The microstructure and boundary phases of in-situ ... · 1. Introduction Silicon nitride (Si 3 N 4)...

Materials Science and Engineering A254 (1998) 242–252

The microstructure and boundary phases of in-situ reinforcedsilicon nitride

Mingqi Liu, Sia Nemat-Nasser *Center of Excellence for Ad6anced Materials, Department of Applied Mechanics and Engineering Sciences, Uni6ersity of California, San Diego,

La Jolla, CA 92093-0416, USA

Received 29 November 1997; received in revised form 13 March 1998

Abstract

The microstructure of an in-situ reinforced silicon nitride, gas pressure sintered with La2O3, Y2O3, and SrO additives and thenheat treated, is examined with X-ray diffraction, SEM, and high-resolution TEM. Two crystalline rare-earth apatite phases,La5Si3O12N and Y5Si3O12N, are identified at the grain pockets and at the two-grain boundaries. The thickness of the crystallinephases at the two-grain boundaries is approximately 1.7 nm, in compliance with the suggested equilibrium intergranular spacing.A glassy phase is also present at the grain pockets and at the two-grain boundaries due to incomplete crystallization of theboundary phases. The thickness of the amorphous phase at the two-grain boundaries varies from 0.7 to 3.0 nm, suggesting thatcompositional inhomogeneities exist in these areas. Based on the microstructural observations, the structures of the crystallineboundary phases, the equilibrium intergranular film thickness, and the mechanisms causing incomplete recrystallization of theglassy phase in the in-situ reinforced silicon nitride are discussed. © 1998 Published by Elsevier Science S.A. All rights reserved.

Keywords: Microstructure; Boundary phases; In-situ reinforced silicon nitride

1. Introduction

Silicon nitride (Si3N4) materials are one class ofpromising structural materials for high-temperature ap-plications. High resistance to thermal shock, as well ashigh strength, high fracture toughness, and high resis-tance to chemical attack, make silicon nitrides suitablematerials for turbine engine components operating un-der extreme conditions. Significant advances in develop-ing silicon nitrides for gas turbine applications havebeen achieved in the past few years. A silicon nitridehaving a pronounced acicular microstructure, known asthe in-situ reinforced (ISR) Si3N4, shows significantlyimproved properties including creep resistance, fracturetoughness, and strength consistency [1–4]. The superiormechanical properties of ISR silicon nitride are mainlyattributed to the in-situ grown b-Si3N4 grains with highaspect ratios, as well as their unique grain-boundarystructure.

It is known that the bonding between Si and N isbasically covalent. The solid-state diffusion is very slow,thus preventing densification of the silicon nitride.However, if a liquid phase is introduced into the systemduring sintering, high densities are achieved. This liquidphase is introduced by means of oxide sintering addi-tives, which form a eutectic liquid with the oxidizedsurface layers of silicon nitride powders. After sintering,this liquid phase is usually retained in the glassy inter-granular phases which cause a deterioration in themechanical properties of the material, particularly atelevated temperatures. Five approaches for reducingthe detrimental influence of the glassy phase have beensuggested: (i) transient liquid-phase sintering whereconstituents of the glassy phase are absorbed into thesilicon nitride grains by solid solution [5], (ii) post-sin-tering heat treatments to alter the glassy phase compo-sition and structure [6–8], (iii) increasing therefractoriness of the boundary phase [9], (iv) usingdensification processes such as gas pressure sinteringwhich requires less liquid phase [10], and (v) crystalliz-ing the grain-boundary phase by properly selectingsintering additives and/or heat treatment [11,12]. The

* Corresponding author. Tel.: +1 619 5344772; fax: +1 6195342727; e-mail: [email protected]

0921-5093/98/ - see front matter © 1998 Published by Elsevier Science S.A. All rights reserved.

PII S0921-5093(98)00679-0

M. Liu, S. Nemat-Nasser / Materials Science and Engineering A254 (1998) 242–252 243

last four approaches have been used extensively in therecent development of high-temperature silicon nitrideceramics, including the ISR silicon nitride materialinvestigated in the present study.

In the last two decades, there has been growingrecognition that the properties of Si3N4 materials aresignificantly influenced by the grain-boundary phasespresent [13–19]. Since a basic understanding of theintergranular phase and structures is essential if im-provements in the performance of silicon nitride are tobe achieved, extensive research has been done on thecharacterization of the grain-boundary phases of siliconnitride materials. By adding rare-earth oxide additivesto silicon nitrides, Thomas et al. [15,20–23] have founda family of RE2Si2O7 crystalline grain-boundary phasesin silicon nitrides which were pressureless-sintered orsintered under relatively low pressures, where RE repre-sents Y, Sm, Gd, Dy, Er, and Yb. On the other hand,a new group of silicon lanthanide oxynitrides of thegeneral formula Ln5Si3O12N, called rare-earth apatiteor H-phase structure, have been identified by Hamon etal. [24] and Rae et al. [25] in mixtures of silicon nitrideand lanthanide oxides hot-pressed at relatively highpressures, where Ln can be Y [25–27], La [24,28,29], Sr[30], Nd, Sm, Gd, and Yb [24].

The purpose of the present research was to study themicrostructure of an in-situ reinforced silicon nitridewith transmission electron microscopy (TEM). The em-phasis of the research, in particular, is placed on thestructural characterization of grain-boundary phases byhigh-resolution electron imaging and electron diffrac-tion in order to gain information on their crucial role ininfluencing the materials’ high-temperature mechanicalproperties. This initial work is currently being followedby systematic mechanical testing of this material atrelatively high temperatures.

2. Experimental procedure

The in-situ reinforced silicon nitride (AS800) used inthe present study is produced by AlliedSignalAerospace, NJ. Silicon nitride powders are sintered at atemperature above 1750°C in a nitrogen atmospherewith a gas pressure of about 20.7 MPa (3000 psi). Thesintering aids include La2O3, Y2O3, and a small amountof SrO. After sintering, a post-sintering heat treatmentis performed. The material thus obtained has a 4-pointbend strength of about 500 MPa (72 ksi) at 1400°C anda room-temperature fracture toughness of about 8 MPam−2 [2].

The as-received material was in the form of a billet,with dimensions of approximately 20 cm×12 cm×2cm. The material was first examined with a Cambridge360 scanning electron microscope (SEM). Silicon, lan-thanum, yttrium, and strontium peaks were identified in

the energy dispersive X-ray spectroscopic (EDS) spec-trum obtained from the polished surface of the mate-rial, and the lanthanum peaks are found to besignificantly stronger than the yttrium and strontiumpeaks (Fig. 1), because La2O3 is the major sinteringadditive. An unetched sample was also analyzed by anXDS-2000 X-ray diffractometer, with a scan rate of 1°min−1. The diffractometer is equipped with a computeranalyzer and a line focus Cu tube operated at 40 kVand 40 mA, and has been calibrated using an externalsilicon standard with a lattice constant of 0.54306 nm at25°C. Fig. 2 is an X-ray diffraction pattern from thematerial, showing large peaks from b-Si3N4, as well assmall peaks from some other crystalline phases. Ananalysis of these small peaks indicates that they belongto either La5Si3O12N or Y5Si3O12N, suggesting thatcrystalline boundary phases may have formed duringsintering and/or post-sintering heat treatment. No a-Si3N4 peaks are observed in the diffraction pattern,confirming a full transformation from a-Si3N4 to b-Si3N4 during high-temperature sintering.

Detailed examination of the in-situ reinforced siliconnitride was performed on a JEOL ARM-1000 highresolution TEM, operated at 800 kV, at the NationalCenter for Electron Microscopy (NCEM), LawrenceBerkeley Laboratory, University of California, Berke-ley, CA. The microscope has a point-to-point resolutionof 1.6 A and a specimen bi-axial tilting capability of940°. The high TEM operating voltage not only im-proves the resolution and reduces ionization damage,but also allows increased penetration [31]. TEM sam-ples were prepared from the central region of the billet.Standard 3 mm discs were ground down to approxi-mately 100 mm in thickness, polished on both sides toremove the residual stresses, dimpled to near-opticaltransparency in the central region, and finally ion-

Fig. 1. EDS spectrum from the in-situ reinforced silicon nitride,revealing peaks of Si, La, Y, and Sr. The Au peaks are from thecoating material.

M. Liu, S. Nemat-Nasser / Materials Science and Engineering A254 (1998) 242–252244

Fig. 2. X-ray diffraction pattern from a unetched specimen, revealing the presence of three crystalline structures: (A) b-Si3N4, (B) La5Si3O12N,and (C) Y5Si3O12N.

milled to produce large areas sufficiently thin for TEMexamination.

The TEM operation conditions for determining theatomic configuration in the vicinity of the ceramicgrain-boundaries have been discussed by Clarke andThomas [32,33]. Accordingly, the boundary has to betilted to an edge-on position to the incident beam.Furthermore, the interface should be flat with a lowdensity of interfacial steps, and the grains on either sideof the interface should be aligned in an orientationsuitable for structure imaging or at least for formingone set of lattice fringes. The structure of the grain-boundary phase was determined by using the selected-area electron diffraction (SAED) technique on aJOEL-2000FX TEM.

3. Results

Fig. 3(a) is a SEM micrograph of the polished sur-face of the silicon nitride specimen. The material con-tains approximately 90% Si3N4 grains and 10% secondphase at the grain boundaries and the grain pockets(multi-grain junctions). The grain structure can beclearly seen after the specimen is etched at 400°C inmolten NaOH for 4–6 min. As shown in Fig. 3(b), theAS800 silicon nitride consists primarily of acicularSi3N4 grains. The acicular microstructure of the mate-rial is believed to result from anisotropic growth of thehexagonal Si3N4 grains to maximize the area of the

low-energy (100) prismatic planes. Because of the highsintering temperature, Si3N4 grains with various grainwidths and lengths are formed. The average grain widthwas determined to be approximately 0.8 mm and theaspect ratio is greater than 4. Fig. 4 is a SEM frac-tograph of the material subjected to high-strain-rateshear, revealing a typical interlocking microstructure,which is highly resistant to deformation. Under com-pressive loads, the acicular grains interlock, resulting inincreased creep resistance by inhibiting grain-boundarysliding. Under tensile stresses, in addition to limitinggrain-boundary sliding, the elongated grains improvethe stress rupture properties by bridging the microc-racks [3].

A similar microstructure is observed in low-magnifi-cation TEM micrographs, as shown in Fig. 5(a) and(b). The microstructure is characterized by low-contrastSi3N4 grains and a high-contrast boundary phase at thegrain pockets. Strain contours, resulting from the resid-ual stresses within the grains and the contact stressbetween neighboring grains, are observed in the grain-boundary vicinity. Like other silicon nitride materials,some dislocations are found in the specimen, mainly inthe large grains. SAED patterns from more than 50individual grains reveal no evidence of the existence ofa-Si3N4 grains, which is consistent with the XRD re-sults (Fig. 2). Therefore, a complete transformationfrom a-Si3N4 to b-Si3N4 has been achieved by theprocessing conditions employed. No second phase par-ticles are identified in the microstructure, indicating


that all the sintering additives were dissolved in theeutectic liquid during the sintering process.

Fig. 6 shows a typical two-grain boundary betweentwo grain pockets. The silicon nitride grains are seen atbright contrast, while the grain-boundary phase ap-pears at dark contrast with near-uniform intensity. Thegrain boundary appears to be straight and well-defined,having a very small film thickness. High-resolutionTEM images and SAED patterns from the boundaryareas of the material reveal the presence of an amor-phous or glassy phase and fully-crystallized phases, aswill be illustrated below.

Fig. 7 is a high-resolution electron micrograph takenat a planar grain-boundary region, showing an amor-phous phase at a two-grain boundary. The incidentbeam is along the �012� and �001� directions of thetwo neighboring b-Si3N4 grains, respectively. The (100)and (110) planes are found to be the dominant grainboundary of the b-Si3N4 grains, due to their relatively

Fig. 4. SEM fractograph of the material subjected to high-strain-rateshear. The Si3N4 grains show high aspect ratios and form an inter-locking structure.

low energy when compared with other crystallineplanes of silicon nitride. The amorphous phase appearsto form along a straight grain boundary and shows aslightly light contrast with respect to the b-Si3N4

grains. The thickness of the boundary phase is approx-imately 1.8 nm, which is typical for amorphous grain-boundary phases in silicon nitrides and is in fairly goodagreement with the equilibrium intergranular distance(on the order of 1 nm) between two straight Si3N4 grainboundaries [34].

Fig. 8(a) is a high-resolution TEM image showing acrystalline boundary phase present along a straightgrain boundary between two b-Si3N4 grains, at an areaapproximately 0.1 mm from the grain pocket. The d-spacings of the lattice fringes of the two b-Si3N4 grainsare 0.66 and 0.27 nm, respectively, corresponding to the(100) and (101) fringes of b-Si3N4. The grain-boundaryphase shown in Fig. 8(a) exhibits sharp lattice fringes,which are significantly different from those of theneighboring Si3N4 grains regarding d-spacing and ori-entation, indicating that a crystalline phase has beenformed at the two-grain boundary. The fringes of thecrystalline boundary phase orient approximately 45°with respect to the b-Si3N4 (100) planes and approxi-mately 80° with respect to the b-Si3N4 (101) planes.There is a well-defined boundary between the crys-talline grain-boundary phase and the neighboring Si3N4

grains. The fringe spacing of the crystalline grain-boundary phase was determined to be 0.24 nm, whichdoes not match any d-spacing found in silicon nitride(b-Si3N4 or a-Si3N4). The thickness of the crystallineboundary phase is approximately 1.6 nm, which isslightly smaller than the thickness of the amorphousgrain-boundary phase (1.8 nm) shown in Fig. 7, but stillmatches well with the equilibrium intergranular dis-tance in Si3N4.

Fig. 3. SEM micrographs of the polished surfaces of the specimen. (a)The unetched surface. The boundary phase (white) is extensively seenat the grain pockets. (b) The etched surface. The boundary phase hasbeen etched away, leaving acicular Si3N4 grains.


A crystalline grain-boundary phase with a thicknessother than 1.6 nm is also identified. Fig. 8(b) is ahigh-resolution TEM image from another boundaryregion of the specimen, showing a crystalline phasepresent along a straight two-grain boundary, at an areanear the grain pocket. Compared with the crystallinephase shown in Fig. 8(a), the crystalline boundaryphase in Fig. 8(b) shows a similar fringe orientationwith respect to the neighboring b-Si3N4 grains, but aslightly larger fringe spacing (0.26 nm). The thicknessof the crystalline boundary phase in Fig. 8(b) wasdetermined to be approximately 1.9 nm, which also isin compliance with the equilibrium intergranular spac-ing. Kleebe et al. [35] have noted that the intergranularspacing in silicon nitride is independent of the orienta-tion of the neighboring Si3N4 grains, but varies with thecomposition of the boundary phases. If the boundary-phase composition is indeed the dominant factor con-trolling the film thickness of a two-grain boundary, thehigh-resolution TEM images shown in Fig. 8(a) and (b)suggest that there could be more than one crystalline

Fig. 6. TEM micrograph showing a straight two-grain boundarybetween two grain pockets.

grain-boundary phase formed in the in-situ reinforcedsilicon nitride.

During high-resolution TEM examination, it wasnoted that the crystalline boundary phase is often seenat two-grain boundaries near the multi-grain pockets,whereas the boundary phase located at the middleportion of a two-grain boundary is often, if not always,an amorphous phase. The observation suggests that thecrystallization of the boundary phase most likely ini-tiates at the grain pockets where the majority of thesintering aids are accumulated. Therefore, a detailedinvestigation was focused on these areas in order toverify the presence of crystalline boundary phase in thematerial.

Fig. 9 is a high-resolution TEM image of a three-grain pocket, showing evidence of a crystallineboundary phase present in the pocket. The latticefringes of a nearby grain shown in this micrograph are

Fig. 5. Low-magnification TEM micrographs of the specimen. MostSi3N4 grains are relatively ‘clean’. No second phase particles areobserved.

Fig. 7. High-resolution TEM micrograph from a two-grain boundaryarea. The b-Si3N4 grains are separated by an amorphous phase witha thickness of approximately 1.8 nm.


Fig. 8. High-resolution TEM images revealing the presence of crys-talline grain-boundary phase at the two-grain boundaries. The d-spacings of the crystalline boundary phase are determined to be 0.24nm in (a) and 0.26 nm in (b). The thickness of the boundary phase isapproximately 1.6 nm in (a) and is about 1.9 nm in (b). It is notedthat the lattice fringes of the boundary phase always orient approxi-mately 45° with respect to the (100) planes of the neighboringb-Si3N4 grains.

presence of a thin glassy layer at the grain pockets,between crystalline RE2Si2O7 boundary phases andSi3N4 grains. The formation of the thin glassy interfa-cial layer is attributed to the impurities present in thestarting Si3N4 powders and sintering additives.

The structures of the crystalline boundary phases arealso identified by using an SAED technique. The cam-era length of the microscope was carefully calibratedfor each period of operation, by using the standarddiffraction patterns obtained from the central regionsof the b-Si3N4 grains. Due to the small intergranulararea, the amount of the crystalline phase segregated ata two-grain boundary is not sufficient to produce atwo-dimensional electron diffraction pattern for struc-tural analysis. Most SAED patterns obtained at thetwo-grain boundaries only show one-dimensional reflec-tion spots, which are not enough to positively identify agiven crystal structure because the d-spacing deter-mined from these spots could match several possiblecrystalline structures. This is illustrated in Fig. 10,which is an SAED pattern from a two-grain boundarywhere a crystalline boundary phase is identified byhigh-resolution imaging. It shows a strong �012� stan-dard diffraction pattern from one b-Si3N4 grain and aline of sharp (100) spots from the neighboring b-Si3N4

grain. In addition, some diffraction spots with signifi-cantly lower intensity (arrows) are observed around theorigin. The d-spacing of these spots was determined tobe 0.548 nm, which does not belong to Si3N4. Althoughthis d-spacing matches the (011( ) spacing of La5Si3O12N,it also matches several other possible compounds.Therefore, a detailed SAED study is focused on thetwo-dimensional diffraction patterns obtained from the

Fig. 9. A crystalline boundary phase present at a grain pocket,showing the two-dimensional lattice fringes. The fringe spacings ofthe boundary phase are 0.32 and 0.33 nm, respectively, correspondingto the (210) and (11( 2) d-spacings of La5Si3O12N. A residual glassyphase with a thickness of few molecular layers is found at theinterface between the crystalline boundary phase and the b-Si3N4

grain.

b-Si3N4 (100) fringes which are parallel to the grainboundary, as often seen in silicon nitride. Theboundary phase shows two sets of lattice fringes, ap-proximately perpendicular to each other. The fringespacings of the boundary phase were determined to be0.32 nm (horizontal fringes) and 0.33 nm (verticalfringes), respectively. The d-spacings and the orienta-tions of the boundary phase fringes do not match thoseof either b-Si3N4 or a-Si3N4. Instead, they are in goodagreement with the (210) and (11( 2) lattice planes (d1=0.317 nm; d2=0.334 nm; and the inter-plane angle=86°) of silicon lanthanum oxynitride (La5Si3O12N). It isinteresting to note that a residual glassy phase, which isa few molecular layers thick, is formed at the interfacebetween the crystalline boundary phase and the b-Si3N4

grain. Thomas and cowokers [21,23] also noted the


Fig. 10. SAED pattern from a two-grain boundary area. It consists ofa �012� standard diffraction pattern from a b-Si3N4 grain, a line ofsharp (100) spots from the neighboring b-Si3N4 grain, and some lowintensity spots (arrows) from the crystalline boundary phase.

also suggest that the crystalline boundary phases at thegrain pocket are single crystals, formed by large-scalegrowth of a single nucleus. This is consistent with thecrystallization behavior of the RE2Si2O7 boundaryphases observed at the grain pockets [22]. Based on thetwo crystalline structures identified at the grain pockets,the lattice fringes of the crystalline phases at the two-grain boundaries as shown in Fig. 8(a) and (b) are mostlikely the (103) fringes of La5Si3O12N and the (21( 2)fringes of Y5Si3O12N, respectively. Based on X-ray dif-fraction spectrum, high-resolution TEM images, andSAED patterns, it is therefore concluded that crys-talline grain-boundary phases of La5Si3O12N andY5Si3O12N have been formed in the in-situ reinforcedsilicon nitride examined in the present study.

Even though SrO is one of the sintering additivesused in preparing the subject material, no evidence of acrystalline silicon strontium oxynitride (Sr5Si3O12N)phase is found at the grain boundaries of AS800 silicon

relatively large grain pocket areas where a crystallineboundary phase has been identified by high-resolutionTEM imaging.

Fig. 11(a) is an SAED pattern from a grain pocketsurrounded by three b-Si3N4 grains. Although the two-dimensional diffraction spots from this area show arelatively lower intensity when compared with thosefrom the b-Si3N4 grains, they are relatively sharp,indicating that the material at the grain pocket has beenwell crystallized. The diffraction pattern shown in Fig.11(a) consists of two sets of spots, perpendicular toeach other. The spots along the horizontal directionhave a d-spacing of 0.838 nm, while the spots along thevertical direction have a d-spacing of 0.402 nm. Theycorrespond to the reflections from the (100) and (1( 21( )planes of La5Si3O12N. The direction of the incidentbeam is thus determined to be the �012� direction ofLa5Si3O12N. Based on an analysis of the SAED pat-terns from the specimen, La5Si3O12N is found to be thedominant crystalline boundary phase at the grain pock-ets, presumably due to the high content of La2O3 in thestarting powders. Another crystalline boundary phaseidentified in the material is silicon yttrium oxynitride(Y5Si3O12N), as shown in Fig. 11(b). The SAED pat-tern shown in Fig. 11(b) is from a relatively large grainpocket surrounded by four b-Si3N4 grains. Two d-spac-ings are identified from this two-dimensional diffractionpattern. One is 0.474 nm (spots along the horizontaldirection) and the other is 0.390 nm. The inter-spotangle is approximately 66°. The d-spacings and spotorientation determined from Fig. 11(b) matches wellwith the �123� standard pattern from Y5Si3O12N, whilethe three spots nearest to the origin are reflections fromthe (21( 0), (1( 21( ), and (111( ) crystal planes of Y5Si3O12N.The diffraction patterns shown in Fig. 11(a) and (b)

Fig. 11. SAED patterns from the crystallized material at the grainpockets. (a) �012� diffraction pattern of La5Si3O12N. (b) �123�diffraction pattern of Y5Si3O12N.


Fig. 12. High-resolution TEM images of the amorphous phase at thegrain pockets. It is noted that the curved b-Si3N4 grain boundary hasa stronger bonding with the amorphous phase than the straightboundaries.

phous phase at the grain pockets is unknown atpresent. But the contrast of the amorphous boundaryphase does vary from one location to another, suggest-ing a compositional inhomogeneity of the amorphousphase at the grain boundaries.

4. Discussion

Both crystalline boundary phases identified in AS800silicon nitride, La5Si3O12N and Y5Si3O12N, belong to arare-earth apatite family with a hexagonal crystal struc-ture. Their lattice constants are also very similar. ForLa5Si3O12N, ao is 0.9684 nm and co is 0.7275 nm. ForY5Si3O12N, the corresponding parameters are 0.9436and 0.6822 nm, respectively. The most widely studiedapatite structure is fluorapatite, Ca5P3O12F. Previousstudies have shown that substitution of rare-earthcations into the fluorapatite structure would not alterthe symmetry of the unit cell but would result in aslight change of the lattice parameters. For rare-earthapatites, the lattice parameters usually increase withincreasing ‘ionic’ radii of the cation [30]. Fig. 13 is aschematic packing drawing of the rare-earth apatitestructure viewed along the co axis [36]. In this structure,an oxygen site in the SiO4 tetrahedron is replaced by anitrogen atom and each oxygen or nitrogen atom issurrounded by six lanthanum (or yttrium) ions. Electri-cal neutrality is maintained by producing a vacancy inevery second Si–O–N tetrahedron. The formation ofcrystallization phases instead of amorphous phases atsome grain boundary areas is considered as one of theimportant reasons resulting in the superior strength ofAS800 silicon nitride at elevated temperatures.Y5Si3O12N, for example, is found to be stable up toabout 1750°C [26], which is significantly higher than themelting temperatures of most glassy phases identified insilicon nitride.

nitride. This is probably due to the fact that the contentof SrO is the lowest among the three sintering additivesused in the material. Therefore, the crystallineSr5Si3O12N boundary phase, if formed, would be hardlydetectable. It is also possible that strontium is onlypresent in the amorphous phase.

Our observations have shown that the second phasespresent in the two-grain boundary regions could beeither crystalline phases or amorphous phases. Thesame is true for the grain pockets. It is noted in thepresent study that not all the grain pockets examinedare filled with crystalline boundary phases. Instead,some of them are still filled with the amorphous phase.The high-resolution TEM images of the amorphousphase at the grain pockets indicate that the interfacebetween the b-Si3N4 grain and the amorphous phaseusually is well-defined along the curved boundaries(Fig. 12(a)) but is less well-defined along the straightboundaries (Fig. 12(b)). The composition of the amor-

Fig. 13. Schematic packing drawing of the rare-earth apatite structureviewed along the co axis. In order of decreasing size, the spheres areO or N, La or Y, and Si.


Fig. 14. High-resolution TEM images show significant variation ofthe amorphous film thickness at the two-grain boundaries. Theintergranular spacing is approximately 0.7 nm in (a) and is about 3.0nm in (b).

analysis of the interfacial energies and the force balancenormal to the grain boundary, the presence of anequilibrium intergranular film thickness is the result oftwo competing interaction forces: (1) an attractive vander Waals force between the grains on either side of theboundary which tends to thin the film, and (2) arepulsive force, due to the structure of the boundaryphase, against the van der Waals force. An equilibriumfilm thickness on the order of 1 nm is proposed basedupon the short-range balance of these two interactionforces. Furthermore, it is concluded that for a givenceramic material, the final film thickness is ultimatelydependent on the composition of the starting liquid.This is confirmed by the different film thicknesses of theLa5Si3O12N and Y5Si3O12N crystalline boundary phases(Fig. 8(a) and (b)) observed in the present study. There-fore, the wide fluctuation of the intergranular amor-phous film thickness observed in AS800 silicon nitrideindicates that there exist significant compositional inho-mogeneities at the boundaries of b-Si3N4 grains. This ismost likely due to the limited viscosity of the liquidphase in these regions, as well as the presence ofimpurities from the starting powders. Based upon thehigh-resolution images obtained in the present study, itis thus suggested that there may be several amorphousboundary phases with different compositions present atthe two-grain boundaries of AS800 silicon nitride.

During the processing of AS800 silicon nitride, at-tempts have been made to eliminate the glassy phase atthe grain boundaries, such as adding rare-earth oxidessuitable to form crystalline phases, using gas pressuresintering which requires less liquid phase, and perform-ing post-sintering heat treatments to assist crystalliza-tion at the grain boundaries. However, the resultsobtained in the present study still show the presence ofan amorphous phase at the grain boundaries and thegrain pockets (Figs. 7, 12 and 14), suggesting thatcomplete recrystallization of boundary phases in siliconnitride is hard to achieve. Such a limited recrystalliza-tion at the grain boundaries of ceramics is attributed totwo factors. The first one is the difficulty for thecrystalline phase to nucleate at the grain boundaries.Bernard-Granger et al. [37] have found that nucleationof the boundary phase in silicon nitride takes place insome discrete places at the interface between theboundary phase and Si3N4 grains. Once created, nucleigrow into the boundary phase by displacement of well-faceted crystallization fronts. Therefore, nucleation isheterogeneous and is not an easy step in the crystalliza-tion process. For the crystallization to initiate, thenucleation needs, at the same time, local compositionfluctuations, with respect to the average composition, inorder to reach locally the composition of a crystallinephase, and probably a propitious orientation with re-spect to the neighboring silicon nitride grains to favoran ‘epitaxial’ nucleation and growth.

The thickness of the crystalline phases at the two-grain boundaries is relatively stable, varying from 1.5 to1.9 nm, with an average value of approximately 1.7 nm.The thickness of the amorphous phases at the two-grainboundaries, on the other hand, shows variation over awide range. The high resolution TEM images shown inFig. 14(a) and (b) illustrate two extreme cases observedin the present study. In Fig. 14(a), the thickness of theamorphous film is approximately equal to the (100)d-spacing of b-Si3N4, i.e. about 0.7 nm. In Fig. 14(b),on the other hand, the amorphous phase shows athickness greater than 3.0 nm. The variation of theamorphous phase thickness observed in AS800 siliconnitride suggests that the amorphous boundary-phasechemistry may vary from one area to another. In thepioneering work on ceramic intergranular spacing,Clark [34] indicates that thin intergranular films shouldmaintain an equilibrium thickness at the two-grainboundary of polycrystalline ceramics. According to an


The second barrier to the crystallization of theboundary phase in ceramics is due to the increase of thestrain energy during phase transformation. As sug-gested by Raj and Lange [38], complete crystallizationof small amounts of the glassy phase segregated at thegrain boundaries is much more difficult than crystalliza-tion of a bulk material of the same composition. This isbecause the glassy phase segregated at the grainboundaries is constrained from flowing over distanceslarger than the grain size, since the channels throughwhich the liquid must move are quite small, usually onthe order of a few nanometers, as can be seen from thehigh resolution TEM images shown in Figs. 7, 8, 12(b)and Fig. 14. The initial glassy phases resulting from thesintering aids, therefore, are essentially confined withinrigid ceramic crystals. As a result, the glassy phase hasto sustain hydrostatic stress (shear stresses are unlikelybecause they will be relaxed by the fluid flow within theintergranular region). Any change in the volume uponcrystallization will set up hydrostatic stresses so that thestrain energy associated with it must be accounted forin the crystallization process. In general, volume changeduring the phase transformation from a glassy to acrystalline phase in a confined region will give rise tostrain energy which would oppose the transformation.The strain energy introduced depends upon the differ-ence in the specific volumes of the crystalline and theglassy phases, the bulk moduli of the two phases, andthe size of the crystalline phase formed. Stress andstrain energy calculations performed by Raj and Langehave shown that the strain energy becomes quite largewhen the crystalline phase grows to a size comparableto the glass pocket. Therefore, there are always condi-tions when the glassy phase may crystallize only par-tially, or not crystallize at all. In most instances, theelastic constraint on the intergranular phase will lead tolimited crystallization of the grain-boundary phase.

5. Conclusions

The following conclusions can be drawn from thepresent study on the microstructure of the in-situ rein-forced silicon nitride: (1) Crystalline rare-earth apatiteboundary phases (La5Si3O12N and Y5Si3O12N) havebeen identified at the grain pockets, and along thetwo-grain boundaries at areas near the pockets. (2) Thethickness of the crystalline phases at the two-grainboundaries is approximately 1.7 nm, confirming thatthere is an equilibrium film thickness for the secondphase at the two-grain boundary. (3) The glassy phaseis also observed at the grain pockets and at the two-grain boundaries, indicating that the crystallization ofthe boundary phases is incomplete. (4) The thickness ofthe amorphous phase at the two-grain boundariesvaries from 0.7 to 3.0 nm, suggesting that significant

compositional inhomogeneities of the boundary phasesexist in these areas. (6) The superior high-temperaturestrength of the material is believed to result from thepartial crystallization of the boundary phases.

Acknowledgements

This work was supported by the Army ResearchOffice (ARO) contract DAAL 04-95-1-0609 andTuskegee (ARO) contract DAAH04-95-1-0369 to theUniversity of California, San Diego. This researchmade use of the University of California, San Diego’sX-ray/SEM facilities at the Scripps Institute ofOceanography’s Analytical Facility. The authors wouldlike to thank Mr John Pollinger at AlliedSignal forproviding the sample material and the National Centerfor Electron Microscopy (NCEM), Lawrence BerkeleyLaboratory, University of California, Berkeley, for theuse of the JEOL ARM-1000 high-resolution electronmicroscope. The authors would also like to thank thereviewer of Materials Science & Engineering A forhis/her comments and suggestions, which have beenpartially integrated into the final version of thismanuscript.

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