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Acta Materialia 57 (2009) 2607–2612

Continuous in situ functionally graded silicon nitride materials

Manuel Belmonte *, Jesus Gonzalez-Julian, Pilar Miranzo, Maria Isabel Osendi

Institute of Ceramics and Glass, CSIC, Campus de Cantoblanco, Kelsen 5, Madrid 28049, Spain

Received 29 October 2008; received in revised form 8 January 2009; accepted 12 January 2009Available online 25 March 2009

Abstract

Functionally graded materials can enhance the performance of components under demanding operating conditions, although thedevelopment of residual stresses across the gradient and scaling up the manufacturing process to mass production present some limita-tions. To overcome these problems, we present a one-step approach for processing continuous functionally graded silicon nitride (Si3N4)materials from a sole homogeneous powder composition, using spark plasma sintering as a densification technique. Through the controlof the temperature profile within the compact, specimens with a continuous variation of a- and b-phase content, as well as grain size,were achieved. A continuous gradation of mechanical properties, in particular, hardness and toughness, were measured in these speci-mens. This approach offers unprecedented opportunities to design custom-made Si3N4 components by taking advantage of the partic-ularities of field-assisted sintering methods.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Functionally graded materials; Ceramics; Spark plasma sintering; Hardness; Toughness

1. Introduction

The concept of functionally graded materials (FGMs),first proposed in the mid-1980s in Japan [1], is based onthe development of compositional and/or microstructuralgradients within the material leading towards graded prop-erties in one, two or three dimensions [2,3]. Compared tomonolithic and composite materials, these structures offerthe possibility of improving performance of componentsunder demanding technological applications, e.g. in fuse-lage surfaces of spacecraft, thermal and environmental bar-rier coatings for turbines and engines, super-hard cuttingtools or artificial bones [2,4,5]. There are numerous meth-ods to process FGMs [3,4,6]; the main challenge is toreduce residual stress concentrations through the gradientin order to enhance the reliability of these materials.

It is well known that silicon nitride (Si3N4) materialspossess good thermomechanical and tribological proper-ties, and hence these ceramics are suitable for use inwear-resistant technological applications, such as valves

1359-6454/$36.00 � 2009 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2009.01.043

* Corresponding author. Tel.: +34 917355863; fax: +34 917355843.E-mail address: mbelmonte@icv.csic.es (M. Belmonte).

in diesel engines, bearings, sealing rings, and metal cuttingand shaping tools [7]. One important advantage of Si3N4

materials is the ability to tailor their microstructures and,consequently, properties, by controlling the a ? b phasetransformation that occurs during sintering at high temper-ature of the mostly a-phase original Si3N4 powders [8]. Inthis sense, the growth of large, elongated b-Si3N4 grainsof high aspect ratio produces an in situ toughening mecha-nism [9,10], although the material becomes less hard due tothe decreasing a-phase content. The spatial control of thisphase transformation leading to phase gradients in thespecimens would open great technological possibilities,such as high toughness and hardness values in a singlespecimen. There is little information available regardinggraded Si3N4 materials [11,12], or their derivative SiAlONceramics [13,14]. Some of these FGMs consist of layeredsystems which were step-wise fabricated by stacking eithertwo a-Si3N4 powders with different mean particle sizes [11]or tape casting a number of layers with gradual variationsof the a/b SiAlON ratio [14]. The Si3N4 bilayer systemdeveloped by Lee et al. [11] improved the contact damageresistance of Si3N4 ceramics thanks to a hard coating offine a-phase-rich grains on a soft substrate of coarse b-rich

rights reserved.

Fig. 1. Schematic diagram of the SPS graphite die and punch set-up usedto produce Si3N4 FGMs.

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grains, while the graded SiAlON [14] exhibited a steppedgradient in composition and hardness. Other strategies toimprove the contact damage resistance of Si3N4-basedmaterials have included infiltration of a Si3N4 specimenby a silicon oxynitride glass to develop near-surface elasticmodulus gradients [12], or, alternatively, packing a com-pact of b-SiAlON into a-SiAlON powder, which resultsin the formation of a dense a-SiAlON layer during the sin-tering process [13]. In both cases, the surface gradientsextended inside the specimen up to a maximum of 400 lm.

However, the design of layered or stepped materials orsurface gradients present some disadvantages, such as thecomplexity of scaling up the process to mass production,the difficulty in scaling up the FGMs to large specimens,or the nucleation of stresses at interfaces that would com-promise the reliability of the component under operatingconditions. Therefore, to overcome those difficulties, newapproaches must be considered and, in this context, wepresent here the fabrication of continuous in situ function-ally graded Si3N4 material from a sole homogeneous start-ing powder using the spark plasma sintering (SPS)technique [15,16]. This technique is a pressure-assistedpulsed direct current sintering process that generates tem-perature gradients within the powder compact using aproper experimental set-up [17–19], which has beenrecently utilized to make functionally graded boron car-bide–aluminium composites [19]. In the present work, byvarying the contact sections between plungers and die inthe SPS machine, we can gain control over the degree ofa ? b phase transformation in the Si3N4 specimen, andhence can tailor its microstructure and properties.

2. Materials and methods

A homogeneous powder mixture containing a-Si3N4

(SN-E10 grade, UBE Industries) and, as sintering addi-tives, 2 wt.% of Al2O3 (SM8, Baikowski Chimie) and5 wt.% of Y2O3 (Grade C, H.C. Starck GmbH & Co.)was prepared by ball milling for 24 h in ethanol, usingnylon balls. Afterwards, the slurry was dried using a rotaryevaporator and sieved through a 63 lm mesh.

Sintering tests were performed in SPS equipment (SPS-510CE, SPS Syntex Inc.): 4 g of powder mixture was placedin a graphite die, which was surrounded by a porous car-bon felt insulator to reduce the radiation heat losses fromthe outer die wall surface. In addition, the die was linedwith graphite foil to prevent reaction of the sample withthe die, and two graphite foil discs were inserted betweenthe punches and the die to ensure electrical contact.Fig. 1 shows a scheme of the SPS graphite punch set-upused to produce the temperature gradient across the com-pacted Si3N4 powders. The effective current intensity perarea unit, and hence the temperature, was varied by mod-ifying the sections of the punches that are in contact withthe graphite spacers and with the electrodes. Thus, whilethe upper punch diameter was 20 mm, the bottom onewas completely enclosed in the die, which rested on one

spacer. Therefore, in this situation the bottom section con-tact was fixed by the external die diameter, which was50 mm. This asymmetric configuration of the graphitecomponents was used to perform SPS tests at three differ-ent set point temperatures (1550, 1600 and 1650 �C), keep-ing constant the rest of the sintering conditions, i.e. heatingrate of 133 �C min�1 up to 1400 �C, and then 50 �C min�1

up to the maximum temperature; holding stage at the max-imum temperature for 5 min; uniaxial pressure of 50 MPaapplied during the whole thermal cycle; and vacuum pres-sure of 6 Pa. Temperature was controlled with a pyrometerfocused on a hole drilled in the middle of the external diesurface through half the thickness of the die wall. In addi-tion, the pyrometer was aligned with the top side of thebottom punch, which allows the temperature of the bottomside of the sintered specimens to be measured. The sinteredspecimens were labelled as FGM-1550, FGM-1600 andFGM-1650.

Density was measured by the water immersion method.Crystalline phases and a/b transformation degree [20] weredetermined by X-ray diffraction (XRD, Bruker D5000, Sie-mens) analysis of the bulk specimens. The samples were cutinto three small pieces to perform XRD analyses at the cen-tre and near the edge of the specimens. Diffraction patternsalong the cross-sections of the specimens were recorded byremoving around 250 lm from the top side by gentle grind-ing; this was sequentially repeated through the specimenthickness.

The microstructure of the specimens was observed byscanning electron microscopy (SEM, S-4700, Hitachi) onpolished and CF4/5 vol.% O2 plasma-etched surfaces at apower of 100 W for 25 s. Quantitative analysis of themicrostructure was done using imaging analysis tools;more than 1000 features were measured on the SEM micro-graphs to determine Si3N4 grain diameter and apparent

M. Belmonte et al. / Acta Materialia 57 (2009) 2607–2612 2609

aspect ratio distributions. To estimate the porosity, an areaof 300 lm2 was analyzed for each specimen.

Hardness (H) and fracture toughness (KIC) were deter-mined by Vickers indentation methods [21], using a testingmachine provided with a hardness measurement unit(Zwick/Roell, Zhu 2.5). Indentation loads of 49 N wereperformed and stage was moved in steps of 250 lm todetermine the hardness evolution along the cross-sectionsof the FGM specimens. On the other hand, indentationloads at 490 N were done to estimate fracture toughnessat both the top and bottom surfaces of FGMs, as well asalong the cross-sections. H and KIC values represent anaverage of at least three and five indentations, respectively.

3. Results and discussion

Cross-section views of Si3N4 specimens are shown inFig. 2. Two areas are clearly distinguished in the specimenssintered at temperatures 61600 �C: a thicker, dark-greyone in which Si3N4 was fully densified (3.23 g cm�3), anda thinner porous white one. As the set temperatureincreased from 1550 to 1600 �C, the densification boundaryadvanced towards the bottom surface. Finally, the samplesintered at the highest SPS temperature (1650 �C) exhibiteda homogeneous dark-grey colour along its cross-sectionassociated with its complete densification. Therefore, atfirst glance, these results confirmed a temperature gradientwithin the specimens, decreasing from top to bottom. Fur-thermore, the slightly curved densification boundaryrevealed little difference in the radial temperature distribu-tion: the temperature was higher in the specimen core inagreement with the model prediction for non-conductingsamples [22,23].

XRD analyses were initially performed at the top andbottom surfaces of each specimen (Table 1). They showedthe presence of a- and b-Si3N4 as sole crystalline phases,with different a/b ratios depending on the surface (top or

Fig. 2. Low-magnification cross-section views of the different FGMspecimens. The densification boundary is clearly observed for the FGM-1550 and FGM-1600 samples. Specimens dimensions of 4 mm thicknessand 20 mm diameter.

bottom) and the sintering temperature. As expected, anincrease in the SPS temperature, once the material was den-sified, led to higher degree of a ? b phase transformation.In this sense, top dense surfaces exhibited low a-phase con-tents ranging from 12 to 4 vol.% (Table 1), whereas at thebottom surfaces the phase transformation depended on thedensification of the material. Thus, porous white surfaces(FGM-1550 and FGM-1600 specimens) remained nearlyuntransformed (85 vol.% of a-phase compared to95 vol.% of the starting powders), while dense FGM-1650surface contained 61 vol.% of a-phase.

The development of continuously graded structures wasconfirmed by XRD analyses performed on the cross-sec-tions of the specimens. Fig. 3 represents the a-phase linecontours for FGM-1550 and FGM-1650 samples showingthe continuous phase gradation. The two-dimensional con-tour plots are initially flat, as seen at the bottom of thespecimens, and, like the top side, present some curvatureas temperature rises. From that point, the centre exhibiteda slightly larger transformation ratio than the edges. As theb-phase boundary moved towards the top surface, the dif-ferences in a-phase content between the edge and the centreincreased, but never exceeded 10 vol.% (Fig. 3). Thisbehaviour is related to a radial temperature gradient closeto the top punch. Anselmi-Tamburini et al. [22] observedradial temperature gradients in SPS samples, especiallyfor non-conducting materials such as alumina, whichbecame larger as the current increased. The present asym-metric configuration would lead to a curved temperaturedistribution at the top of the sample due to the upperpunch [22], whereas at the bottom a flattened distributionwould occur. Tiwari et al. [23], using thermoelectrical sim-ulations, found a strong dependence of the radial tempera-ture distribution on both power input and thermalconductivity of the sample. According to their simulationsfor the same sample diameter, high Joule heat densities andmoderate thermal conductivity materials, around 20–40 W m�1 K�1, like the present Si3N4 [24], a radial temper-ature difference from centre to edge of �10–20 �C can beestimated. This value is in good agreement with the slightexperimental differences in a-phase content.

Field emission scanning electron microscopy observa-tions (Fig. 4) of the specimens show obvious microstruc-tural differences within each material along their cross-section. FGM-1550 specimen (Fig. 4a) has about6.5 vol.% of porosity (determined by image analysis meth-ods) close to the a-phase-rich bottom end, having noporosity from a height of �600 lm up to the top. Themicrostructure of the porous region has a narrow distribu-tion of equiaxed Si3N4 grains of 180 nm average grain size,D (Table 1), similar to the original Si3N4 crystallite size, i.e.�200 nm [25]. The growth of b-elongated grains graduallyprogressed in the fully dense region towards the top sur-face, leading to a bimodal microstructure of large elon-gated grains embedded in a matrix of equiaxed smallergrains. In fact, significant changes in D (�110%, from180 to 380 nm) and average aspect ratio, AR (�60%, from

Table 1a-phase content, morphological characteristics (average grain size, D; and average aspect ratio, AR) and mechanical parameters (hardness, H; and fracturetoughness, KIC) for the different FGM specimens, in both top (T) and bottom (B) surfaces. H and KIC values represent an average of at least three and fiveindentations, respectively.

Material a-Phase (vol.%) D (nm) AR H (49 N) (GPa) KIC (490 N) (MPa m1/2)

T B T B T B T B T B

FGM-1550 12 85 380 180 2.3 1.4 15.7 ± 0.3 16.2 ± 0.2 5.2 ± 0.1 –FGM-1600 10 85 440 180 2.2 1.4 16.0 ± 0.1 18.7 ± 0.2 5.8 ± 0.2 –FGM-1650 4 61 500 200 2.3 1.5 15.4 ± 0.3 20.4 ± 0.2 5.7 ± 0.1 3.5 ± 0.2

Fig. 3. XRD two-dimensional contour plots representing the a-phasecontents along the cross-sections of (a) FGM-1550 and (b) FGM-1650specimens.

Fig. 4. SEM micrographs of polished and plasma-etched (a) FGM-1550,(b) FGM-1600 and (c) FGM-1650 specimens. The images were taken atincreasing distances (1000 lm steps) from the top surface.

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1.4 to 2.3), were quantified (Table 1). FGM-1600 andFGM-1650 samples (Fig. 4b and c, respectively) exhibiteda similar microstructural trend as FGM-1550, i.e. the mate-rials become more bimodal towards the top surfaces,although FGM-1600 and FGM-1650 exhibit a larger graincoarsening. This phenomenon is explained because the den-sification process progressed more in these specimens. Inthis sense, FGM-1600 presented less porosity (1.1 vol.%)than FGM-1550, achieving full densification at �300 lmfrom the base of the specimen, while FGM-1650 was densealong the whole sample. In both specimens, graded micro-structures were achieved, increasing D from �200 nm to440 nm (120%) and 500 nm (150%) for FGM-1600 andFGM-1650, respectively, with a similar increase in ARfrom �1.5 to �2.3 (�50%).

From these results, temperature differences along theaxial direction of �150 and �50 �C for FGM-1550 and

FGM-1650, respectively, were estimated as follows: (i)the temperature at the bottom surfaces is given by thepyrometer reading and (ii) the top surfaces of both gradedstructures reached 1700 �C, as Si3N4 samples sintered atthat temperature for 5 min with a symmetric positioningof the punches showed a-phase contents below 4%. Thehighest temperature gradient of the FGM-1550 could beassociated with the porosity remaining near the bottomsurface (Fig. 4a), which would produce a low thermal con-ductivity layer. In fact, as the a-phase content at the first�1000 lm from the bottom was similar to that measuredat the bottom surface of FGM-1650 (60%, Fig. 3), a tem-perature of �1650 �C can be estimated at that position,representing a temperature increase of 100 �C in the first1000 lm.

The type of graded microstructures developed in thepresent work can readily be attained by SPS as this tech-nique yields dense Si3N4 materials with almost negligible

Fig. 5. (a) Hardness evolution along the thickness of the FGM-1550 andFGM-1650 specimens. (b) Toughness values assessed along the cross-section of FGM-1650 representing the average values (j symbols), as wellas the data for the different crack path orientations, as depicted in theoptical micrograph in (c) (J, N and . symbols for X, Y top and Y bottomdirections, respectively). KIC measured at the top and the bottom surfaces(circled) are also plotted in (b). The indentation shown in (c) wasperformed at 1000 lm from the top surface. H and KIC values represent anaverage of at least three and five indentations, respectively.

M. Belmonte et al. / Acta Materialia 57 (2009) 2607–2612 2611

phase transformation or grain growth [25–27]. Despitethese experimental results, the sintering mechanisms hereare still under debate. Plasma generation between the par-ticles during the pulsed current stage [28], which wouldpromote particle surface cleaning and enhancement ofmass transport [29], was initially claimed, but it has notbeen confirmed for non-conducting powders. Morerecently, Shen et al. [30] proposed a dynamic ripeningmechanism based on the enhanced motion of charged spe-cies by the electric field that promotes diffusion andhomogenization of the formed liquid. This improvedhomogenization of the liquid would explained the densifi-cation with scarce transformation observed here forSi3N4 FGMs, and may be a consequence of better wettingof the particles by the liquid phase due to interfacial charg-ing phenomena.

Presumably the development of graded microstructuresshould promote gradients in the mechanical properties ofFGMs (Fig. 5). The hardness (H) evolution from top tobottom surfaces for FGM-1550 and FGM-1650 specimensis plotted in Fig. 5a. Both materials exhibited a continuousgradient in hardness that increased with the a-phase con-tent. In the case of dense FGM-1650 sample, H values ran-ged from 15.4 up to 20.4 GPa. Compared with the reportedhardness values for a-rich Si3N4 materials, typically in therange of 18–20 GPa [8,31], the result for FGM-1650 isquite remarkable taking into account its high b-phase con-tent (�40 vol.%). This improvement in hardness is due tothe benefits of having submicron, instead of micron-sized,grain structures [32]. On the other hand, the abruptdecrease in hardness observed for FGM-1550 beyond the3000 lm from the top surface can be explained by theincreased porosity in this material. The hardness at zeroporosity, H0, can be assessed using the following equation[33]:

HH 0

¼ ð1� hÞ2ebh;

where H is the measured hardness, h is the porosity of thespecimen measured by image analysis methods, and b is aconstant parameter equal to 2.6. Substituting the adequatevalues in the above expression gives a number of 22 GPafor H0, which seems reasonable for such a a-rich densematerial (85%) of fine grain size (�200 nm). In bothFGMs, the difference in hardness values between the twosides was greater than 30%.

Differences in fracture toughness values, KIC, measuredat the bottom and top surfaces were clearly observed forFGM-1650 (Table 1), ranging from 3.5 to 5.7 MPa m½,which means an increase in KIC higher than 60%. In thecase of FGM-1550 and FGM-1600 specimens, the presenceof pores and the stresses induced by the high a-phase con-tent precluded correct indentation crack measurement atthe bottom surfaces. However, taking into account thecomparable KIC values assessed for those samples at thetop surfaces, compared to that of FGM-1650, a similartrend in toughness should also be expected.

Fracture toughness values along the cross-section ofFGM-1650 specimen are collected in Fig. 5b, and they wereindependently assessed for different crack path orientationsas depicted in the optical micrograph of Fig. 5. At firstglance, two main issues can be established. The first oneis the confirmation of a continuous toughness gradientalong the thickness of the specimen (see average plot inFig. 5b); as expected, the variations in fracture toughness

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showed a reverse trend to that of hardness. The secondimportant issue is the significant difference in crack exten-sion and, therefore, in KIC, for X and Y crack path direc-tions (Fig. 5b and c). Moreover, cracks in the Y directionare always shorter, i.e. higher KIC, when propagatingtowards the top surface (Y top) than towards the bottomone (Y bottom). These differences in KIC behaviour areexplained by the microstructure evolution within the spec-imen. When cracks extend towards b-rich top zones withelongated coarser grains (Table 1 and Fig. 4c), typicaltoughening mechanisms such as pull-out, crack deflectionor crack bridging are developed, which improve crackpropagation resistance (maximum KIC value was6.2 MPa m½ for a crack growing in the Y top direction at1000 lm from top surface; see Fig. 5b). Accordingly, crackspropagating towards the bottom surface, with equiaxedfine-grained microstructures, showed lower KIC values(Table 1 and Fig. 5b). Furthermore, the axial pressureapplied during SPS induced a certain alignment of the elon-gated grains perpendicular to the pressing direction,favouring the development of toughening mechanisms inthe Y direction instead of the X.

As has been demonstrated, Si3N4 materials with gradedmicrostructures and properties can be obtained by control-ling the temperature profile within the compact by adjust-ing the SPS die set-up. This original method opens newopportunities for developing a range of Si3N4 components.In this sense, contact damage resistance could be enhancedby gradually increasing the b-phase content towards thecontact surface [11]. Furthermore, materials with localizedsuperplastic deformation [34], where the grain growth ofnanostructures is locally restrained, might be also devel-oped. Finally, carbon nanotubes could be grown intograded porous structures supported mechanically by adense substrate of similar composition [35], allowing theiruse as catalysts or membranes.

4. Summary

Continuous in situ functionally graded Si3N4 materialshave been fabricated from a sole homogeneous startingpowder using the SPS technique. The approach is basedon developing a temperature gradient within the startingpowder compact, which allowed a close tailoring of theSi3N4 microstructure. Materials with continuous gradientsin a/b phase contents and grain sizes were readily obtained,and continuous variation in the mechanical response ofthese structures was evidenced as well.

Acknowledgements

Thanks are given to A. de Pablos (ICV, Spain) for assis-tance with SPS and XRD tests. Funding for this work was

provided by Spanish Ministry of Science and Innovation(MICINN) and Spanish National Research Council(CSIC) through projects MAT2006-7118, HA2007-0083and CSIC200660I191. J. Gonzalez-Julian acknowledgesthe financial support of the JAE (CSIC) fellowshipProgram.

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