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Nb(Si,C,N) composite materials densified by spark plasma sintering
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ScienceDirect Journal of the European Ceramic Society xxx (2015) xxx–xxx
Nb(Si,C,N) composite materials densified by spark plasma sintering Martin Seifert a , Zhijian Shen b , Walter Krenkel a , Günter Motz a,∗ a
b
University of Bayreuth, Ceramic Materials Engineering (CME), D-95440 Bayreuth, Germany Stockholm University, Department of Materials and Environmental Chemistry, S-106 91 Stockholm, Sweden
Abstract Nb2 N/Nb5 Si3 /Nb ceramic–metal and Nb(C,N)/Nb5 Si(3+x) Cx ceramic like composites were densified by spark plasma sintering (SPS). The precursor multiphase powders were synthesized by solid-state reactions of a polysilazane precursor with niobium powder addition. By varying the amount of reactive species of both precursor and niobium the phase composition and structure of the synthesized multiphase powders can be tailored from metal dominated to mainly ceramic like. The results demonstrated that the SPS approach leads to highly dense samples at 1600 ◦ C with an applied uniaxial pressure of 100 MPa. Furthermore, XRD measurements and EBSD analysis proofed that SPS is a feasible method to retain the original phase compositions and grain sizes within the multiphase powders during sintering. The evaluation of the Vickers indents revealed both a dependency of the measured values on phase composition and on residual porosity. © 2015 Published by Elsevier Ltd. Keywords: Polymer derived ceramics; Spark-plasma sintering; Composite material; Reactive pyrolysis; Refractory metal
1. Introduction In general, advanced ceramics provide exceptional properties like high hardness, resistance to wear, high strength to weight ratio as well as high stiffness and good creep resistance. These materials are lack, however, of plastic deformation at ambient temperatures thus fail normally in a catastrophic manner. Ceramic–metal composites combine the benefits of both ceramics and metals by exhibiting increased fracture toughness because of the ductile deformation behaviour of the metal phases involved. In order to investigate new material combinations for ceramic–metal or ceramic composites for high performance applications niobium and it’s related compounds seem to be natural candidates due to their specific properties. Niobium and binary Nb(C,N) ceramic or intermetallic Nb–Si phases demonstrate high melting points above 2000 ◦ C, high hardness as well as high stiffness and strength at comparatively low densities (7–8 g/cm3 ) [1–7]. Furthermore, as it has been reported in the literature it is possible to improve material properties like strength
∗
Corresponding author. Tel.: +49 921 555505; fax: +49 921 555502. E-mail addresses: [email protected] (M. Seifert), [email protected] (Z. Shen), [email protected] (W. Krenkel), [email protected] (G. Motz).
or hardness of steels by the dispersion of Nb2 N or NbCx Ny phases [8,9]. Especially in situ formed Nb/Nb5 Si3 compounds were intensively investigated because this material combination exhibit very good creep behaviour and high strength at 1300√◦ C and a room temperature fracture-toughness up to 21 MPa m [10–12]. Furthermore, by the dispersion of NbC within in situ Nb/Nb5 Si3 composites crack deflection and branching occur [13]. In terms of the high melting points of Nb and related ceramic or intermetallic phases the processing of dense composites e.g. by liquid metal infiltration of porous ceramic preforms is very complex and difficult to accomplish. The powder metallurgical approach in combination with different sintering techniques like reactive hot pressing (HP), hot isostatic pressing (HIP) or spark plasma sintering (SPS) require elaborate powder preparation steps like mechanical alloying (MA) or self-combustion high temperature synthesis (SHS) to produce sinterable powders [14–17]. Recent work on ceramic composites derived from a carbon rich polysilazane filled with Niobium particles demonstrated the formation of silicide, carbide and nitride phases at the Nb/SiCN interphase via solid-state reaction during precursor pyrolysis at 1600 ◦ C in argon atmosphere [18]. By varying the amount of the reactive educts within the starting composition it is possible to control phase formation and to tailor the microstructure.
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.02.005 0955-2219/© 2015 Published by Elsevier Ltd.
Please cite this article in press as: Seifert M, et al. Nb(Si,C,N) composite materials densified by spark plasma sintering. J Eur Ceram Soc (2015), http://dx.doi.org/10.1016/j.jeurceramsoc.2015.02.005
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ARTICLE IN PRESS M. Seifert et al. / Journal of the European Ceramic Society xxx (2015) xxx–xxx Table 1 Composition and density of the starting powder mixtures. Powder
Nb amount
7Nb3P 6Nb4P 5Nb5P 4Nb6P
Fig. 1. Flow chart for the manufacturing of Nb(C,N)/Nb5 Si3 composites.
By increasing the amount of polymer within the starting compositions an interconnected particle network with a core-rim structure is formed accompanied by formation of porosity values up to 50%. The aim of this work is to obtain dense Nb(Si,C,N) composite materials by additive-free sintering of multiphase powders using SPS technique. 2. Experimental procedure 2.1. Powder synthesis The Nb(Si,C,N) multiphase powders used for this work were synthesized via solid-state reaction of Nb powder and the polysilazane precursor HTTS during pyrolysis. The polymer results from chemical crosslinking of the commercially available liquid polysilazane HTT 1800 (Fa. Clariant Advanced Materials GmbH, Sulzbach, Germany). The crosslinking process is described in detail elsewhere [21]. In order to produce the precursor powder, HTTS is dissolved in pentane and 3 wt.% of the radical initiator dicumylperoxide (DCP 98%, Sigma Aldrich, St. Louis, USA) was added to the solution. By adding DCP the crosslinking temperature of HTTS could be decreased to 150 ◦ C. Subsequently, the solvent was removed with rotary evaporation. The drying process results in brittle HTTS powder agglomerates which can be sieved to a particle size smaller than 63 m. Fig. 1 shows the scheme for manufacturing the composite materials. The starting mixtures containing niobium (Hauner Metallische Werkstoffe, Roethenbach, Germany, 99.8%, <45 m) and HTTS powder were homogenized by dry mixing and subsequent ball milling for 6 h. The resulting mixtures were filled in a heatable stainless steel cylinder for thermal crosslinking of the PCSZ at 160 ◦ C in air and for shaping of cylindrical specimen. Pyrolysis of the cured green bodies in an Al2 O3 tube furnace (Thermal Technology, Santa Rosa, USA) at 1600 ◦ C in Argon (purity 5.0) atmosphere with a heating rate of 3 K/min and annealing for 3 h at maximum temperature lead to specimen with a highly porous and weak bounded particle network with porosities up to 50–70 vol.%. For this reason it is
Powder density after pyrolysis
(vol.%)
(wt.%)
(g/cm3 )
70 60 50 40
94.8 92.1 88.6 83.9
8.11 7.91 7.69 7.43
simple to grind up the porous specimens by hand in a mortar to prepare the multiphase powders with a particle size smaller than 63 m suitable for the SPS process. The composition of the powder mixtures is listed in Table 1. According to the previous work the region between 60 and 50 vol.% of Nb within the starting composition is of utmost interest because it involves the transition from phase compositions dominated by residual metallic niobium to compositions solely containing intermetallic or ceramic phases. In order to investigate the influence on both, phase composition and mechanical properties, the amount of Nb was varied from 70 to 40 vol.%. The designation of the powders corresponds to the composition of the starting powder mixtures. The sintered samples are designated by the affix–Sx while x represents the respective sinter run. 2.2. SPS-process Spark plasma sintering was performed in a SPS facility (Dr. Sinter 2050, Sumitomo Coal Mining Co. Ltd., Tokyo, Japan) under medium vacuum of 6 Pa at 1600 ◦ C. For each experiment a powder volume of 0.2 cm3 was filled into a graphite die with an inner diameter of 12 mm lined with graphite foil. The temperature was first increased to 600 ◦ C in 1 min and after a setting time of 2 min it was raised again to the pre-set sintering temperature by applying different heating rates. The parameters used for different sinter runs are listed in Table 2. During the second heating step the temperature was monitored by a radiation pyrometer focused on the out surface of the graphite die. To ensure reproducibility the maximum pressure of 100 MPa was applied during the setting time at 600 ◦ C and kept constantly at this value for entire period of sintering. Dwell times at the sintering temperature were set to 3 and 5 min for investigating the influence on densification and grain growth. It is expected that the unique sintering mechanisms of SPS facilitated by the applicable high heating and cooling rates will Table 2 Different sintering parameters used for the experiments. Run
Heating rate (K/min)
Pressure (MPa)
Sintering temperature (◦ C)
Holding time (min)
S1 S2 S3 S4
200 200 100 100
100 100 100 100
1400 1600 1600 1600
3 3 3 5
Please cite this article in press as: Seifert M, et al. Nb(Si,C,N) composite materials densified by spark plasma sintering. J Eur Ceram Soc (2015), http://dx.doi.org/10.1016/j.jeurceramsoc.2015.02.005
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support densification of the synthesized powders [19] and retain metastable phases, as well as prevent coarsening or changes in the phase compositions of the generated Nb–Si and Nb(C,N) phases. The linear changes of the sample and of the die during sintering were recorded with a dilatometer integrated in the SPS facility. The densification curves were derived from the shrinkage curves considering the relative densities of the sintered samples. The shrinkage curves were corrected for the thermal expansion of the graphite die in order to obtain the real shrinkage curves of the multiphase powders. Eq. (1) concludes the approach of using the normalized data for calculating the relative densities during sintering. ρi = ρs
Is Is + L s − L i
(1)
In this equation ρi is the intermediate density at sintering, ρS is the relative density after sintering, lS is the sample height after sintering. LS and Li are defined as the punch positions during and after sintering. The density of the sintered samples ρS was calculated by measuring the relative density with the Archimedes method and divided by the theoretical density of the milled multiphase powders measured with He-Pycnometry. In order to obtain information about the sintering kinetics the linear shrinkage rate, defined as dl/dt, was plotted as a function of time.
2.3. Characterization methods In order to measure the true density of the synthesized powders He-Pycnometry measurements (Accupyc 1330, Micromeritics, Germany) were carried out. The density and open porosity of the sintered samples were determined using the Archimedes’ method. Before investigating the microstructure, phase composition and hardness, the sintered samples were polished with diamond slurries down to 1 m finish. Crystalline phases were detected by X-ray diffraction analysis of the grinded powders and the polished surfaces of the sintered bodies (D8 ADVANCE, Bruker AXS, Karlsruhe, Germany) using monochromated CuK␣ radiation. The amount of residual Nb was determined by quantitative Rietveld refinement using the evaluation software TOPAS (V4.2, Bruker AXS, Karlsruhe, Germany) and PDF-4+ 2012 structural database for crystallographic information. Weighted Scanning electron microscopy (SEM, JSM6400, Jeol, Japan) was used to characterize the structure and morphology of the particles as well as the microstructure of the sintered samples. Electron back-scattered diffraction analysis (EBSD) was performed by using a field emission SEM (Leo Gemini 1530, Zeiss, Oberkochen, Germany) equipped with Oxford Instruments INCA energy 400 EDS and HKL Nordlys EBSD detectors to yield information about phase distribution of the composites. The Vickers hardness of the composites was determined with the Vickers indentation technique under a load of 98 N for 15 s.
Fig. 2. XRD patterns of the Nb(Si,C,N) powders synthesized at 1600 ◦ C in Ar – Atmosphere (␣: Nb5 Si3 , : Nb2 N, ♦: Nb, : Nb2 C, ␥: Nb5 Si4 C, 䊉:NbC, : Nb2 CN).
3. Results and discussion 3.1. Multiphase powders The pyrolysis of HTTS involves several thermally induced reactions comprising the transition from polymer into an amorphous SiCN ceramic at temperatures between 400–1000 ◦ C, the formation of Si-N, Si-C nano-domains and a free carbon phase at T < 1400 ◦ C, as well as the crystallization of these Si-N and Si-C environments at T > 1400 ◦ C [22–26]. In the presence of reactive metal like Nb the crystallization is inhibited because of reactions at the Nb/SiCN interphase at temperatures exceeding 1000 ◦ C. The mechanism of such reaction and the phase formation were described in detail elsewhere [18]. Fig. 2 shows the results of the XRD-measurements of the synthesized multiphase powders after pyrolysis at 1600 ◦ C in argon atmosphere. For a niobium content of 94.8 wt.% (70 vol.%) within the starting composition the synthesized powder particles are composed of residual metallic Nb, metal rich -Nb2 N and ␣-Nb5 Si3 phase. The amount of residual niobium determined with qualitative Rietveld refinement is 31 vol.% for 7Nb3P (Rwp = 14.32, GOF = 1.53) and 11.0 vol.% for 6Nb4P (Rwp = 16.24, GOF = 1.57). By increasing the amount of HTTS gradually to 60 vol.% the amount of metallic niobium is completely converted by the reaction with SiCN. Furthermore, a low temperature distorted Nb2 C phase (-Fe2 N prototype) is formed [27]. The enhanced amount of carbon also supports the formation of metastable Nb5 Si4 C which is the ␥-modification (Mn5 Si3 prototype) of tetragonal ␣-Nb5 Si3 with carbon atoms located at octahedral sites of the Nb–Si lattice. Nb2 N is in equilibrium with Nb5 Si3 . Murakami et al. reached the same results by investigating Nb–Si–N composites prepared with spark plasma sintering [28]. Cubic Nb2 CN was detected within the 4Nb6P powder and it may be a result of the reaction of Nb2 N with
Please cite this article in press as: Seifert M, et al. Nb(Si,C,N) composite materials densified by spark plasma sintering. J Eur Ceram Soc (2015), http://dx.doi.org/10.1016/j.jeurceramsoc.2015.02.005
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Fig. 3. Nb(Si,C,N) multiphase powders synthesized via solid state reaction of different Nb/HTTS mixtures: (a) 7Nb3P, (b) 6Nb4P, (c) 5Nb5P, (d) 4Nb6P.
carbon phases or enhanced carbon diffusion. No crystalline SiC or Si3 N4 phases, typically resulting from the crystallization of the amorphous SiCN phase, could be detected. It can be assumed that the amount of HTTS is completely converted during reaction by forming Nb(Si,C,N) phases. EBSD measurements carried out in the previous work proofed already that when the added amount of Nb was 50 vol.% or less particles with a core-rim structure were formed, which composed of intermetallic Nb5 Si3 and Nb5 Si4 C phases at the rim and ceramic Nb(C,N) phases at the particle core. Fig. 3 displays examples of the synthesized particles. It is obvious that particle morphology is strongly dependent on the phase composition of the particles. According to the previous work the structures at the niobium particle surfaces of 7Nb3P (Fig. 3a) can be ascribed to Nb2 N and Nb5 Si3 phases. The particle surfaces of the powders 5Nb5P and 4Nb6P (Fig. 3c,d) are composed of
coarse-grained Nb5 Si3 phase and fine-grained Nb5 Si4 C, respectively. 3.2. Densification As related to the pyrolysis temperature selected for powder synthesis the sintering trials were carried out at a maximum sintering temperature of 1600 ◦ C to avoid the phase transformations. Different heating rates of 100 and 200 K/min and holding times of 3 and 5 min were chosen in order to analyse the densification behaviour and the sintering kinetics. The first sintering experiments were carried out at 1400 ◦ C and 1600 ◦ C with an applied heating rate of 200 K/min to investigate the influence of the sintering temperature on the densification behaviour of the multiphase powders. The densification curves shown in Fig. 4(left) revealed that at the intermediate stage of sintering
Fig. 4. Relative densities of the specimens as a function of sintering time at a maximum sintering temperature of 1400 ◦ C (left) and 1600 ◦ C (right), applied heating rate is 200 K/min in both cases.
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the densification behaviour of the powders is comparable in the range of 1000–1400 ◦ C. After reaching of the maximum sintering temperature the densification was interrupted suddenly. By raising the sintering temperature from 1400 to 1600 ◦ C an enhanced densification could be observed at the final stage of sintering. While the samples with the composition of 7Nb3P and 6Nb4P exhibited relative densities of 99.2% and 99.1% after sintering at 1600 ◦ C (S2) the ceramic-like powders 5Nb5P and 4Nb6P showed poor densification. Taking into account that there is a residual Nbamount within the powders 7Nb3P and 6Nb4P the enhanced densification may be a result of plastic deformation of the metallic Nb-phase. Eriksson et al. [20] had the same observations while investigating the densification of (Ti)x (TiB2 )1−x ceramic–metal composites. They showed that due to plastic deformation of the Ti particles the densification was increased and therefore sintering temperatures could be reduced. In contrast, 5Nb5P and 4Nb6P showed decreased sintering activity. It is expected, that the silicide phases on the particle surface counteracts the densification at lower temperatures. Chen et al. [16] obtained while investigating the sintering behaviour of Nb/Nb5 Si3 composites fabricated with the SPS technique that there is need for temperatures above a critical sintering temperature of 1400 ◦ C to achieve highly dense composites. The sinter run S2 revealed comparable results. A general observation of the SPS-experiments is that it is possible to remove residual porosity by lowering the heating rates and prolonging the sintering time, respectively. A drawback of decreased heating rates is enhanced grain growth. To overcome this problem the heating rate was gradually decreased to 100 K/min. The evaluation of the sintering curves plotted in Fig. 5 clarifies that decreasing the heating rate and prolonging the sintering time to at least 18 min lead to highly dense composites with densities for 5Nb5P-S4 of 98.8% and 4Nb6P-S4 97.2%. The compositions of 7Nb3P and 6Nb4P exhibited negligible differences of the relative densities between 3 min and 5 min of holding time. The shrinkage rates of the different powder mixtures revealed information about the sintering kinetics. It is obvious that sintering activity of the ceramic–metal powders (7Nb6P and 6Nb4P) is increased compared to 5Nb5P and 4Nb6P. The maximum densification of 7Nb3P and 6Nb4P took place at temperatures of 1300–1400 ◦ C whereas for the powders with the silicide phases at the particle surface (5Nb5P and 4Nb6P) the maximum densification was observed at 1500 ◦ C. The holding time at maximum sintering temperatures had a distinct influence on densification of the ceramic like Nb(C,N)/Nb5 Si3+x Cx powders (5Nb5P und 4Nb6P). According to the ternary Nb–Si–C phase diagram reported in the literature [32] Nb5 Si4 C is in the solid-state region and no liquid phases should form at these temperatures under inert conditions in a conventional furnace. The melting point of Nb5 Si3 [33] is also higher than the chosen sintering temperature. But it should be noted that during the SPS process the local temperature inside the sample might be higher than what measured. Particularly in the present case when the sample is electrical conductive. As a result, phases which form at a temperature higher than 1600 ◦ C could be also occur locally at the surface of the multiphase particles. Therefore, we assumed that processes like the
5
Fig. 5. Relative densities and shrinkage rates of the different powder mixtures parameter setup S4.
formation of a viscous grain boundary phase or enhanced grain boundary sliding are the dominating densification mechanisms during sintering. But further work is demanded to confirm these effects. The chemical reactions between the multiphase particles or crystallization effects occurring during SPS densification will be discussed afterwards in Section 3.3. The fractured surfaces of the composites sintered with S4 setup are illustrated in Fig. 6. The ceramic–metal composites 7Nb3P-S4 and 6Nb4P-S4 are highly dense. The cleavage-like areas on the fractured surface of 7Nb3P-S4 correspond to ductile-ruptured niobium particles. In contrast, with increasing amount of Nb–Si and Nb(C,N) phases for samples 5Nb5P-S4 and 4Nb6P-S4 the fractured surfaces appear as for brittle materials and ceramics, respectively. No delamination between core and rim is visible and it was possible to retain the origin structure of the particles. The amount of residual niobium within 6Nb4P-S4 seems to have less influence on the fracture mode. 3.3. Phase composition and microstructure The XRD patterns of the samples densified at 1600 ◦ C and 5 min holding time (S4) shown in Fig. 7 reveal slight changes compared to the XRD results presented in Fig. 2. It can be noted that samples 7Nb3P-S4 and 6Nb4P-S4 show peak broadening effects of the Nb2 N and residual Nb-phase. Sample 5Nb5P-S4 exhibited a strong decrease of Nb5Si3 phase from 35.4 wt.%
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Fig. 6. Fractured surfaces of composites after sintering at 1600 ◦ C for 5 min: (a) 7Nb3P-S4, (b) 6Nb4P-S4, (c) 5Nb5P-S4, (d) 4Nb6P-S4.
to 1.6 wt.% by forming Nb5Si4C. Especially, sample 4Nb6PS4 revealed slight changes of the peak intensities. Hence, the amount of Nb2 CN and NbC is increased while Nb2 N declines after the SPS process. Quantitative Rietveld analysis showed that the amount of Nb2 N phase is reduced from 15.5 to 9.6 wt.% after sintering while the amount of metastable Nb5Si4 C remained almost unchanged. Preliminary results on Nb(Si,C,N) composites derived by Nb-particle loaded polyorganosilazanes [18] revealed no detectable changes in the chemical composition of the composite materials after pyrolysis for 30 h at 1600 ◦ C in
Fig. 7. XRD-pattern of composites densified at 1600 ◦ C for 5 min (S4) (␣: Nb5 Si3 , : Nb2 N, ♦: Nb, : Nb2 C, ␥: Nb5 Si4 C, 䊉: NbC, : Nb2 CN).
Argon. Based on the occurrence of Nb5 Si4 C phase and increased amounts of Nb2 CN as well as NbC we assume that an enhanced amount of carbon is incorporated in all of the samples by the reaction with the graphite parts, i.e. of the die and the punches during sintering. Therefore, chemical reactions among the particles during SPS densification cannot be assigned properly. The results of the XRD measurements and the refined lattice parameters were used for determining each phase while running the EBSD analysis. Fig. 8 reveals the results of the analysis exemplarily for the ceramic–metal composite 7Nb3P-S4 and the composite ceramic 4Nb6P-S4. The phase distribution map of 7Nb3P-S4 (Fig. 8b) exhibits coarse-grained Nb5 Si3 with a mean grain size diameter of 1.2 m and Nb2 N phases embedded in a niobium matrix with a size of 14.4 m. The value for the diameter of the Nb2 N grains is 6.8 m. The black dots visible at the grain boundaries in Fig. 8b could be, on the one hand, very small Nb2 N or Nb5 Si3 grains formed in grain boundaries, or originate, on the other hand, from pseudosymmetries or non-indexed structures caused by a shift in lattice parameters of the related Nb2 N or Nb5 Si3 phases due to carbon atoms on interstices. As expected the structure of sample 4Nb6P-S4 is completely different by exhibiting the already mentioned core-rim structure with ceramic Nb(C,N) phases covered by fine-grained Nb5 Si4 C phases with a grain size of 0.9 m (Fig. 8d). Compared to composite 7Nb3P-S4 the size of Nb2 N has decreased to 1.8 m. As it is reported in the literature NbC can inhibit grain growth during pressure assisted sintering [34,35]. The diameter of Nb2 CN and NbC grains was determined to an average value of 0.7 m and 0.8 m, respectively.
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Fig. 8. Results of the EBSD measurement: band contrast maps of 7Nb3P-S4 (a) and 4Nb6P-S4 (b) with the corresponding phase distribution (c and d) with Nb5 Si4 C (4Nb6P-S4) or Nb5Si3 (7Nb3P-S4), Nb2 N, Nb2 CN, NbC and Nb. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
It can be concluded that holding times of at least 5 min at 1600 ◦ C are required to reach relative densities of more than 97%. 3.4. Mechanical properties The Vickers hardness diagram shown in Fig. 9 reveals that there is a dependency of sample hardness on both the phase composition and the sintering parameters.
By decreasing the amount of residual Nb an increase in sample hardness from 8.7 ± 0.1 GPa to 11.3 ± 0.3 GPa could be observed. This effect can be directly related to the reduction of the amount of plastically deformable Nb-matrix phase. Heian et al. [29] observed similar behaviour at Nb/Nb5 Si3 functionally graded composites. By decreasing the Nb-content from 30 vol.% to at least 3 vol.% the Vickers-hardness of the composites was increased from 7.3 GPa to 12 GPa. The highest hardness values of 12.8 ± 0.3 GPa and 12.2 ± 0.4 GPa were measured for the ceramic-like composites 5Nb5P and 4Nb6P. An influence of the core-shell structure or a difference in hardness in relation to Nb5 Si3 and Nb5 Si4 C could not be observed. According to the literature [30,31] there are some approaches to describe the hardness as a function of porosity with the mathematical relationship of Eq. (2). H = H0 · e−αP
Fig. 9. Composite hardness plotted in dependency on composition and sinter parameters.
(2)
Via least-squares fits it is possible to obtain hardness values (H0 ) for the pore free composites and the material specific constant α. The results of these fits are plotted in Fig. 10. A correlation of the curves to the measured hardness values clearly showed a dependency on the porosity. The values for the hardness of the sintered bodies of 7Nb3P-S4 and-S4 6Nb4P are in good agreement with the calculated values of 8.9 ± 0.1 GPa and 11.5 ± 0.3 for the pore-free composites. The values of for 5Nb5P-S4 and 4Nb6P-S4 (14.0 ± 0.4 GPa and 15.4 ± 0.4 GPa) deviates due to the remaining porosity. Furthermore, at porosity values smaller than 20% the influence of phase composition on hardness becomes apparent.
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References
Fig. 10. Fitted profiles for the composite hardness as a function of the sample porosity.
4. Conclusion Nb2 N/Nb5 Si3 /Nb ceramic–metal and Nb(C,N)/Nb5 Si(3+x) Cx ceramic composites were densified by spark plasma sintering (SPS). The required multiphase powders were successfully synthesized via reactive pyrolysis of polysilazane HTTS and Nb by using starting powder materials. It was confirmed that phase composition and particle structure is influenced by the amount of reactive species within the starting composition. SPS is verified as a feasible process for manufacturing dense composites and for retaining the structure and phase composition of the starting powders. Due to plastic deformation the ceramic–metal composites reached relative densities of >99% at 1600 ◦ C after sintering for 3 min with an applied heating rate of 200 K/min. Because of the core-shell structure of the ceramic like particles holding time up to 5 min at a maximum heating rate of 100 K/min lead to final sample densities of >97% th. D. In terms of the densification behaviour of the ceramic like powders with Nb5 Si3 or Nb5 Si4 C phases at the particle surface, processes like the formation of a viscous grain boundary phase or enhanced grain boundary sliding are supposed to be the dominating mechanisms during sintering. To provide relative densities of more than 94% a critical sintering temperature of 1400 ◦ C is necessary for the ceramic–metal composites (7Nb3P-Sx, 6Nb4P-Sx) and 1600 ◦ C for the ceramic composites (5Nb5P-Sx, 4Nb6P-Sx). The results demonstrate that dense Nb2 N/Nb5Si3/Nb ceramic–metal composites as well as Nb(C,N)/Nb5Si(3 + x)Cx ceramic composites can be produced by the SPS process, in which the phase composition of the starting powders is preserved. The Vickers-hardness measurements revealed a strong dependency on the phase composition and microstructure as well as the residual porosity of the composites. Acknowledgements This work was financially supported by the DFG graduate school 1229 “Stable and metastable multiphase systems at high application temperatures” and the European commission Marie
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Nb(Si,C,N) composite materials densified by spark plasma sintering Martin Seifert a , Zhijian Shen b , Walter Krenkel a , Günter Motz a,∗ a
b
University of Bayreuth, Ceramic Materials Engineering (CME), D-95440 Bayreuth, Germany Stockholm University, Department of Materials and Environmental Chemistry, S-106 91 Stockholm, Sweden
Abstract Nb2 N/Nb5 Si3 /Nb ceramic–metal and Nb(C,N)/Nb5 Si(3+x) Cx ceramic like composites were densified by spark plasma sintering (SPS). The precursor multiphase powders were synthesized by solid-state reactions of a polysilazane precursor with niobium powder addition. By varying the amount of reactive species of both precursor and niobium the phase composition and structure of the synthesized multiphase powders can be tailored from metal dominated to mainly ceramic like. The results demonstrated that the SPS approach leads to highly dense samples at 1600 ◦ C with an applied uniaxial pressure of 100 MPa. Furthermore, XRD measurements and EBSD analysis proofed that SPS is a feasible method to retain the original phase compositions and grain sizes within the multiphase powders during sintering. The evaluation of the Vickers indents revealed both a dependency of the measured values on phase composition and on residual porosity. © 2015 Published by Elsevier Ltd. Keywords: Polymer derived ceramics; Spark-plasma sintering; Composite material; Reactive pyrolysis; Refractory metal
1. Introduction In general, advanced ceramics provide exceptional properties like high hardness, resistance to wear, high strength to weight ratio as well as high stiffness and good creep resistance. These materials are lack, however, of plastic deformation at ambient temperatures thus fail normally in a catastrophic manner. Ceramic–metal composites combine the benefits of both ceramics and metals by exhibiting increased fracture toughness because of the ductile deformation behaviour of the metal phases involved. In order to investigate new material combinations for ceramic–metal or ceramic composites for high performance applications niobium and it’s related compounds seem to be natural candidates due to their specific properties. Niobium and binary Nb(C,N) ceramic or intermetallic Nb–Si phases demonstrate high melting points above 2000 ◦ C, high hardness as well as high stiffness and strength at comparatively low densities (7–8 g/cm3 ) [1–7]. Furthermore, as it has been reported in the literature it is possible to improve material properties like strength
∗
Corresponding author. Tel.: +49 921 555505; fax: +49 921 555502. E-mail addresses: [email protected] (M. Seifert), [email protected] (Z. Shen), [email protected] (W. Krenkel), [email protected] (G. Motz).
or hardness of steels by the dispersion of Nb2 N or NbCx Ny phases [8,9]. Especially in situ formed Nb/Nb5 Si3 compounds were intensively investigated because this material combination exhibit very good creep behaviour and high strength at 1300√◦ C and a room temperature fracture-toughness up to 21 MPa m [10–12]. Furthermore, by the dispersion of NbC within in situ Nb/Nb5 Si3 composites crack deflection and branching occur [13]. In terms of the high melting points of Nb and related ceramic or intermetallic phases the processing of dense composites e.g. by liquid metal infiltration of porous ceramic preforms is very complex and difficult to accomplish. The powder metallurgical approach in combination with different sintering techniques like reactive hot pressing (HP), hot isostatic pressing (HIP) or spark plasma sintering (SPS) require elaborate powder preparation steps like mechanical alloying (MA) or self-combustion high temperature synthesis (SHS) to produce sinterable powders [14–17]. Recent work on ceramic composites derived from a carbon rich polysilazane filled with Niobium particles demonstrated the formation of silicide, carbide and nitride phases at the Nb/SiCN interphase via solid-state reaction during precursor pyrolysis at 1600 ◦ C in argon atmosphere [18]. By varying the amount of the reactive educts within the starting composition it is possible to control phase formation and to tailor the microstructure.
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.02.005 0955-2219/© 2015 Published by Elsevier Ltd.
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Nb amount
7Nb3P 6Nb4P 5Nb5P 4Nb6P
Fig. 1. Flow chart for the manufacturing of Nb(C,N)/Nb5 Si3 composites.
By increasing the amount of polymer within the starting compositions an interconnected particle network with a core-rim structure is formed accompanied by formation of porosity values up to 50%. The aim of this work is to obtain dense Nb(Si,C,N) composite materials by additive-free sintering of multiphase powders using SPS technique. 2. Experimental procedure 2.1. Powder synthesis The Nb(Si,C,N) multiphase powders used for this work were synthesized via solid-state reaction of Nb powder and the polysilazane precursor HTTS during pyrolysis. The polymer results from chemical crosslinking of the commercially available liquid polysilazane HTT 1800 (Fa. Clariant Advanced Materials GmbH, Sulzbach, Germany). The crosslinking process is described in detail elsewhere [21]. In order to produce the precursor powder, HTTS is dissolved in pentane and 3 wt.% of the radical initiator dicumylperoxide (DCP 98%, Sigma Aldrich, St. Louis, USA) was added to the solution. By adding DCP the crosslinking temperature of HTTS could be decreased to 150 ◦ C. Subsequently, the solvent was removed with rotary evaporation. The drying process results in brittle HTTS powder agglomerates which can be sieved to a particle size smaller than 63 m. Fig. 1 shows the scheme for manufacturing the composite materials. The starting mixtures containing niobium (Hauner Metallische Werkstoffe, Roethenbach, Germany, 99.8%, <45 m) and HTTS powder were homogenized by dry mixing and subsequent ball milling for 6 h. The resulting mixtures were filled in a heatable stainless steel cylinder for thermal crosslinking of the PCSZ at 160 ◦ C in air and for shaping of cylindrical specimen. Pyrolysis of the cured green bodies in an Al2 O3 tube furnace (Thermal Technology, Santa Rosa, USA) at 1600 ◦ C in Argon (purity 5.0) atmosphere with a heating rate of 3 K/min and annealing for 3 h at maximum temperature lead to specimen with a highly porous and weak bounded particle network with porosities up to 50–70 vol.%. For this reason it is
Powder density after pyrolysis
(vol.%)
(wt.%)
(g/cm3 )
70 60 50 40
94.8 92.1 88.6 83.9
8.11 7.91 7.69 7.43
simple to grind up the porous specimens by hand in a mortar to prepare the multiphase powders with a particle size smaller than 63 m suitable for the SPS process. The composition of the powder mixtures is listed in Table 1. According to the previous work the region between 60 and 50 vol.% of Nb within the starting composition is of utmost interest because it involves the transition from phase compositions dominated by residual metallic niobium to compositions solely containing intermetallic or ceramic phases. In order to investigate the influence on both, phase composition and mechanical properties, the amount of Nb was varied from 70 to 40 vol.%. The designation of the powders corresponds to the composition of the starting powder mixtures. The sintered samples are designated by the affix–Sx while x represents the respective sinter run. 2.2. SPS-process Spark plasma sintering was performed in a SPS facility (Dr. Sinter 2050, Sumitomo Coal Mining Co. Ltd., Tokyo, Japan) under medium vacuum of 6 Pa at 1600 ◦ C. For each experiment a powder volume of 0.2 cm3 was filled into a graphite die with an inner diameter of 12 mm lined with graphite foil. The temperature was first increased to 600 ◦ C in 1 min and after a setting time of 2 min it was raised again to the pre-set sintering temperature by applying different heating rates. The parameters used for different sinter runs are listed in Table 2. During the second heating step the temperature was monitored by a radiation pyrometer focused on the out surface of the graphite die. To ensure reproducibility the maximum pressure of 100 MPa was applied during the setting time at 600 ◦ C and kept constantly at this value for entire period of sintering. Dwell times at the sintering temperature were set to 3 and 5 min for investigating the influence on densification and grain growth. It is expected that the unique sintering mechanisms of SPS facilitated by the applicable high heating and cooling rates will Table 2 Different sintering parameters used for the experiments. Run
Heating rate (K/min)
Pressure (MPa)
Sintering temperature (◦ C)
Holding time (min)
S1 S2 S3 S4
200 200 100 100
100 100 100 100
1400 1600 1600 1600
3 3 3 5
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support densification of the synthesized powders [19] and retain metastable phases, as well as prevent coarsening or changes in the phase compositions of the generated Nb–Si and Nb(C,N) phases. The linear changes of the sample and of the die during sintering were recorded with a dilatometer integrated in the SPS facility. The densification curves were derived from the shrinkage curves considering the relative densities of the sintered samples. The shrinkage curves were corrected for the thermal expansion of the graphite die in order to obtain the real shrinkage curves of the multiphase powders. Eq. (1) concludes the approach of using the normalized data for calculating the relative densities during sintering. ρi = ρs
Is Is + L s − L i
(1)
In this equation ρi is the intermediate density at sintering, ρS is the relative density after sintering, lS is the sample height after sintering. LS and Li are defined as the punch positions during and after sintering. The density of the sintered samples ρS was calculated by measuring the relative density with the Archimedes method and divided by the theoretical density of the milled multiphase powders measured with He-Pycnometry. In order to obtain information about the sintering kinetics the linear shrinkage rate, defined as dl/dt, was plotted as a function of time.
2.3. Characterization methods In order to measure the true density of the synthesized powders He-Pycnometry measurements (Accupyc 1330, Micromeritics, Germany) were carried out. The density and open porosity of the sintered samples were determined using the Archimedes’ method. Before investigating the microstructure, phase composition and hardness, the sintered samples were polished with diamond slurries down to 1 m finish. Crystalline phases were detected by X-ray diffraction analysis of the grinded powders and the polished surfaces of the sintered bodies (D8 ADVANCE, Bruker AXS, Karlsruhe, Germany) using monochromated CuK␣ radiation. The amount of residual Nb was determined by quantitative Rietveld refinement using the evaluation software TOPAS (V4.2, Bruker AXS, Karlsruhe, Germany) and PDF-4+ 2012 structural database for crystallographic information. Weighted Scanning electron microscopy (SEM, JSM6400, Jeol, Japan) was used to characterize the structure and morphology of the particles as well as the microstructure of the sintered samples. Electron back-scattered diffraction analysis (EBSD) was performed by using a field emission SEM (Leo Gemini 1530, Zeiss, Oberkochen, Germany) equipped with Oxford Instruments INCA energy 400 EDS and HKL Nordlys EBSD detectors to yield information about phase distribution of the composites. The Vickers hardness of the composites was determined with the Vickers indentation technique under a load of 98 N for 15 s.
Fig. 2. XRD patterns of the Nb(Si,C,N) powders synthesized at 1600 ◦ C in Ar – Atmosphere (␣: Nb5 Si3 , : Nb2 N, ♦: Nb, : Nb2 C, ␥: Nb5 Si4 C, 䊉:NbC, : Nb2 CN).
3. Results and discussion 3.1. Multiphase powders The pyrolysis of HTTS involves several thermally induced reactions comprising the transition from polymer into an amorphous SiCN ceramic at temperatures between 400–1000 ◦ C, the formation of Si-N, Si-C nano-domains and a free carbon phase at T < 1400 ◦ C, as well as the crystallization of these Si-N and Si-C environments at T > 1400 ◦ C [22–26]. In the presence of reactive metal like Nb the crystallization is inhibited because of reactions at the Nb/SiCN interphase at temperatures exceeding 1000 ◦ C. The mechanism of such reaction and the phase formation were described in detail elsewhere [18]. Fig. 2 shows the results of the XRD-measurements of the synthesized multiphase powders after pyrolysis at 1600 ◦ C in argon atmosphere. For a niobium content of 94.8 wt.% (70 vol.%) within the starting composition the synthesized powder particles are composed of residual metallic Nb, metal rich -Nb2 N and ␣-Nb5 Si3 phase. The amount of residual niobium determined with qualitative Rietveld refinement is 31 vol.% for 7Nb3P (Rwp = 14.32, GOF = 1.53) and 11.0 vol.% for 6Nb4P (Rwp = 16.24, GOF = 1.57). By increasing the amount of HTTS gradually to 60 vol.% the amount of metallic niobium is completely converted by the reaction with SiCN. Furthermore, a low temperature distorted Nb2 C phase (-Fe2 N prototype) is formed [27]. The enhanced amount of carbon also supports the formation of metastable Nb5 Si4 C which is the ␥-modification (Mn5 Si3 prototype) of tetragonal ␣-Nb5 Si3 with carbon atoms located at octahedral sites of the Nb–Si lattice. Nb2 N is in equilibrium with Nb5 Si3 . Murakami et al. reached the same results by investigating Nb–Si–N composites prepared with spark plasma sintering [28]. Cubic Nb2 CN was detected within the 4Nb6P powder and it may be a result of the reaction of Nb2 N with
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Fig. 3. Nb(Si,C,N) multiphase powders synthesized via solid state reaction of different Nb/HTTS mixtures: (a) 7Nb3P, (b) 6Nb4P, (c) 5Nb5P, (d) 4Nb6P.
carbon phases or enhanced carbon diffusion. No crystalline SiC or Si3 N4 phases, typically resulting from the crystallization of the amorphous SiCN phase, could be detected. It can be assumed that the amount of HTTS is completely converted during reaction by forming Nb(Si,C,N) phases. EBSD measurements carried out in the previous work proofed already that when the added amount of Nb was 50 vol.% or less particles with a core-rim structure were formed, which composed of intermetallic Nb5 Si3 and Nb5 Si4 C phases at the rim and ceramic Nb(C,N) phases at the particle core. Fig. 3 displays examples of the synthesized particles. It is obvious that particle morphology is strongly dependent on the phase composition of the particles. According to the previous work the structures at the niobium particle surfaces of 7Nb3P (Fig. 3a) can be ascribed to Nb2 N and Nb5 Si3 phases. The particle surfaces of the powders 5Nb5P and 4Nb6P (Fig. 3c,d) are composed of
coarse-grained Nb5 Si3 phase and fine-grained Nb5 Si4 C, respectively. 3.2. Densification As related to the pyrolysis temperature selected for powder synthesis the sintering trials were carried out at a maximum sintering temperature of 1600 ◦ C to avoid the phase transformations. Different heating rates of 100 and 200 K/min and holding times of 3 and 5 min were chosen in order to analyse the densification behaviour and the sintering kinetics. The first sintering experiments were carried out at 1400 ◦ C and 1600 ◦ C with an applied heating rate of 200 K/min to investigate the influence of the sintering temperature on the densification behaviour of the multiphase powders. The densification curves shown in Fig. 4(left) revealed that at the intermediate stage of sintering
Fig. 4. Relative densities of the specimens as a function of sintering time at a maximum sintering temperature of 1400 ◦ C (left) and 1600 ◦ C (right), applied heating rate is 200 K/min in both cases.
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the densification behaviour of the powders is comparable in the range of 1000–1400 ◦ C. After reaching of the maximum sintering temperature the densification was interrupted suddenly. By raising the sintering temperature from 1400 to 1600 ◦ C an enhanced densification could be observed at the final stage of sintering. While the samples with the composition of 7Nb3P and 6Nb4P exhibited relative densities of 99.2% and 99.1% after sintering at 1600 ◦ C (S2) the ceramic-like powders 5Nb5P and 4Nb6P showed poor densification. Taking into account that there is a residual Nbamount within the powders 7Nb3P and 6Nb4P the enhanced densification may be a result of plastic deformation of the metallic Nb-phase. Eriksson et al. [20] had the same observations while investigating the densification of (Ti)x (TiB2 )1−x ceramic–metal composites. They showed that due to plastic deformation of the Ti particles the densification was increased and therefore sintering temperatures could be reduced. In contrast, 5Nb5P and 4Nb6P showed decreased sintering activity. It is expected, that the silicide phases on the particle surface counteracts the densification at lower temperatures. Chen et al. [16] obtained while investigating the sintering behaviour of Nb/Nb5 Si3 composites fabricated with the SPS technique that there is need for temperatures above a critical sintering temperature of 1400 ◦ C to achieve highly dense composites. The sinter run S2 revealed comparable results. A general observation of the SPS-experiments is that it is possible to remove residual porosity by lowering the heating rates and prolonging the sintering time, respectively. A drawback of decreased heating rates is enhanced grain growth. To overcome this problem the heating rate was gradually decreased to 100 K/min. The evaluation of the sintering curves plotted in Fig. 5 clarifies that decreasing the heating rate and prolonging the sintering time to at least 18 min lead to highly dense composites with densities for 5Nb5P-S4 of 98.8% and 4Nb6P-S4 97.2%. The compositions of 7Nb3P and 6Nb4P exhibited negligible differences of the relative densities between 3 min and 5 min of holding time. The shrinkage rates of the different powder mixtures revealed information about the sintering kinetics. It is obvious that sintering activity of the ceramic–metal powders (7Nb6P and 6Nb4P) is increased compared to 5Nb5P and 4Nb6P. The maximum densification of 7Nb3P and 6Nb4P took place at temperatures of 1300–1400 ◦ C whereas for the powders with the silicide phases at the particle surface (5Nb5P and 4Nb6P) the maximum densification was observed at 1500 ◦ C. The holding time at maximum sintering temperatures had a distinct influence on densification of the ceramic like Nb(C,N)/Nb5 Si3+x Cx powders (5Nb5P und 4Nb6P). According to the ternary Nb–Si–C phase diagram reported in the literature [32] Nb5 Si4 C is in the solid-state region and no liquid phases should form at these temperatures under inert conditions in a conventional furnace. The melting point of Nb5 Si3 [33] is also higher than the chosen sintering temperature. But it should be noted that during the SPS process the local temperature inside the sample might be higher than what measured. Particularly in the present case when the sample is electrical conductive. As a result, phases which form at a temperature higher than 1600 ◦ C could be also occur locally at the surface of the multiphase particles. Therefore, we assumed that processes like the
5
Fig. 5. Relative densities and shrinkage rates of the different powder mixtures parameter setup S4.
formation of a viscous grain boundary phase or enhanced grain boundary sliding are the dominating densification mechanisms during sintering. But further work is demanded to confirm these effects. The chemical reactions between the multiphase particles or crystallization effects occurring during SPS densification will be discussed afterwards in Section 3.3. The fractured surfaces of the composites sintered with S4 setup are illustrated in Fig. 6. The ceramic–metal composites 7Nb3P-S4 and 6Nb4P-S4 are highly dense. The cleavage-like areas on the fractured surface of 7Nb3P-S4 correspond to ductile-ruptured niobium particles. In contrast, with increasing amount of Nb–Si and Nb(C,N) phases for samples 5Nb5P-S4 and 4Nb6P-S4 the fractured surfaces appear as for brittle materials and ceramics, respectively. No delamination between core and rim is visible and it was possible to retain the origin structure of the particles. The amount of residual niobium within 6Nb4P-S4 seems to have less influence on the fracture mode. 3.3. Phase composition and microstructure The XRD patterns of the samples densified at 1600 ◦ C and 5 min holding time (S4) shown in Fig. 7 reveal slight changes compared to the XRD results presented in Fig. 2. It can be noted that samples 7Nb3P-S4 and 6Nb4P-S4 show peak broadening effects of the Nb2 N and residual Nb-phase. Sample 5Nb5P-S4 exhibited a strong decrease of Nb5Si3 phase from 35.4 wt.%
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Fig. 6. Fractured surfaces of composites after sintering at 1600 ◦ C for 5 min: (a) 7Nb3P-S4, (b) 6Nb4P-S4, (c) 5Nb5P-S4, (d) 4Nb6P-S4.
to 1.6 wt.% by forming Nb5Si4C. Especially, sample 4Nb6PS4 revealed slight changes of the peak intensities. Hence, the amount of Nb2 CN and NbC is increased while Nb2 N declines after the SPS process. Quantitative Rietveld analysis showed that the amount of Nb2 N phase is reduced from 15.5 to 9.6 wt.% after sintering while the amount of metastable Nb5Si4 C remained almost unchanged. Preliminary results on Nb(Si,C,N) composites derived by Nb-particle loaded polyorganosilazanes [18] revealed no detectable changes in the chemical composition of the composite materials after pyrolysis for 30 h at 1600 ◦ C in
Fig. 7. XRD-pattern of composites densified at 1600 ◦ C for 5 min (S4) (␣: Nb5 Si3 , : Nb2 N, ♦: Nb, : Nb2 C, ␥: Nb5 Si4 C, 䊉: NbC, : Nb2 CN).
Argon. Based on the occurrence of Nb5 Si4 C phase and increased amounts of Nb2 CN as well as NbC we assume that an enhanced amount of carbon is incorporated in all of the samples by the reaction with the graphite parts, i.e. of the die and the punches during sintering. Therefore, chemical reactions among the particles during SPS densification cannot be assigned properly. The results of the XRD measurements and the refined lattice parameters were used for determining each phase while running the EBSD analysis. Fig. 8 reveals the results of the analysis exemplarily for the ceramic–metal composite 7Nb3P-S4 and the composite ceramic 4Nb6P-S4. The phase distribution map of 7Nb3P-S4 (Fig. 8b) exhibits coarse-grained Nb5 Si3 with a mean grain size diameter of 1.2 m and Nb2 N phases embedded in a niobium matrix with a size of 14.4 m. The value for the diameter of the Nb2 N grains is 6.8 m. The black dots visible at the grain boundaries in Fig. 8b could be, on the one hand, very small Nb2 N or Nb5 Si3 grains formed in grain boundaries, or originate, on the other hand, from pseudosymmetries or non-indexed structures caused by a shift in lattice parameters of the related Nb2 N or Nb5 Si3 phases due to carbon atoms on interstices. As expected the structure of sample 4Nb6P-S4 is completely different by exhibiting the already mentioned core-rim structure with ceramic Nb(C,N) phases covered by fine-grained Nb5 Si4 C phases with a grain size of 0.9 m (Fig. 8d). Compared to composite 7Nb3P-S4 the size of Nb2 N has decreased to 1.8 m. As it is reported in the literature NbC can inhibit grain growth during pressure assisted sintering [34,35]. The diameter of Nb2 CN and NbC grains was determined to an average value of 0.7 m and 0.8 m, respectively.
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Fig. 8. Results of the EBSD measurement: band contrast maps of 7Nb3P-S4 (a) and 4Nb6P-S4 (b) with the corresponding phase distribution (c and d) with Nb5 Si4 C (4Nb6P-S4) or Nb5Si3 (7Nb3P-S4), Nb2 N, Nb2 CN, NbC and Nb. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
It can be concluded that holding times of at least 5 min at 1600 ◦ C are required to reach relative densities of more than 97%. 3.4. Mechanical properties The Vickers hardness diagram shown in Fig. 9 reveals that there is a dependency of sample hardness on both the phase composition and the sintering parameters.
By decreasing the amount of residual Nb an increase in sample hardness from 8.7 ± 0.1 GPa to 11.3 ± 0.3 GPa could be observed. This effect can be directly related to the reduction of the amount of plastically deformable Nb-matrix phase. Heian et al. [29] observed similar behaviour at Nb/Nb5 Si3 functionally graded composites. By decreasing the Nb-content from 30 vol.% to at least 3 vol.% the Vickers-hardness of the composites was increased from 7.3 GPa to 12 GPa. The highest hardness values of 12.8 ± 0.3 GPa and 12.2 ± 0.4 GPa were measured for the ceramic-like composites 5Nb5P and 4Nb6P. An influence of the core-shell structure or a difference in hardness in relation to Nb5 Si3 and Nb5 Si4 C could not be observed. According to the literature [30,31] there are some approaches to describe the hardness as a function of porosity with the mathematical relationship of Eq. (2). H = H0 · e−αP
Fig. 9. Composite hardness plotted in dependency on composition and sinter parameters.
(2)
Via least-squares fits it is possible to obtain hardness values (H0 ) for the pore free composites and the material specific constant α. The results of these fits are plotted in Fig. 10. A correlation of the curves to the measured hardness values clearly showed a dependency on the porosity. The values for the hardness of the sintered bodies of 7Nb3P-S4 and-S4 6Nb4P are in good agreement with the calculated values of 8.9 ± 0.1 GPa and 11.5 ± 0.3 for the pore-free composites. The values of for 5Nb5P-S4 and 4Nb6P-S4 (14.0 ± 0.4 GPa and 15.4 ± 0.4 GPa) deviates due to the remaining porosity. Furthermore, at porosity values smaller than 20% the influence of phase composition on hardness becomes apparent.
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References
Fig. 10. Fitted profiles for the composite hardness as a function of the sample porosity.
4. Conclusion Nb2 N/Nb5 Si3 /Nb ceramic–metal and Nb(C,N)/Nb5 Si(3+x) Cx ceramic composites were densified by spark plasma sintering (SPS). The required multiphase powders were successfully synthesized via reactive pyrolysis of polysilazane HTTS and Nb by using starting powder materials. It was confirmed that phase composition and particle structure is influenced by the amount of reactive species within the starting composition. SPS is verified as a feasible process for manufacturing dense composites and for retaining the structure and phase composition of the starting powders. Due to plastic deformation the ceramic–metal composites reached relative densities of >99% at 1600 ◦ C after sintering for 3 min with an applied heating rate of 200 K/min. Because of the core-shell structure of the ceramic like particles holding time up to 5 min at a maximum heating rate of 100 K/min lead to final sample densities of >97% th. D. In terms of the densification behaviour of the ceramic like powders with Nb5 Si3 or Nb5 Si4 C phases at the particle surface, processes like the formation of a viscous grain boundary phase or enhanced grain boundary sliding are supposed to be the dominating mechanisms during sintering. To provide relative densities of more than 94% a critical sintering temperature of 1400 ◦ C is necessary for the ceramic–metal composites (7Nb3P-Sx, 6Nb4P-Sx) and 1600 ◦ C for the ceramic composites (5Nb5P-Sx, 4Nb6P-Sx). The results demonstrate that dense Nb2 N/Nb5Si3/Nb ceramic–metal composites as well as Nb(C,N)/Nb5Si(3 + x)Cx ceramic composites can be produced by the SPS process, in which the phase composition of the starting powders is preserved. The Vickers-hardness measurements revealed a strong dependency on the phase composition and microstructure as well as the residual porosity of the composites. Acknowledgements This work was financially supported by the DFG graduate school 1229 “Stable and metastable multiphase systems at high application temperatures” and the European commission Marie
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