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Materials Science and Engineering A 488 (2008) 475–481

Microstructure and properties of translucent Mg–sialonceramics prepared by spark plasma sintering

Y. Xiong, Z.Y. Fu ∗, H. Wang, Y.C. Wang, J.Y. Zhang, Q.J. ZhangState Key Lab of Advanced Technology for Materials Synthesis and Processing,Wuhan University of Technology, Wuhan 430070, People’s Republic of China

Received 6 August 2007; received in revised form 12 November 2007; accepted 14 November 2007

bstract

Translucent Mg–sialon ceramics were prepared using spark plasma sintering (SPS) �-Si3N4 powders with 9 wt% AlN and 3 wt% MgO asintering additives. Microstructural observations indicate that the optical and mechanical properties of Mg–sialon ceramics are affected by theensity and �′:�′-phase ratio in sintered bodies, which are tailored by controlling the content of formed liquid phase and optimizing the parametersf spark plasma sintering in present study. The material is toughened by the existence of a small amount of �′-sialon. The reason that the two-phase

omposite does not greatly compromise optical property could be attributed to the fine equiaxed microstructures and low content of �′-phase.ranslucent Mg–sialon ceramics achieving 66.4% of maximum transmittance, 21.4 ± 0.3 GPa hardness, and 6.1 ± 0.1 MPa m1/2 fracture toughnessere prepared by spark plasma sintering at 1850 ◦C for 5 min.2007 Elsevier B.V. All rights reserved.

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eywords: Translucent Mg–sialon ceramics; Spark plasma sintering; Microstru

. Introduction

Sialon ceramics are Si3N4-based solid solutions. The twoajor phases are �′ and �′, which are isostructural with �-

nd �-Si3N4, respectively. The �′-sialons that have the generalormula MxSi12−m−nAlm+nOxN16−n are produced by replacingm + n) (Si–N) bonds with m (Al–N) and n (Al–O) bonds in the-Si3N4 structure. Electrical neutrality in the solid solution isaintained by adding M cations to the interstitial space, where, the stabilizer, is one of the cations Li, Mg, Ca, Y and most

are-earth elements [1,2]. One of the most important advantagesf �′-sialon is that the stabilizing elements can be incorporatednto �′-sialon matrix, and the sintering of these materials occurshrough the formation of transient liquid phase which results ineramics with reduced amounts of residual glassy phase [2]. Fur-hermore, �′-sialon exhibits higher hardness and better thermal

hock resistance than �′-sialon.

Researches on sialon ceramics have been mostly concen-rated on the structural applications because of their excellent

∗ Corresponding author. Tel.: +86 27 87662983; fax: +86 27 87215421.E-mail addresses: xiongyan1980@hotmail.com (Y. Xiong),

yfu@uchit.edu.cn (Z.Y. Fu).

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921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2007.11.041

; Phase transformation

echanical properties. In recent years, there are several reports3–9] (as shown in Table 1) regarding the optical proper-ies of sialon ceramics, primarily concerning �′-sialon. Its considered that the application field of sialon ceramicsould be broadened, for example as optical window, if they

an be prepared with high transparency [8]. However, mostf transparent/translucent sialon ceramics were fabricated byot pressing with hour(s) soaking and doped by rare earthxides. Spark plasma sintering (SPS) is a newly developedintering technique, mostly known for its rapid densificationrocess, which enables us not only to shorten the densifi-ation process but also control the obtained microstructures.ecently, the authors have first reported the preparation of

ranslucent Mg–sialon ceramic using spark plasma sinteringf �-Si3N4 powder with AlN and MgO as sintering additives10]. In the present paper, microstructures and microstructuralffects on the properties of Mg–sialon ceramics are investi-ated.

. Experimental procedure

The mixture of �-Si3N4 powder (SE-10 grade, ≥95% �-hase, Ube Industries, Tokyo, Japan) 9 wt% AlN (F grade,okuyama Co., Soda, Japan), and 3 wt% MgO (Shanghai Chem-

476 Y. Xiong et al. / Materials Science and Engineering A 488 (2008) 475–481

Table 1Phase and properties of literature-reported translucent sialon ceramics

Stabilizing element Sintering conditions Phase Properties Ref.

Maximumtransmittance

(%) Hardness(GPa)

Fracture toughness(MPa m1/2)

– 1700–1800 ◦C sintered under apressure of 100–300 kg cm−2

�′-Sialon 40 (0.65-mm-thick) – – [3]

Yb Hot-pressed at 1750 ◦C for 4 h Yb-�′-sialon – – – [4]Nd, Ce, Yb Capsule-free HIP sintering – – – – [2]Y Hot-pressed at 1800 ◦C for 1 h

with a pressure of 40 MPaY-�′-sialon – 21.12 2.64 [5]

Y-�′/�′-sialon 35 (0.5-mm-thick) 19.46 4.97Lu Hot-pressed at 1950 ◦C for 2 h

with a pressure of 40 MPa in a0.9 MPa nitrogen atmosphere

Lu-�′-sialon >70 (0.5-mm-thick) 19.27 ± 0.76 2.55 ± 0.11 [6]

Dy Hot-pressed at 1900 ◦C for 1 hwith a pressure of 20 MPa

Dy-�′-sialon 65.2 (1.0-mm-thick) 20 5.1 [7]

Dy, Y SPS-sintered at 1700 ◦C with apressure of 50 MPa

(Dy, Y)-�′-sialon 64 (0.5-mm-thick) – – [8]

Gd Hot-pressed at 1800 ◦C for 1 hwith a pressure of 25 MPa

Gd-�′-sialon ∼65 (0.7-mm-thick) – – [9]

M 66.4 (0.5-mm-thick) 21.4 ± 0.3 6.1 ± 0.1 Our work

iawagamitpows

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Mg-�′/�′-sialon

cal, Shanghai, China) were ball-milled for 24 h using ethanols the mixing medium. After dried in vacuum, 3.5 g mixturesere poured into the graphite molds and sintered in a SPS

pparatus (Model 1050, Sumitomo Coal Mining Co. Ltd., Kana-awa, Japan) at a heating rate of 100 ◦C/min under a nitrogentmosphere of −0.03 MPa. The temperature during SPS waseasured by an infrared pyrometer focused on a small hole

n graphite mold. A pressure of 30 MPa was applied fromhe beginning of sintering. After 5 min soaking time, sam-les of 20 mm in diameter and 3–4 mm in thickness werebtained. For comparison, green sample of the same compositeas hot-pressed at 1900 ◦C for 60 min under 30 MPa pres-

ure.Relative densities of the sintered bodies were measured by

rchimedes method. Phase identification was conducted by X-ay diffractometry (XRD; Model D/MAX-RBX, Rigaku Co.,okyo, Japan) with Cu K� radiation and Si as an internal stan-ard at a speed of 0.5◦/min. The software UnitCell was used forhe calculation of lattice parameter. The x value of �′ phases

xSi12−(m+n)Al(m+n)OnN16−n was calculated from the meanalues of xa and xc obtained in the following equations [11]:

(nm) = 0.775 + 0.0156xa (1)

(nm) = 0.562 + 0.0162xc (2)

he z value of �′-sialon phases Si6−zAlzOzN8−z was determinedrom the mean value of za and zc obtained in the followingquations [12]:

(A ) = 7.603 + 0.0297za (3)

(A ) = 2.907 + 0.0225zc (4)

The volume fraction of �′ phase in the sintered bodiesas determined by the formula developed by Gazzara and

Fig. 1. XRD patterns of (a) mixture powder and (b) sintered translucentMg–sialon ceramics.

Y. Xiong et al. / Materials Science and Engineering A 488 (2008) 475–481 477

Fig. 2. (a) Scanning electronic micrograph and element distribution of (b) O, (c) Mg, (d) Al and (e) Si in translucent Mg–sialon ceramics.

Table 2Calculated lattice parameters of green � powder and 1850 ◦C SPS-sintered sample

Cell parameters �′-Phase �′-Phase

a (A) c (A) Cell volume (A3) a (A) c (A) Cell volume (A3)

Green powder 7.7532 5.6194 292.5403 – – –Sintered body 7.7901 5.6545 297.1739 7.6281 2.9233 147.3103

x = 0.24 z = 0.78

Table 3Selected properties of prepared translucence Mg–sialon ceramics

Sintering parameters Relativedensity (%)

�′-Phasecontent (vol%)

Maximumtransmittance (%)

Hardness (GPa) Fracture toughness(MPa m1/2)

1800 ◦C SPS-sintered for 5 min 98.3 90.6 55.2 19.6 ± 0.6 5.6 ± 0.31850 ◦C SPS-sintered for 5 min 99.5 86.6 66.4 21.4 ± 0.3 6.1 ± 0.11870 ◦C SPS-sintered for 5 min 99.5 83.7 60.1 19.9 ± 0.4 6.5 ± 0.21900 ◦C HP-sintered for 60 min 98.2 75.3 – – –

4 and Engineering A 488 (2008) 475–481

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78 Y. Xiong et al. / Materials Science

essier [13]. Microstructures were observed by scanning elec-ron microscopy (SEM; Model 5610, JEOL, Tokyo, Japan)nd transmission electron microscopy (TEM; Model 2010,EOL, Tokyo, Japan) attached with an energy dispersive X-ay microanalyzer (EDX, Oxford). The sintered bodies wereliced to 0.5 mm in thickness and polished for optical transmit-ance measurement by Fourier transform infrared spectroscopyFTIR; Model Nexus, ThermoNicolet Nexus, Waltham, USA) in.5–5.5 �m wavelength range. Vickers hardness was conductednder a load of 98 N for 15 s and calculated using the followingquation [14]:

v10 = 0.47P

a2 (5)

racture toughness (KIC) was determined at room temperaturey indentation-fracture under the same load and calculated con-ulting the equation proposed by Evans and Charles [15] andiiharra et al. [16] as following:

KICΦ

(Ha1/2)= 0.15k

( c

a

)−3/2(6)

here Φ (≈3) and k(= 3.2) are two dimensionless constants, Hs the hardness, a is the half-diagonal of the Vickers indent andis the radius of the surface crack.

. Results

.1. Phase identification

The molar percent of Si3N4, AlN and MgO in the greenixture are 68.1 mol%, 23.8 mol% and 8.1 mol%, respectively.ccording to the phase diagram of MgO–Si3N4–AlN systemeasured by Kuang et al. [17], single �′-sialon cannot be

btained from the composite of Si3N4–MgO: 3AlN and theesultants from the present green component should be locatedt the two-phase (�′ + �′) or three-phase region (�′ + �′ + AlN).he XRD patterns of green mixture powder and the sample sin-

ered at 1850 ◦C for 5 min are shown in Fig. 1. It is difficult tond the traces of AlN and MgO after sintering. The peaks of �′-ialon are rather strong compared with those of �′-sialon, whichndicates that the content of �′-sialon is rather low.

The calculated lattice parameters of green powder andhe sample SPS-sintered at 1850 ◦C for 5 min are listed inable 2. The enlargement in the lattice parameters of the sin-

ered body confirms that the sintered body consists mainlyf Mg0.24Si11.28Al0.72O0.24N15.76 with a small amount ofi5.22Al0.78O0.78N7.22. According to the formula of Mg2+-tabilized �′-sialon, the composite should correspond tog0.42Si10.74Al1.26O0.42N15.58. However, the calculated value

f x is 0.24, which could be attributed to the rapid densifica-

ion of SPS (the whole sintering process was finished within0 min). The remnant Mg and Al elements are abundant in somerain boundary and triple junctions, as shown in the scanninglectronic image micrograph (Fig. 2).

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ig. 3. Scanning electronic image micrograph of 1850 ◦C sintered Mg–sialoneramics.

.2. Microstructures

SEM micrograph of fracture surface for 1850 ◦C sinteredg–sialon is shown in Fig. 3. Transgranular and intergranu-

ar fracture can be observed. Some grains were pull out, fromhich could be deduced good fracture toughness of the sample

mproved by grain pull-out mechanism [2]. TEM micrographsf 1850 ◦C sintered Mg–sialon ceramics are shown in Fig. 4. Ashown in Fig. 4(a), the sample consists of fine equiaxed �′-phaseith small content of elongated grains of low aspect ratio. No

xtreme grain growth occurs and the average grain size is smallerhan 1.0 �m, which could be attributed to the rapid densifica-ion of SPS process. From Fig. 4(b), it can be observed that theormation of elongated grain by solution-reprecipitation mech-nism that small �-Si3N4 grain dissolves in liquid ponds andeposits at c-axis direction [18]. Few secondary phases can alsoe observed at triple grain junctions as shown in Fig. 4(c). Theattice fringe micrographs (Fig. 4(d and e)) show that intergran-lar glassy phase is hardly observed between the sialon grainshile snatchy lattice fringes can be observed in the triple junc-

ion area, which indicates that the junction is filled with mixturesf amorphous and crystallographic phases. However, the XRDesult shows no other crystallographic phase, which could bettributed to the low content. The existence of lattice fringes inunction area could be explained by the formation of MgAlSiN3n MgO–Si3N4–AlN system with the presence of liquid phaseshrough the following reactions [17]:

MgO + Si3N4 → Mg2SiO4 + 2MgSiN2 (7)

gSiN2 + AlN → MgAlSiN3 (8)

.3. Optical and mechanical properties

The appearance of translucent Mg–sialon slices of 0.5 mm inhickness is shown in Fig. 5. It can be observed that the sam-

le sintered at 1850 ◦C showed the best optical transmittance.he slice is tan in color. Characters under the slice are clearly

egible through the sample. The transmittances against wave-ength curves of the three samples are shown in Fig. 6. Although

Y. Xiong et al. / Materials Science and Engineering A 488 (2008) 475–481 479

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ig. 4. TEM micrographs of (a) low magnified image, (b) formation of elongatee) triple grain junction in 1850 ◦C sintered Mg–sialon ceramics.

he shape of three curves is similar, the transmittance of the

ample sintered at 1850 ◦C is the highest in all the wavelengthange. The transmittance gradually increases with the increasef wavelength and achieves the maximum value of 66.4% at.47 �m wavelength. With the further increase of wavelength

wtcr

ins, (c) triple junction area, and lattice fringe images of (d) grain boundary and

fter 4.5 �m, transmittance decreases sharply to zero at 5 �m

avelength. The transmittance value keeps higher than 50% in

he range of 2.0–4.7 �m wavelength, which indicates good opti-al transmittance of prepared Mg–sialon in medium infraredange.

480 Y. Xiong et al. / Materials Science and Engineering A 488 (2008) 475–481

Fig. 5. Appearance of (a) 1800 ◦C, (b) 1850 ◦C and (c) 1870 ◦C sinte

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ig. 6. Optical transmittance curves of different temperatures sintered translu-ent Mg–sialon ceramic slices (of 0.5 mm thickness) in 1.5–5.5 �m wavelengthegion.

The bending strength value of sintered translucent Mg–sialoneramics was not obtained because of the limitation of sampleize. Selected properties of prepared Mg–sialon ceramics arehown in Table 3. The sample sintered at 1800 ◦C shows theowest hardness and fracture toughness due to its lowest density.or the sample sintered at 1850 ◦C, high content of �′-sialonhase results in the hardness as high as 21.4 ± 0.3 GPa. Theracture toughness value calculated from indentation method isatisfying 6.1 ± 0.1 MPa m1/2, which can be attributed to thexistence of a small amount of �′-sialon whose grain possessesn elongated morphology.

. Discussions

To obtain polycrystalline ceramics with good optical trans-arency, the sintered body should, first of all, be highly densifiedo eliminate light scattering by micropores. And the grain bound-ry phases, which can scatter light due to unmatched refractivesndex with the matrix, should be avoided. The sintered bodys prefered to be composed of a single phase, i.e. �-sialon, ifoth �′ and �′ phases present, the different refractive indicest grain boundaries will increase light scattering. The opticalransmittance of the translucent Mg–sialon ceramics is mainly

nfluenced by the density and �′ content of the sintered body,hich explains the best optical property of 1850 ◦C sintered

ample. Meanwhile, the material should be strong enough toeet requirements of severe working environments, for instance

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red translucent Mg–sialon ceramic slices of 0.5 mm thickness.

issile dome materials. Furthermore, transmittance can also bemproved by decreasing wall thickness. As shown in Table 1, lit-ratures regarding optical property of sialon ceramics primarilyoncern single �′-sialon phase with fine exquiaxed microstru-ures. Those materials possess high hardness, but relative lowracture toughness, commonly less than 3 MPa m1/2, which mayimit application of those materials. Dy-�′-sialon prepared byu et al. [7] achieved the highest fracture toughness value of.1 MPa m1/2 in reported literatures, however, a strong absorb-ng band occurs at ∼2.75 �m wavelength in Dy-�′-sialon andt ∼3.5 �m wavelength in (Dy, Y)-�′-sialon [8]. Jones et al. [5]oubled the fracture toughness of single-phase �′-material byroducing two phase �′/�′ composite sialons. However, the opti-al transmittance is greatly lost due to the high �′-phase content.n present study, Mg–sialon ceramic achieves satisfying opticalnd mechanical properties, which could be attributed to the fol-owing reason. (1) The content of �′-phase is low, which does notreatly compromise optical property of two-phase �′/�′ com-osite. (2) The material consists of mainly fine equiaxed grainsith an average grain size less than 1 �m. Light transmittance

n polycrystalline materials increases with decreasing grain size19]. Duplex �′/�′-sialon ceramics offer possibilities of tailoringhe microstructure by varying the �′:�′-phase ratio, for exam-le, equiaxed �′-sialon grains can be matched with elongated′-grains to form toughened composites, and consequently theroperties of the final product can be improved, because theyombine the high fracture toughness of �′-sialon with the goodardness and optical transmittance of �′-sialon.

Four possibilities of forming sialon solid solutions presenthemselves as follows:

-Si3N4 → �′-sialon (9)

-Si3N4 → �′-sialon (10)

-Si3N4 → �′-sialon (11)

-Si3N4 → �′-sialon (12)

The development of duplex microstructures appears to beependent on the rate of nucleation of �′-sialon [20]. �-Si3N4taring powders lead to faster Si3N4 → �′/�′-sialon transforma-ions than �-Si3N4 staring powders, because of a higher forceor the less stable � powders [21]. Moreover, transformations

hat occur by solution-reprecipitation require a new phaseucleated from the liquid phases (as shown in Fig. 3(b)),hich are formed from reactions between induced sintering

dditives and oxide layers on starting particles surface. It is an

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Y. Xiong et al. / Materials Science

ffective way to tailor the final obtained structures of the ratiof �′:�′-phase through controlling the amount of liquid phases.he results from Sung et al. [22] have shown that an increasingmount of AlN could retard the phase transformation andower the density of the sintered bodies due to the decreasingontent of formed liquid phases in Si3N4–AlN–MgO system.hree weight percent of MgO and 9 wt% AlN were found

o be the optimal addition content in our work. However, theensification of Mg–sialon ceramics by conventional methods isather difficult due to the decrease in the amount of liquid phase.he sample prepared by hot-pressing achieves only 98.2% of

he theoretical density and 75.3 vol% �′-phase content. Longeroaking time and higher sintering temperature could bring outigher density and greater phase transformation. In presenttudy, SPS technique is applied to harmonize the conflict. SPSs mostly known for its rapid densification process, hence highintering density and �′-phase content could be obtained byhortening soaking time under high temperature.

. Conclusions

1) Translucent Mg–sialon ceramics achieving 66.4% ofmaximum infrared transmittance, 21.4 ± 0.3 GPa hardness,and 6.1 ± 0.1 MPa m1/2 fracture toughness were preparedby spark plasma sintering �-Si3N4 powder at 1850 ◦Cfor 5 min with 9 wt% AlN and 3 wt% MgO as sinteringadditives. Improved fracture toughness could be attributedto the existence of a small amount of �′-sialon withelongated morphology. The reason that the two-phasematerial does not greatly compromise optical propertycould be attributed to the fine equiaxed microstructures andlow content of �′-phase.

2) The optical and mechanical properties of translucent

Mg–sialon ceramics are affected by the �′:�′-phase ratioand the density in sintered bodies, which were tailoredby controlling the content of formed liquid phase andoptimizing sintering parameters.

[[[[

ngineering A 488 (2008) 475–481 481

3) SPS technique is an efficient densification method forthe fabrication of translucnet Mg–sialon. High sinteringdensity and �′-phase content could be obtained from itsrapid consolidation process with shorter soaking time underhigh temperature.

cknowledgments

The authors thank for financially support by Ministry of Edu-ation of China (No. 704034, PCSIRT0644) and the Nationalatural Science Foundation of China (No. 50772081).

eferences

[1] F.L. Riley, J. Am. Ceram. Soc. 83 (2000) 245.[2] H. Mandal, J. Euro. Ceram. Soc. 19 (1999) 2349.[3] M. Mitomo, Y. Moriyoshi, T. Sakai, et al., J. Mater. Sci. 1 (1980) 25.[4] B.S.M. karunaratne, R.J. Lumby, M.H. Lewis, J. Mater. Res. 11 (1996)

2790.[5] M.I. Jones, H. Hyuga, K. Hirao, J. Am. Ceram. Soc. 86 (2003) 520.[6] M.I. Jones, H. Hyuga, K. Hirao, et al., J. Am. Ceram. Soc. 87 (2004)

714.[7] X.L. Su, P.L. Wang, W.W. Chen, et al., J. Am. Ceram. Soc. 87 (2004)

730.[8] X.L. Su, P.L. Wang, W.W. Chen, et al., J. Mater. Sci. 39 (2004) 6257.[9] W.W. Chen, X.L. Su, P.L. Wang, et al., J. Am. Ceram. Soc. 88 (2005)

2304.10] Y. Xiong, Z.Y. Fu, H. Wang, et al., J. Am. Ceram. Soc. 90 (2007) 1647.11] Z.-J. Shen, T. Ekstrom, N. Nygren, J. Am. Ceram. Soc. 79 (1996) 721.12] T. Ekstrom, P.O. Kall, M. Nygren, et al., J. Euro. Ceram. Soc. 24 (1998)

1853.13] C.P. Gazarra, D.R. Messier, Ceram. Bull. 56 (1977) 777.14] S. Kurama, M. Herrmann, H. Mandal, J. Eur. Ceram. Soc. 22 (2002) 109.15] A.G. Evans, E.A. Charles, J. Am. Ceram. Soc. 59 (1976) 371.16] K. Niihara, R. Morena, D.P.H. Hasselman, J. Mater. Sci. Lett. 1 (1982) 13.17] S.F. Kuang, Z.K. Huang, W.Y. Sun, et al., J. Mater. Sci. Lett. 9 (1990) 69.

19] R. Apetz, M.P.B. van Bruggen, J. Am. Ceram. Soc. 86 (2003) 480.20] J. Kim, A. Rosenflanze, I.W. Chen, J. Am. Ceram. Soc. 12 (2000) 1819.21] A. Rosenflanze, I.W. Chen, J. Eur. Ceram. Soc. 19 (1999) 2337.22] R.J. Sung, T. Kusunose, T. Nakayama, et al., Ceram. Trans. 165 (2005) 15.