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Carbon influence on fracture toughness of niobium carbides
Journal of the European Ceramic Society 39 (2019) 5167–5173
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Original Article
Carbon influence on fracture toughness of niobium carbides a
b
b
T c,d
Xingyuan Zhao , Maanas Togaru , Qianying Guo , Christopher R. Weinberger , ⁎ Leslie Lambersona, , Gregory B. Thompsonb a
Department of Mechanical Engineering, Colorado School of Mines, Golden, CO, 80401, United States Department of Metallurgical and Materials Engineering, University of Alabama, Tuscaloosa, AL, 35487, United States Department of Mechanical Engineering, Colorado State University, Fort Collins, CO, 80523, United States d School of Advanced Materials Discovery, Colorado State University, Fort Collins, CO, 80523, United States b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Niobium carbides Fracture toughness Hardness Knoop indentation Crack propagation Microstructure
This paper explores the fracture behavior of niobium carbides of varying compositions between NbC1.0 and NbC0.5. The surface crack in flexure (SCF) method was used to evaluate the fracture toughness as a function of carbon concentration. Additionally, hardness measurements were conducted with a Knoop indenter, and X-ray diffraction (XRD) and scanning electron microscopy (SEM) were used to identify the phase content and microstructures. As the carbon content decreased, the hardness increased from 8 GPa for NbC1.0 to 12 GPa for NbC0.5 and the fracture toughness decreased from 2.5 MPa m to 0.44 MPa m . Notably, the NbC0.67 sample exhibited a secondary precipitate lath-like microstructure with the laths indexed to β-Nb2C and a KIC near 2 MPa m . Though similar lath like structures in tantalum carbides have been reported to yield a KIC of approximately 15 MPa m , the laths in these two materials have fundamentally different structures where bonding in the former is comprised of β-Nb2C and the latter of ζ-Ta4C3-x. This results in the observed different fracture properties, which can be explained through concepts of microstructural toughening.
1. Introduction The group VB transition metal carbides (TMCs) comprise a collection of ultrahigh temperature ceramics (UHTCs) that display excellent stability in extreme environments [1]. They are often characterized by their extraordinary melting points, which are mostly above 3000 °C, high hardness, strength, and wear resistance. As a result, these materials have found applications for bearings, cutting tools, nuclear cladding, and thermal protection systems [2–4]. As would be expected, these properties are linked to the crystal structure of the carbide phase, with its structure dictated by the carbon-to-metal ratio. With changes in this ratio, diverse and wide phase stability fields develop, as shown in Fig. 1, whose structures are largely based on close packing of the metal atoms with carbon atoms filling the octahedral interstices. These structures exhibit a mixture of covalent, metallic, and ionic bonding between the transition metal and carbon atoms [3]. In the mixed phase fields, the two present phases often have a preferred orientation relationship with each other such that particular morphologies precipitate, enabling the microstructure to have a profound impact on the carbide’s strength and fracture properties [5–9]. This microstructure-driven property has garnered considerable
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attention for the tantalum carbides. The precipitation of the zeta phase, ζ-Ta4C3-x, results in either a crisscross pattern of nanometer thick lathes within equiaxed TaC grains or a series of parallel lathes when it precipitates from Ta2C grains which causes it, under sinter processing, to form acicular grains [7]. These precipitation events are driven by the close-packed orientation relationship of the zeta phase with either the four variants of {111} planes in the rock-salt B1 TaC structure or the single {0001} close-packed plane from the trigonal Ta2C structure [7,12,13]. This orientation relationship is driven by associated low interfacial energy between the zeta and matrix phases, a direct consequence of the orientation relationships, and the low thermodynamic driving forces for coarsening [14,15]. When a high volume fraction of the zeta phase is present, the fracture toughness in the tantalum carbides has been shown to increase significantly. Hackett el al. [5] presented a comprehensive fracture investigation of single and mixed-phase tantalum carbides by hot pressing them at 1800 °C for 120 min with C/Ta ratios of 0.6, 0.7, 0.8, 0.9 and 1.0, and evaluating the toughness using a single-edge-precracked beam (SEPB) method [16]. In the single phase TaC composition between a C/Ta ratio of 0.8 to 1.0, the fracture toughness decreased (ranging from 5.3 to 3.8 MPa m ) whereas hardness increased (ranging from 13.5 to
Corresponding author at: Mechanical Engineering Department, Colorado School of Mines, 1500 Illinois Street, Golden, CO 80401, United States. E-mail address: [email protected] (L. Lamberson).
https://doi.org/10.1016/j.jeurceramsoc.2019.08.022 Received 18 June 2019; Accepted 17 August 2019 Available online 18 August 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
Journal of the European Ceramic Society 39 (2019) 5167–5173
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Fig. 1. Phase diagram of (a) tantalum carbides [10] and (b) niobium carbide, with the tick marks indicating overall compositions investigated [11].
be drawn from this prior work in the tantalum carbides is if other group VB TMCs exhibit similar fracture toughness values. To answer this question and provide additional insight into the fracture behavior of these materials, we turn our attention to the niobium-carbon system which shares many phases with the tantalum carbide system and hence may share similar microstructures and toughness. A comparison of the tantalum-carbon and niobium-carbon phase diagrams can be seen in Fig. 1. Since both metals are in the same periodic table family, we expect similar phases to form, which they do and provide initial evidence that niobium carbides could have comparable microstructure and toughness behavior as similarly processed tantalum carbides. Specifically, both carbides share a B1 (rocksalt) monocarbide phase which can retain its equilibrium structure with significant loss of carbon. Upon cooling from the substoichiometric high temperature B1 phase field, both carbides precipitate a M6C5 vacancy ordered rocksalt phase. However, this vacancy ordered phase is rarely reported because of the difficulties associated with the ordering of the carbon vacancies. With further carbon loss, the tantalum carbides precipitate the ζ-Ta4C3-x phase, with DFT work by Yu et al. [21] and Weinberger et al. [22] confirming its relative phase stability. In Gusev’s phase diagram shown in Fig. 1, the niobium carbide’s ζ-Nb4C3-x is listed as only a high temperature phase existing between 1294 °C (1592 K) to 1552 °C (1850 K). However, the presence of the zeta phase on this phase diagram is taken from the work of Wiesenberger et al. [23] who showed ζ-Nb4C3-x being stable below 1575 °C and Weinberger and Thompson [24] have pointed out that, based on the literature, this phase should be stable at low temperatures and that its lower decomposition temperature prevents it from being observed in many samples that are processed at high temperatures. When it was present, Wiesenberger et al. [23] described that the ζ phase precipitates in a preferred orientation in β -Nb2C phase as needle shape grains or in between NbC and β -Nb2C phase as bands, which would be similar to the aforementioned tantalum carbide based laths. The phase diagram of Gusev predicts a vacancy ordered Nb3C2 phase, but that has yet to be reported experimentally [24] and would not exist if the zeta phase is stable at low temperatures. The Nb2C phase is reported to have three polytypes which are a function of temperature [24–28] whereas the Ta2C has two polytypes, a lower temperature trigonal α-Ta2C and a hexagonal L’3 β-Ta2C structure. Though the phase diagram by Gusev correctly shows the orderdisorder transition for Ta2C, Fig. 1a, the niobium carbide diagram, Fig. 1b, only shows the β-Nb2C phase. The correct phase sequence for Nb2C has a high temperature γ phase which has the L’3 structure, an intermediate temperature trigonal β-Nb2C phase (which has the ε-Fe2N prototype structure), and a low temperature α-Nb2C phase which is orthorhombic [25] with all three structures in agreement with DFT simulations [21,29]. Though there are some subtle differences in the
20 GPa) with a decreasing C/Ta ratio. Meanwhile the mean grain size decreased slightly from 7.75 to 5.82 μm. For the two-phase mixture composition at a C/Ta ratio of 0.7 (55.4 wt% TaC and 44.6 wt% of ζTa4C3) and 0.6 (21 wt% Ta2C and 79 wt% of ζ-Ta4C3), the fracture toughness exhibited a significant increase from 7.9 MPa m to 12.7 MPa m , respectively with hardness value at 0.6 being 9.2 GPa. Note that the wt. % given in the parenthesis are the equilibrium fractions based on a tie line from the phase diagram reported by Gusev [11] and serves only as a guide to the reader to the expected equilibrium fraction of the zeta phase. More often than not in these two-phase fields, a metastable mixture of three phases of TaCx, Ta4C3-x, and Ta2C exists because of the experimental challenges of explicitly forming only two phases [17]. This high fracture toughness result was later confirmed by Liming et al. [6] who reported 13.8 MPa m in a TaC0.7 sample that was sparkplasma-sintered (SPS) at 1700 °C. These authors further showed that when the sintering temperature was increased from 1800 °C and 1900 °C, the thickness of ζ-Ta4C3 grains increased while the fracture toughness decreased suggesting a microstructure contribution to the mechanical response. Motivated by these reports, Sygnatowicz et al. [8] examined the fracture toughness, fracture strength, resistance curves and hardness of tantalum carbides with high weight contents of ζTa4C3-x. These authors noted an indirect relationship between fracture toughness and hardness, with a direct relationship between fracture toughness and grain size (and hence an indirect relationship between grain size and hardness). The tantalum carbides exhibited a rising Rcurve behavior which is attributed largely to grain bridging, facilitated by easy cleavage along basal planes of the zeta phase. The idea of weak fracture along the basal planes is further supported by density functional theory (DFT) calculations [18] which revealed that the metalmetal planes in the zeta phase laths are indeed quite weak compared to the metal-carbon planes in the zeta phase and in TaC. Based on this microstructure evidence, Schulz et al. [9] examined the crack propagation in a tantalum carbide TaC0.67 which had a high volume content of the ζ-Ta4C3-x phase using micro- and nano-indentation relative to the indent and lath orientations. The indentation results showed long, straight cracks parallel to the zeta phase laths and short cracks normal to the laths, indicating crack propagation is easy along the basal plane of the zeta phase but not normal to it. This would be in agreement with the aforementioned DFT simulations [18]. Similar type of basal cracking behavior has been reported in the MAX phases Ti3AlC2 and Ti3SiC2 systems [19,20] suggesting synergy between these studies where links between microstructure and properties have been drawn. Collectively, these findings have now suggested that nanoscale precipitation of the zeta phase laths contributes to the high fracture toughness on the tantalum carbides. An outstanding question that can 5168
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2.2. Microstructure characterization The microstructure of the four compositions was imaged by scanning electron microscopy (SEM) using a JEOL 7000 FE with electron backscattered diffraction (EBSD) for phase and grain size analysis. The secondary and/or backscattered electron micrographs were used to estimate the porosity in the samples by taking a relative ratio of the collective pore area to the total area of the scanned image in order to calculate a void fraction. At least three images from different areas of the sample were used to calculate overall porosity, adjusting the magnification depending on the feature sizes to resolve the pores within a user defined field of view. As EBSD can be difficult to discern the subtle phase differences in the metal-rich niobium carbides [30], X-ray diffraction was also performed to further confirm the phase structures for these phases. The XRD was performed on a Bruker D8 Discover diffractometer operated under 40 keV and 35 mA. The XRD scans also enabled the volume fraction of the phases to be computed as previously done in procedures reported in [17,30].
2.3. Fracture characterization The HIP-ed carbide billets were electro-discharged machined into beam-shape specimens with dimensions of approximately 4 mm × 5 mm × 45 mm, and four to five specimens were tested at each composition. To remove any detrimental surface formations from machining, all samples were hand polished using 600 grit SiC paper. Fracture toughness was measured using the surface crack in flexure (SCF) method following ASTM standard C1421 [16]. Each specimen was pre-cracked using a Knoop indenter for 12 s at 25 N on a Shimadzu AG-IC 50 kN load frame with a 1.5 ° tilt. The indentation impression was created on the 3 mm x 45 mm surface, in order to minimize material removal necessary to avoid residual stresses [31,32]. The impression was measured under a microscope (as shown for NbC0.67 in Fig. 2), and subsequently the surface was polished to a 30 μm finish. The specimen was then tested in a three-point bend fixture on a AG-IC load frame with a 50 kN load cell at a loading rate of 0.5 mm/min. The upper fixture was equipped with an Epsilon LVDT deflectometer to track the displacement. Force at fracture was determined from the instantaneous drop of the load. A cotton pad was placed between the lower fixture supports to prevent post-fracture damage. After failure, fracture surfaces were imaged using SEM to determine the pre-crack size. If no stable crack extension was found, the Knoop impression length was assumed to be the critical crack length [16]. The SCF method was chosen over the SENB (single edge notch bend) method due to the limitation of manufactured sample sizes; however it has been shown that the SENB and SCF techniques show fairly good agreement between experimentally measured and estimated critical flaw sizes, and are suitable for obtaining reliable fracture toughness measurements for hard metals [33].
Fig. 2. Schematic of samples used for fracture investigations. (i) Specimen size (not to scale) showing orientation of the indention impression (ii) Micrograph of the Knoop indentation impression prior to fracture testing for a NbC0.67 specimen.
phases between the niobium and tantalum carbides, the fact that they share a majority of similar phases suggests they consequently also share similar microstructural features. For example, in a recent report by Smith et al. [30], the precipitation of the β-Nb2C phase within NbCx grains resulted in nanometer thick lath-like morphology as those seen in the zeta phase containing tantalum carbides. In this work, we begin an investigation to determine how reproducible the high toughness may be in these transition metal carbides that share such similarities. Instead of examining the potential for toughening from the ζ-Nb4C3-x phase, which can be difficult to process because of its low decomposition temperature, we have instead focused on the easier to form β-Nb2C phase. This particular phase presents an interesting case study because the phase can form laths similar to ζTa4C3-x [30] however, unlike the zeta phase, the β-Nb2C phase does not have basal planes comprised predominantly of metal-metal bonds, but carbon-metal bonds across all of its basal planes. This will allow us to investigate if the microstructure is necessary or sufficient for the high fracture toughness.
2. Materials and methods 2.1. Specimen fabrication
2.4. Deformation and crack propagation
The niobium carbide billets were fabricated by mixing commercial Nb (ABCR GmbH & Co, < 45 μm, 99.8%) and NbC (ABCR GmbH & Co, < 1.1 μm, 99.7%) powder at various ratios with C/Nb composition ratios of 1.0, 0.75, 0.67, and 0.5. Mixture of powders were hot isostatic pressed (HIP) in an argon environment at 205 MPa at 1800 °C for 1 h, and cooled in a canning furnace at 27 °C /min to room temperature before removal [30]. Thermo-combustion analysis (LECO) provided by a certified commercial vendor (Northern Analytical) confirmed the carbon content for each sample, with the oxygen levels < 0.308 ± 0.015% wt.% for all samples. The average starting powder sizes of Nb and NbC was 8.25 μm and 0.67 μm, respectively [30].
Deformation and crack propagation of the niobium carbides were also studied using Knoop indentation similar to the study performed by Schultz et al [9]. The Knoop indentation was also used to calculate material hardness using ASTM C1326 [34]. Specimens were machinepolished to maintain flatness using 1 μm diamond paste, and then polished with aqueous 0.05 μm silica slurry for 4 h on Buehler VibroMet polishing machine to produce a clean surface finish. The polished surfaces were indented with the Knoop indenter at 49 N for 20 s using a AG-IC load frame equipped with a 5 kN load cell. Each sample was indented at 3 locations spread out on the polished surface, and the resulting indentation impressions were imaged via SEM.
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3. Results & discussion The XRD scans for the niobium carbides are shown in Fig. 3 with the various SEM micrographs for each composition shown in Fig. 4. According to the XRD scans, the NbC1.0 and NbC0.75 were confirmed to contain the single phase B1 rocksalt structure. The presence of the B1 structure would be expected for the NbC1.0 as it was sintered using only the NbC starting powders, whereas NbC0.75 was a mixture of both NbC and Nb powders. Upon a closer inspection of Fig. 1b, at the HIP maximum processing temperature and at this sub-stoichiometry, the B1 structure can still accommodate the depletion of carbon from its sublattice and thus a single B1 structure is present. Upon cooling from the HIP temperature, we would expect that the ζ-Nb4C3-x and Nb6C5 phases to precipitate out of the NbCx matrix. The fact that we did not observe either phase indicates that the NbCx phase is quenched in and the cooling was too fast to either precipitate out the zeta phase or to allow carbon vacancy ordering in the rocksalt structure. Both the NbC1.0 and NbC0.75 samples revealed equiaxed grains with averages sizes of 1.2 ± 0.4 μm and 23.5 ± 8.5 μm, respectively, in Fig. 4(a)-(b), with the grain size differences in the NbC0.75 attributed to the metal powder addition to the monocarbide powders, with similar powder mixing and final grain size variations seen in other hot isostatic pressed carbides [7,30].
Fig. 3. XRD scan of the various niobium carbide samples.
Fig. 4. SEM micrographs for (a) NbC1.0 (b) NbC0.75 (c) NbC0.67 and (d) NbC0.5. Within each image (i) represents a backscattered micrograph with the dark box being a magnified image of the region to the right of (i). (ii) represents the magnified image from (i). (iii) electron backscattered diffraction image revealing the texture of each grain by the inverted pole coordination. (iv) the grain size distribution from (iii). 5170
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1.5 ± 0.1 MPa m with the hardness increasing to 11.5 ± 2.7 GPa. The NbC0.67 sample recovered the fracture toughness slightly to 2.0 ± 0.3 MPa m with its hardness also slightly increasing to 13.5 ± 1.3 GPa. This sample, as with the others, retained the equiaxed NbCx grains but now has Nb2C laths within it (as shown in Fig. 4(c)-ii). Note this sample also has the lowest porosity at 0.9%. Finally, the NbC0.5 is mostly composed of equiaxed β-Nb2C but, as noted above, some distinct regions also revealed small NbC grains with β-Nb2C laths. This sample’s fracture toughness significantly decreased to 0.4 ± 0.2 MPa m with the hardness only slightly decreased to 11.8 ± 1.4 GPa. Reviewing the hardness trends for the niobium carbides as shown in Table 1, the loss of carbon results in an increase in these values. Though the NbC0.67 sample has the highest average value, it falls within the statistical errors of the NbC0.75 and NbC0.5 samples. The grain size continually increased with decreasing carbon content, but the hardness did not statistically change over these grain sizes, suggesting that grain size effect does not significantly contribute to the hardness values. Niobium carbides are known to exhibit an anomalous hardness rise between NbC1.0 and NbC0.86 [36,37], whose origin is debated [38–41], but also has a decrease with subsequent carbon loss. The hardness trends noted here are likely a convolution of this anomalous hardness rise and the porosity decrease observed with carbon loss. The fracture toughness values exhibit an indirect relationship to the hardness and generally decrease with carbon loss. The only exception being the modest increase noted for the NbC0.67 composition, but it is still lower than that of the NbC1.0 sample and likely caused by the reduction in porosity instead of the microstructure of the phases. The low fracture toughness value for the NbC0.67, while having a lath-like microstructure, does not show the exceptional toughness observed in the tantalum carbides processed at the same temperature and carbon composition Consequently, there does not appear to be any trend of the fracture toughness relative to grain size, suggesting the expected result that grain size does not influence fracture toughness. The low fracture toughness of NbC0.5, as compared to the other niobium carbides, may be a result of the substantial number of microcracks observed in the sample. Some of the microcracks are of similar size as the initial indent used to initiate fracture, which could influence the calculated fracture toughness. The ability to precipitate in the niobium carbides a similar lath microstructure as observed in the tantalum carbides, but with a chemically and crystallographic distinct phase does provide an opportunity to better understand the microstructure influence of the laths on fracture toughness in the transition metal carbides. In this case, the β-Nb2C is a crystal structure that is based on an hexagonal close-packed (HCP) ordering of the metal atoms and the carbon atoms ordering in the octahedra interstices, whereas the zeta phase is comprised of alternating face centered cubic (FCC)/HCP stacking sequences that are a result of carbon depleted stacking faults. The fact that the laths in the NbC0.67 samples did not offer the same toughening value observed in the tantalum carbides raises the question of why not; or alternatively, what is different between the two cases? The answer to this question has been proposed by Yu et al. in their theoretical investigation into the influence of bonding in the group IVB transition metal carbides [18]. Notably, they point out that the depletion of carbon planes associated with the change in stacking sequence creates preferred cleavage planes in the zeta phase (or between any carbon depleted stacking fault). This preferred cleavage plane, and the complex microstructure, causes crack meandering, crack deflection and most importantly grain bridging, resulting in an increase in fracture toughness. That fracture behavior is evident in the tantalum carbon system when the zeta phase is present [8,9]. Consequently the crystallographic structure of the lath appears to play a significant role in the toughening mechanism, and is not present in the niobium carbides as the β-Nb2C phase does not have carbon depleted planes. Its structure has ordered carbon atoms on all the basal planes and thus there is a
The XRD scan of the NbC0.67 sample, which is predicted to be in the NbCx and β-Nb2C phase field during its maximum HIP condition, confirms the existence of both of these phases. The lack of the Nb6C5, Nb3C2, or ζ-Nb4C3-x phases in the NbC0.67 sample, as predicted by the phase diagram of Fig. 1b and the work of Wiesenberger, is again attributed to kinetic limitations or the stability of the zeta phase, none of which are present at the high temperature processing condition. The microstructure contained equiaxed grains with an average size of 44.9 ± 22.5 μm. Within these grains a series of nanometer scale Nb2C laths can be seen, via the arrows in Fig. 4c-ii serving as a guide to the eye. In Fig. 4c-ii, a magnified image of a few grains clearly reveal the laths which were then grey scale threshold to yield an estimate of the phase fraction, determined to be approximately 26% based on a ratio of lath surface coverage over the entire image surface. Note that the EBSD images, 4c(iii) only shows limited number of laths due to the limitations of the EBSD method in detecting the misorientation on the scale of the laths. Finally the NbC0.5 sample is processed in the single Nb2C phase field. The XRD of this sample revealed the β-Nb2C phase, with a small fraction of NbC. The presence of NbCx is attributed to incomplete homogenization and reaction between the NbC and Nb starting metal powders. This can be seen in the lower enlarged SEM micrograph, Fig. 4(d)-ii. In this grain, distinct Nb2C laths within the NbCx grains is present. As noted in the recent paper by Smith et al.[30], as the carbon diffuses out of NbC and reacts with the Nb metal, it forms the larger, equiaxed Nb2C grains. As the monocarbide grains become depleted in carbon at the HIP temperature, Fig. 1b, the carbon deficient monocarbide precipitates the hemi-carbide yielding the lath-like microstructure within those grains. Since the NbC was the smaller starting powder, the incomplete homogenization leads to these isolated multiphase regions in the consolidated microstructure. Based on the XRD intensity ratio [35], the NbC comprises of approximately 17% phase fraction. For this sample, the average equiaxed grain size is 126.3 ± 68.8 μm. The larger grain size with decreasing carbon content is again attributed to the larger starting Nb metal powder as compared to the smaller NbC powder, along with the faster diffusivity of metal for grain growth as compared to a carbide. In addition, it is clear that the NbC0.5 sample has a substantial number of microcracks, most likely a result of the shrinkage associated with NbC grains that loose carbon during processing. The tabulation of the measured mechanical properties and relevant microstructure features are given in Table 1. As shown in Table 1, the porosity of the billets decreased once the metal-rich carbides precipitated. Since the HIP process temperature was constant for all billets, and the metal-rich carbides have slightly lower melting temperatures than the monocarbide, the homologous sintering temperature would be higher for these compositions and likely facilitated more densification for the same absolute temperature and time enabling the billets then to achieve a lower porosity. Furthermore, the local exothermic release that would occur by the reaction of the carbon with the metal to form these phases would also contribute to enhancing densification. The average fracture toughness value of 2.5 ± 0.6 MPa m for the NbC1.0 was found to be the highest among all four compositions whereas its Knoop hardness, 8.4 ± 3.2 GPa, was the lowest value among all the compositions. With loss of carbon but a retention of the same B1 phase, the fracture toughness of the NbC0.75 sample decreased to
Table 1 Various experimental results of the niobium carbides.
KIC (MPa m ) HK (GPa) Grain Size (μm) Porosity (%)
NbC1.0
NbC0.75
NbC0.67
NbC0.5
2.5 ± 0.5 8.4 ± 3.2 1.2 ± 0.4 5.9
1.5 ± 0.1 11.5 ± 2.7 23.5 ± 8.5 6.8
2.0 ± 0.3 13.5 ± 1.3 44.9 ± 22.5 0.9
0.4 ± 0.2 11.8 ± 1.4 126.3 ± 68.8 3.4
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However, the NbC0.67 samples, with β-Nb2C laths, had a fracture toughness near 2 MPa m . The lack of change in toughness highlights the importance of local metal-metal bonding that occurs in carbon depleted planes of the zeta phase, but is absent in the β-Nb2C phase. It would be interesting to further see if the ζ-Nb4C3-x phase can create a microstructure that will toughen niobium carbide composites, although we hypothesize that the processing of such large monolithic ceramics will be difficult due to the speculated lower processing temperature (ζNb4C3-x’s decomposition temperature is 1575 °C) for this phase, which will increase porosity in the microstructure. Acknowledgements C.R. Weinberger and G.B. Thompson recognize Air Force Office of Scientific Research grant FA9550-15-1-0217andFA9550-15-1-0095, respectively. L. Lamberson gratefully acknowledges partial support from the NSF CAREER Grant No. 1751989. References [1] W.G. 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Shetty, [ζ]-Ta4C3−x: A High Fracture Toughness Carbide with Rising-Crack-Growth-Resistance (R-Curve) Behavior. American Ceramic Society, J. Am. Ceram. Soc. 98 (8) (2015) 2601. [9] B.C. Schulz, et al., Experimental investigation into the crack propagation in multiphase tantalum carbide ceramics, Mater. Sci. Eng. A 695 (2017) 315–321. [10] A.I. Gusev, A.S. Kurlov, V.N. Lipatnikov, Atomic and vacancy ordering in carbide ζTa4C3−x (0.28⩽x⩽0.40) and phase equilibria in the Ta–C system, J. Solid State Chem. 180 (11) (2007) 3234–3246. [11] A.I. Gusev, Order–disorder transformations and phase equilibria in strongly nonstoichiometric compounds, Physics-Uspekhi 43 (1) (2000) 1–37. [12] D.J. Rowcliffe, G. Thomas, Structure of non-stoichiometric TaC, Mater. Sci. Eng. 18 (2) (1975) 231–238. [13] M. Lewis, J. Billingham, P. Bell, Nonstoichiometry in Ceramic Compounds, Univ. of Warwick, Coventry, Eng., 1973. [14] C.R. Weinberger, et al., A computational investigation into the microstructures and stability of the zeta phase in transition metal carbides and nitrides, Adv. Appl. Ceram. 117 (sup1) (2018) s26–s33. [15] H. Yu, et al., A model for understanding the formation energies of nanolamellar phases in transition metal carbides and nitrides, Model. Simul. Mater. Sci. Eng. 24 (5) (2016) 55004. [16] ASTM, ASTM Standard C 1421-01b: Standard Test Method for Determination of Fracture Toughness of Advanced Ceramics at Ambient Temperature, 100 Barr Harbor Dr., PO Box C 700, West Conshohocken, PA, 2019 19428. [17] R.A. Morris, et al., Variations in tantalum carbide microstructures with changing carbon content, Int. J. Appl. Ceram. Technol. 10 (3) (2013) 540–551. [18] H. Yu, G.B. Thompson, C.R. Weinberger, The role of chemistry and bonding in regulating fracture in multiphase transition metal carbides and nitrides, Extreme Mech. Lett. 17 (2017) 1–6. [19] N.V. Tzenov, M.W. Barsoum, Synthesis and characterization of Ti3AlC2, J. Am. Ceram. Soc. 83 (4) (2000) 825–832. [20] T. El-Raghy, et al., Damage mechanisms around hardness indentations in Ti3SiC2, J. Am. Ceram. Soc. 80 (2) (1997) 513–516. [21] X.-X. Yu, C.R. Weinberger, G.B. Thompson, Ab initio investigations of the phase stability in group IVB and VB transition metal carbides, Comput. Mater. Sci. 112 (2016) 318–326. [22] C.R. Weinberger, G.B. Thompson, A computational search for the zeta phase in the tantalum carbides, J. Am. Ceram. Soc. 102 (3) (2019) 1454–1462. [23] H. Wiesenberger, W. Lengauer, P. Ettmayer, Reactive diffusion and phase equilibria in the V–C, Nb–C, Ta–C and Ta–N systems, Acta Mater. 46 (2) (1998) 651–666. [24] C.R. Weinberger, G.B. Thompson, Review of phase stability in the group IVB and VB transition‐metal carbides, J. Am. Ceram. Soc. 101 (10) (2018) 4401–4424. [25] T. Epicier, P. Convert, High temperature powder neutron diffraction studies of structural transformations in transition metal hemicarbides M2C1− x: I. Compounds of the Vth group (V2C, Nb2C, Ta2C), Phys. B Condens. Matter 156 (1989) 41–43. [26] B. Lönnberg, T. Lundström, Thermal expansion and phase analytical studies of Nb2C, J. Less Common Met. 113 (2) (1985) 261–268.
Fig. 5. Crack propagation initiated from the Knoop indent on a NbC0.75 sample. Note how the crack travels relatively unimpeded, with the crack propagating through the laths (dashed circled region).
reduced preference for fracture. Consequently, there is likely a much more homogenous fracture energy in β-Nb2C as compared to the zeta phase. The experimental report here provides evidence for supporting the hypothesis of Yu et al. [18] that the metal-metal bonding across the close-packed planes is critical in creating the high fracture toughness and provides further insight that not only nanoscale laths that create a torturous path are required, but the crystallography of those laths must create the localized metal-metal bonding within its unit cell to control the fracture. This notion can be confirmed by SEM micrographs of the crack propagation in the NbC0.67 sample. Fig. 5 reveals intergranular cracking with the cracks propagating easily through the laths, even with the cracks oriented normal to the laths (dashed circle region). In the work of Schultz et al.[9] such propagation in ζ-Ta4C3-x stunted crack growth. This finding provides the clearest evidence that the β-Nb2C laths provide minimal resistance to micro-cracking, and therefore contributes little to the fracture toughness as evidenced by its low value. 4. Conclusion In this work, we examined the fracture toughness and hardness of a series of niobium carbides at various carbon contents of C/Nb 1.0, 0.75, 0.67, and 0.5. The NbC1.0 and NbC0.75 samples composed equiaxed B1 rocksalt structures whereas NbC0.67 was an equiaxed NbCx matrix with laths of β-Nb2C. Finally equiaxed NbC0.5 was fabricated, but not all of the NbC converted to Nb2C revealing smaller grains of NbC with laths of Nb2C. As the carbon content decreased, the grains grew in size from approximately 1 μm to 120 μm. The grain growth is contributed to not only the larger starting Nb metal powder as compared to NbC used in mixing the composition of the billets, but also the easier diffusivity of the metal as compared to the carbide at the HIP temperature of 1800 °C. The hardness increased from approximately 8 GPa from NbC1.0 to 12 GPa for NbC0.5, with the fracture toughness decreasing from approximately 2.5 MPa m to 0.4 MPa m over the same range. The NbC0.67 samples comprised of a similar microstructure to the high fracture toughness tantalum carbides at similar compositions. The tantalum carbides have a lath based microstructure associated with the presence of the Ta4C3-x phase. These laths, comprised of carbon depleted stacking faults, provide toughening that results in fracture toughness values as high as 15 MPa m observed experimentally. 5172
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Advanced Ceramics, ASTM International, West Conshohocken, PA, 2019. [35] B. Warren, X-ray Diffraction, Dover, New York, 1990, p. 253. [36] H. Holleck, Material selection for hard coatings, J. Vac. Sci. Technol. A Vac. Surf. Films 4 (6) (1986) 2661–2669. [37] I. Vinitskii, Relation between the properties of monocarbides of groups IV–V transition metals and their carbon content, Powder Metall. Met. Ceram. 11 (6) (1972) 488–493. [38] X.-X. Yu, G.B. Thompson, C.R. Weinberger, Influence of carbon vacancy formation on the elastic constants and hardening mechanisms in transition metal carbides, J. Eur. Ceram. Soc. 35 (1) (2015) 95–103. [39] S.-H. Jhi, et al., Vacancy hardening and softening in transition metal carbides and nitrides, Phys. Rev. Lett. 86 (15) (2001) 3348. [40] G. Morgan, M. Lewis, Hardness anisotropy in niobium carbide, J. Mater. Sci. 9 (3) (1974) 349–358. [41] D. Rowcliffe, G. Hollox, Hardness anisotropy, deformation mechanisms and brittleto-ductile transition in carbide, J. Mater. Sci. 6 (10) (1971) 1270–1276.
[27] B. Vishwanadh, et al., Synthesis and phase transformation mechanism of Nb2C carbide phases, J. Alloys Compd. 671 (2016) 424–434. [28] S. Alyamovskii, et al., Orthorhombic β′ phases of niobium and vanadium carbides, Russ. J. Inorg. Chem. 12 (1967) 301. [29] L. Wu, et al., The phase stability and mechanical properties of Nb–C system: Using first-principles calculations and nano-indentation, J. Alloys Compd. 561 (2013) 220–227. [30] C.J. Smith, C.R. Weinberger, G.B. Thompson, Phase stability and microstructural formations in the niobium carbides, J. Eur. Ceram. Soc. 38 (15) (2018) 4850–4866. [31] J.J. Petrovic, et al., Effects of residual stresses on fracture from controlled surface flaws, J. Am. Ceram. Soc. 59 (3–4) (1976) 177–178. [32] J.J. Petrovic, Effect of indenter geometry on controlled-surface-Flaw fracture toughness, J. Am. Ceram. Soc. 66 (4) (1983) 277–283. [33] Y. Torres, D. Casellas, M. Anglada, L. Llanes, Fracture toughness evalution of hardmetals: influence of testing procedure, Int. J. Refract. Metals Hard Mater. 19 (1) (2001) 27–34. [34] ASTM, ASTM C1326 Standard Test Method for Knoop Indentation Hardness of
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Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Original Article
Carbon influence on fracture toughness of niobium carbides a
b
b
T c,d
Xingyuan Zhao , Maanas Togaru , Qianying Guo , Christopher R. Weinberger , ⁎ Leslie Lambersona, , Gregory B. Thompsonb a
Department of Mechanical Engineering, Colorado School of Mines, Golden, CO, 80401, United States Department of Metallurgical and Materials Engineering, University of Alabama, Tuscaloosa, AL, 35487, United States Department of Mechanical Engineering, Colorado State University, Fort Collins, CO, 80523, United States d School of Advanced Materials Discovery, Colorado State University, Fort Collins, CO, 80523, United States b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Niobium carbides Fracture toughness Hardness Knoop indentation Crack propagation Microstructure
This paper explores the fracture behavior of niobium carbides of varying compositions between NbC1.0 and NbC0.5. The surface crack in flexure (SCF) method was used to evaluate the fracture toughness as a function of carbon concentration. Additionally, hardness measurements were conducted with a Knoop indenter, and X-ray diffraction (XRD) and scanning electron microscopy (SEM) were used to identify the phase content and microstructures. As the carbon content decreased, the hardness increased from 8 GPa for NbC1.0 to 12 GPa for NbC0.5 and the fracture toughness decreased from 2.5 MPa m to 0.44 MPa m . Notably, the NbC0.67 sample exhibited a secondary precipitate lath-like microstructure with the laths indexed to β-Nb2C and a KIC near 2 MPa m . Though similar lath like structures in tantalum carbides have been reported to yield a KIC of approximately 15 MPa m , the laths in these two materials have fundamentally different structures where bonding in the former is comprised of β-Nb2C and the latter of ζ-Ta4C3-x. This results in the observed different fracture properties, which can be explained through concepts of microstructural toughening.
1. Introduction The group VB transition metal carbides (TMCs) comprise a collection of ultrahigh temperature ceramics (UHTCs) that display excellent stability in extreme environments [1]. They are often characterized by their extraordinary melting points, which are mostly above 3000 °C, high hardness, strength, and wear resistance. As a result, these materials have found applications for bearings, cutting tools, nuclear cladding, and thermal protection systems [2–4]. As would be expected, these properties are linked to the crystal structure of the carbide phase, with its structure dictated by the carbon-to-metal ratio. With changes in this ratio, diverse and wide phase stability fields develop, as shown in Fig. 1, whose structures are largely based on close packing of the metal atoms with carbon atoms filling the octahedral interstices. These structures exhibit a mixture of covalent, metallic, and ionic bonding between the transition metal and carbon atoms [3]. In the mixed phase fields, the two present phases often have a preferred orientation relationship with each other such that particular morphologies precipitate, enabling the microstructure to have a profound impact on the carbide’s strength and fracture properties [5–9]. This microstructure-driven property has garnered considerable
⁎
attention for the tantalum carbides. The precipitation of the zeta phase, ζ-Ta4C3-x, results in either a crisscross pattern of nanometer thick lathes within equiaxed TaC grains or a series of parallel lathes when it precipitates from Ta2C grains which causes it, under sinter processing, to form acicular grains [7]. These precipitation events are driven by the close-packed orientation relationship of the zeta phase with either the four variants of {111} planes in the rock-salt B1 TaC structure or the single {0001} close-packed plane from the trigonal Ta2C structure [7,12,13]. This orientation relationship is driven by associated low interfacial energy between the zeta and matrix phases, a direct consequence of the orientation relationships, and the low thermodynamic driving forces for coarsening [14,15]. When a high volume fraction of the zeta phase is present, the fracture toughness in the tantalum carbides has been shown to increase significantly. Hackett el al. [5] presented a comprehensive fracture investigation of single and mixed-phase tantalum carbides by hot pressing them at 1800 °C for 120 min with C/Ta ratios of 0.6, 0.7, 0.8, 0.9 and 1.0, and evaluating the toughness using a single-edge-precracked beam (SEPB) method [16]. In the single phase TaC composition between a C/Ta ratio of 0.8 to 1.0, the fracture toughness decreased (ranging from 5.3 to 3.8 MPa m ) whereas hardness increased (ranging from 13.5 to
Corresponding author at: Mechanical Engineering Department, Colorado School of Mines, 1500 Illinois Street, Golden, CO 80401, United States. E-mail address: [email protected] (L. Lamberson).
https://doi.org/10.1016/j.jeurceramsoc.2019.08.022 Received 18 June 2019; Accepted 17 August 2019 Available online 18 August 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. Phase diagram of (a) tantalum carbides [10] and (b) niobium carbide, with the tick marks indicating overall compositions investigated [11].
be drawn from this prior work in the tantalum carbides is if other group VB TMCs exhibit similar fracture toughness values. To answer this question and provide additional insight into the fracture behavior of these materials, we turn our attention to the niobium-carbon system which shares many phases with the tantalum carbide system and hence may share similar microstructures and toughness. A comparison of the tantalum-carbon and niobium-carbon phase diagrams can be seen in Fig. 1. Since both metals are in the same periodic table family, we expect similar phases to form, which they do and provide initial evidence that niobium carbides could have comparable microstructure and toughness behavior as similarly processed tantalum carbides. Specifically, both carbides share a B1 (rocksalt) monocarbide phase which can retain its equilibrium structure with significant loss of carbon. Upon cooling from the substoichiometric high temperature B1 phase field, both carbides precipitate a M6C5 vacancy ordered rocksalt phase. However, this vacancy ordered phase is rarely reported because of the difficulties associated with the ordering of the carbon vacancies. With further carbon loss, the tantalum carbides precipitate the ζ-Ta4C3-x phase, with DFT work by Yu et al. [21] and Weinberger et al. [22] confirming its relative phase stability. In Gusev’s phase diagram shown in Fig. 1, the niobium carbide’s ζ-Nb4C3-x is listed as only a high temperature phase existing between 1294 °C (1592 K) to 1552 °C (1850 K). However, the presence of the zeta phase on this phase diagram is taken from the work of Wiesenberger et al. [23] who showed ζ-Nb4C3-x being stable below 1575 °C and Weinberger and Thompson [24] have pointed out that, based on the literature, this phase should be stable at low temperatures and that its lower decomposition temperature prevents it from being observed in many samples that are processed at high temperatures. When it was present, Wiesenberger et al. [23] described that the ζ phase precipitates in a preferred orientation in β -Nb2C phase as needle shape grains or in between NbC and β -Nb2C phase as bands, which would be similar to the aforementioned tantalum carbide based laths. The phase diagram of Gusev predicts a vacancy ordered Nb3C2 phase, but that has yet to be reported experimentally [24] and would not exist if the zeta phase is stable at low temperatures. The Nb2C phase is reported to have three polytypes which are a function of temperature [24–28] whereas the Ta2C has two polytypes, a lower temperature trigonal α-Ta2C and a hexagonal L’3 β-Ta2C structure. Though the phase diagram by Gusev correctly shows the orderdisorder transition for Ta2C, Fig. 1a, the niobium carbide diagram, Fig. 1b, only shows the β-Nb2C phase. The correct phase sequence for Nb2C has a high temperature γ phase which has the L’3 structure, an intermediate temperature trigonal β-Nb2C phase (which has the ε-Fe2N prototype structure), and a low temperature α-Nb2C phase which is orthorhombic [25] with all three structures in agreement with DFT simulations [21,29]. Though there are some subtle differences in the
20 GPa) with a decreasing C/Ta ratio. Meanwhile the mean grain size decreased slightly from 7.75 to 5.82 μm. For the two-phase mixture composition at a C/Ta ratio of 0.7 (55.4 wt% TaC and 44.6 wt% of ζTa4C3) and 0.6 (21 wt% Ta2C and 79 wt% of ζ-Ta4C3), the fracture toughness exhibited a significant increase from 7.9 MPa m to 12.7 MPa m , respectively with hardness value at 0.6 being 9.2 GPa. Note that the wt. % given in the parenthesis are the equilibrium fractions based on a tie line from the phase diagram reported by Gusev [11] and serves only as a guide to the reader to the expected equilibrium fraction of the zeta phase. More often than not in these two-phase fields, a metastable mixture of three phases of TaCx, Ta4C3-x, and Ta2C exists because of the experimental challenges of explicitly forming only two phases [17]. This high fracture toughness result was later confirmed by Liming et al. [6] who reported 13.8 MPa m in a TaC0.7 sample that was sparkplasma-sintered (SPS) at 1700 °C. These authors further showed that when the sintering temperature was increased from 1800 °C and 1900 °C, the thickness of ζ-Ta4C3 grains increased while the fracture toughness decreased suggesting a microstructure contribution to the mechanical response. Motivated by these reports, Sygnatowicz et al. [8] examined the fracture toughness, fracture strength, resistance curves and hardness of tantalum carbides with high weight contents of ζTa4C3-x. These authors noted an indirect relationship between fracture toughness and hardness, with a direct relationship between fracture toughness and grain size (and hence an indirect relationship between grain size and hardness). The tantalum carbides exhibited a rising Rcurve behavior which is attributed largely to grain bridging, facilitated by easy cleavage along basal planes of the zeta phase. The idea of weak fracture along the basal planes is further supported by density functional theory (DFT) calculations [18] which revealed that the metalmetal planes in the zeta phase laths are indeed quite weak compared to the metal-carbon planes in the zeta phase and in TaC. Based on this microstructure evidence, Schulz et al. [9] examined the crack propagation in a tantalum carbide TaC0.67 which had a high volume content of the ζ-Ta4C3-x phase using micro- and nano-indentation relative to the indent and lath orientations. The indentation results showed long, straight cracks parallel to the zeta phase laths and short cracks normal to the laths, indicating crack propagation is easy along the basal plane of the zeta phase but not normal to it. This would be in agreement with the aforementioned DFT simulations [18]. Similar type of basal cracking behavior has been reported in the MAX phases Ti3AlC2 and Ti3SiC2 systems [19,20] suggesting synergy between these studies where links between microstructure and properties have been drawn. Collectively, these findings have now suggested that nanoscale precipitation of the zeta phase laths contributes to the high fracture toughness on the tantalum carbides. An outstanding question that can 5168
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2.2. Microstructure characterization The microstructure of the four compositions was imaged by scanning electron microscopy (SEM) using a JEOL 7000 FE with electron backscattered diffraction (EBSD) for phase and grain size analysis. The secondary and/or backscattered electron micrographs were used to estimate the porosity in the samples by taking a relative ratio of the collective pore area to the total area of the scanned image in order to calculate a void fraction. At least three images from different areas of the sample were used to calculate overall porosity, adjusting the magnification depending on the feature sizes to resolve the pores within a user defined field of view. As EBSD can be difficult to discern the subtle phase differences in the metal-rich niobium carbides [30], X-ray diffraction was also performed to further confirm the phase structures for these phases. The XRD was performed on a Bruker D8 Discover diffractometer operated under 40 keV and 35 mA. The XRD scans also enabled the volume fraction of the phases to be computed as previously done in procedures reported in [17,30].
2.3. Fracture characterization The HIP-ed carbide billets were electro-discharged machined into beam-shape specimens with dimensions of approximately 4 mm × 5 mm × 45 mm, and four to five specimens were tested at each composition. To remove any detrimental surface formations from machining, all samples were hand polished using 600 grit SiC paper. Fracture toughness was measured using the surface crack in flexure (SCF) method following ASTM standard C1421 [16]. Each specimen was pre-cracked using a Knoop indenter for 12 s at 25 N on a Shimadzu AG-IC 50 kN load frame with a 1.5 ° tilt. The indentation impression was created on the 3 mm x 45 mm surface, in order to minimize material removal necessary to avoid residual stresses [31,32]. The impression was measured under a microscope (as shown for NbC0.67 in Fig. 2), and subsequently the surface was polished to a 30 μm finish. The specimen was then tested in a three-point bend fixture on a AG-IC load frame with a 50 kN load cell at a loading rate of 0.5 mm/min. The upper fixture was equipped with an Epsilon LVDT deflectometer to track the displacement. Force at fracture was determined from the instantaneous drop of the load. A cotton pad was placed between the lower fixture supports to prevent post-fracture damage. After failure, fracture surfaces were imaged using SEM to determine the pre-crack size. If no stable crack extension was found, the Knoop impression length was assumed to be the critical crack length [16]. The SCF method was chosen over the SENB (single edge notch bend) method due to the limitation of manufactured sample sizes; however it has been shown that the SENB and SCF techniques show fairly good agreement between experimentally measured and estimated critical flaw sizes, and are suitable for obtaining reliable fracture toughness measurements for hard metals [33].
Fig. 2. Schematic of samples used for fracture investigations. (i) Specimen size (not to scale) showing orientation of the indention impression (ii) Micrograph of the Knoop indentation impression prior to fracture testing for a NbC0.67 specimen.
phases between the niobium and tantalum carbides, the fact that they share a majority of similar phases suggests they consequently also share similar microstructural features. For example, in a recent report by Smith et al. [30], the precipitation of the β-Nb2C phase within NbCx grains resulted in nanometer thick lath-like morphology as those seen in the zeta phase containing tantalum carbides. In this work, we begin an investigation to determine how reproducible the high toughness may be in these transition metal carbides that share such similarities. Instead of examining the potential for toughening from the ζ-Nb4C3-x phase, which can be difficult to process because of its low decomposition temperature, we have instead focused on the easier to form β-Nb2C phase. This particular phase presents an interesting case study because the phase can form laths similar to ζTa4C3-x [30] however, unlike the zeta phase, the β-Nb2C phase does not have basal planes comprised predominantly of metal-metal bonds, but carbon-metal bonds across all of its basal planes. This will allow us to investigate if the microstructure is necessary or sufficient for the high fracture toughness.
2. Materials and methods 2.1. Specimen fabrication
2.4. Deformation and crack propagation
The niobium carbide billets were fabricated by mixing commercial Nb (ABCR GmbH & Co, < 45 μm, 99.8%) and NbC (ABCR GmbH & Co, < 1.1 μm, 99.7%) powder at various ratios with C/Nb composition ratios of 1.0, 0.75, 0.67, and 0.5. Mixture of powders were hot isostatic pressed (HIP) in an argon environment at 205 MPa at 1800 °C for 1 h, and cooled in a canning furnace at 27 °C /min to room temperature before removal [30]. Thermo-combustion analysis (LECO) provided by a certified commercial vendor (Northern Analytical) confirmed the carbon content for each sample, with the oxygen levels < 0.308 ± 0.015% wt.% for all samples. The average starting powder sizes of Nb and NbC was 8.25 μm and 0.67 μm, respectively [30].
Deformation and crack propagation of the niobium carbides were also studied using Knoop indentation similar to the study performed by Schultz et al [9]. The Knoop indentation was also used to calculate material hardness using ASTM C1326 [34]. Specimens were machinepolished to maintain flatness using 1 μm diamond paste, and then polished with aqueous 0.05 μm silica slurry for 4 h on Buehler VibroMet polishing machine to produce a clean surface finish. The polished surfaces were indented with the Knoop indenter at 49 N for 20 s using a AG-IC load frame equipped with a 5 kN load cell. Each sample was indented at 3 locations spread out on the polished surface, and the resulting indentation impressions were imaged via SEM.
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3. Results & discussion The XRD scans for the niobium carbides are shown in Fig. 3 with the various SEM micrographs for each composition shown in Fig. 4. According to the XRD scans, the NbC1.0 and NbC0.75 were confirmed to contain the single phase B1 rocksalt structure. The presence of the B1 structure would be expected for the NbC1.0 as it was sintered using only the NbC starting powders, whereas NbC0.75 was a mixture of both NbC and Nb powders. Upon a closer inspection of Fig. 1b, at the HIP maximum processing temperature and at this sub-stoichiometry, the B1 structure can still accommodate the depletion of carbon from its sublattice and thus a single B1 structure is present. Upon cooling from the HIP temperature, we would expect that the ζ-Nb4C3-x and Nb6C5 phases to precipitate out of the NbCx matrix. The fact that we did not observe either phase indicates that the NbCx phase is quenched in and the cooling was too fast to either precipitate out the zeta phase or to allow carbon vacancy ordering in the rocksalt structure. Both the NbC1.0 and NbC0.75 samples revealed equiaxed grains with averages sizes of 1.2 ± 0.4 μm and 23.5 ± 8.5 μm, respectively, in Fig. 4(a)-(b), with the grain size differences in the NbC0.75 attributed to the metal powder addition to the monocarbide powders, with similar powder mixing and final grain size variations seen in other hot isostatic pressed carbides [7,30].
Fig. 3. XRD scan of the various niobium carbide samples.
Fig. 4. SEM micrographs for (a) NbC1.0 (b) NbC0.75 (c) NbC0.67 and (d) NbC0.5. Within each image (i) represents a backscattered micrograph with the dark box being a magnified image of the region to the right of (i). (ii) represents the magnified image from (i). (iii) electron backscattered diffraction image revealing the texture of each grain by the inverted pole coordination. (iv) the grain size distribution from (iii). 5170
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1.5 ± 0.1 MPa m with the hardness increasing to 11.5 ± 2.7 GPa. The NbC0.67 sample recovered the fracture toughness slightly to 2.0 ± 0.3 MPa m with its hardness also slightly increasing to 13.5 ± 1.3 GPa. This sample, as with the others, retained the equiaxed NbCx grains but now has Nb2C laths within it (as shown in Fig. 4(c)-ii). Note this sample also has the lowest porosity at 0.9%. Finally, the NbC0.5 is mostly composed of equiaxed β-Nb2C but, as noted above, some distinct regions also revealed small NbC grains with β-Nb2C laths. This sample’s fracture toughness significantly decreased to 0.4 ± 0.2 MPa m with the hardness only slightly decreased to 11.8 ± 1.4 GPa. Reviewing the hardness trends for the niobium carbides as shown in Table 1, the loss of carbon results in an increase in these values. Though the NbC0.67 sample has the highest average value, it falls within the statistical errors of the NbC0.75 and NbC0.5 samples. The grain size continually increased with decreasing carbon content, but the hardness did not statistically change over these grain sizes, suggesting that grain size effect does not significantly contribute to the hardness values. Niobium carbides are known to exhibit an anomalous hardness rise between NbC1.0 and NbC0.86 [36,37], whose origin is debated [38–41], but also has a decrease with subsequent carbon loss. The hardness trends noted here are likely a convolution of this anomalous hardness rise and the porosity decrease observed with carbon loss. The fracture toughness values exhibit an indirect relationship to the hardness and generally decrease with carbon loss. The only exception being the modest increase noted for the NbC0.67 composition, but it is still lower than that of the NbC1.0 sample and likely caused by the reduction in porosity instead of the microstructure of the phases. The low fracture toughness value for the NbC0.67, while having a lath-like microstructure, does not show the exceptional toughness observed in the tantalum carbides processed at the same temperature and carbon composition Consequently, there does not appear to be any trend of the fracture toughness relative to grain size, suggesting the expected result that grain size does not influence fracture toughness. The low fracture toughness of NbC0.5, as compared to the other niobium carbides, may be a result of the substantial number of microcracks observed in the sample. Some of the microcracks are of similar size as the initial indent used to initiate fracture, which could influence the calculated fracture toughness. The ability to precipitate in the niobium carbides a similar lath microstructure as observed in the tantalum carbides, but with a chemically and crystallographic distinct phase does provide an opportunity to better understand the microstructure influence of the laths on fracture toughness in the transition metal carbides. In this case, the β-Nb2C is a crystal structure that is based on an hexagonal close-packed (HCP) ordering of the metal atoms and the carbon atoms ordering in the octahedra interstices, whereas the zeta phase is comprised of alternating face centered cubic (FCC)/HCP stacking sequences that are a result of carbon depleted stacking faults. The fact that the laths in the NbC0.67 samples did not offer the same toughening value observed in the tantalum carbides raises the question of why not; or alternatively, what is different between the two cases? The answer to this question has been proposed by Yu et al. in their theoretical investigation into the influence of bonding in the group IVB transition metal carbides [18]. Notably, they point out that the depletion of carbon planes associated with the change in stacking sequence creates preferred cleavage planes in the zeta phase (or between any carbon depleted stacking fault). This preferred cleavage plane, and the complex microstructure, causes crack meandering, crack deflection and most importantly grain bridging, resulting in an increase in fracture toughness. That fracture behavior is evident in the tantalum carbon system when the zeta phase is present [8,9]. Consequently the crystallographic structure of the lath appears to play a significant role in the toughening mechanism, and is not present in the niobium carbides as the β-Nb2C phase does not have carbon depleted planes. Its structure has ordered carbon atoms on all the basal planes and thus there is a
The XRD scan of the NbC0.67 sample, which is predicted to be in the NbCx and β-Nb2C phase field during its maximum HIP condition, confirms the existence of both of these phases. The lack of the Nb6C5, Nb3C2, or ζ-Nb4C3-x phases in the NbC0.67 sample, as predicted by the phase diagram of Fig. 1b and the work of Wiesenberger, is again attributed to kinetic limitations or the stability of the zeta phase, none of which are present at the high temperature processing condition. The microstructure contained equiaxed grains with an average size of 44.9 ± 22.5 μm. Within these grains a series of nanometer scale Nb2C laths can be seen, via the arrows in Fig. 4c-ii serving as a guide to the eye. In Fig. 4c-ii, a magnified image of a few grains clearly reveal the laths which were then grey scale threshold to yield an estimate of the phase fraction, determined to be approximately 26% based on a ratio of lath surface coverage over the entire image surface. Note that the EBSD images, 4c(iii) only shows limited number of laths due to the limitations of the EBSD method in detecting the misorientation on the scale of the laths. Finally the NbC0.5 sample is processed in the single Nb2C phase field. The XRD of this sample revealed the β-Nb2C phase, with a small fraction of NbC. The presence of NbCx is attributed to incomplete homogenization and reaction between the NbC and Nb starting metal powders. This can be seen in the lower enlarged SEM micrograph, Fig. 4(d)-ii. In this grain, distinct Nb2C laths within the NbCx grains is present. As noted in the recent paper by Smith et al.[30], as the carbon diffuses out of NbC and reacts with the Nb metal, it forms the larger, equiaxed Nb2C grains. As the monocarbide grains become depleted in carbon at the HIP temperature, Fig. 1b, the carbon deficient monocarbide precipitates the hemi-carbide yielding the lath-like microstructure within those grains. Since the NbC was the smaller starting powder, the incomplete homogenization leads to these isolated multiphase regions in the consolidated microstructure. Based on the XRD intensity ratio [35], the NbC comprises of approximately 17% phase fraction. For this sample, the average equiaxed grain size is 126.3 ± 68.8 μm. The larger grain size with decreasing carbon content is again attributed to the larger starting Nb metal powder as compared to the smaller NbC powder, along with the faster diffusivity of metal for grain growth as compared to a carbide. In addition, it is clear that the NbC0.5 sample has a substantial number of microcracks, most likely a result of the shrinkage associated with NbC grains that loose carbon during processing. The tabulation of the measured mechanical properties and relevant microstructure features are given in Table 1. As shown in Table 1, the porosity of the billets decreased once the metal-rich carbides precipitated. Since the HIP process temperature was constant for all billets, and the metal-rich carbides have slightly lower melting temperatures than the monocarbide, the homologous sintering temperature would be higher for these compositions and likely facilitated more densification for the same absolute temperature and time enabling the billets then to achieve a lower porosity. Furthermore, the local exothermic release that would occur by the reaction of the carbon with the metal to form these phases would also contribute to enhancing densification. The average fracture toughness value of 2.5 ± 0.6 MPa m for the NbC1.0 was found to be the highest among all four compositions whereas its Knoop hardness, 8.4 ± 3.2 GPa, was the lowest value among all the compositions. With loss of carbon but a retention of the same B1 phase, the fracture toughness of the NbC0.75 sample decreased to
Table 1 Various experimental results of the niobium carbides.
KIC (MPa m ) HK (GPa) Grain Size (μm) Porosity (%)
NbC1.0
NbC0.75
NbC0.67
NbC0.5
2.5 ± 0.5 8.4 ± 3.2 1.2 ± 0.4 5.9
1.5 ± 0.1 11.5 ± 2.7 23.5 ± 8.5 6.8
2.0 ± 0.3 13.5 ± 1.3 44.9 ± 22.5 0.9
0.4 ± 0.2 11.8 ± 1.4 126.3 ± 68.8 3.4
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However, the NbC0.67 samples, with β-Nb2C laths, had a fracture toughness near 2 MPa m . The lack of change in toughness highlights the importance of local metal-metal bonding that occurs in carbon depleted planes of the zeta phase, but is absent in the β-Nb2C phase. It would be interesting to further see if the ζ-Nb4C3-x phase can create a microstructure that will toughen niobium carbide composites, although we hypothesize that the processing of such large monolithic ceramics will be difficult due to the speculated lower processing temperature (ζNb4C3-x’s decomposition temperature is 1575 °C) for this phase, which will increase porosity in the microstructure. Acknowledgements C.R. Weinberger and G.B. Thompson recognize Air Force Office of Scientific Research grant FA9550-15-1-0217andFA9550-15-1-0095, respectively. L. Lamberson gratefully acknowledges partial support from the NSF CAREER Grant No. 1751989. References [1] W.G. Fahrenholtz, et al., Refractory Diborides of Zirconium and Hafnium, J. Am. Ceram. Soc. 90 (5) (2007) 1347–1364. [2] D.J. Rowcliffe, Plastic deformation of transition metal carbides, Deformation of Ceramic Materials II, Springer, Boston, MA, 1984, pp. 49–71. [3] E. Wuchina, et al., UHTCs: Ultra-High Temperature Ceramic Materials for Extreme Environment Applications vol. 16, (2007), pp. 30–36. [4] L.E. Toth, Transition Metal Carbides and Nitrides 7 Academic Press, New York, 1971. [5] K. Hackett, et al., Phase constitution and mechanical properties of carbides in the Ta-C system, J. Am. Ceram. Soc. 92 (10) (2009) 2404–2407. [6] L. Limeng, et al., Microstructure and mechanical properties of spark plasma sintered TaC0.7 ceramics, J. Am. Ceram. Soc. 93 (10) (2010) 2945–2947. [7] R.A. Morris, et al., Microstructural formations and phase transformation pathways in hot isostatically pressed tantalum carbides, Acta Mater. 60 (1) (2012) 139–148. [8] M. Sygnatowicz, R.A. Cutler, D.K. Shetty, [ζ]-Ta4C3−x: A High Fracture Toughness Carbide with Rising-Crack-Growth-Resistance (R-Curve) Behavior. American Ceramic Society, J. Am. Ceram. Soc. 98 (8) (2015) 2601. [9] B.C. Schulz, et al., Experimental investigation into the crack propagation in multiphase tantalum carbide ceramics, Mater. Sci. Eng. A 695 (2017) 315–321. [10] A.I. Gusev, A.S. Kurlov, V.N. Lipatnikov, Atomic and vacancy ordering in carbide ζTa4C3−x (0.28⩽x⩽0.40) and phase equilibria in the Ta–C system, J. Solid State Chem. 180 (11) (2007) 3234–3246. [11] A.I. Gusev, Order–disorder transformations and phase equilibria in strongly nonstoichiometric compounds, Physics-Uspekhi 43 (1) (2000) 1–37. [12] D.J. Rowcliffe, G. Thomas, Structure of non-stoichiometric TaC, Mater. Sci. Eng. 18 (2) (1975) 231–238. [13] M. Lewis, J. Billingham, P. Bell, Nonstoichiometry in Ceramic Compounds, Univ. of Warwick, Coventry, Eng., 1973. [14] C.R. 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Fig. 5. Crack propagation initiated from the Knoop indent on a NbC0.75 sample. Note how the crack travels relatively unimpeded, with the crack propagating through the laths (dashed circled region).
reduced preference for fracture. Consequently, there is likely a much more homogenous fracture energy in β-Nb2C as compared to the zeta phase. The experimental report here provides evidence for supporting the hypothesis of Yu et al. [18] that the metal-metal bonding across the close-packed planes is critical in creating the high fracture toughness and provides further insight that not only nanoscale laths that create a torturous path are required, but the crystallography of those laths must create the localized metal-metal bonding within its unit cell to control the fracture. This notion can be confirmed by SEM micrographs of the crack propagation in the NbC0.67 sample. Fig. 5 reveals intergranular cracking with the cracks propagating easily through the laths, even with the cracks oriented normal to the laths (dashed circle region). In the work of Schultz et al.[9] such propagation in ζ-Ta4C3-x stunted crack growth. This finding provides the clearest evidence that the β-Nb2C laths provide minimal resistance to micro-cracking, and therefore contributes little to the fracture toughness as evidenced by its low value. 4. Conclusion In this work, we examined the fracture toughness and hardness of a series of niobium carbides at various carbon contents of C/Nb 1.0, 0.75, 0.67, and 0.5. The NbC1.0 and NbC0.75 samples composed equiaxed B1 rocksalt structures whereas NbC0.67 was an equiaxed NbCx matrix with laths of β-Nb2C. Finally equiaxed NbC0.5 was fabricated, but not all of the NbC converted to Nb2C revealing smaller grains of NbC with laths of Nb2C. As the carbon content decreased, the grains grew in size from approximately 1 μm to 120 μm. The grain growth is contributed to not only the larger starting Nb metal powder as compared to NbC used in mixing the composition of the billets, but also the easier diffusivity of the metal as compared to the carbide at the HIP temperature of 1800 °C. The hardness increased from approximately 8 GPa from NbC1.0 to 12 GPa for NbC0.5, with the fracture toughness decreasing from approximately 2.5 MPa m to 0.4 MPa m over the same range. The NbC0.67 samples comprised of a similar microstructure to the high fracture toughness tantalum carbides at similar compositions. The tantalum carbides have a lath based microstructure associated with the presence of the Ta4C3-x phase. These laths, comprised of carbon depleted stacking faults, provide toughening that results in fracture toughness values as high as 15 MPa m observed experimentally. 5172
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[27] B. Vishwanadh, et al., Synthesis and phase transformation mechanism of Nb2C carbide phases, J. Alloys Compd. 671 (2016) 424–434. [28] S. Alyamovskii, et al., Orthorhombic β′ phases of niobium and vanadium carbides, Russ. J. Inorg. Chem. 12 (1967) 301. [29] L. Wu, et al., The phase stability and mechanical properties of Nb–C system: Using first-principles calculations and nano-indentation, J. Alloys Compd. 561 (2013) 220–227. [30] C.J. Smith, C.R. Weinberger, G.B. Thompson, Phase stability and microstructural formations in the niobium carbides, J. Eur. Ceram. Soc. 38 (15) (2018) 4850–4866. [31] J.J. Petrovic, et al., Effects of residual stresses on fracture from controlled surface flaws, J. Am. Ceram. Soc. 59 (3–4) (1976) 177–178. [32] J.J. Petrovic, Effect of indenter geometry on controlled-surface-Flaw fracture toughness, J. Am. Ceram. Soc. 66 (4) (1983) 277–283. [33] Y. Torres, D. Casellas, M. Anglada, L. Llanes, Fracture toughness evalution of hardmetals: influence of testing procedure, Int. J. Refract. Metals Hard Mater. 19 (1) (2001) 27–34. [34] ASTM, ASTM C1326 Standard Test Method for Knoop Indentation Hardness of
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