Formation of an a-SiAlON Layer on b-SiAlON and Its Effect onMechanical Properties

Xin Jiang,†,‡ Yong-Kee Baek,§ Sung-Min Lee,† and Suk-Joong L. Kang*,†,¶

Center for Interface Science and Engineering of Materials (CISEM), Korea Advanced Institute of Science andTechnology (KAIST), Taejon 305-701, Korea

Agency for Defense Development, Taejon 305-600, Korea

We have developed a technique forin situ formation of ana-SiAlON layer on sintered b-SiAlON. The technique con-sists of packing a powder compact ofb-SiAlON composi-tion with a-SiAlON composition powder, presintering, andsintering. Using this technique, it is possible to control thethickness of the layer by changing the presintering condi-tions during heating to a sintering temperature. The intro-duction of an a-SiAlON layer increases the surface hard-ness and improves the wear and oxidation resistance, witha moderate reduction of flexural strength. The surfacemodification of b-SiAlON may provide opportunities ofwider application of b-SiAlON to wear- and oxidation-resistant components.

I. Introduction

SIALON ceramics are the solid solutions ofa-Si3N4 or b-Si3N4 with oxygen, aluminum, and some other metal ions

(M). The general formulae are MpSi12−(m+n)Alm+nOnN16−n fora-SiAlON and Si6−zAl zOzN8−z for b-SiAlON, where the sub-scriptsp, m, n,andz are variables within their respective solu-bility range.1 In general,a-SiAlON ceramics are characterizedby their equiaxed grains with high hardness and good wear andoxidation resistance, whereasb-SiAlON ceramics are charac-terized by their elongated grains with good flexural strengthand fracture toughness.2 To combine advantages in the physicalproperties of both SiAlONs,a/b SiAlON composites have longbeen studied.3–6 The microstructures of the SiAlON compos-ites generally consist ofa- and b-SiAlON grains uniformlydistributed in materials, and their properties exhibit intermedi-ate values, between those ofa- andb-SiAlON. For example, incomparison to pureb-SiAlON, the a/b-SiAlON compositeshows improved hardness and oxidation resistance with a re-duction of flexural strength and fracture toughness. Theseproperties vary almost linearly with the formation of addedSiAlON.2

Surface modification is another technique to improve theproperties of ceramics.7–9Surface modification sometimes pro-vides opportunities to introduce surface compressive stressesthat result from the difference in the thermal expansion coef-ficients between a surface layer and the bulk. Some previous

results have shown that the surface nitridation of dense SiCcould improve surface-sensitive mechanical properties becauseof the thermally induced compressive stresses in the Si3N4layer.10–12 So far, however, the introduction of compressivestresses at the surface of SiAlON materials has rarely beenreported. Instead, the deposition of a second phase at the Si3N4surface was proposed to improve oxidation resistance.13,14Ion-assisted deposition of a platinum layer on dense Si3N4 im-proved high-temperature stability in oxidizing atmospheres.14

However, such a deposition process may have limitations, be-cause of high costs and/or the weak bonding strength of thecoatings.

In the present investigation, we have developed a processthat allows structural modification ofb-SiAlON at the surfaceto form ana-SiAlON layer. The surface-modifiedb-SiAlON isexpected to exhibit higher hardness at the specimen surface andbetter wear and oxidation resistance thanb-SiAlON while re-taining the bulk toughness ofb-SiAlON.

II. Experimental Procedure

The starting materials were Si3N4 (SN-E10, Ube IndustriesCo., Tokyo, Japan), Al2O3 (AKP-50, Sumitomo Chemicals,Tokyo, Japan), AlN (C-Grade, H. C. Stark, Berlin, Germany),and Y2O3 (Shin-etsu Chemical Co., Tokyo, Japan). Three typesof b-SiAlON specimens were prepared: 91.5Si3N4–2.4AlN–6.1Al2O3, 83.1Si3N4–4.9AlN–12.0Al2O3, and 74.6Si3N4–7.3AlN–18.1Al2O3 (by weight percent), withz 4 0.5, 1, and1.5, respectively, in Si6−zAl zOzN8−z. To enhance the densifica-tion of the specimens, an extra 5 wt% of Y2O3 was added toeach specimen. Proportioned powders were attrition-milled inethyl alcohol with ZrO2 balls for 1 h. After drying and sieving,the powders ofb-SiAlON composition were compacted intopellets of 4 mm × 16 mm × 16 mm and 4 mm × 5 mm × 40 mm(by pressing uniaxially with very low pressure) and then iso-statically pressed under 200 MPa.

For the sintering ofb-SiAlON compacts, two types of pack-ing powders were used: BN (Bornitrid-s, ElektroschmelzwerkKempten?GmbH, Munchen, Germany) anda-SiAlON powder(composition of 73.6Si3N4–6.3Y2O3–6.9Al2O3–13.2AlN(wt%)), which corresponded to Y0.33Si9.3Al2.7O1.7N14.3. Thea-SiAlON composition powder was prepared by using attritionmilling in ethyl alcohol with ZrO2 balls for 1 h, as in the caseof b-SiAlON composition powder.

Depending on the packing powder, two types of thermalcycles were adopted for the sintering ofb-SiAlON. In the caseof BN powder packing, the specimen was heated to 1750°C ata rate of 20 K/min and held at the temperature for 1 h in 1 atmof nitrogen gas. In the case ofa-SiAlON composition powderpacking, the specimen was isothermally presintered at 1600°Cfor ∼1 min (hereafter denoted as 0 h), 3 h, or 6 h and thensintered at 1750°C for 1 h in 1 atm ofnitrogen gas. Throughseveral preliminary experiments, the presintering was found tobe essential to form ana-SiAlON layer with an appreciablethickness on ab-SiAlON body.

I-W. Chen—contributing editor

Manuscript No. 191068. Received April 18, 1997; approved October 18, 1997.Supported by the Korea Science and Engineering Foundation, through the Center

for Interface Science and Engineering of Materials, and also by the Agency forDefense Development.

*Member, American Ceramic Society.†Center for Interface Science and Engineering of Materials.‡On leave from the Shanghai Institute of Ceramics, Chinese Academy of Sciences,

Shanghai 200050, China. Supported by CISEM, KAIST, and Science and TechnologyPolicy Institute (STEPI) during his research stay in Korea.

§Agency for Defense Development.¶Author to whom correspondence should be addressed.

J. Am. Ceram. Soc., 81 [7] 1907–12 (1998)Journal

1907

The crystalline phases present in the specimens were iden-tified via X-ray diffractometry (XRD) using CuKa. The rela-tive amount ofa-SiAlON, a/(a+b), was calculated by the for-mula proposed by Gazzara and Messier:15

a~%! =I102~a! + I210~a!

I102~a! + I210~a! + I101~b! + I210~b!

2 100 (1)

where I102(a) is the diffraction intensity of the (102) plane ofa-SiAlON. The microstructures of sintered specimens wereobserved via scanning electron microscopy (SEM) after etch-ing the specimens in molten NaOH at 400°C for∼1 min. Thecation composition of the specimens on their etched cross-sectional plane was determined via energy-dispersive X-rayspectroscopy (EDX).

The sintered rectangular specimens were ground to 2.5 mm ×3.5 mm × 25 mm for the four-point bend test. During thesample preparation, however, the tensile surface was main-tained as-sintered and just slightly polished with diamond pasteof 1 mm size, to keep the structure of the surface layer. Somesintered square specimens with dimensions of∼3 mm ×12 mm × 12 mm were slightly polished for the hardness, tough-ness, and oxidation tests. Other square specimens were groundand polished to disks 6.3 mm in diameter and 1.6 mm thickfor the wear test. The wear surface was also maintainedas-sintered.

The flexural strength was measured by using the four-pointbend test with an outer span of 12.5 mm and an inner span of

6.25 mm. Four or five samples were tested for each condition.The hardness and toughness were measured by the indentionmethod with a Vickers diamond indentor and a load of 2 kg.After the indentation test, the fracture toughnessKc was ob-tained from the equation16

Kc = 0.018S E

HvD1/2S P

c3/2D (2)

whereE is the Young’s modulus,Hv the Vickers indentationhardness,P the indentation load, andc the half-crack length.The value ofE was assumed to be 300 GPa for the calculation.The sliding-wear test was conducted using a ball-on-three-flatconfiguration.17 An alumina ball 12.7 mm in diameter wasrotated on three disk samples that were aligned with their sur-face normals in tetrahedral coordination relative to the rotationaxis of the ball. The rotation speed of the ball was 100 rpm. Aweight of 70 kg was loaded on the alumina ball to apply thenormal force of 280 N on each disk. The entire specimenassembly was immersed in paraffin oil during the test.

The oxidation experiment was performed in an alumina tubefurnace at temperatures of 1200°, 1300°, and 1350°C in air.The test samples were placed on two platinum wires that weresupported by an alumina plate, to prevent contact between thealumina plate and the specimens during the oxidation treat-ment, and heated up to the oxidation temperatures at a heatingrate of 45 K/min. The weight change of the specimens wasmeasured by a microbalance with a resolution of 10−4 g.

Fig. 1. SEM micrographs of (a)b-SiAlON sintered at 1750°C for 1 h in BN packing powder and (b and c)b-SiAlON presintered at 1600°C for3 h and then sintered at 1750°C for 1 h in a-SiAlON packing powder ((b) center region and (c) surface).

1908 Journal of the American Ceramic Society—Jiang et al. Vol. 81, No. 7

III. Results and Discussion

(1) Formation of a-SiAlON on b-SiAlONDuring the sintering at 1750°C for 1 h, the powder compacts

were densified to∼99% of theoretical density. Figure 1 showsthe SEM micrographs of the sinteredb-SiAlON specimens(z 4 1) with BN powder (Fig. 1(a)) and witha-SiAlON com-position powder (Figs. 1(b and c)) packing. The microstruc-tures consist of the typical SiAlON grains and a glassy matrix.However, the shape of the grains seems to be somewhat dif-ferent between the specimens. In particular, the shape of thegrains at the surface of theb-SiAlON specimen packed witha-SiAlON powder and presintered at 1600°C for 3 h ismostlyequiaxed, implying that the grains area-SiAlON. However, thegrain shape in the interior of the same specimen (Fig. 1(b)) ismore elongated, similar to that in the specimen packed with BN(Fig. 1(a)).

The XRD patterns in Fig. 2 confirm that the grains at thesurface of the specimen packed witha-SiAlON powder aremostly a-SiAlON, whereas those in the bulk of the specimenareb-SiAlON. Therefore, ana-SiAlON layer seems to form atthe surface of theb-SiAlON specimen when using thea-SiAlON packing powder.

The thickness of thea-SiAlON layer increased as the presin-tering time at 1600°C increased, as shown in Fig. 3. However,the layer thickness did not increase as the sintering time at1750°C increased. In addition, when a fully denseb-SiAlONspecimen was heat-treated with thea-SiAlON packing powder,no detectablea-SiAlON phase formed at the surface, even after9 h at a temperature of 1750°C. These results indicate that the

formation of ana-SiAlON layer does not result from a directchemical reaction between theb-SiAlON specimen and thea-SiAlON packing powder but rather from a mass transferfrom the packing powder to the compact around the presinter-ing temperature.

Figure 4 plots the variations of cation composition with thedistance from the surface on a cross-sectional plane of thespecimen presintered at 1600°C for 3 h and then sintered at1750°C for 1 h. Yttrium and aluminum contents are higher atthe specimen surface in thea-SiAlON layer than inside thespecimen. Because the yttrium content in a liquid formed in thepacking powder is higher than that in theb-SiAlON compact,this result confirms again the transfer of the cations from thepacking powder into the layer. The mass transfer may occur vialiquid flow and/or via gas-phase transport when a liquid phaseforms in the packing powder. Because the size of pore capil-laries in the compact must be much smaller than that in thepacking powder, a liquid formed in the packing powder isexpected to penetrate into theb-SiAlON compact. The pen-etration depth of the liquid should then be determined bypresintering time, as observed in Fig. 3. However, the liquidpenetration may be very limited, because of the high viscosityof the liquid and its reaction with solid particles in the compact.Vapor-phase transport may also be operative, if the vapor pres-sure of the liquid is high enough to cause the mass transfer.However, gas-phase transport would be insignificant, even un-der an appreciable vapor pressure, because the transport shouldoccur in both directions—from the packing powder to the com-pact and vice versa.

Because the chemical composition of the material trans-ported from the packing powder is in equilibrium witha-SiAlON, a-SiAlON grains form during presintering and sin-tering, which results in the formation of ana-SiAlON layer atthe b-SiAlON surface, as shown in Fig. 1(c). This result issimilar to that observed previously, where the transformationof a-SiAlON to b-SiAlON occurred via the penetration of aliquid whose composition is in equilibrium withb-SiAlON.18

When b-SiAlON compacts withz values of 0.5 and 1.5packed with thea-SiAlON packing powder were presinteredand sintered,a+b anda+12H surface layers were respectivelyformed. According to the phase equilibria of the Si3N4–Al2O3–Y2O3 system,19 many subsolidus compatibility tetrahedra areavailable. During the material transport from the packing pow-der, the location of the final composition of theb-SiAlONsurface is determined by the chemical composition of packingpowder and by theb-SiAlON composition itself. The forma-tion of 12H SiAlON at the surface ofb-SiAlON with a z valueof 1.5 again confirms that 12H is in equilibrium witha-SiAlON at an alumina-rich terminal.20

Fig. 2. XRD patterns of (a) the surface and (b) the bulk ofb-SiAlONpresintered at 1600°C for 3 h and then sintered at 1750°C for 1 h ina-SiAlON packing powder.

Fig. 3. Measured fraction ofa-SiAlON versus distance from surfaceof the specimens presintered at 1600°C for 0, 3, and 6 h, and thensintered at 1750°C for 1 h in a-SiAlON packing powder.

July 1998 Formation of ana-SiAlON Layer onb-SiAlON and Its Effect on Mechanical Properties 1909

(2) Mechanical PropertiesFigure 5 plots the measured hardness and fracture toughness

with distance from the surface of the modifiedb-SiAlONspecimen presintered at 1600°C for 3 h. Because thea-SiAlONlayer was thin (<300mm), a light indentation load of 2 kg wasused for the measurement. The hardness increases graduallyfrom the bulk toward the surface, whereas the fracture tough-ness decreases. This result occurs, of course, from the increasein a-SiAlON content at the surface. Even though the microin-dentation fracture toughness at the surface region is lower thanthat of the bulk, the bulk fracture toughness of the material isbelieved to be the same as that ofb-SiAlON.

To evaluate the effect of ana-SiAlON layer on the flexuralstrength ofb-SiAlON, a-SiAlON specimens were also pre-pared by sintering compacts of thea-SiAlON packing powder(73.6Si3N4–6.3Y2O3–6.9Al2O3–13.2AlN (wt%)) at 1750°C for1 h in 1 atm of nitrogen gas. The measured average flexuralstrengths ofa-SiAlON, b-SiAlON, and modifiedb-SiAlON(presintered at 1600°C for 3 h) with ana-SiAlON layer were316, 622, and 476 MPa, respectively, with the maximum scat-

ter of the data of (+28, −18), (+100, −53), and (+59, −84),respectively. The flexural strength of the modifiedb-SiAlONis lower than that ofb-SiAlON but is obviously higher thanthat of a-SiAlON. The reduction of the flexural strength ofb-SiAlON with the formation of ana-SiAlON layer may bemainly due to the introduction of tensile stresses in thea-SiAlON layer, because the thermal expansion coefficient ofa-SiAlON (aa-SiAlON 4 3.7 × 10−6/°C)21 is higher than that ofb-SiAlON (ab-SiAlON 4 3.1 × 10−6/°C).21 The stress (s) in-troduced at the surface of the layer may be estimated by usingANSYS 5.2 code22 and a well-known equation:23

s =Ec~ac − as!DT

~1 − nc!F2Shc

hsDSEc

EsD ~1 − ns!

~1 − nc!+ 1G−1

(3)

whereEc andEs are the Young’s moduli of the coated layer andthe substrate, respectively;a is the mean thermal expansioncoefficient,n the Poisson’s ratio,hc the thickness of the coatedlayer, hs the thickness of the substrate material, andDT thetemperature difference below the solidus of the SiAlON. In the

Fig. 5. Variation of (j) Vickers hardness and (d) indentation fracture toughness with distance from the surface of the specimen presintered at1600°C for 3 h and then sintered at 1750°C for 1 h in a-SiAlON packing powder.

Fig. 4. Variation of cation concentrations (silicon, aluminum, and yttrium) with distance from the surface of a specimen corresponding to that inFig. 3, presintered at 1600°C for 3 h and then sintered at 1750°C for 1 h in a-SiAlON packing powder.

1910 Journal of the American Ceramic Society—Jiang et al. Vol. 81, No. 7

calculation, the fraction ofa-SiAlON at the surface was takento be 100% and assumed to decrease linearly to 0% at 250mmfrom the surface, following the measureda-SiAlON fraction inFig. 3. Assuming values of 300 GPa for the Young’s modulusof SiAlON,24 0.25 for Poisson’s ratio,24 and 1300°C for thetemperature difference below the solidus of our SiAlON, themaximum tensile stresss is calculated to be∼220 MPa. Thisvalue is in good agreement with the reduction of flexuralstrength, when we take into account inaccuracies in the dataused and the relaxation of the stresses in the material. There-fore, the difference in the thermal expansion coefficients mustbe the major cause for the reduction of flexural strength.

Table I lists the measured wear scar diameter of theb-SiAlON and modifiedb-SiAlON (presintered at 1600°C for 3and 6 h) specimens. For a given sliding time, the wear scardiameters of the modifiedb-SiAlONs are smaller than that ofthe b-SiAlON, which indicates that the wear resistance ofb-SiAlON is much improved by introducing ana-SiAlON layerat its surface. The resistance also seems to increase as thethickness ofa-SiAlON layer increases for our specimens.Therefore, the wear resistance can be optimized for the specificpurpose of the application ofb-SiAlON.

(3) Oxidation BehaviorFigure 6 shows the measured weight gain per unit surface

area with annealing time at various temperatures in air. Thetested specimens werea-SiAlON, b-SiAlON, and modifiedb-SiAlON (presintered at 1600°C for 3 h). The weight gain ofthe surface-modified specimen is higher than that ofa-SiAlONbut is considerably lower than that ofb-SiAlON. Thus, thea-SiAlON layer onb-SiAlON seems to improve the oxidationresistance ofb-SiAlON.

Figure 7 shows an SEM micrograph of a surface-modifiedb-SiAlON specimen that has been oxidized at a temperature of1350°C for 10 h. At the surface of the specimen, an oxidized

layer is formed and parallel cracks are introduced between theoxide scale and the bulk matrix. The oxide scale with an un-even surface seems to contain small crystals and pores. AnXRD analysis identified the crystalline phases that are presentin the oxide scale as yttrium disilicate (Y2Si2O7) and a-cristobalite. An EDX analysis revealed that the white areaswith irregular shape are Y2Si2O7 and the small grains with darkcontrast area-cristobalite. The oxide scale formed at the sur-face of a-SiAlON or surface-modifiedb-SiAlON containsmore glass and less Y2Si2O7 than that formed at the surface ofb-SiAlON. The number of cracks formed in the oxidized layerof the surface-modifiedb-SiAlON are much less than those oftheb-SiAlON. This result may be attributed to the formation ofa smaller amount of Y2Si2O7 in the surface-modifiedb-SiAlON than inb-SiAlON. In fact, the formation of Y2Si2O7considerably increases the specific volume of oxidizedSiAlON.25

IV. Conclusions

When sintering powder compacts ofb-SiAlON compositionwithin a packing powder ofa-SiAlON composition, a densea-SiAlON layer formed at the surface of sinteredb-SiAlON. A

Table I. Measured Wear Scar Diameter ofb-SiAlON andSurface-Modified b-SiAlON Specimens

Specimen

Measured wear diameter (mm)

Slidingtime 4 80 min

Slidingtime 4 240 min

b-SiAlON 596 772Modified b-SiAlON

(presintered for 3 h) 524 694Modified b-SiAlON

(presintered for 6 h) 508 653

Fig. 6. Measured weight gain of various SiAlON specimens with oxidation time in air at 1200°, 1300°, and 1350°C; the surface-modifiedb-SiAlON was presintered at 1600°C for 3 h and then sintered at 1750°C for 1 h in a-SiAlON packing powder.

Fig. 7. SEM micrograph of the oxide scale formed in surface-modified b-SiAlON after annealing at 1350°C for 10 h in air.

July 1998 Formation of ana-SiAlON Layer onb-SiAlON and Its Effect on Mechanical Properties 1911

glass phase that formed in the packing powder seems to havepenetrated into theb-SiAlON powder compact during heatingto the sintering temperature and induced the formation ofa-SiAlON grains at the surface. The thickness of thea-SiAlONlayer that formed increased as the holding time increased at atemperature below the sintering temperature.

The formation of ana-SiAlON layer increased the surfacehardness and improved the wear and oxidation resistance ofb-SiAlON. Compared with these benefits, the reduction offlexural strength may be acceptable; the flexural strength of thesurface-modifiedb-SiAlON showed an intermediate value be-tween those ofb-SiAlON and a-SiAlON. Therefore, the de-veloped technique ofin situ modification of theb-SiAlONsurface toa-SiAlON may be applied to the preparation ofsintered components that have high demands on wear and oxi-dation resistance.

Acknowledgments: The authors are grateful to Dr. Seong-Jai Cho forthe wear test and to Ms. Won-Kyung Choi for the stress analysis.

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