Sintered porous cermets based on TiB2 and TiB2–TiC–Mo2C

Materials Chemistry and Physics 67 (2001) 226–233

Sintered porous cermets based on TiB2 and TiB2–TiC–Mo2C

Manmohan Singha, K.N. Raib, G.S. Upadhyayab,∗a Materials Science Programme, Indian Institute of Technology, Kanpur, India

b Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur, India

Abstract

TiB2 powder, with different binders (Ni and Ni/Mn), after milling were cold compacted (300 MPa) and sintered in H2 at 1300 and 1350◦Cfor 1 h. To improve the sintering behaviour, TiC/Mo2C alloy carbide was added and the milled charge along with the same binders (Niand Ni/Mn) was cold compacted and sintered under similar conditions. Sintered density, porosity, transverse rupture strength (TRS), grainsize and lattice parameter of binder and hard phases were measured. Better densification was observed with Ni/Mn binder as comparedto Ni binder for either hard phase based systems. Maximum value of TRS was noted for TiB2–TiC–Mo2C–40 wt.% Ni/Mn cermet. Meltexudation was observed for either hard phase based systems with Ni binder. © 2001 Published by Elsevier Science B.V.

1. Introduction

TiB2 based cermets have received much attention fortheir potential applications because of their superior ther-momechanical properties under various conditions. Thedensification of single phase TiB2 ceramic is complicatedby two characteristics of this compound. The high meltingpoint of 2980◦C requires sintering temperature of the orderof 1800–2300◦C to obtain more than 95% of theoreticaldensity. Such a high temperature results in a rapid mass dif-fusion and permits the attainment of high densities withinacceptable sintering times between minutes and hours. Thisrapid mass transport, however, is combined with a tremen-dous grain growth of faceted crystals. Secondly, the hexa-gonal structure of TiB2 results in marked thermal expansionanisotropy. The expansion difference is between 37 and 42%,between 25 and 930◦C and this difference increases as thetemperature exceeds 930◦C. The expansion anisotropy pro-duces considerable internal stress during cooling and gen-erates micro cracking when the grain size is above a criticalvalue of approximately 20mm. The micro cracking occurs inthe grains and at the grain boundaries with a resultant degra-dation of macroscopic mechanical properties [1–4]. To mini-mize the above mentioned problems in sintering of TiB2, itis worthwhile to search cermets based on this ceramic.

During liquid phase sintering of TiB2 with suitablemetallic binders, the latter should wet and partially dissolveand reprecipitate the refractory compound during sintering.However, the formation of brittle and comparatively soft

∗ Corresponding author.E-mail address:[email protected] (G.S. Upadhyaya).

binary or ternary borides must be inhibited because theyconsume a good part of ductile binder. The relatively lowsinterability of TiB2 based cermet systems can be attributedto the fact that the metallic binders do not form sufficientliquid phase and easily change into borides during sinteringby the reaction of TiB2 with impurities such as free boronand boron oxides [5].

Since TiB2 is stable in the presence of TiC [6] and thelatter has good structural and thermodynamic compatibilitywith TiB2 [7], it was proposed to introduce TiC as a ternaryadditive. Among refractory borides and carbides, the heatof formation of the former is greater than the latter con-firming higher degree of covalent bond which causes betterwetting of refractory carbide by transition metal melts ascompared to borides. This is also confirmed from the factthat refractory carbides in contrast to borides possess ho-mogeneity ranges. This peculiarity in the chemical bondingis important in developing cermet systems. Among refrac-tory carbides, TiC–Mo2C system exhibits better wettabilitywith nickel as compared to TiC [8]. Keeping this in mind,both these carbides, i.e. TiC and Mo2C have been added instraight TiB2–Ni cermets. In addition, the binder modifica-tion is also called for in order to lower the melting point ofpure nickel by alloying. For this, the selection of manganesewas made. Such manipulation in both hard phase and binderphase is desirable in order to achieve a good design of theresultant cermet system.

2. Experimental procedure

The starting material used to fabricate green bodies werecommercially available TiB2, TiC–Mo2C, Ni, and Ni/Mn

0254-0584/01/$ – see front matter © 2001 Published by Elsevier Science B.V.PII: S0254-0584(00)00428-4

M. Singh et al. / Materials Chemistry and Physics 67 (2001) 226–233 227

Table 1Characteristics of different powders

Powder Manufacturer Chemical analysis /particle size Apparentdensity (g cm−3)

Flow rate(s (50 g)−1)

TiB2 Hermann C. Starck, Berlin B: 30.3, C: 0.3, O: 0.25, N: 0.39, Ti: balance;average particle size (FSSS): 9.5mm

1.65 Non-free flowing

TiC–Mo2C Treibacher Chemische Werke, Austria Ctotal: 14.18, Cfree: 0.05, Fe: 0.05, O: 0.13,Mo2C: 39.9, TiC: 60. 1; average particle size(FSSS): 4.8mm

2.36 Non-free flowing

Ni SF-300, Sherrit Gordon Mines, Canada H2: 0.266, Co: 0.186, Cu: 0.002, Fe: 0.009,S: 0.039, C: 0.017, SiO2: 0.002, Ni: balance;sieve analysis: 0–30mm — 31.9%, 30–40mm— 65.4%

3.43 23

Ni/Mn BSA Metal Powders, Birmingham Ni: 39.0, Mn: balance; sieve analysis:−53mm— 0.1%,−53+ 45mm — 2.3%,−45+ 38mm— 7.5%,−38mm— 90.1%

1.76 12.2

powders. The characteristics of the different powders aregiven in Table 1. Weighed quantity of powder was manu-ally mixed followed by wet ball milling in acetone for 10 hin “Fritsch Pulverisette 5” centrifugal type ball mill usingWC lined ball containing 19.95 mmφ WC balls. Two massper cent micronized wax powder was added as a binder toimprove the green strength. The powder mixture slurrieswere dried at room temperature. The dried powder mixturewas later granulated manually. Rectangular parallelepipedshaped green compact of size 25.1 mm × 8.1 mm andthickness approximately 2–5 mm were prepared from themilled powder in a single acting hydraulic press at 300 MPapressure.

Sintering was performed in a silicon carbide resistanceheated tubular furnace in dry H2 atmosphere (dew point−35◦C). Dewaxing was done at 400◦C for 2 h with heatingrate 2◦C min−1. The sintering temperatures selected were1300 and 1350◦C, respectively. After dewaxing, heating ratewas increased to 8◦C min−1 and sintering was done for 1 h atselected temperature. Densities of green compacts were cal-culated from the mass and physical dimension measurementsof the sample. For the measurement of sintered density, xy-lene impregnation/water immersion method was followed.

The transverse rupture strength (TRS) of sintered speci-mens was evaluated in the three point bending test usinga fixture containing 5 mmφ WC rollers having 15 mmspan (ASTM specifications B 406-76). Load was appliedin the 10 T capacity MTS testing machine with a crossheadspeed 0.2 mm min−1. The sintered compacts were groundon diamond wheel, and polished on STRUERS polisherwith DP cloth (type DOR) using diamond paste (2.5mm).The polished samples were etched by using HNO3 (10 ml)and HCl (10 ml). The scanning micrographs and frac-tographs of sintered samples were taken using JEOL, 840Ascanning electron microscope. X-ray diffraction studyof sintered compacts was carried out with the help ofISODEBYEFLEX-2002 X-ray diffractometer using Cu Karadiation.

3. Results

3.1. TiB2 based cermets

Fig. 1 shows the effect of Ni binder content on the sintereddensity, from which it is clear that the sintered density ofthe cermets gradually increases with increase in the bindercontent. Sintered density is slightly more after higher tem-perature sintering compared to 1300◦C except for 10% Nibinder. Fig. 2 shows the effect of Ni/Mn binder content onthe sintered density. The sintered density increases slightlywith increase in Ni/Mn binder content and is maximumat 40% Ni/Mn binder for 1350◦C sintering. With increase

Fig. 1. Variation of sintered density and % porosity of TiB2 based cermetswith nickel binder content.

228 M. Singh et al. / Materials Chemistry and Physics 67 (2001) 226–233

Fig. 2. Variation of sintered density and % porosity of TiB2 based cermetswith nickel–manganese binder content.

in sintering temperature, the sintered density, in general,increases.

The difference between sintering behaviour of TiB2–Niand TiB2–Ni/Mn cermets is characterized by sweating (exu-dation) phenomena in case of TiB2 cermets containing30 and 40 wt.% Ni binders. Very small spherical particleswere found on the surface of the sweated samples. How-ever, no such sweating phenomena was observed in case ofTiB2–Ni/Mn cermets.

Fig. 3 shows the effect of Ni binder content on the TRSof the cermets. It is clear that TRS decreases with the in-crease in Ni binder content in case of either sintering tem-perature. After higher temperature sintering, i.e. 1350◦C,TRS was found to decrease except for 30 and 40 wt.% Nibinder as compared to 1300◦C. Significant variation in TRSis observed as the Ni content increases from 10 to 30 wt.%irrespective of sintering temperature. The effect of Ni/Mnbinder content on the TRS of the cermets is shown in Fig. 4.After 1300◦C sintering, the TRS value increases with in-crease in Ni/Mn content from 10 to 20 wt.%, then slightlychanges at 30 wt.%. A still further increase in Ni/Mn contentdecreases the TRS values of the cermets. With increase insintering temperature, improvement in TRS of the cermets,in general, is observed.

Fig. 5 shows the scanning photomicrographs of fracturedsurfaces of TRS test pieces of some typical TiB2–Ni andTiB2–Ni/Mn cermets. It is evident that brittle mode of frac-ture for either type of cermet is present. Fig. 6 shows thescanning photomicrographs of some of the TiB2–Ni and

Fig. 3. Variation of TRS of TiB2 based cermets with nickel binder content.

TiB2–Ni/Mn cermets. Hard phase TiB2 has an irregular geo-metrical shaped grains, while the binder forms a liquid phasearound the hard phase. Comparison of the microstructuresof TiB2–Ni and TiB2–Ni/Mn cermets confirms a good wet-ting in the latter system. With increase in binder content ineither system better densification occurs, which is related toenhanced diffusion, as evidenced from the microstructures(Fig. 6). Grain size variation of boride phase for some of theTiB2–Ni and TiB2–Ni/Mn cermets is shown in Table 2, fromwhich it is evident that with increase in amount of binder,the grain size of boride phase decreases in case of Ni binder,while the reverse is true in case of Ni/Mn binder.

A typical X-ray diffraction scan for TiB2–10% Ni cermetis shown in Fig. 7. Lattice parameters of binder and TiB2

Fig. 4. Variation of TRS of TiB2 based cermets with nickel–manganesebinder content.


Fig. 5. SEM fractographs of TiB2–Ni and TiB2–Ni/Mn cermets sinteredat 1350◦C: (a) TiB2–10 wt.% Ni; (b) TiB2–10 wt.% Ni/Mn.

phases were calculated for cermets sintered at 1350◦C andthe values are shown in Table 3. With increase in Ni bindercontent, the lattice parameter of binder increases and thec/aratio of TiB2 phase slightly decreases. In case of cermetswith Ni/Mn binder, the reverse is true.

3.2. TiB2–TiC–Mo2C based cermets

The effect of Ni binder content on sintered density of50TiB2–50(TiC+ Mo2C) based cermets is shown in Fig. 8,

Table 2Variation of TiB2 grain size with Ni and Ni/Mn binder in TiB2 andTiB2–TiC–Mo2C based cermets sintered at 1350◦C

Baseceramic

Binder(wt.%)

Grain size forNi containingcermet (mm)

Grain size forNi/Mn containingcermet (mm)

TiB2 10 18.7 5.640 16.5 9.9

TiB2–TiC–Mo2C 10 7.0 6.340 6.5 6.9

Fig. 6. Scanning electron micrographs of TiB2–Ni and TiB2–Ni/Mn cer-mets sintered at 1350◦C: (a) TiB2–10 wt.% Ni; (b) TiB2–10 wt.% Ni/Mn.

from which it is evident that the sintered density of thecermets increases with increase in Ni binder content af-ter sintering at either temperature. Increasing the sinteringtemperature slightly increases the sintered density. The %sintered porosity is invariably maximum for 10 wt.% Nibinder composition.

Fig. 9 shows the effect of Ni/Mn binder content on sin-tered density and porosity of boride–carbide based cermets.Sintered density increases with increase in Ni/Mn bindercontent after either temperature sintering. Increase in sinter-ing temperature increases the sintered density. For both thesintering temperatures, cermets with minimum binder con-

Table 3Lattice parameters of binder and TiB2 phases in TiB2 based cermets(sintered at 1350◦C)

% Binder Binder (nm) TiB2 (nm)

10 Ni 0.3520 a = 0.3030, c = 0.3221, c/a = 1.06340 Ni 0.3527 a = 0.3036, c = 0.3226, c/a = 1.062610 Ni/Mn 0.3522 a = 0.3030, c = 0.3217, c/a = 1.061740 Ni/Mn 0.3515 a = 0.3028, c = 0.3226, c/a = 1.0654


Fig. 7. X-ray diffraction scan for TiB2–10% Ni sintered cermet.

tent were found to have maximum porosity. The differencebetween cermets with these two binder systems is appar-ent in sweating response during liquid phase sintering, suchthat 40 wt.% Ni containing cermet exhibits sweating at ei-ther sintering temperature, while Ni/Mn containing cermetdoes not show any such behaviour.

Figs. 10 and 11 show the effect of Ni and Ni/Mn binderon TRS of the cermets such that the value increases withincrease in Ni or Ni/Mn binder content after either tempera-ture sintering. An improvement in TRS value of the cermetshas been observed with increase in sintering temperature. Asignificant variation in TRS is observed after 1350◦C whenthe binder content was increased from 20 to 40 wt.%. Fig. 12shows the fractographs of 50TiB2–50(TiC + Mo2C)–Niand 50TiB2–50(TiC + Mo2C)–Ni/Mn cermets containing10 wt.% binder. In all the cases, the mode of fracture was

Fig. 8. Variation of sintered density and % porosity of TiB2–TiC–Mo2Cbased cermets with nickel binder content.

brittle. SEM photographs of boride–carbide based cermetscontaining 40 wt.% binder are shown in Fig. 13. It is inter-esting to note that in such cermets the hard phase grains arenot that much faceted as in case of TiB2 based systems. Bycomparing the photographs of cermets with Ni and Ni/Mnbinder, it is clear that there has been relatively good wettingin the latter system.

Grain size of boride phase for some of the cermets withNi and Ni/Mn binders are given in Table 2. It is evident thatwith increase in Ni binder grain size gets slightly reduced,but with increase in Ni/Mn binder content the situation isreverse. Further, there is insignificant difference in the grainsize of the hard phase with either binder system.

Lattice parameter variation of TiB2 and binder phasesin this class of sintered cermets (Table 4) reveal that withincrease in Ni binder the lattice parameter of the binder

Fig. 9. Variation of sintered density and % porosity of TiB2–TiC–Mo2Cbased cermets with nickel–manganese binder content.


Fig. 10. Variation of TRS of TiB2–TiC–Mo2C based cermets with nickelbinder content.

Fig. 11. Variation of TRS of TiB2–TiC–Mo2C based cermets withnickel–manganese binder content.

increases slightly, whereas thec/a ratio of TiB2 phase in-creases. In case of cermets with Ni/Mn binder, slight de-crease in the lattice parameter of binder phase andc/a ratioof TiB2 are observed. In case of Ni/Mn sintered cermets,

Table 4Lattice parameters of binder and TiB2 phases in TiB2–TiC–Mo2C basedcermets (sintered at 1350◦C)

% Binder Binder (nm) TiB2 (nm)

10 Ni 0.3523 a = 0.3037, c = 0.3212, c/a = 1.057640 Ni 0.3528 a = 0.3035, c = 0.3230, c/a = 1.064210 Ni/Mn 0.3527 a = 0.3030, c = 0.3235, c/a = 1.06840 Ni/Mn 0.3525 a = 0.3027, c = 0.3227, c/a = 1.066

Fig. 12. SEM fractographs of TiB2–TiC–Mo2C–Ni and TiB2–TiC–Mo2C–Ni/Mn cermets sintered at 1350◦C: (a) TiB2–TiC–Mo2C–10 wt.%Ni; (b) TiB2–TiC–Mo2C–10 wt.% Ni/Mn.

thec/a ratio of TiB2 is greater than in case of Ni containingcermets.

4. Discussion

The results on the effect of different binder additionon both types of ceramic systems are schematically sum-marized in Table 5. In case of TiB2–Ni system, the lowsolubility ratio seems to be the reason for increase in %

Table 5Role of increase in binder content on the sintering behaviour and propertiesof TiB2 and TiB2/TiC/Mo2C based cermets

Properties TiB2 basedcermetsa

TiB2/TiC/Mo2Cbased cermetsa

% Porosity ↑ ⇓? ↓ ⇓TRS ↓ ⇑? ↑ ⇑Grain size ↓ ⇑ Insignificant variationLattice parameter of binder↑ ⇓ ↑ ⇓Binder exudation ↑ Nil ↑ Nil

a ↑ — Ni binder; ⇑ — Ni/Mn binder; ? — no uniform trend.


Fig. 13. Scanning electron micrographs of TiB2–TiC–Mo2C–Niand TiB2–TiC–Mo2C–Ni/Mn cermets sintered at 1350◦C: (a)TiB2–TiC–Mo2C–40 wt.% Ni; (b) TiB2–TiC–Mo2C–40 wt.% Ni/Mn.

sintered porosity with increase in nickel binder, in spite ofthe small wetting angle of nickel with TiB2 [8]. In addition,the possibility of intermediate compound (Ni20.3Ti2.7B6,the tau phase) [2] might also inhibit the melt spreading andsubsequent sintering.

Though the exudation of nickel occurs in both TiB2–Niand TiB2–TiC–Mo2C–Ni cermets (for higher bindercontent), comparatively lesser amount of exudation inTiB2–TiC–Mo2C–Ni cermet confirms that nickel is rela-tively good binder for this system. As nickel dissolves TiCbetter than TiB2, TiB2–TiC–Mo2C based cermets achievebetter densification. Besides the physicochemical aspect, inthe refractory boride/carbide based systems due to the goodstructural and thermodynamic compatibility of TiC withTiB2, a coherence is established between the most denselypacked lattice planes of TiC and TiB2 [7], which promotesdensification. The favourable interfacial match in TiB2 andTiC is assumed to encourage a high mobility of atomsacross the interface, leading to betterment in properties.

The replacement of nickel binder with nickel–manganeseenhances densification due to its better wetting and solubi-

lity ratio compared to nickel. Because of lower melting pointof nickel–manganese alloy (1020◦C) compared to nickel(1453◦C), a significant reduction in viscosity results in abetter densification. Due to the above stated facts, there isnot much difference in the densification and microstructurecharacteristics of cermets containing Ni/Mn binder.

Increasing the binder content enhances the amount of liq-uid phase which facilitates in grain growth of TiB2 hardphase by dissolution–reprecipitation process which is con-firmed by experimental data (Table 2). Smaller grain size ofTiB2 hard phase in the carbide based cermets as comparedto boride based cermets is due to the presence of refractorycarbides which inhibit grain growth of TiB2. The insignif-icant variation in TiB2 grain size in such cermets appearsto be due to the preferential dissolution of TiC/Mo2C inthe binder in comparison to TiB2 based ones. It is interest-ing to note that the lattice dilation of binder phase (Ni orNi/Mn) in case of TiB2 based cermets is more as comparedto TiB2–TiC–Mo2C based cermets. This suggests that thepresence of refractory carbides prevents the dissolution ki-netics of TiB2 in the binder. Further detailed investigationin the direction is needed.

In case of TiB2 based cermets, the decrease in TRS withincrease in binder content is directly related to the increasedporosity. On the other hand, in case of TiB2–TiC–Mo2Cbased cermets better TRS value is attributed to multipleslip systems possessed by TiC [7] and its better densifi-cation as compared to the TiB2 based ones. Though thedifference in % sintered porosity of these two different re-fractory hard phase based systems is not very significant,TiB2–TiC–Mo2C based cermets show positive sinteredproperties as the sintering temperature/nickel binder contentare increased.

5. Conclusions

Considering the results and discussion of the present in-vestigation, the following conclusions can be drawn.1. The % sintered porosity of TiB2–Ni cermet decreases

with increase in binder content but reverse is true forTiB2–Ni/Mn system. With increase in temperature, the% sintered porosity generally decreases in both theabove systems. In case of TiB2–TiC–Mo2C–Ni andTiB2–TiC–Mo2C–Ni/Mn cermets, % sintered poro-sity generally decreases with increase in binder content/sintering temperature.

2. TRS of TiB2–Ni system generally decreases with in-crease in binder amount/sintering temperature but re-verse is true for TiB2–Ni/Mn system. In case of TiB2–TiC–Mo2C–Ni and TiB2–TiC–Mo2C–Ni/Mn systems,TRS generally increases with binder amount/sinteringtemperature.

3. TiB2–Ni cermets generally exhibit a relatively large TiB2grain size as compared to TiB2–Ni/Mn cermets.


References

[1] G. Petzow, R. Telle, New developments in the field of refractoryhard metals based on cemented boride, in: S. Somiya (Ed.),Advanced Ceramics, Terra Scientific Publishing Company, Tokyo,1987, pp. 131–143.

[2] P.S. Sklad, C.S. Yust, Characterisation of TiB2–Ni ceramicsby transmission and analytical electron microscopy, in: R.K.Viswanadham, D.J. Rowcliffe, J. Gurland (Eds.), Science of HardMaterials, Plenum Press, New York, 1981, pp. 911–930.

[3] Z.T. Zakhariev, M.S. Ivanova, I. Serebriakova, Hard materials basedon cemented TiB2–WC–Co Alloys, Z. Metallkd. 85 (11) (1994)801–803.

[4] V.J. Tennery, C.B. Finch, C.S. Yust, G.W. Clark, Structure propertycorrelation for TiB2 ceramics densified using active liquid metals,

in: R.K. Viswanadham, D.J. Rowcliffe, J. Gurland (Eds.), Scienceof Hard Materials, Plenum Press, New York, 1981, pp. 891–909.

[5] D.G. Ahn, A. Kawasaki, R. Watanabe, Microstructure and mechanicalproperties of TiB2–W cermets prepared by HIP, Mater. Trans. JIM37 (5) (1996) 1078–1083.

[6] P. Schwarzkopf, R. Kieffer, Refractory Hard Metals, Macmillan, NewYork, 1953, pp. 276–285.

[7] T.J. Davies, A.A. Ogwu, Characterisation of composite of TiC+TiB2 bonded with nickel based binder alloy: mechanical andmicrostructural properties, Powder Metall. 38 (1) (1995) 39–44.

[8] G.S. Upadhyaya, Sintering aspects of non-oxide ceramics, in: C.Ganguly, S.K. Roy, P.R. Roy (Eds.), Advanced Ceramics, Trans TechPublications, Zurich, 1990, pp. 205–225.

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Sintered porous cermets based on TiB2 and TiB2–TiC–Mo2C