Development of Fe NbC Cermet From Powder Obtained by Self Propagating High Temperature Synthesis...

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Development of FeNbC cermet from powderobtained by self-propagating hightemperature synthesis

P. G. Esteban* and E. Gordo

The present work studies the development of a 30Fe70NbC (wt-%) cermet by powder metallurgy

techniques. The base powder was obtained by the process self-propagating high temperature

synthesis, and this powder was modified by mechanical milling using two different approaches:

comminution milling to reduce particle size (without modifying powder microstructure) and high

energy milling to reduce both particle and reinforcement size in order to improve powder

microstructure. Both modified powders were compacted by uniaxial and cold isostatic pressing.

The compacts were sintered in vacuum at different temperatures and times, and characterised by

several techniques, including the determination of density and hardness, study by XRD to

determine the constituents after sintering, and a complete microstructural analysis by SEM. The

present study reveals that high energy milling permits to obtain enhanced materials when

compared with materials sintered from comminuted powder.

Keywords: Cermet, FeNbC, Iron based composite materials, High energy milling, Self-propagating high temperature synthesis, Uniaxial pressing, Coldisostatic pressing

IntroductionDevelopment of iron matrix composites with highceramic reinforcement contents opens an interestingway for science and technology. The technical motiva-tion for these materials begins in the interest of theindustries of cemented carbides and cermets to sub-stitute totally or partially either the cobalt or nickelmatrix, because Co and Ni are expensive materials andalso considered hazardous for health and environment.Also, iron matrix materials are attractive because oftheir abundance (that reduces costs) and recyclingability.1,2 Iron has the possibility of being heat treated3

and it can be used in a broad range of applications,especially in the case of high temperature resistantmatrixes. From a scientist point of view, it is a lowinvestigated knowledge area because most studies onmetal matrix composites (MMCs) are based on non-ironmatrix composites.4,5 However, although iron has a lowsolubility with ceramic phases, the risk in an iron matrixmaterial with ceramic reinforcement is the possiblereactions between components that can happen duringmanufacturing, promoting ceramic dissolution andreprecipitation in brittle phases.

There are many published studies on ceramic reinforcediron matrix composites, but most of them study materialswith low ceramic contents, up to 15 wt-%, because of thedifficulties found to increase reinforcement contents

without decreasing mechanical properties.68 One of thesedifficulties is to achieve an homogeneous reinforcementdispersion in the metal matrix and a right bond betweenmatrix and reinforcement. Therefore, it could be adesirable option to start from a composite powder formedby a reinforcement embedded in a metal matrix. Highenergy milling9 has been used in many studies to createthis type of composite powder.10,11 However, the presentwork starts from a FeNbC composite powder obtainedby self-propagating high temperature synthesis1214 (SHS)process and the particles of this powder are formed byNbC embedded in an iron matrix. The high carbidecontent of this powder leads to consolidation difficultiesusing conventional powder metallurgy techniques.Uniaxial compaction and cold isostatic pressing (CIP)are compared in the present study in order to show thedifferent properties of final materials depending on theconsolidation technique employed.

The main objective of the present work is thedevelopment of an iron based composite material withhigh content of reinforcement: the composition of thepowder is 30 wt-%Fe (matrix) and 70 wt-%NbC (rein-forcement). The initial powders were obtained by SHSand then processed by powder metallurgy techniques.This material is conceived to replace cemented carbidesin some tool applications.

In order to achieve this main objective, the research isapproached through the following partial steps:

(i) characterisation of the initial powder

(ii) performing a comminution milling to reduceparticle size and a high energy milling toenhance powder microstructure

Department of Materials Science and Engineering, University Carlos III ofMadrid, Spain

*Corresponding author, email [email protected]

2006 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteReceived 24 October 2005; accepted 11 February 2006DOI 10.1179/174329006X102825 Powder Metallurgy 2006 VOL 49 NO 2 153

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(iii) powder pressing by uniaxial and cold isostaticpressing to determine differences in the finalmaterials

(iv) sintering study to evaluate densities, hardnessand dimensional shrinkage of the samplesobtained by the different milling processes andcompaction methods employed

(v) analysis of the microstructures obtained in thedifferent conditions.

Experimental procedureThe base material is a composite powder with iron asmatrix and NbC as reinforcement, with a maximumparticle size of 300 mm. This powder is manufacturedby the SHS process in London & ScandinavianMetallurgical Co. Ltd. (LSM, England). The composi-tion of the powder is given in Table 1.

Starting from initial powder, Fig. 1 shows the guide-lines followed in the present work. First of all, initialpowder was characterised and preliminary experimentswere performed to evaluate its capability of beingpressed and sintered. As a result, it was decided tocarry out a milling process of the powder in order toadjust it. Therefore, it was mechanically milled in twoways, a comminution milling to reduce particle size anda high energy milling in order to enhance powdermicrostructure. Therefore, the present work covers thestudy of the three types of powder, as shown in Fig. 1.

Comminution milling was performed in a centrifugalball mill (Fritsch Pulverisette 6), with a powder/ballsweight ratio of 1 : 1, at 550 rev min21 rotating speed andin air atmosphere. The diameter of the balls is 10 mm,and both the balls and the container are made of amartensitic stainless steel. High energy milling wasperformed in a planetary ball mill (Fritsch Pulverisette5), with a powder/balls weight ratio of 1 : 10, at400 rev min21 rotary speed, in air atmosphere and with10 mm diameter stainless steel martensitic balls.

In order to study both milling processes, samples ofpowder were taken every 2 h during milling to evaluate

changes and determine the optimum milling time for thetwo milling processes.

The characterisation of the powder was performedwith these techniques: X-ray diffraction (XRD), scan-ning electron microscopy (SEM), carbon content analy-sis and particle size laser analysis.

X-ray diffraction was used to determine compositionand amorphisation of the powders during the millingprocesses and to identify the possible new phases formedin the process. Diffraction was carried out in a PhilipsXPert diffractometer customised with these parameters:40 kV, 40 mA, Cu radiation l51.5405 and countingtime 0.6 s.

Morphological and microstructural analyses wereperformed with the aid of SEM in a Philips XL 30microscope, provided with EDX to take semiquantita-tive composition analysis.

Carbon content was controlled to study decarburisa-tion of the powder during the milling processes. Thisanalysis was performed in a LECO CS-200 equipment.

Two different powder compacting methods were used:uniaxial pressing at 700 MPa and cold isostatic pressingat 2500 bar (250 MPa); the latter maintained for 5 min.Green density was calculated from the ratio of the massand the volume of the green compacts. Green densityvalues can show differences between the two pressingtechniques and its influence on the density of sinteredsamples.

Differential thermal analysis (DTA) was carriedout to find out changes in materials during theheating and sintering stages. DTA were performedunder the same conditions as sintering. This analysis isuseful to determine the possible presence of a liquidphase during sintering. The equipment used was aNetzsch STA 409 C.

Sintering was carried out in a high vacuum furnace(Carbolite HVT, 1024 mbar). Sintering temperature wasvaried in the range of 13001450uC and two differentsintering times were used: 30 and 60 min. Heating andcooling speeds were adjusted to 5uC min21.After sintering, density of samples was measured by

the Archimedes method. The values obtained are shownas relative densities calculated as a percentage of thetheoretical value, which was obtained by the rule ofphases, resulting in a value of 7.77 g cm23.

Hardness was also measuresd using the Vickersmethod with 30 kg load (HV30).

Also, XRD analysis was performed to evaluatechanges in compounds after sintering. All sinteredsamples were analysed by SEM in order to establish alink between properties and microstructures.

Results and discussion

Characterisation of powdersInitial powder has a maximum particle size of 300 mm,with an average size of 5 mm, which reveals that thepowder is brittle (owing to its high content in NbC). Themeasured carbon content is 7.26 wt-%, finding out anincrease of this content while analysing small particles(which match with fragmented carbides). The morphol-ogy and microstructure of the initial powder are shownin Figs. 2 and 3 respectively.

The microstructure shown in Fig. 3 is the key tounderstand the main motivation of the present study:

Table 1 Composition of initial powder, wt-%

Nb Fe C Al O Ti Si N Mn

61.25 28.97 7.49 0.64 0.61 0.55 0.26 0.12 0.11

1 Experimental procedure scheme

Esteban and Gordo FeNbC cermet from powder by self-propagating high temperature synthesis

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each particle is itself a composite material. But there aretwo factors that seem not to be convenient in advance:maximum particle size is very large to achieve goodcompaction and sintering and carbide size is too large toensure good properties in sintered materials. For thesereasons, the present work studies two different types ofmilling to change the characteristics of the initialpowder. On one hand, a comminution milling wascarried out to reduce particle size without changingmicrostructure and on the other hand, a high energymilling was performed, which can change both particlesize and microstructure, decreasing reinforcement size,promoting an homogeneous distribution of carbides andmaking an energy activation of the powder to promotesintering.

Powder microstructure after comminution milling isshown in Fig. 4. The maximum particle size is 100 mm,with an average size of 4 mm and a 7.35 wt-% carboncontent (note that there is not variation with regard tothe original powder). Microstructure has not beenmodified, containing the same large carbides as theoriginal powder. Figure 5 shows a particle of powderafter high energy milling for 6 h, in which a decrease inboth particle and carbide sizes can be seen in ahomogeneous microstructure that has changed withregard to the original powder. High energy powder hada 10 mm average particle size, with a 6.33 wt-% carboncontent, which in fact reveals decarburising after highenergy milling.

The average particle size and wt-% carbon content forthe three types of powder are shown in Table 2.

There were no changes in the number of constituentsof the powder after millings, as revealed by the XRDanalyses of the both milled powders and original powder(see Fig. 6). However, the XRD analysis shows adifference between high energy powder and the twoother samples. Although the three samples have the

2 Initial powder morphology

3 Initial powder microstructure

Table 2 Average particle size and carbon content, wt-%

Initialpowder

Comminutedpowder

High energypowder

Average particle size, mm 5 4 10Carbon content, wt-% 7.26 7.35 6.33

4 Microstructure of comminuted powder

5 Microstructure of high energy milled powder

6 X-ray diffraction patterns of original powder, comminu-

ted powder and high energy powder


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same composition, peaks of high energy powder areshorter and thicker than those of the two others,showing the amorphisation of high energy powder,activated to promote sintering.

CompactionPressing of initial powder was found difficult, leading toobtain weak green compacts and porous sinteredsamples with low hardness. These results reveal theneed of decreasing powder particle size. Preliminaryexperiments showed that green compacts made fromcomminuted powder were stronger than those madefrom initial powder, therefore it was found necessary toperform the comminution milling.

Both comminuted and high energy milled powderswere uniaxially and isostatically pressed as described inthe Experimental procedure and in Fig. 7. The values ofthe green densities obtained are shown in Table 3.

Related to compaction methods, uniaxial compactsachieve higher green densities than isostatic samples,which is due to the lower pressure employed in the latter.Related to the type of powder, samples made fromcomminuted powder reach higher green densities thansamples made from high energy powder, as it can beexpected because high energy powders are strainhardened after the milling process.7

Sintered materialsMaterials obtained from comminuted powder

Figure 8 shows relative densities for both uniaxial andisostatic samples made from comminuted powder,sintered for 30 and 60 min. For the same sinteringtimes, uniaxial samples show higher densities than theisostatic samples. Also, samples sintered for 60 minreach higher densities than samples sintered for 30 min.Maximum densities are reached by uniaxial samplessintered for 60 min while isostatic samples sintered for30 min have the lower density values.

The same behaviour has been found for hardnessvalues. Figure 9 shows that maximum hardness isachieved by uniaxial samples sintered for 60 min. Interms of density and hardness, uniaxial samples sinteredfor 30 min are equivalent to isostatic samples sinteredfor 60 min.

These properties are related to sample microstruc-tures, revealing porous microstructures of those samplessintered below 1400uC. Also, samples sintered for30 min present higher porosity than samples sinteredfor 60 min. Figure 10 shows the microstructure of anuniaxial sample from comminuted powder sintered at1350uC for 60 min. The left image shows the highporosity encountered while the right image shows the

7 Scheme of different powders and compaction methods

employed

Table 3 Relative density of green compacts

Uniaxialcompacts, %

Isostaticcompacts, %

Comminuted powder 70 67High energy powder 61 58

8 Relative densities of uniaxial and isostatic samples

made from comminuted powder, sintered in vacuum

for 30 and 60 min

9 Measured hardness of uniaxial and isostatic samples

made from comminuted powder sintered for 30 and

60 min

10 Microstructure of uniaxial sample made from com-

minuted powder sintered at 1350uC for 60 min: leftimage in secondary electron (SE) and right image in

backscattered electron (BSE)



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two phases corresponding to Fe (grey contrast) andNbC (bright contrast).

At the temperature of 1400uC, porosity is reducedconsiderably, probably owing to the formation of aliquid phase during sintering. Figure 11 shows thesection of an uniaxial sample from comminuted powdersintered at 1400uC for 60 min. It presents two regions,corresponding to the outer and the inner regions of thesample. The outer region (left side of the section) showslower porosity than the inner one (right side), which is aporous region.

Analyses in the outer region have revealed thebeginning of an intermetallic formation, identified asFe2Nb. A detail image of this formation at 1425uC isshown in Fig. 12.

The intermetallic phase is more extensive at highertemperatures and sintering times, forming continuousregions which makes the material more fragile. Thisintermetallic phase is present in smaller amount inisostatic samples than in uniaxial samples. The amountof intermetallic phase seems to be related to the amountof residual porosity in the sample. The intermetallicformation hypothetically begins when there is no longerexpectation of porosity reduction in the sample. Then,after 60 min sintering, isostatic samples have stillcapability of reducing their porosity while uniaxialsamples reach their minimum porosity, beginning theformation of the intermetallic.

This intermetallic phase is identified as e in the FeNbphase diagram15 of Fig. 13, which could appear duringthe solidification of the liquid phase formed duringsintering at temperatures over 1400uC. Decarburisationof the liquid phase (probably induced by oxidation)could force the liquid phase to solidify by the FeNbdiagram instead of the FeNbC one.The identification of the intermetallic phase was

confirmed by XRD and its composition was analysedby EDX (see Table 4). This composition is in the rangeof the e phase in the FeNb phase diagram.

Materials obtained from high energy milled powder

Next results can demonstrate that high energy materialsachieve enhanced properties with respect to comminu-tion materials. These properties can be related to theimproved microstructure of high energy materials.Figure 14 shows the density curves of comminuted

and high energy samples sintered for 60 min. Bothuniaxial and isostatic samples from high energy milled

11 Secondary electron image of section of uniaxial sam-

ple sintered at 1400uC for 60 min

12 Backscattered electron image of formation of Fe2Nb

intermetallic phase under sintering conditions of

1425uC and 60 min

13 FeNb phase diagram15

Table 4 Intermetallic EDX analysis

Element wt-% at.-%

SiK 8.18 17.81NbL 41.97 27.62FeK 49.85 54.57Total 100 100

14 Relative density comparison between comminuted

and high energy (HE) samples sintered for 60 min



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powders reach higher density values than those madefrom comminuted powder. Uniaxial samples obtainbetter properties than isostatic samples, as it happenedwith the materials made by comminution material.

Hardness is much higher in those materials made fromhigh energy powder than in those made from commin-uted powder. Figure 15 shows that comminutionmaterials reach a maximum value near 500 HV30 whilehigh energy samples obtain a maximum value of800 HV30 at 1450uC.The explanation to these differences in properties can

again be explained by the microstructure of the samples.Comparing the microstructures of samples from com-minuted powder and high energy powder, it can be seenthat the latter present finer carbides, less porosity and amore homogeneous microstructure than the samplesfrom comminuted powder. These differences can beanalysed comparing Fig. 16 (isostatic comminutionsample sintered at 1425uC for 60 min) and Fig. 17(isostatic high energy sample sintered at 1400uC for30 min). Note that both images are taken at the samemagnification. These results show the improvementobtained by high energy milling the initial powder, asit was the main objective of this milling.

Figure 18 shows a uniaxial sample, where the mediumcontrast matches with the intermetallic phase. Figure 19shows an isostatic sample, where the main phases are Fe(dark contrast) and NbC (white contrast), and there is a

little amount of intermetallic phase (grey contrast).Looking into the detail microstructure of high energysamples in Fig. 18, it can be observed that high energysamples have more intermetallic amount in theirmicrostructure than comminution samples. Also, thepresence of intermetallic phase is clearly higher inuniaxial samples than in isostatic samples. This seemsto be the reason that leads uniaxial samples to havemuch more hardness than isostatic samples.

Although hardness is higher because of the presenceof this intermetallic phase, it could also cause the

15 Hardness comparison between comminution and high

energy samples sintered for 60 min

16 Microstructure of isostatic sample made from com-

minuted powder sintered at 1425uC for 60 min: leftimage in SE and right image in BSE

17 Microstructure of isostatic sample made from high

energy milling powder, sintered at 1400uC for 30 min:left image in SE and right in BSE

18 Microstructure of high energy uniaxially pressed

material, sintered at 1350uC for 60 min

19 Microstructure of high energy isostatically pressed

material, sintered at 1375uC for 60 min



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fragilisation of the material. This amount of interme-tallic phase in uniaxial samples is possibly related toporosity as it was explained in comminution materials.When a near minimum porosity state is reached, theformation of the intermetallic phase seems to begin. Thisformation succeeds at lower temperatures and times inuniaxial samples than in isostatic materials. Thisintermetallic phase was identified by XRD as shown inFig. 20.In high energy materials uniaxially pressed, the

intermetallic formation starts at lower temperatures incomparison with comminution samples. Because thehigh energy materials are energy activated, the liquidphase could appear at temperatures lower than 1400uCand the solidification of this phase would follow the FeNb phase diagram. Also, before forming the liquidphase, this energy activation could promote diffusion ofcarbon from the NbC to the nearby iron, forming a FeNb system in decarburised regions. Decarburising wasalso shown during the milling process, which is also animportant factor that contributes to the intermetallicformation.

Conclusions1. It has been possible to achieve consolidation and

sintering of the initial FeNbC powder. However, it wasneeded to process the powder by high energy milling inorder to enhance material properties.2. Both comminution and high energy milling pro-

cesses lead to obtain two different microstructures ofpowders, which translate into different sintering beha-viour, different microstructures and different finalproperties of the sintered materials. Materials madefrom high energy milled powders reach higher values of

density and hardness than those materials made fromcomminuted powder.

3. There is a presence of a Fe2Nb intermetallic phasewhich hardens the material, but could fragilise thematerial. The presence of this intermetallic phase ishigher in uniaxial samples than in isostatic samples.Also, this intermetallic phase is more extensive in highenergy materials than in comminution materials.

4. Uniaxial pressed samples reach higher densitiesand hardness than isostatic samples. Besides, uniaxialsamples have more amount of intermetallic phase.

Acknowledgements

The authors wish to thank the staff members of theDepartment of Materials Science and Engineering at theUniversity Carlos III of Madrid for their support andtheir enthusiasm during the making of the present work.Also especial thanks to the European PowderMetallurgy Association (EPMA) for the award givento the Master Thesis related to the present work, whichwas presented in the PM2004 Congress celebrated atVienna.

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20 X-ray diffraction analysis of sintered and uniaxially

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