Chapter Eight

Introduction and Background on Transition Metal Carbides


8.1 Technological Uses of Transition Metal Carbides
8.2 Structures of Transition Metal Carbides
8.3 Synthesis and Characterization
8.4 Tour of the Group IV-VI Carbides
8.5 Goals of the Part II Project
8.6 References

8.1 Technological Uses of Transition Metal Carbides

8.1.1 Structural/refractory uses1,2

The carbides of the transition metals in Groups IV - VI have extremely high melting points (Table 8.1) and are therefore referred to collectively as the “refractory carbides.” In addition to their stability at high temperatures, these compounds are extremely hard (Table 8.2), finding industrial use in cutting tools and wear-resistant parts. Their hardness is retained to very high temperatures, and they have low chemical reactivity – they are attacked only by concentrated acid or base in the presence of oxidizing agents at room temperature, and retain good corrosion resistance to high temperatures. The refractory carbides are strong, with Young’s modulus values – a measure of elastic deformation resistance – rivaling those of SiC at room temperature. In addition, they have good thermal shock resistance and good thermal conductivity, permitting heat to be drawn away from the working surface of the tool. This gives them a benefit over other refractory materials, which do not conduct heat so well. (Table 8.3)

Tungsten carbide, WC, is the most commonly used for fabrication as “cemented carbide” tools for cutting steel, in which the carbide is bonded in a metal matrix, usually cobalt. Cobalt is used because it wets the carbide particles and therefore behaves as a good binder without having significant ability to dissolve the carbide, so that the carbide is left pure in the bound form. However, pure WC-Co cemented carbides tend to weld locally with the steel being cut. TiC, TaC, and NbC are often used in conjunction with WC because TiC locally forms a layer of TiO2 or TiO23 which protects the tool from wear, and TaC and NbC raise the melting temperature and oxidation resistance of the tool.

For high-temperature applications, the carbides are used as pure-material sintered parts or in a Co/Mo/W/carbide sintered composite. They outperform the standard alloys and so-called “superalloys” in such applications as rocket nozzles and jet engine parts, where erosion resistance at temperatures of 2500 oC and up is crucial. TiC1-x and VC1-x in particular maintain high strengths up to 1800 oC and therefore can be used as high temperature structural materials, provided that internal and surface flaws, such as stress cracks and pores introduced during fabrication and sintering, are removed. Such defects lead to a high room-temperature brittleness; plastic flow relieves internal stresses caused by defects and leads to reduced brittleness at high temperatures. Plastic deformation occurs particularly via a mechanism of dislocation glide along {111} planes.

It must be noted that there is some variation in the literature with respect to the reports of assorted mechanical and thermodynamic properties of the refractory carbides. The transition metal carbides show a range of nonstoichiometries and possibilities for vacancy ordering, so the precise phases being tested for a given property are often unclear. Furthermore, small concentrations of oxygen present as metal oxide are famously difficult to remove – or even detect – and can be expected to affect the properties of the material.

Mixed-metal carbides have been examined for their melting point and hot-hardness behavior as well. The hardness arc-cast or zone-melted samples of (Ta0.8Hf0.2)C1+x was tested (by indentation using diamond or B4C tips) and compared with that of the similarly-prepared pure-metal carbides TaC1-x and HfC1+x over the temperature range 800 - 2000 oC.4 In all cases the hot hardness decreased with increasing temperature. For temperatures up to 1400 oC the hardness order increased as HfC1+x < (Ta0.8Hf0.2)C1+x < TaC1-x, but over 1400 oC the mixed carbide began to outperform the tantalum carbide. Hardnesses varied from 600 kg/mm2 to 150 kg/mm2 over the temperature range tested.

The melting points of the mixed-metal carbides outperform those of the pure-metal carbides, as well. Samples prepared by vacuum reduction of the mixed oxide powders at 2000 oC, followed by sintering at 2200 oC and 2500 oC in good vacuum were tested for their melting behavior.5 8TaC.ZrC and 4TaC.HfC had melting points 3890 oC and 3990 oC, respectively, somewhat higher than those of the pure-metal carbides (measured at 3470, 3750, and 3840 oC for Zr, Hf, and Ta carbides, respectively). The higher melting points in the mixed-metal carbides was attributed to composition changes due to the selective evaporation of carbon during melting.

Table 8.1: Melting points of the carbides of Groups IV-VI
Metal Name
Metal MP (oC)2
Carbide (MC) MP (oC)*2
Carbide (MC) MP (oC)**6



1810 (Cr3C2)



* Values listed in Toth (see Ref. 2)
** Values listed in Kirk-Othmer Encyclopedia of Chemical Technology (See Ref. 6)
Table 8.2: Properties of Group IV - VI Carbides2
Lattice (Å)
Young’s Modulus x 106 psi
Micro-hardness (kg/mm2)
Therm. Exp. Coef. x 10-6
a: 11.47,
b: 5.545,
c: 2.830
a: 7.244, b: 6.004, c: 5.199
4.9//a, 8.2//c
a: 2.906, c: 2.837
2100 (basal plane)
5.0//a, 4.2//c

Table 8.3: Properties of other Refractory Materials6
Melting Point (oC)
Microhardness (kg/mm2)
2300 dec
C, diamond
3800 dec

8.1.2 Catalytic uses

In addition to the technological uses of transition metal carbides which exploit their high hardness and stability at high temperatures, certain carbides have been examined for their catalytic properties in a number of reactions. This is in addition to their being potential supports for more traditional catalytic materials (Ni, Pt, Rh, etc.) due to their high heat stability.

The observation has been made that tungsten was active as a catalyst and showed good selectivity toward xylene formation during the isomerization and hydrogenolysis of 1,1,3-trimethylcyclopentane, but only after an induction period, as is also characteristic of platinum and paladium.7 Base transition metals do not behave this way. The behavior was explained by invoking the formation of tungsten carbide on the surface of the metal; WC is similar to Pt in selectivity for neopentane isomerization as well.8 Mo2C has been found to behave similarly to Ru in CO-H2 reactions.9 These are merely two examples: in fact, carbides of the Group IV-VI metals have been studied for their activity in oxidation, hydrogenation/dehydrogenation, isomerization, hydrogenolysis, and CO-H2 reactions, and in many cases have been found to rival the performance of the less economic Group VIII metals. While the refractory carbides do not show high activity for oxidation reactions (for example, the rate of H2 oxidation follows the order metal >> carbide > oxide for Group V and VI metals10, and the rate of NH3 oxidation over refractory carbides is lower than that over Group VIII metals11), they are as active as the transition metals themselves for hydrogenation and dehydrogenation reactions. In isomerization reactions, WC, Pt, and Ir are unique in their high activity and selectivity.12 However, it is not necessary for the refractory carbides to be more or even equally active in catalyzing given reactions compared with the noble metals, because the lower cost of the carbides will in many cases offset the losses in catalytic activity.

The Group IV-VI metal carbides are, however, difficult to prepare as high purity (e.g. free of surface oxide or graphitic and amorphous carbon contamination) and high surface area powders. There have been a number of approaches to synthesizing the carbides in a way that produces a high specific surface area, including carburization of a spray of the oxide powder using a CH4/H2 gas mixture,13 reaction of the metal oxide vapors with activated carbon,14 ultrasound irradiation of the metal carbonyl,15 carburization of a precursor deposited onto a support such as alumina,16,17 and similar approaches.18 The problem of surface contamination has been addressed as well, with most solutions centering on the idea of activation of the surface of the carbide via a thermal treatment prior to catalytic use. The thermal treatment includes heating in vacuum, which has been found to activate TaC, TiC, and WC for hydrogenation of ethylene,19,20 or reduction in flowing hydrogen gas.14,21 These methods, of course, are for removing surface oxygen; surface carbon (amorphous or graphitic) is difficult to remove.

The catalytic uses of the transition metal carbides are mentioned in the context of this project because the low-temperature precursor methods reported in Chapter 9 have the potential for yielding high-surface area carbides possibly useful for catalytic applications. However, the structural and refractory characteristics of the carbides were the focus of the project reported in Part II of this thesis, and no exploration of the catalytic potential, even characterization as basic as finding the surface area via BET analysis, was considered.

8.2 Structures of Transition Metal Carbides

Most of the transition metal monocarbides form in the B1 (NaCl) structure, fcc metal with carbon occupying the octahedral interstitial sites. The shortest M-M distance is about 30% greater in the B1 carbide than in the pure metal for the Group IV and V carbides, but drops to less than 10% greater for the Group VI or VII carbides.18 At 100% site occupancy, the stoichiometry of the carbide is MC1.0, though this situation is rarely realized. The concentration and ordering, if any, of the vacancies that result from a nonstoichiometric M-C ratio have a great effect on the thermodynamic, mechanical, electronic, and magnetic properties of the metal carbides; however, the details of these effects are a matter of some debate in the literature, due to the difficulties inherent in synthesizing pure compounds and in measuring the exact details of the crystal structure of a given sample. The metal carbides share many characteristics with the metals themselves, having a plastic deformation like the fcc metals which, while lowering the high-temperature hardness, protects parts fabricated from the carbides from catastrophic failure in response to stresses.

Most of the Group IV-V metal carbides conform to the Hägg rules, which were developed empirically to “predict” the structures of the transition metal borides, carbides, halides, and nitrides.22 The structure adopted by the metal carbide is determined by the ratio of the radius of the nonmetal atom (rx) to that of the metal atom (rm). For r = rx / rm < 0.59, the structures adopted are the simple A1, A2, A3, and hexagonal lattices. For r > 0.59, more complex structures form to prevent the expansion of the lattice – which would have been required for the simple structure to accommodate the large nonmetal atom – from taking the metal atoms beyond the distance for favorable metal-metal interactions. The monocarbides take on an fcc metal lattice with carbon atoms on the octahedral interstitial sites, while random occupation of half of the Oh sites in M2C or M3C leads to the L3’ (anti-NiAs) structure, and carbon occupation of the trigonal prismatic sites in the hcp lattice formed by the tungsten atoms leads to the CdI2 structure.2 Non-Hägg structures are known among the carbides as well, a major example being Cr2C3. Specific phases available to each metal will be discussed in Section 8.4 below.

One much-noted feature of the structure of the transition metal carbides is that the lattice adopted by the metal in the carbide is never that of the parent metal. That is to say, if the metal has an hcp lattice, its carbide has the metal on an fcc lattice; the fcc parent metal occupies a non-cubic lattice in the daughter carbide; and bcc parent metals have both fcc and hcp lattices available in the daughter carbides. This has been explained using Engel-Brewer theory of metals,23,24 in which the structure adopted by a metal or alloy depends on the s-p electron count.18 With increasing s-p electron count the metal structure progresses from bcc to hcp to fcc across the transition series. Likewise, the Group IV and V metal carbides MC form in the B1 structure rather than a hexagonal form because the incompletely filled bands of the host metals can accommodate a high ratio of sp-electron-rich carbon to metal. In Group VI the stoichiometry M2C occurs often, while Groups VII and VIII, when they form carbides at all, take on metal-rich stoichiometries M3C and M4C, consistent with an attempt to avoid filling antibonding levels in the metal bands.25

The nature of the bonding in the monocarbides is a matter of some debate, although all agree that the simple ionic model (M+C- or M-C+) is not consistent with the properties of the carbides. Ionic materials will not typically slip on the {111} planes due to the strong repulsive (coulombic) interactions across the shear plane in the half-glide position; instead, they will slip on the {110} or {100} planes.2 The catalytic behavior of WC is similar to that of Pt, as noted in Section 8.1.2, and this is explained according to the addition of the carbon valence electrons to those of tungsten to give a count equal to that of platinum. However, the exact direction of electron transfer is the subject of some controversy. X-ray photelectron spectroscopy (XPS) and electronegativity considerations suggest simple M --> C electron donation, but the XPS data is questionable due to the possibility for backdonation or screening effects.25 Furthermore, simple M ‡ C donation would result in ionic compounds such as the alkali and alkaline earth metal carbides. However, these materials are resistors with low optical transparency or reflectivity, and readily hydrolyze to the metal oxide and hydrocarbon. In contrast, the transition metal carbides are conductors with a shiny metallic and colored appearance and are hydrolytically stable. Band occupation suggests C --> M electron transfer. Carbon appears to combine its sp electrons with the metal spd bands, leading to the similarity between the crystal structures, reactivities, and catalytic activities of the refractory carbides and the Group VIII metals. Such a donation scheme may be used to explain, if crudely, the trend in melting point maxima, which occurs in Group VI for the metals, Group V for the carbides, and Group IV for the analogous mononitrides: the maxima may be associated with the half-filled d shell.25

The Group IV-V carbides are able to form continuous solid solutions with each other over a wide range of compositions, but are only partially miscible with the carbides of the Group VI - VIII metals.18 Ceramic materials of mixed C, N, and O composition are also common, with oxycarbides of definite stoichiometry having been reported for a number of the early transition metals. Even “pure” metal carbides tend to contain small amounts of dissolved oxygen.

8.3 Synthesis and Characterization

The distinction between “traditional” methods and “precursor” methods made in this section is admittedly somewhat arbitrary, since any starting material besides the elements themselves could be considered to be “precursors.” However, the separation has been made so that simple starting materials are included with “traditional” methods, and starting materials requiring some high degree of design and synthetic effort have been designated “precursor” methods.

8.3.1 Traditional Methods2

The usual method of preparing polycrystalline transition metal carbides on the research scale entails the direct reaction of metal or metal hydride powders with carbon. Pure materials with a homogenous composition are difficult to achieve, however, requiring high-purity gases or good vacuum in combination with very high tempertures. The methods are summarized in Equations 8.3-8.6.

(8.3a) M + C --> MC
(8.3b) MH + C --> MC + H2
Direct reaction by melting or sintering the starting material with carbon in vacuum or inert atmosphere.

(8.4) MxOy + C --> MC + CO
Reaction of the metal oxide and excess carbon in inert or reducing atmosphere.

(8.5a) M + CxH2x+y --> MC + H2
(8.5b) M + CO --> MC + O2
Reaction of the metal with a carburizing gas.

(8.6a) MCln + CxH2x+y --> MC + HCl + (CqHr)
(8.6b) M(CO)n + H2 --> MC + (CO, CO2, H2, H2O)
Reaction of the metal halide or carbonyl vapor with hydrogen.

The reaction temperatures for synthesis of the refractory carbides according to Equations 8.3a, b are listed in Table 8.5. However, most carbide systems must be heated for several hours at temperatures over 2000 oC to ensure compositional homogeneity. With good vacuum, such treatment also will remove oxygen contamination from many of the metal carbides. Reports regarding the range of temperatures required to form the carbides according to other reaction schemes give variable numbers.

Table 8.5: Reaction temperatures for direct formation of metal carbides2
Temperature Range (oC)
8.3a, b
1700 - 2100
8.3a, b
1800 - 2200
8.3a, b
1900 - 2300
1100 - 1200
1300 - 1400
1300 - 1500
1400 - 1800
1200 - 1400
1400 - 1600

8.3.2 Precursor Methods

Many of the reported precursor methods for making high-temperature carbides focus on the synthesis of alpha- or beta-SiC rather than on the carbides of the early transition metals. In addition the solid-state reaction between silica and carbon (Acheson process),6 beta-SiC has been made by pyrolysis of high-Si rice hulls (a carbothermal reduction of SiO2 collected by the rice plant and locallized in the hulls). This process is commercially practiced. It has also been synthesized, more notably, by the pyrolysis of polymers such as poly(dimethylsilane) or poly(phenylmethylsilane).26 These polymers are synthesized via the Wurtz-type coupling of appropriate chlorosilanes, e.g. coupling in the presence of sodium or potassium metal (M) with loss of MCl driving the polymerization. Linear polycarbosilanes produce nearly zero ceramic yield: branching on the silicon backbone is crucial. The presence of Si-H and Si-vinyl functionalities lead to the potential for hydrosilation cross-linking and a still-higher ceramic yield.

Similar work has been done to give organometallic precursors for the transition metal carbides. This work is dogged by the relatively few compounds of the Group IV - V metals which are stable, have a low C/M ratio (so that excess free carbon is avoided), and which do not contain oxygen. Those which do fulfill these criteria tend to be sublimable, leading to a low ceramic yield, since the opportunities for polymerization or cross-linking of these materials are few. Seyferth and Tracy attempted to couple dimethyltitanocene with methyldichlorosilane in the presence of sodium metal to produce a copolymer precursor for (Si, Ti)C, but little of the product was soluble in toluene and the ceramic yield was low.3 Dichlorobis(2,4-pentanedionato)titanium(IV) was reacted similarly with methyldichlorosilane in the presence of Na, but the products were soluble only in THF and contained chlorine. Finally, dimethylmetalloocene, which decomposes in ultraviolet light to Cp2M and 2 .CH3 (M=Ti, Zr, Hf), was used with the polysilane made by the Würtz coupling of HMeSiCl2 to produce air-sensitive metal-containing polysilanes with ceramic yields ranging from 44% (Hf) to approximately 80% (Ti). The polymers were at least sparingly soluble in aromatic or ethereal solvents, sometimes more soluble, depending on the specific reaction conditions.

Unpublished research performed in the Seyferth labs polymerized zirconium and hafnium metallocene methacrylates for use as precursors for ZrC and HfC. The precursors were solids or viscous polymers. Liquid precursors were obtained when substituted metallocene dichlorides were polymerized with n-butyllithium and a diyne. These gave ceramic yields near 60%, but the process suffered from the inability to scale it up.

Most other research involving organometallic precursors for the refractory carbides has centered on the small-molecule, monomeric metallocene compounds for use with various forms of chemical vapor deposition of films of these the carbide materials.

Precursors for the refractory carbides have been made from the metal alkoxides transesterified with polyhydroxyl alcohols and mixed with phenolic resins or furfuryl alcohol to provide excess carbon. In most of these cases the precursor was a solid, though some soluble precursor polymers resulted. Past work along this line will be discussed in more detail in Chapter 9.

8.3.3 Characterization

Complete characterization of the refractory carbides is difficult. Since many of the important mechanical and catalytic properties are sensitive to a number of factors which tend to vary widely among samples, there is a variation in literature reports regarding measurement of these properties. These factors include (1) the crystal structure and lattice parameters, including the presence of vacancy ordering; (2) the chemical composition, including not only the overall carbon-to-metal ratio present in the bulk sample, but the amount of free carbon versus combined (lattice) carbon; (3) the impurity concentration, particularly that of oxygen; (4) the overall defect structure, including grain size, dislocations, and porosity; and (5) the sample homogeneity.

The crystal structure can be found and the lattice parameters measured by X-ray and neutron diffraction experiments. Sample homogeneity may be deduced from X-ray diffraction (XRD) data by noting the sharpness of the splitting the Kalpha-1 and Kalpha-2 lines, and crystalline impurity phases present in concentrations over approximately 5% may often be detected using XRD. Even large oxygen impurity concentrations, however, may escape notice due to the formation of a solid solution with the carbide. The effects of medium-to-large oxygen impurity concentrations on lattice parameters are not well-described. Very careful XRD analysis may also detect ordering of carbon atoms or vacancies in the structures of the carbides of the lighter metals, when such ordering is accompanied by a distortion of the lattice. The specific details of the ordering would not be easy to determine even using modelling techniques such as Rietveld analysis;27 ordering without an accompanying symmetry or lattice parameter change would only be detectable using a complex Fourier analysis of the XRD peak intensities.28 Neutron or electron diffraction experiments, on the other hand, would give unambiguous information regarding the ordering of carbon atoms on the interstitial sites, however, electron diffraction sample preparation is made difficult by the need to make a thin section of the brittle and possibly porous carbide material, and neutron diffraction is not immediately available to most researchers.

Chemical analysis of the bulk material primarily yields information about the quantity of carbon in the sample. The total carbon content may be found by heating the carbide in a stream of oxygen, completely converting it to the metal oxide. The CO2 generated is captured and measured by weight or by the conductivity of a combusted CO2-O2 mixture in a commercial Leco unit. This process is complicated by the high temperatures tolerated by the refractory carbides, since it relies on all of the carbon’s being driven off as CO2. Diffusion of carbon and/or oxygen through the metal carbide and/or oxide is slow even at very high temperatures. Free carbon (graphite) may be determined by dissolving the carbide in a mixture of hydrofluoric and nitric acids (warning: the dissolution reaction is violent and exothermic). The unbound carbon does not dissolve; it is collected, washed, and analyzed as CO2.

Oxygen analysis is extremely difficult because once the solid solution M(C,O) has formed the oxygen is nearly impossible to remove. This is particularly true in the case of nonstoichiometric carbides MC1-x, where empty octahedral interstitial sites are very inviting to the stray oxygen atom. Vacuum-fusion techniques, e.g. heating the refractory carbide to 2400-2800 oC in a graphite mold or platinum bath under vacuum, followed by determination of the oxygen removed as COx, are the major methods available, but are often only partially successful. This is particularly true of the carbides of the Group IV metals. Neutron activation analysis, in which oxygen is activated by the reaction 16O(n,p)16N, then monitoring the 6.1 and 7.1 MeV gamma-radiation from 16N, has proven useful.

Sample homogeneity and local compositions may be measured by microprobe analysis, in which the X-ray emission of elements on irradiation with an electron beam, as in an electron microscope, is measured and correlated with the concentration of the element. The microprobe must be equipped with a thin detector window to give good quantitative information regarding light elements (Z< 10), however. Electron energy loss spectroscopy (EELS), on the other hand, can sensitively detect elements in the Z < 10 region.29 Both of these methods are extremely locallized, however, and multiple areas of several samples must be examined for conclusions regarding the bulk material to be validly made.

Vacancy concentration is likewise difficult to measure due to the ambiguities in chemical composition measurements. Where the composition is well-known, comparison of X-ray density with a physically-measured density will give a good estimate of the vacancy concentration. Likewise, density measurements can give information regarding larger-scale defects such as porosity. Electron microscopy can give reasonable estimates of the grain size.

8.4 Tour of the Group IV - V Carbides

The material in Section 8.4 is primarily based on E.K. Storms’ The Refractory Carbides.1

8.4.1 General Trends

Approximate preparation of the transition metal carbides is straightforward, but ensuring a given stoichiometry and purity against oxygen contamination is famously difficult. Variations in the quantity of vacancies on the carbon (or, less frequently, the metal) lattice, as well as variations in the amount of dissolved oxygen, lead to a wide range of claims regarding even basic thermodyanmic, mechanical, and electromagnetic data for the early transition metal carbides. Removal of the oxycarbide phases, which can be considered to be solid solutions of MO and MC, depends on the partial pressure of CO over the sample to be purified. At the high-carbon end of the stoichiometric range, Equation 8.7 leads to additional dissolved oxygen in the lattice if CO is removed. At the low-carbon end, Equation 8.8 occurs independently of the CO partial pressure. Thus if excess graphite is present, a high CO partial pressure leads to more nearly stoichiometric carbide as Equation 8.7 is forced left; but the success of this approach in purifying a given metal carbide depends on the stability of the oxycarbide, the annealing temperature and time, and, clearly, the partial pressure of carbon monoxide.

(8.7) MxOy + C (dissolved in MC) = MO (dissolved in MC) + CO(g)
(8.8) MxOy + M (dissolved in MC) = MO (dissolved in MC)

Group IV carbides are difficult to purify without melting; heating up to 2000 oC will result in increased oxygen contamination if the vacuum is not better than 10-6 torr. Moreover, as noted in the previous section, few straightforward chemical methods exist for finding the oxygen level in the Group IV carbides MC1-x bulk materials; none are reliable. The Group V and VI carbides purify readily at temperatures over 1800 oC.

Slow diffusion of carbon in all of the refractory carbides results in stoichiometry gradients which are difficult to detect in bulk materials but which may compromise the material strength, hardness, and high-temperature behavior. The lattice parameter and the sharpness of the XRD pattern can give some rough indication of the homogeneity, however.

Group IV metals tend to form in a single cubic phase with a limiting stoichiometry near MC1.0, but which normally varies from MC0.5 to MC0.97 depending on the synthesis conditions. Group V metals form an M2C phase in addition to the monocarbide. The composition range of the M2C phase is narrow at room temperature, with decomposition into the cubic phase plus liquid at high temperatures. The V-C system has a cubic phase extending to MC0.88, while NbC and TaC approach MC1.0 and melt congruently even at carbon-deficient stoichiometries. The Group VI metals have a more complex M-C phase diagram, forming a number of distinct compositions. The chromium carbides behave uniquely, while the Mo-C and W-C systems have common features, with the MC and M2C phases stable at high temperatures. (Table 8.4) The trends in melting points indicate that the Group IV, V, and lower two Group VI metals have strong M-C and M-M bonds, distinct from the Groups IA - IIIB metals, which form acetylenic M-C bonds, and from the Groups VIII - IIB metals, which form unstable carbides, if at all.

Removal of bound carbon causes the lattice parameter to decrease in most of the refractory carbides, but to increase in TiC and ZrC. The behavior of HfC on decrease of the lattice carbon-to-metal ratio is uncertain due to the variation in this behavior among reports. For several metal carbides, the variation of lattice parameter with carbon content is linear. Removal of carbon from the lattice also is associated with reductions in hardness, at least for the Group IV carbides. Dissolved oxygen lowers the lattice parameter in Group IV carbides, while it increases it in Group V carbides, and has an uncertain effect for carbides of Group VI metals. The effect of oxygen contamination on mechanical properties is not clearly reported in the literature.

The refractory carbides show high chemical resistance but will react under certain conditions. At high temperatures, the high-carbon compositions form hydrocarbons in the presence of hydrogen. The reactions with oxygen have been indicated above (Equations 8.7, 8.8), and are complex. The carbides will form the nitrides at high temperatures and in the presence of N2, NH3, or N2/H2 mixtures; however, the cubic carbides and nitrides are completely miscible.

8.4.2 Group IV

The conversion of TiO2 to TiC occurs via the intermediates Ti3O5, Ti2O3, and TiO in the temperature range 1000 - 1500 oC. The carbide closest to TiC1.0 forms at 1600-1700 oC under 1-10 torr of CO. Titanium hydride and carbon form TiC1.0-x after 1 hour in vacuum at 1200 oC. TiC has also been formed by heating a tungsten wire or carbon filament in an atmosphere of TiCl4, H2, and hydrocarbon, or by reaction of CaC2, TiCl4, and H2 at 800 oC (CaC2 and CaCl2 are removed by washing with water after the reaction is complete). The last traces of oxygen are difficult to remove, and have a significant effect on the material properties. Later heating may recontaminate even a “pure” sample of titanium carbide if the vacuum is not sufficiently good; a nonprotecting, non-adherent TiO2 (anatase) layer forms at about 450 oC. The Ti-C system has one cubic compound of formula TiC, although other phases have occasionally been claimed. TiC is metallic and grey, and is stable to most concentrated acids or bases. It will dissolve completely in HNO3 and combinations of HNO3 with HCl (aqua regia), HF, and H2SO4.

The reduction of ZrO2 proceeds via Zr2O3 and ZrO to the carbide between 950-1200 oC. It has also been formed using ZrH plus graphite or from ZrCl4 in the presence of hydrogen and hydrocarbon vapor. Attempts to remove oxygen completely are generally unsuccessful except under melting conditions. “Pure” ZrC heated at temperatures under 1800 oC tends to gather oxygen up to several percent. The Zr-C system contains one cubic compound, ZrC, and the lattice parameter varies with oxygen contamination noticibly at levels of 1000 ppm. ZrC is somewhat more susceptible to acid attack than is TiC and oxidizes rapidly above 500 oC.

HfO2 forms an oxycarbide of constant composition between 1743 - 2033 oC and under 70-1000 torr of CO, with Hf2O3 forming at 1000-1200 oC and the HfC-HfO solid solution between 1300-1800 oC. HfC forms a carbon-deficient lattice between 1800 - 2000 oC, but can be made stoichiometric by repeated heating at 1900 oC. HfC has also been formed from HfCl4 + H2 + CH4 in the presence of a hot tungsten wire and from hafnium hydride and carbon. It is one the most difficult carbides to rid of its oxygen, only becoming “pure” when melted or heated at temperatures in excess of 2500 oC in good vacuum. The Hf-C system has one cubic phase HfC, but the composition can range to a low of HfC0.52. Its melting point increases with increasing carbon content up to a maximum, then HfC forms a eutectic with carbon. There has been no study of lattice parameter variation with either oxygen or carbon content.

8.4.3 Group V

Vanadium carbide has been formed by heating V2O5 or V2O3 with carbon for two hours at 1800 oC in 1-10 torr of CO. V2O5 begins reacting with carbon at 435 oC, and the oxygen concentration is relatively easily reduced to below detectable limits by higher-temperature treatments or by the reaction of vanadium metal or hydride with carbon. Loss of vanadium at high temperatures and low carbon content presents a difficulty, but nearly-saturated VC can be heated to 2000 oC without loss of vanadium. The carbide has been made by treatment of VCl4 in an atmosphere of hydrogen and methane at 1500 - 2000 oC. The two main phases are cubic VC, available over the range VC0.78-VC1.0, and the hexagonal beta-V2C, available for C/V atom ratios of 0.4-0.5 between approximately 1500 - 2000 oC, but presenting a very narrow range of stable compositions at room temperature. Reports of V5C and V4C3 have been discredited. VC will react with dry HCl gas at 750 oC to produce VCl2, methane, and hydrogen, and has a high rate of oxidation in air, with powdered VC and V2C being attacked slowly by air even at room temperature. V2C is soluble in hot 50% HCl solution, leaving a carbon residue, but VC is inert under these conditions; both of the vanadium carbides are attacked by concentrated nitric, sulfuric, and perchloric acids.

Niobium (V) oxide begins to react with carbon at 675 oC, forming NbO2 and NbCx below 1200 oC and forming an NbCxOy solid solution between 1450 and 1500 oC. Pure NbC is accessible by heating the metal and carbon powders directly, but high temperatures and heating times are required to complete the reaction and drive off oxygen and nitrogen contaminants. The conditions are made less stringent by the presence of an H2 atmosphere. NbCl5 heated in the presence of hydrogen and methane forms the pure carbide at 900 -1000 oC. The Nb-C system has the cubic phase NbC and two crystal forms (alpha, room temperature, with a very narrow composition range and beta, existing between 2300-3000 oC and over the C/Nb ratio range 0.4-0.5) of Nb2C. A third, zeta phase of Nb2C (C/Nb range 0.5-0.7) has been suggested to exist on the basis of a single weak powder pattern line, but has not been verified. The lattice parameter increases as the C/Nb atom ratio approaches unity, and is increased as well by the presence of oxygen and nitrogen. NbC is inert even to boiling aqua regia but will dissolve in HNO3/HF mixtures; it is severely corroded in air at temperatures above 1100 oC. Its color ranges from grey (NbC0.9) to lavender (NbC0.99).

The tantalum carbide system is relatively easy to free of oxygen impurities, but due to the slow rate of carbon diffusion it tends to have inhomogeneities in its bulk composition. Evaporation of carbon at temperatures above 2400 oC renders the use of high temperatures to establish a uniform composition problematic. Reaction from the elements in vacuum begins at approximately 1000 oC but is slow to reach completion. Use of hydrogen or methane atmospheres increases the reaction rate but requires a post-synthesis vacuum annealling step to remove dissolved hydrogen. TaC cannot be made from TaCl5 in a hydrocarbon/hydrogen atmosphere due to the formation of metallic tantalum, but has been made with varying success by heating Ta wire in methane. Arc-melting tends to produce C-deficient, inhomogenous carbides. The Ta-C system has the cubic phase TaC and a hexagonal compound Ta2C (actually the C6 anti-CdI2 structure type due to ordering of the carbon atoms) with an α-β transition near 2000 oC. A ζ phase has been claimed at 2000 - 3000 oC over the C/Ta atom ratio 0.70-0.75 but its existence remains a point of debate. The cubic phase persists over a wide temperature and composition range, possibly as low as TaC0.58 and verified down to at least TaC0.74. The composition TaC0.89 is the highest-melting substance known. Very small amounts of carbon (Ta64C) result in a tetragonal distortion to the normally bcc Ta parent lattice. The TaC lattice varies linearly with composition, with the equation C/Ta = -25.641 + 5.9757a. It is grey and metallic in appearance up to about TaC0.85, then becomes increasingly brown with rising carbon content until the golden TaC0.99 is reached. It is the most acid-stable of the refractory carbides, dissolving in a nitric/hydrofluoric acid mix, and reacts with pure oxygen above 800 oC. Loss of carbon results from lower-temperature reactions with oxygen in air.

Table 8.4: Range of Melting Points for Group IV-VI Carbides1
Element MP (oC)
Maximum MP (oC)
MP in the presence of C (oC)
Atom Ratio at Max MP

8.5 Goals of the Part II Project

The specific goals of Part II of this thesis were the synthesis of a industrially-useful liquid precursor for tantalum and hafnium carbides, TaC and HfC. As indicated in Section 8.3.2 above, most precursor methods have resulted in solid or at best soluble polymeric materials. Liquid precursors were preferred in this project so that removal of solvent during fabrication of HfC or TaC monoliths could be avoided. The liquid precursors were to be used in conjunction with proprietary technology owned by Foster-Miller, Inc., of Waltham, Massachusetts, to fabricate high-performance parts of solid TaC or HfC. Factors of concern included (1) ease of synthesis, particularly of scale-up; (2) ease of handling (e.g. room-temperature stability, reasonable air-stability); (3) clean decomposition to the target carbide without undue contamination by metal oxide or by excess graphitic carbon; (4) low volatility – e.g. the material should decompose in situ during pyrolysis, rather than distill; (5) low gas production on decomposition, so that bubbles in the solid would not become points of structural weakness; (6) low viscosity over a temperature range near or just above room temperature (e.g. the precursor was to be a liquid or low-melting solid). Cost was an addtional concern but was left to be addressed in later phases of the project.

Early work on the project focused on the use of organometalic precursors for TaC and for HfC. The main idea behind this approach was that the precursor was best formed by putting the M-C bonds in place during the synthesis, eliminating any need to remove heteroatom bonds (particularly oxygen) from the metal center via carbothermal reduction during the pyrolysis step. Large, waxy groups or smaller, symmetry-destroying groups, would enforce the liquid form of the precursor. Efforts first focused on using CpTaCl4 or Cp2TaCl2 as starting materials, and using organolithium or Grignard reagents to add alkyne groups based on 1-hexyne, trimethylsilylacetylene, and triethylsilylacetylene by a salt metathesis; however, the products consistently decomposed, possibly due to reduction of the tantalum by the organolithium or -magnesium reagent. A similar approach using hafnocene and substituted hafnocenes (e.g. bis(methylcyclopentadienyl)hafnium dichloride, bis(vinyldimethylsilylcyclopentadienyl) hafnium dichloride) gave liquid precursors which pyrolyzed to HfC, and which were obtained in pure form. Stirring pentakis(dimethylamido)tantalum with alkynes, with loss of ≥ 2 dimethylamine groups, gave liquid precursors which pyrolyzed to TaC but which defied purification and characterization attempts. Thus this work showed some success, but was dogged by the sensitivity of the organometallic compounds made. The organometallic approach has not been discussed in detail in this Thesis.

Practical concerns of scalability were a major factor in the abandonment of this approach in favor of a sol-gel-type route in which the alkoxide in alcohol solution would be hydrolyzed to form the precursor, which would be reduced carbothermally during pyrolysis. This route was superior in ease of synthesis, scale-up, handling, and cost, but was found to give a too-low mass-of-carbide-per-gram-of-precursor. It also formed a glassy, bubble-ridden solid which was high in excess carbon. The work is discussed in Chapter 9, and has been continued by researchers at Foster-Miller, Inc.

It should be noted by the reader that the work described in Chapter 9 involved a practical challenge to the academic chemist. Traditional methods of purifying chemical materials for characterization include primarily crystallization and distillation or sublimation. However, the materials synthesized in these chapters were designed not to crystallize, and were expected to decompose on heating rather than to distill. Therefore, purification for the sake of getting NMR, MS, or elemental analytic evidence of the structure claimed was rarely possible. Production of a formulation with reproducible properties and pyrolysis behavior was given priority over isolation and characterization of an exact precursor molecule.

8.6 References

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