morphology on the example of powder B. Under these milling conditions, the molybdenum particles were plastically deformed and show a flat shape. Most of the hard brittle TiC particles were fragmented and afterwards pressed in to the softer molybdenum particles. With further increasing of milling energy, e.g. with increasing rotating speed, the original particle shapes got lost (powder C). Here both kinds of particles were cold welded and form agglomerates (Fig. 1c). After milling, all powders were investigated by XRD measurements. The phase analysis show that the powders consist only of the two phases molybdenum and titanium carbide. However, the plastic deformation during milling leads to strong microstructural
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
3
metals. Hence, the molybdenum particles in powder C are strongly strain hardened. The dislocation densities are not influenced by the TiC content. The deviations between powders B, D and E are inside the error range of XRD measurement. 3.2. Densification
changes especially inside the ductile molybdenum particles. The increase of the BPR and the rotating speed leads to higher crystal defect concentration which is connected to a strong broadening of the XRD peaks. In case of molybdenum, the crystallite size drops down drastically from more than 1000 nm in non-milled molybdenum powders to 40–48 nm in the milled powders A, B, D, E and to 24 nm in the milled powder C. Simultaneously, the dislocation densities in the powders increase. Fig. 2 depicts the strong influence of the milling conditions on dislocation densities. With increasing milling energy the dislocation density increases and reaches a maximum value of more than 10 12 cm − 2 in powders, milled at BPR 10:1 and 300 rpm. Normally, such high values are measured in strong cold worked
Fig. 3 shows the relative densities of all sintered composites in comparison to sintered pure molybdenum (not ball milled). Sintering at 1800 °C leads to relative densities of 97% or 98%, respectively. These values are a little bit higher than those of bulk molybdenum. Therefore, ball powder milling improves the densification process. The combination of plastic deformation, cold welding and fracturing during milling increases the surface area and that leads to an activation of sintering processes. At the lower sintering temperature, the relative density depends slightly on milling conditions and TiC fraction. However, the calculation of relative densities is uncertain, because it was not possible to calculate the theoretical densities of the composites exactly. As explained more in detail in paragraph 3.4, interdiffusion between molybdenum and TiC occurs during sintering. A core-rim structure is formed where the rim consist of (Ti1 −x Mox)Cy. It was not possible to estimate precisely the volume fraction of TiC and (Ti1 −x Mox)Cy and the chemical composition of (Ti1 −x Mox)Cy. Additionally, same composites contain a small fraction of Mo2C, which also influence the theoretical density. For these reasons, it was decided to calculate the relative densities based on the powder composition. By our best knowledge, the relative densities presented in Fig. 3 are lower bounds. In conventional powder metallurgy, where the sintering process starts with a green body, sintering shrinkages range from about zero to 25% [20]. Due to the uniaxial pressing at the SPS, the shrinkage occurred mainly in the pressing direction. This linear intrinsic shrinkage can be calculated from the displacement of the pressing piston which is recorded during whole sintering process. For exact conclusions from these records, the experimental data have to be corrected by subtraction of the contribution of piston displacement, which was caused by compression during applied pressure before sintering starts. This is necessary, because this first compaction depends on the applied pressure as well as on the die filling behavior of powder, and it varied in every batch. Furthermore, the thermal expansion of specimen, die and graphite sheets must be subtracted. For this correction, some baselines were determined experimentally with the same powders and sintering conditions. The procedure is described in more detail in ref. [13]. Nearly S-shaped curves were observed, if the calculated linear shrinkage is represented as a function of time (Fig. 4). This indicates that the SPS process is mainly thermal activated. Only at the beginning of sintering, linear dependence occurs with a sharp transition at about 950 °C. This transition temperature corresponds to 0.4 Tm of pure molybdenum, which is the temperature where appreciable volume diffusion starts in solid molybdenum materials. The volume
Fig. 2. Dislocation densities inside molybdenum powder particles after ball milling.
Fig. 3. Relative densities of all Mo–TiC composites in comparison to pure bulk molybdenum sintered by SPS.
Fig. 1. Morphology of Mo–25 vol.% TiC powders after ball milling under different conditions (SEM, BSE contrast); a) powder A: BPR 5:1, 200 rpm; b) powder B: BPR 10:1, 200 rpm; c) powder C: BPR 10:1, 300 rpm.
4
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
Fig. 5. SEM micrographs (BSE contrast; bright: molybdenum, dark: TiC with core-rimstructure) of Mo–TiC composites, spark plasma sintered at 1800 °C. a) specimen B1800, b) specimen C-1800. Fig. 4. Densification behavior of Mo–TiC composites depending on ball milling conditions (on the top) and on TiC content (on the bottom).
diffusion supports the particle bonding and the densification during sintering which leads to a strong rise of shrinkage above this temperature up to the end of heating up. In contrast, the shrinkage is very low during the short isothermal sintering time of 3 min. Therefore, the densification of the composites takes place mainly during heating up, and the process corresponds to the intermediate stage sintering where neighboring sinter necks grow sufficiently large to overlap [20]. Overall, linear shrinkages up to 40% were achieved, which is higher than the shrinkage of conventional sintering. This is caused by the lower initial powder compaction at SPS. The shrinkage rate is influenced by the ball milling parameters, e.g. by the defect densities inside the powder particles. The increase of milling energy shifts the curves to shorter sintering times. Therefore, the sintering activity of the milled powders rises up with increasing defect densities. The TiC fraction influences the sintering rate as well. The curves shift to shorter times for lower particle contents. Obviously, the small TiC particles, pressed into the molybdenum surface, decrease the contact areas between molybdenum and prevent the growth of molybdenum particle bonds.
particles can be observed in the specimen C-1800 (Fig. 5b). The microstructures of composites with lower TiC fraction look like similar to the microstructures of specimen B-1800 (Fig. 5a). All specimens are macroscopic crack free and have very small porosities which correspond to the measured densities. The TiC particles are cohesively bonded in the molybdenum matrix and a typical core-rim structure was formed. According to the investigation of Cédat et al. [5], the rim consist of (Ti1 −x Mox)Cy which was formed by diffusion of molybdenum into TiC during sintering. The same effect can be observed in TiC particle-reinforced tungsten composites (see chapter 1). In both cases the chemical composition changes continuously in the rim and the bonding between matrix and carbide is improved. The microhardness of all composites is presented in Fig. 6 in comparison to pure SPS molybdenum. It can be shown that the microhardness depends on both, the TiC particle fraction and the milling parameters. The composites produced from the strongest deformed powder C are of the highest hardness which is four times higher than in pure molybdenum. It can be concluded that this strong hardening is caused by a combination of the particle reinforcement and the strain hardening of molybdenum during ball milling. Finally, the dependence of the hardness on the milling conditions is a further
3.3. Microstructure and microhardness of sintered composites Fig. 5 illustrates the microstructure of composites made from the high energy milled powders B and C after sintering at 1800 °C. The SEM micrograph was obtained using backscattered electrons allowing an exact separation between the molybdenum and titanium carbide phases. The shapes of molybdenum grains inside the composites inherit the shape of the plastically deformed molybdenum particles after milling (compare Fig. 2). This indicates that neither recrystallization nor grain growth occur in the molybdenum phase at the chosen sintering conditions. This is in contradiction to the consolidation behavior of pure molybdenum powders (particle size 3–5 μm) showing a significant grain growth at these sintering temperatures. Here, a mean chord length of 19 ± 12 μm was estimated after sintering at 1800 °C [13]. A more homogeneous distribution of the fine TiC
Fig. 6. Results of micro hardness measurements of all SPS Mo–TiC composites.
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
5
argument that no recrystallization or grain growth occurs during sintering.
3.4. Defect structures and phase compositions after sintering After sintering, the integral breath of molybdenum peaks is reduced compared to that of as-milled powders (see for example Fig. 9, XRD patterns (a) and (b)). The dislocation density decreases to an order of magnitude while the crystallite size increases about sixfold. Obviously, recovery processes occur during sintering. Some of the stored internal strain energy is relieved by dislocation motion. The number of dislocations is slightly reduced by annihilation or by moving to grain boundaries. In addition, dislocations rearrange themselves into complex networks and the distortion of lattice planes goes back. However, there is no evidence for recrystallization. In this case, the dislocation density might be reduced dramatically. It can be concluded that the TiC particles also prevent the motion of high-angle grain boundaries as a requirement for recrystallization. After sintering, a new phase was detected by XRD measurements in the most composites (see Fig. 7). The phase was identified as hexagonal β-Mo2C. The occurrence of the Mo2C phase is independent of the sintering temperature, but it depends on milling conditions. The phase was not detected in the composites made from powder A. Probably, the formation of Mo2C is supported either by the higher concentration of crystal defects, which accelerated the carbon diffusion, or by the stronger fragmentation of TiC in the powerful milled powders. The Mo2C phase occurs also in TiC particle-reinforcement molybdenum sheets after sintering, forging and rolling [2] and in the molybdenum alloy TZM [21]. Furthermore, this carbide type was detected in tungsten composites as W2C [10–12]. SEM was used to look for Mo2C precipitates in our SPS composites. Mo–25vol.% TiC specimens C-1800 (with Mo2C) and A-1800 (without Mo2C) were compared. For this purpose, cross sections of both composites were etched by Murakami's reagent. Due to the different chemical potential of the phases, their surfaces were etched differently. Fig. 8 shows scanning electron micrographs of cross sections after etching. The etching reagent roughened the surface of molybdenum grains strongly. In contrast, TiC particles and the core-rim structure are attacked marginally. This creates a strong topographic contrast. Fig. 8a shows a micrograph of the composite C-1800. A cluster of TiC particles is represented, where the (Ti1 −xMox)Cy rims contact each other. Some minute particles, marked with a white circle, exist inside the (Ti1 −xMox)Cy grains. In contrast, no such particles appear in the composite A-1800 (Fig. 8b). For this reason, we assume that the observed particles are Mo2C. Based on the investigations in W-ZrC composites [10], it can be concluded that the Mo2C phase is also precipitated from carbon supersaturated (Ti1 −xMox)Cy solid solution during cooling. Additionally, thin light plates are visible in the (Ti1 −xMox)Cy phase of the C-1800 composite (Fig. 8a). Probably, these plates are low angle
Fig. 7. XRD spectra of SPS Mo–TiC composites sintered at 1600 °C. (•—TiC peaks, ∇— Mo2C peaks, not marked—Mo peaks).
Fig. 8. SEM micrographs (BSE contrast) of etched Mo–TiC composites: a) cluster of (Ti1 −xMox)Cy with precipitates (specimen C-1800), b) the surrounding of a TiC particle without precipitates (specimen A-1800).
grain boundaries decorated with minute precipitates. Finally, the exact formation of Mo2C must be a subject of a further investigation. 3.5. Tempering behavior After sintering, the composites C-1800 were annealed at 1100, 1300 and 1500 °C for 10 h in hydrogen atmosphere in order to investigate the tempering resistance of microstructure. The stability of the microstructure under thermal loading is very important for the application of these composites. The microstructure and the phase composition of all sintered and annealed specimens show no significant changes up to 1300 °C. Above 1300 °C a minor growth of molybdenum grains inside the composites occurs. Fig. 9 shows the XDR spectra of the powder C and its sintered and annealed composites. It is clearly detectable that the Mo2C carbide occurs only after sintering in the bulk specimens and not inside the powder C. The Mo2C fraction seems to decrease marginally at the annealing temperature of 1500 °C. The line breadth of molybdenum in the composites decreases after sintering at 1800 °C, but annealing does not have further influence on
Fig. 9. XRD spectra of a) ball milled powder C, b) composite C-1800, c) annealed at 1100 °C, d) annealed at 1300 °C, e) annealed at 1500 °C (•—TiC peaks, ▽—Mo2C peaks, Mo peaks not marked).
6
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
Generally, the performed investigations contribute to the understanding of the densification behavior of Mo–TiC composites by spark plasma sintering.
Acknowledgments The authors are grateful for the financial support of this work by Dr. Erich Krüger Stiftung and for the abandonment of the molybdenum powder by Dr. A. Hoffmann, Plansee Metall GmbH Reutte. We say also thank to Dr. G. Leichtfried, Plansee Metall GmbH Reutte, for the hydrogen annealing of the composites. Fig. 10. Results of micro hardness measurements of sintered and annealed Mo–25 vol.% TiC composites.
References the line breadth. Therefore, the defect structures adjusted during ball milling is very stable as well as in the sintered and in the annealed composites. The microhardness, measured after annealing, is constant up to annealing temperature of 1300 °C (see Fig. 10). A slight decrease of hardness is detectable after annealing at 1500 °C which is the consequence of the beginning of grain growth in molybdenum. 4. Conclusions Mo–TiC composites were successful prepared by spark plasma sintering from ball milled powders at sintering temperatures of 1600 and 1800 °C and with pressure of 67 MPa. Relative densities of 97% or 98% were reached. It is assumed that it is possible to reach 100% relative density by means of optimization of the milling conditions (BPR and rotation speed) and so to economize a final hot working step. The chosen ball milling parameters influence strongly the microstructure and the defect structure of powder particles which activates the densification process. In contrast to this, the sinter activity decreases with increasing TiC fraction, because the TiC particles prevent the sintering of molybdenum particles. During sintering, interdiffusion between molybdenum and TiC occurs and a typical core-rim structure can be observed where the rim consists of (Ti1 −x Mox)Cy. Moreover, the hexagonal Mo2C phase was detected by XRD measurements in the most composites after sintering. The occurrence of this phase is also influenced by milling parameters but not by the used sintering temperatures. This phase was not observed in composites made from low milled powders. It can be assumed that the formation of Mo2C is a consequence either of the higher concentration of crystal defects or of the stronger fragmentation of TiC in the powerful milled powders. It is particularly emphasized that neither recrystallization nor grain growth occurs in the molybdenum phase of all composites at the chosen sintering and annealing conditions. Even the lowest TiC content is sufficient to prevent recrystallization. The addition of hard TiC particles to molybdenum leads to a significant hardness increase. As expected, the hardness increases with increasing TiC content. Additionally, the milling conditions influence the hardness values. Therefore, the observed strong hardening is caused by a combination of particle reinforcement and strain hardening of molybdenum. During annealing of 10 h in a hydrogen atmosphere, the hardness stays constant up to 1300 °C. A slight decrease of hardness was detected after annealing at 1500 °C which is the consequence of the beginning of grain growth in molybdenum.
[1] Plansee High Performance Materials, “Molybdenum - Material properties and applications”, The official website of the Plansee Group. http://www.plansee.com/ lib/Molybdenum.pdf (accessed March 15, 2010). [2] Myong Ki Yoo, Hiraoka Y, Kurishita H, Kayano H, Choi J. Recrystallization of TiC dispersion Mo-alloy. Int J Refract Met Hard Mater 1996;14:355–64. [3] Le Flem M, Allemand A, Urvoy S, Cédat D, Rey C. Microstructure and thermal conductivity of Mo–TiC cermets processed by hot isostatic pressing. J Nucl Mater 2008;380:85–92. [4] Cédat D, Libert M, Le Flem M, Fandeur O, Rey C, Clavel M, et al. Experimental characterization and mechanical behaviour modelling of molybdenum–titanium carbide composite for high temperature applications. Int J Refract Met Hard Mater 2009;27:267–73. [5] Cédat D, Rey C, Clavel M, Schmidt JH, Le Flem M, Allemand A. Microstructural characterization of a composite Mo reinforced by 25 at.% TiC. J Nucl Mater 2009;385:533–7. [6] Cédat D. Etude expérimentale et numérique du comportement mécanique ď un composite métal-céramique: MoTiC30vol.%. PhD, Ecole Centrale des Arts et Manufactures, Ecole Centrale Paris 2008. [7] Takida T, Mabuchi M, Nakamura M, Igarashi T, Doi Y, Nagae T. Mechanical properties of a ZrC-dispersed Mo alloy processed by mechanical alloying and spark plasma sintering. Mater Sci Eng A 2000;276:269–72 276. [8] Song GM, Wang YJ, Zhou Y. The mechanical and thermophysical properties of ZrC/W composites at elevated temperatures. Mater Sci Eng A 2002;334:223–32. [9] Song GM, Wang YJ, Zhou Y. Thermomechanical properties of TiC particlereinforced tungsten composites for high temperature applications. Int J Refract Met Hard Mater 2003;21:1–12. [10] Zhang T, Wang Y, Zhou Y, Song G. Effect of temperature gradient in the disk during sintering on microstructure and mechanical properties of ZrCp/W composite. Int J Refract Met Hard Mater 2009;27:126–9. [11] Zhang T, Wang Y, Zhou Y, Song G. Effect of heat treatment on microstructure and mechanical properties of ZrC particles reinforced tungsten-matrix composites. Mater Sci Eng A 2009;512:19–25. [12] Zhang T, Wang Y, Zhou Y, Song G. Effect of ZrC particle size on microstructure and room temperature mechanical properties of ZrCp/W composites. Mater Sci Eng A 2010;527:4021–7. [13] Ohser-Wiedemann R, Martin U, Seifert HJ, Müller A. Densification behaviour of pure molybdenum powder by spark plasma sintering. Int J Refract Met Hard Mater 2010;28:550–7. [14] Ohser-Wiedemann R, Weck Ch, Martin U, Müller A, Seifert H-J. Densification behaviour of molybdenum composites during Spark Plasma Sintering. Proceedings of the Euro PM2010 Conference held in Florence, 10th–14th October 2010, Volume 2. Sintering; 2010. [15] Database PDF-2 of the International Centre for Diffraction Data; 2006. [16] The official website of the BGMN program. http://www.bgmn.de/ (accessed March 29 2011). [17] Krivoglaz MA. X-Ray and Neutron Diffraction in Nonideal Crystals. Springer Berlin; 1996. ISBN 3540505644. [18] Impermeable sintered metal materials and hardmetals—determination of density. (ISO 3369: 1975). German version DIN ISO 3369; 1990. [19] Sintered metal materials, excluding hartmetals - Permeable sintered metal materials - Determination of density, oil content and open porosity. (ISO 2738: 1999). German version EN ISO 2738; 2000. [20] German RM. A-Z of Powder Metallurgy. Elsevier Ltd1 8561 74298; 2005. [21] Mrotzek T, Hoffmann A, Martin U. Hardening mechanisms and recrystallization behaviour of several molybdenum alloys. Int J Refract Met Hard Mater 2006;24: 298–305.
Contents lists available at SciVerse ScienceDirect
Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Spark plasma sintering of TiC particle-reinforced molybdenum composites R. Ohser-Wiedemann a,⁎, Ch. Weck a, U. Martin a, A. Müller a, H.J. Seifert b,⁎ a b
Institute of Materials Science, TU Bergakademie Freiberg, Gustav-Zeuner-Str. 5, 09599 Freiberg, Germany Institute for Applied Materials, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, Bldg. 681, 76344 Eggenstein-Leopoldshafen, Germany
a r t i c l e
i n f o
Article history: Received 19 July 2011 Accepted 1 December 2011 Keywords: Molybdenum Spark plasma sintering Composites Microstructure Recrystallization
a b s t r a c t In order to improve the recrystallization resistance and the mechanical properties of molybdenum, TiC particle-reinforcement composites were sintered by SPS. Powders with TiC contents between 6 and 25 vol.% were prepared by high energy ball milling. All powders were sintered both at 1600 and 1800 °C, some of sintered composites were annealed in hydrogen for 10 h at 1100 up to 1500 °C. The powders and the composites were investigated by scanning electron microscopy and XRD. The microhardness and the density of composites were measured, and the densification behavior was investigated. It turns out that SPS produces Mo–TiC composites, with relative densities higher than 97%. The densification behavior and the microhardness of all bulk specimens depend on both the ball milling conditions of powder preparation and the TiC content. The highest microhardness was obtained in composites containing 25 vol.% TiC sintered from the strongest milled powders. The TiC particles prevent recrystallization and grain growth of molybdenum during sintering and also during annealing up to 10 h at 1300 °C. Interdiffusion between molybdenum and carbide particles leads to a solid solution transition zone consisting of (Ti1 −x Mox)Cy carbide. This diffusion zone improves the bonding between molybdenum matrix and TiC particles. A new phase, the hexagonal Mo2C carbide, was detected by XRD measurements after sintering. Obviously, this phase precipitates during cooling from sintering temperature, if (Ti1 −x Mox)Cy or molybdenum, are supersaturated with carbon. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Due to their high melting point, pure molybdenum and its alloys are used for high temperature applications in a variety of industries. However, application temperature is limited by the onset of recrystallization, which can be constraint by addition of fine disperse distributed particles, like carbides and oxides in the molybdenum alloys TZM (titanium–zirconium–molybdenum) or MHC (molybdenum–hafnium– carbon). These molybdenum alloys contain only small quantities of particles [1]. Another way to improve the recrystallization resistance and the mechanical properties at higher temperatures of refractory metals is the creation of metal matrix composites. Yoo et al. [2] investigated the recrystallization behavior of molybdenum sheets which were mechanical alloyed with 1.0 wt.% TiC and 0.2 wt.% C. At first, the powders were sintered by hot isostatic pressing (HIP) at 1300 °C, 200 MPa and 5 h. After it, the specimens were forged and rolled. Transmission electron microscopy investigations showed that very fine TiC particles are homogeneously distributed in the molybdenum phase and are pinned on dislocations and grain boundaries. The micro hardness increases up to 500 HV. Moreover,
⁎ Corresponding author. Tel.: + 49 3731 39 2647; fax: + 49 3731 39 3657. E-mail address: [email protected] (R. Ohser-Wiedemann). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.12.001
the formation of the molybdenum carbide Mo2C was detected in these composites. After annealing at temperatures between 1400 and 2000 °C for 10 h in flowing H2 atmosphere, grain and particle growth occur and the microhardness decreases. However, some small particles are still pinned on the dislocations and grain boundaries even at annealing temperature of 1800 ° C. Therefore, the authors assume, the recrystallization was not complete up to this temperature. At temperatures higher than 1800 °C a coarse polycrystalline microstructure with pores was observed and the microhardness decreased considerably. Mo–TiC composites with various TiC fractions are used as material for nuclear applications [3–6]. Mo–30 vol.% TiC has been proven useful as a composite with high mechanical strength at temperatures higher than 1000 K and up to 1900 K under fast neutron flux. Mixed Mo–TiC powders were successful compacted by HIP at 1600 °C, 2 h under a pressure of 160 MPa [3–5], but until now the recrystallization behavior of the composites was not investigated. After sintering a core-rime-structure was formed surrounding the TiC particles. The rim consist of the solid solution (Ti1 −x Mox)Cy which has the same cubic structure as TiC. This suggests that diffusion occurs in the carbide during sintering. The formation of Mo2C carbides was not detected in these investigations. Takida et al. [7] successfully consolidated ZrC particle dispersed molybdenum (1.4 vol.% ZrC) by spark plasma sintering. The particles were dispersed and no particle coarsening was observed after
2
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
annealing at 1797 °C for 1 h. The grain growth during annealing was strongly limited and the material exhibit excellent mechanical properties. The authors in ref. [8–12] investigated W-TiC or W-ZrC composites with particle contend up to 40 vol.% by hot pressing (HP) at 2000 or 2200 °C for 1 h. The results show that mainly fine particles prevent the tungsten grain growth. The mechanical properties at room temperature are improved depending on both, particle fraction and particle size. Similar to Mo–TiC composites, core-rime-structures were formed by diffusion of tungsten into TiC or ZrC. The formation of (Ti1 −xWx)Cy in W-TiC composites improve the strengthening effect of TiC [9]. In W-ZrC composites a hexagonal W2C phase with the same crystal structure like Mo2C (space group of both carbides P63/mmc [15]) was detected by XRD investigations. The phase was observed in only such composites where the carbide particle sizes are less than 1.7 μm. Obviously, this phase precipitates from the (Zr1 −xWx)Cy solid solution during cooling down from sintering temperature, because the solubility of tungsten in (Zr1 −xWx)Cy reduces with decreasing temperature [12]. The motivation of the presented work is to prepare Mo–TiC composites with low TiC volume fractions (6, 12.5 and 25%) using the spark plasma sintering process. During the SPS process, the dies and powders are heated by a pulsed direct current. This is in contrast to hot pressing or hot isostatic pressing and offers the possibility to sinter at higher heating rates, lower sintering temperatures and within shorter sintering times. To get a homogeneous TiC particle distribution in molybdenum, the Mo–TiC powder mixtures were produced by high energy ball milling with different milling parameters. Therefore, a further topic of the work is to investigate how the different milling parameters control the microstructure and hardness of sintered composites. Molybdenum, considered as the matrix material, was systematically investigated according to its SPS sintering behavior in ref. [13]. First results of the compaction of Mo–TiC composites were published in ref. [14]. 2. Experimental Pure molybdenum powders (99.95 wt.%, Plansee Metall GmbH, Austria; Fisher number 3 to 5 μm) were blended either with 6, 12.5 or 25 vol.% TiC powders (TiC STD 120, H.C. Starck GmbH, Germany; Fisher number 1 to 1.5 μm). Afterwards, powder mixtures were milled in a planetary ball mill (Fritsch “Pulverisette 6”) under air. Grinding bowls and balls were made of steel; the ball diameter was 10 mm. To prevent heating-up during milling, the milling was interrupted with a break of 15 min after a milling time of 15 min. The total milling time (without breaks) was 5 h. The ball powder ratio (BPR) and the rotating speed have been varied to change the milling energy. Table 1 gives an overview of all varied milling parameters. X-ray diffraction (XRD) spectra of all Mo–TiC powder mixtures and sintered composites were recorded with a powder diffractometer using Cu-Kα radiation and a graphite monochromator in the diffracted beam. The PDF data base [15] was used for qualitative phase analysis. Using the Rietveld-analysis BGMN (version 3.4.27) [16], the crystallite size (domain size) and micro strain values were derived from the integral breadths of the molybdenum peaks. The broadening due to the instrumental aberrations was separated by a
Table 1 Milling conditions and marking of the prepared Mo–TiC powders (total milling time 5 h). Ball powder ratio
Rotation speed
Powder identification marks
5:1 10 : 1 10 : 1
min− 1 200 200 300
Mo–25% TiC A B C
Mo–12.5% TiC – D –
Mo–6% TiC – E –
Monte-Carlo simulation. The dislocation density was roughly estimated from the mean squared strain, where an approach of Krivoglaz [17, page 357 ff.] was used. The approach is applicable only for cubic metals and their alloys, which exhibits high dislocation densities. Moreover, elastic isotropy of the crystals is supposed. The Mo–TiC powders were filled without any additives in cylindrical graphite dies of 20 mm in diameter. The sintering runs in a FCT-HP D 25 spark plasma sinter equipment (FCT Systeme GmbH, Rauenstein, Germany). The compaction of the powder mixtures started with raising the external pressure up to 67 MPa, followed by heating up under vacuum to temperatures of 1600 or 1800 °C. The heating rate amounted to 300 K/min, and the holding time at maximum temperature was 3 min for all batches. In contrast to the sintering by HP [8–12] or HIP [2–5], the whole process time of SPS needs 15 min only. The pulse pattern of the electric current was adjusted on 2:1 (10 ms pulse time, 5 ms pause time). The compaction pressure has been decreased during the cooling step. In the following, a combination of the powder name and the sintering temperature is used as a notation of composites, e.g. “A-1600” for composite sintered at 1600 °C using powder A. The densities of all consolidated bulk specimens were measured using the Archimedes technique [18,19]. The relative densities were calculated based on the theoretical densities of molybdenum (10.28 g cm − 3) and titanium carbide (4.93 g cm − 3). The microstructure was studied by scanning electron microscopy (SEM) at cross sections (sections parallel to the acting force). The preparation steps were grinding and mechanical polishing up to 1 μm diamonds abrasive. Same cross sections were chemical etched using Murakami's reagent (a solution mixture of 10 g sodium hydroxide, 10 g potassium ferricyanide, and 100 ml of distilled water). The hardness of the Mo–TiC composites was measured at the cross sections by a Vickers microhardness tester (LECO Instrumente GmbH). The indentations were made in a pattern consisting of three parallel lines each with nine measuring points. In order to investigate the tempering resistance of microstructure, some sintered composites were annealed at 1100, 1300 and 1500 °C for 10 h in hydrogen atmosphere. After annealing the microhardness was measured again, and XRD spectra were recorded. 3. Experimental results and discussion 3.1. Characterization of ball milled Mo–TiC powders Scanning electron microscopy investigations showed that the original molybdenum powder particles are of spherical shape, some of them include small pores (see also [13]). The particles agglomerate and tend to form chains. The powders contain a fraction of fine particles with sizes b 1 μm. In contrast, the finer TiC powders consist of some big angular particles and a lot of small fractions before milling. Fig. 1 shows the morphology of powder mixtures after different milling conditions. At low BPR of 5:1 and a rotating speed of 200 rpm (powder A), the original shapes of both powders are still observable (Fig. 1a). In contrast, the morphology of powder particles changes if the BPR was increased up to 10:1 (powders B, D and E). Fig. 1b gives an impression of morphology on the example of powder B. Under these milling conditions, the molybdenum particles were plastically deformed and show a flat shape. Most of the hard brittle TiC particles were fragmented and afterwards pressed in to the softer molybdenum particles. With further increasing of milling energy, e.g. with increasing rotating speed, the original particle shapes got lost (powder C). Here both kinds of particles were cold welded and form agglomerates (Fig. 1c). After milling, all powders were investigated by XRD measurements. The phase analysis show that the powders consist only of the two phases molybdenum and titanium carbide. However, the plastic deformation during milling leads to strong microstructural
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
3
metals. Hence, the molybdenum particles in powder C are strongly strain hardened. The dislocation densities are not influenced by the TiC content. The deviations between powders B, D and E are inside the error range of XRD measurement. 3.2. Densification
changes especially inside the ductile molybdenum particles. The increase of the BPR and the rotating speed leads to higher crystal defect concentration which is connected to a strong broadening of the XRD peaks. In case of molybdenum, the crystallite size drops down drastically from more than 1000 nm in non-milled molybdenum powders to 40–48 nm in the milled powders A, B, D, E and to 24 nm in the milled powder C. Simultaneously, the dislocation densities in the powders increase. Fig. 2 depicts the strong influence of the milling conditions on dislocation densities. With increasing milling energy the dislocation density increases and reaches a maximum value of more than 10 12 cm − 2 in powders, milled at BPR 10:1 and 300 rpm. Normally, such high values are measured in strong cold worked
Fig. 3 shows the relative densities of all sintered composites in comparison to sintered pure molybdenum (not ball milled). Sintering at 1800 °C leads to relative densities of 97% or 98%, respectively. These values are a little bit higher than those of bulk molybdenum. Therefore, ball powder milling improves the densification process. The combination of plastic deformation, cold welding and fracturing during milling increases the surface area and that leads to an activation of sintering processes. At the lower sintering temperature, the relative density depends slightly on milling conditions and TiC fraction. However, the calculation of relative densities is uncertain, because it was not possible to calculate the theoretical densities of the composites exactly. As explained more in detail in paragraph 3.4, interdiffusion between molybdenum and TiC occurs during sintering. A core-rim structure is formed where the rim consist of (Ti1 −x Mox)Cy. It was not possible to estimate precisely the volume fraction of TiC and (Ti1 −x Mox)Cy and the chemical composition of (Ti1 −x Mox)Cy. Additionally, same composites contain a small fraction of Mo2C, which also influence the theoretical density. For these reasons, it was decided to calculate the relative densities based on the powder composition. By our best knowledge, the relative densities presented in Fig. 3 are lower bounds. In conventional powder metallurgy, where the sintering process starts with a green body, sintering shrinkages range from about zero to 25% [20]. Due to the uniaxial pressing at the SPS, the shrinkage occurred mainly in the pressing direction. This linear intrinsic shrinkage can be calculated from the displacement of the pressing piston which is recorded during whole sintering process. For exact conclusions from these records, the experimental data have to be corrected by subtraction of the contribution of piston displacement, which was caused by compression during applied pressure before sintering starts. This is necessary, because this first compaction depends on the applied pressure as well as on the die filling behavior of powder, and it varied in every batch. Furthermore, the thermal expansion of specimen, die and graphite sheets must be subtracted. For this correction, some baselines were determined experimentally with the same powders and sintering conditions. The procedure is described in more detail in ref. [13]. Nearly S-shaped curves were observed, if the calculated linear shrinkage is represented as a function of time (Fig. 4). This indicates that the SPS process is mainly thermal activated. Only at the beginning of sintering, linear dependence occurs with a sharp transition at about 950 °C. This transition temperature corresponds to 0.4 Tm of pure molybdenum, which is the temperature where appreciable volume diffusion starts in solid molybdenum materials. The volume
Fig. 2. Dislocation densities inside molybdenum powder particles after ball milling.
Fig. 3. Relative densities of all Mo–TiC composites in comparison to pure bulk molybdenum sintered by SPS.
Fig. 1. Morphology of Mo–25 vol.% TiC powders after ball milling under different conditions (SEM, BSE contrast); a) powder A: BPR 5:1, 200 rpm; b) powder B: BPR 10:1, 200 rpm; c) powder C: BPR 10:1, 300 rpm.
4
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
Fig. 5. SEM micrographs (BSE contrast; bright: molybdenum, dark: TiC with core-rimstructure) of Mo–TiC composites, spark plasma sintered at 1800 °C. a) specimen B1800, b) specimen C-1800. Fig. 4. Densification behavior of Mo–TiC composites depending on ball milling conditions (on the top) and on TiC content (on the bottom).
diffusion supports the particle bonding and the densification during sintering which leads to a strong rise of shrinkage above this temperature up to the end of heating up. In contrast, the shrinkage is very low during the short isothermal sintering time of 3 min. Therefore, the densification of the composites takes place mainly during heating up, and the process corresponds to the intermediate stage sintering where neighboring sinter necks grow sufficiently large to overlap [20]. Overall, linear shrinkages up to 40% were achieved, which is higher than the shrinkage of conventional sintering. This is caused by the lower initial powder compaction at SPS. The shrinkage rate is influenced by the ball milling parameters, e.g. by the defect densities inside the powder particles. The increase of milling energy shifts the curves to shorter sintering times. Therefore, the sintering activity of the milled powders rises up with increasing defect densities. The TiC fraction influences the sintering rate as well. The curves shift to shorter times for lower particle contents. Obviously, the small TiC particles, pressed into the molybdenum surface, decrease the contact areas between molybdenum and prevent the growth of molybdenum particle bonds.
particles can be observed in the specimen C-1800 (Fig. 5b). The microstructures of composites with lower TiC fraction look like similar to the microstructures of specimen B-1800 (Fig. 5a). All specimens are macroscopic crack free and have very small porosities which correspond to the measured densities. The TiC particles are cohesively bonded in the molybdenum matrix and a typical core-rim structure was formed. According to the investigation of Cédat et al. [5], the rim consist of (Ti1 −x Mox)Cy which was formed by diffusion of molybdenum into TiC during sintering. The same effect can be observed in TiC particle-reinforced tungsten composites (see chapter 1). In both cases the chemical composition changes continuously in the rim and the bonding between matrix and carbide is improved. The microhardness of all composites is presented in Fig. 6 in comparison to pure SPS molybdenum. It can be shown that the microhardness depends on both, the TiC particle fraction and the milling parameters. The composites produced from the strongest deformed powder C are of the highest hardness which is four times higher than in pure molybdenum. It can be concluded that this strong hardening is caused by a combination of the particle reinforcement and the strain hardening of molybdenum during ball milling. Finally, the dependence of the hardness on the milling conditions is a further
3.3. Microstructure and microhardness of sintered composites Fig. 5 illustrates the microstructure of composites made from the high energy milled powders B and C after sintering at 1800 °C. The SEM micrograph was obtained using backscattered electrons allowing an exact separation between the molybdenum and titanium carbide phases. The shapes of molybdenum grains inside the composites inherit the shape of the plastically deformed molybdenum particles after milling (compare Fig. 2). This indicates that neither recrystallization nor grain growth occur in the molybdenum phase at the chosen sintering conditions. This is in contradiction to the consolidation behavior of pure molybdenum powders (particle size 3–5 μm) showing a significant grain growth at these sintering temperatures. Here, a mean chord length of 19 ± 12 μm was estimated after sintering at 1800 °C [13]. A more homogeneous distribution of the fine TiC
Fig. 6. Results of micro hardness measurements of all SPS Mo–TiC composites.
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
5
argument that no recrystallization or grain growth occurs during sintering.
3.4. Defect structures and phase compositions after sintering After sintering, the integral breath of molybdenum peaks is reduced compared to that of as-milled powders (see for example Fig. 9, XRD patterns (a) and (b)). The dislocation density decreases to an order of magnitude while the crystallite size increases about sixfold. Obviously, recovery processes occur during sintering. Some of the stored internal strain energy is relieved by dislocation motion. The number of dislocations is slightly reduced by annihilation or by moving to grain boundaries. In addition, dislocations rearrange themselves into complex networks and the distortion of lattice planes goes back. However, there is no evidence for recrystallization. In this case, the dislocation density might be reduced dramatically. It can be concluded that the TiC particles also prevent the motion of high-angle grain boundaries as a requirement for recrystallization. After sintering, a new phase was detected by XRD measurements in the most composites (see Fig. 7). The phase was identified as hexagonal β-Mo2C. The occurrence of the Mo2C phase is independent of the sintering temperature, but it depends on milling conditions. The phase was not detected in the composites made from powder A. Probably, the formation of Mo2C is supported either by the higher concentration of crystal defects, which accelerated the carbon diffusion, or by the stronger fragmentation of TiC in the powerful milled powders. The Mo2C phase occurs also in TiC particle-reinforcement molybdenum sheets after sintering, forging and rolling [2] and in the molybdenum alloy TZM [21]. Furthermore, this carbide type was detected in tungsten composites as W2C [10–12]. SEM was used to look for Mo2C precipitates in our SPS composites. Mo–25vol.% TiC specimens C-1800 (with Mo2C) and A-1800 (without Mo2C) were compared. For this purpose, cross sections of both composites were etched by Murakami's reagent. Due to the different chemical potential of the phases, their surfaces were etched differently. Fig. 8 shows scanning electron micrographs of cross sections after etching. The etching reagent roughened the surface of molybdenum grains strongly. In contrast, TiC particles and the core-rim structure are attacked marginally. This creates a strong topographic contrast. Fig. 8a shows a micrograph of the composite C-1800. A cluster of TiC particles is represented, where the (Ti1 −xMox)Cy rims contact each other. Some minute particles, marked with a white circle, exist inside the (Ti1 −xMox)Cy grains. In contrast, no such particles appear in the composite A-1800 (Fig. 8b). For this reason, we assume that the observed particles are Mo2C. Based on the investigations in W-ZrC composites [10], it can be concluded that the Mo2C phase is also precipitated from carbon supersaturated (Ti1 −xMox)Cy solid solution during cooling. Additionally, thin light plates are visible in the (Ti1 −xMox)Cy phase of the C-1800 composite (Fig. 8a). Probably, these plates are low angle
Fig. 7. XRD spectra of SPS Mo–TiC composites sintered at 1600 °C. (•—TiC peaks, ∇— Mo2C peaks, not marked—Mo peaks).
Fig. 8. SEM micrographs (BSE contrast) of etched Mo–TiC composites: a) cluster of (Ti1 −xMox)Cy with precipitates (specimen C-1800), b) the surrounding of a TiC particle without precipitates (specimen A-1800).
grain boundaries decorated with minute precipitates. Finally, the exact formation of Mo2C must be a subject of a further investigation. 3.5. Tempering behavior After sintering, the composites C-1800 were annealed at 1100, 1300 and 1500 °C for 10 h in hydrogen atmosphere in order to investigate the tempering resistance of microstructure. The stability of the microstructure under thermal loading is very important for the application of these composites. The microstructure and the phase composition of all sintered and annealed specimens show no significant changes up to 1300 °C. Above 1300 °C a minor growth of molybdenum grains inside the composites occurs. Fig. 9 shows the XDR spectra of the powder C and its sintered and annealed composites. It is clearly detectable that the Mo2C carbide occurs only after sintering in the bulk specimens and not inside the powder C. The Mo2C fraction seems to decrease marginally at the annealing temperature of 1500 °C. The line breadth of molybdenum in the composites decreases after sintering at 1800 °C, but annealing does not have further influence on
Fig. 9. XRD spectra of a) ball milled powder C, b) composite C-1800, c) annealed at 1100 °C, d) annealed at 1300 °C, e) annealed at 1500 °C (•—TiC peaks, ▽—Mo2C peaks, Mo peaks not marked).
6
R. Ohser-Wiedemann et al. / Int. Journal of Refractory Metals and Hard Materials 32 (2012) 1–6
Generally, the performed investigations contribute to the understanding of the densification behavior of Mo–TiC composites by spark plasma sintering.
Acknowledgments The authors are grateful for the financial support of this work by Dr. Erich Krüger Stiftung and for the abandonment of the molybdenum powder by Dr. A. Hoffmann, Plansee Metall GmbH Reutte. We say also thank to Dr. G. Leichtfried, Plansee Metall GmbH Reutte, for the hydrogen annealing of the composites. Fig. 10. Results of micro hardness measurements of sintered and annealed Mo–25 vol.% TiC composites.
References the line breadth. Therefore, the defect structures adjusted during ball milling is very stable as well as in the sintered and in the annealed composites. The microhardness, measured after annealing, is constant up to annealing temperature of 1300 °C (see Fig. 10). A slight decrease of hardness is detectable after annealing at 1500 °C which is the consequence of the beginning of grain growth in molybdenum. 4. Conclusions Mo–TiC composites were successful prepared by spark plasma sintering from ball milled powders at sintering temperatures of 1600 and 1800 °C and with pressure of 67 MPa. Relative densities of 97% or 98% were reached. It is assumed that it is possible to reach 100% relative density by means of optimization of the milling conditions (BPR and rotation speed) and so to economize a final hot working step. The chosen ball milling parameters influence strongly the microstructure and the defect structure of powder particles which activates the densification process. In contrast to this, the sinter activity decreases with increasing TiC fraction, because the TiC particles prevent the sintering of molybdenum particles. During sintering, interdiffusion between molybdenum and TiC occurs and a typical core-rim structure can be observed where the rim consists of (Ti1 −x Mox)Cy. Moreover, the hexagonal Mo2C phase was detected by XRD measurements in the most composites after sintering. The occurrence of this phase is also influenced by milling parameters but not by the used sintering temperatures. This phase was not observed in composites made from low milled powders. It can be assumed that the formation of Mo2C is a consequence either of the higher concentration of crystal defects or of the stronger fragmentation of TiC in the powerful milled powders. It is particularly emphasized that neither recrystallization nor grain growth occurs in the molybdenum phase of all composites at the chosen sintering and annealing conditions. Even the lowest TiC content is sufficient to prevent recrystallization. The addition of hard TiC particles to molybdenum leads to a significant hardness increase. As expected, the hardness increases with increasing TiC content. Additionally, the milling conditions influence the hardness values. Therefore, the observed strong hardening is caused by a combination of particle reinforcement and strain hardening of molybdenum. During annealing of 10 h in a hydrogen atmosphere, the hardness stays constant up to 1300 °C. A slight decrease of hardness was detected after annealing at 1500 °C which is the consequence of the beginning of grain growth in molybdenum.
[1] Plansee High Performance Materials, “Molybdenum - Material properties and applications”, The official website of the Plansee Group. http://www.plansee.com/ lib/Molybdenum.pdf (accessed March 15, 2010). [2] Myong Ki Yoo, Hiraoka Y, Kurishita H, Kayano H, Choi J. Recrystallization of TiC dispersion Mo-alloy. Int J Refract Met Hard Mater 1996;14:355–64. [3] Le Flem M, Allemand A, Urvoy S, Cédat D, Rey C. Microstructure and thermal conductivity of Mo–TiC cermets processed by hot isostatic pressing. J Nucl Mater 2008;380:85–92. [4] Cédat D, Libert M, Le Flem M, Fandeur O, Rey C, Clavel M, et al. Experimental characterization and mechanical behaviour modelling of molybdenum–titanium carbide composite for high temperature applications. Int J Refract Met Hard Mater 2009;27:267–73. [5] Cédat D, Rey C, Clavel M, Schmidt JH, Le Flem M, Allemand A. Microstructural characterization of a composite Mo reinforced by 25 at.% TiC. J Nucl Mater 2009;385:533–7. [6] Cédat D. Etude expérimentale et numérique du comportement mécanique ď un composite métal-céramique: MoTiC30vol.%. PhD, Ecole Centrale des Arts et Manufactures, Ecole Centrale Paris 2008. [7] Takida T, Mabuchi M, Nakamura M, Igarashi T, Doi Y, Nagae T. Mechanical properties of a ZrC-dispersed Mo alloy processed by mechanical alloying and spark plasma sintering. Mater Sci Eng A 2000;276:269–72 276. [8] Song GM, Wang YJ, Zhou Y. The mechanical and thermophysical properties of ZrC/W composites at elevated temperatures. Mater Sci Eng A 2002;334:223–32. [9] Song GM, Wang YJ, Zhou Y. Thermomechanical properties of TiC particlereinforced tungsten composites for high temperature applications. Int J Refract Met Hard Mater 2003;21:1–12. [10] Zhang T, Wang Y, Zhou Y, Song G. Effect of temperature gradient in the disk during sintering on microstructure and mechanical properties of ZrCp/W composite. Int J Refract Met Hard Mater 2009;27:126–9. [11] Zhang T, Wang Y, Zhou Y, Song G. Effect of heat treatment on microstructure and mechanical properties of ZrC particles reinforced tungsten-matrix composites. Mater Sci Eng A 2009;512:19–25. [12] Zhang T, Wang Y, Zhou Y, Song G. Effect of ZrC particle size on microstructure and room temperature mechanical properties of ZrCp/W composites. Mater Sci Eng A 2010;527:4021–7. [13] Ohser-Wiedemann R, Martin U, Seifert HJ, Müller A. Densification behaviour of pure molybdenum powder by spark plasma sintering. Int J Refract Met Hard Mater 2010;28:550–7. [14] Ohser-Wiedemann R, Weck Ch, Martin U, Müller A, Seifert H-J. Densification behaviour of molybdenum composites during Spark Plasma Sintering. Proceedings of the Euro PM2010 Conference held in Florence, 10th–14th October 2010, Volume 2. Sintering; 2010. [15] Database PDF-2 of the International Centre for Diffraction Data; 2006. [16] The official website of the BGMN program. http://www.bgmn.de/ (accessed March 29 2011). [17] Krivoglaz MA. X-Ray and Neutron Diffraction in Nonideal Crystals. Springer Berlin; 1996. ISBN 3540505644. [18] Impermeable sintered metal materials and hardmetals—determination of density. (ISO 3369: 1975). German version DIN ISO 3369; 1990. [19] Sintered metal materials, excluding hartmetals - Permeable sintered metal materials - Determination of density, oil content and open porosity. (ISO 2738: 1999). German version EN ISO 2738; 2000. [20] German RM. A-Z of Powder Metallurgy. Elsevier Ltd1 8561 74298; 2005. [21] Mrotzek T, Hoffmann A, Martin U. Hardening mechanisms and recrystallization behaviour of several molybdenum alloys. Int J Refract Met Hard Mater 2006;24: 298–305.