Fabrication of β-sialon nanoceramics by high-energy mechanical milling and spark plasma

sintering

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2005 Nanotechnology 16 1569

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INSTITUTE OF PHYSICS PUBLISHING NANOTECHNOLOGY

Nanotechnology 16 (2005) 1569–1573 doi:10.1088/0957-4484/16/9/027

Fabrication of β-sialon nanoceramics byhigh-energy mechanical milling and sparkplasma sinteringXin Xu, Toshiyuki Nishimura, Naoto Hirosaki, Rong-Jun Xie,Yoshinobu Yamamoto and Hidehiko Tanaka

National Institute for Materials Science, Advanced Materials Laboratory, 1-1 Namiki,Tsukuba, Ibaraki 305-0044, Japan

E-mail: XU.Xin@nims.go.jp

Received 8 October 2004, in final form 2 June 2005Published 1 July 2005Online at stacks.iop.org/Nano/16/1569

AbstractA single-phase β-sialon nanoceramic, Si5AlON7, has been prepared byhigh-energy mechanical milling, followed by spark plasma sintering. Aftermilling, the starting powder mixture (Si3N4, AlN, Al2O3) was mostlytransformed into an amorphous phase that contains a large number ofwell-dispersed nanocrystalline β-Si3N4 particles. Milling promoted themixing and reaction among the starting powders; β-sialon grains with thedesigned composition were formed directly through precipitation from thehomogeneous amorphous phase on the nanocrystalline β-Si3N4 particles.Milling also improved the sintering ability of the starting powders, so thedensification temperature was lowered by about 100 ◦C compared with thatfor as-received powders. Particle rearrangement was the main densificationmechanism for the milled powders. A homogeneous microstructurecomposed of equiaxed β-sialon grains with a diameter of about 50 nm wasobtained after sintering at 1550 ◦C for 5 min.

1. Introduction

Silicon nitride-based ceramics are considered among the mostpromising materials for advanced structural applications, dueto their unique combination of excellent properties such ashigh strength, high hardness, high toughness, good oxidationresistance, good chemical stability, low coefficient of thermalexpansion, and low density. Sialon ceramics, in comparisonwith Si3N4 ceramics, have better high-temperature propertiesand higher hardness due to the smaller amount of intergranularglassy phase [1, 2].

β-sialons, which were the first group developed in thesialon family, are formed by replacing up to two thirds of theSi in the β-Si3N4 by Al, provided that valency compensation ismaintained by the replacement of an equivalent concentrationof N by O, giving a range of β-sialons, Si6−zAlzOzN8−z with0 < z < 4.2. Since the difference between the respectivebond lengths (0.174 nm for Si–N and 0.175 nm for Al–O)is small, the lattice strain is also small and the extent of thereplacement is wide [2]. Many efforts have been made to

enhance the properties of β-sialon ceramics in the past threedecades [3–8].

Nanomaterials have come to constitute a fast growingfield of research in the past 20 years. At the beginning,mainly metals were investigated; later a large amount ofwork was performed on nanoceramic materials [9, 10]. Anincrease in hardness, bending strength, Weibull modulus,and abrasive wear resistance has been observed in thesenanoceramics [11]. Most importantly, the nanoceramicsexhibit excellent superplasticity [12–14], which is veryimportant for cost-effective near-net-shape forming ofcomplex ceramic components.

The preparation of sialon nanoceramics has not beenreported up to now. Failures appear to be connected todifficulties in preparing nanosized powders and densificationwithout fatal grain growth, which also occurs in many ceramicsystems. In this paper, we succeed in preparing β-sialonnanoceramics for the first time. The novel method is basedon high-energy mechanical milling followed by spark plasmasintering [15]. The effects of high-energy mechanical milling

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Figure 1. XRD patterns of BP0 and BP2.

on densification, phase transformation, and microstructuraldevelopment have been investigated.

2. Experimental procedure

Commercially available β-Si3N4 (NP500 grade, DenkiKagaku Kogyo Co., Tokyo, Japan, d50 = 0.5 µm), AlN (GradeF, Tokuyama Soda Co., Tokyo, Japan), and Al2O3 (99.9% pure,Sumitomo Chemical Co., Tokyo, Japan) were used as startingpowders. The starting mixture with nominal compositionSi5AlON7 (z = 1) was mixed in ethanol using silicon nitrideballs for 4 h. The as-received powder mixture (denoted asBP0), about 6.2 g, was then high-energy mechanically milled(Seishin planetary ball mill, PM-1200, Japan) using siliconnitride balls 5 mm in diameter and a silicon nitride pot 350 mlin volume. The ball-to-powder weight ratio was 20:1, millingspeed 475 rpm, and milling time 4 h. The powder mixtureobtained was denoted as BP2.

The powder mixtures were then compacted in a carbondie with a 15 mm inner diameter and sintered using a sparkplasma sintering system (Sumitomo Coal Mining Co. Ltd.,Tokyo, Japan) under a compressive stress of 30 MPa. Theheating and cooling rates were about 300 and 600 ◦C min−1,respectively. All the processing steps were performed underN2 atmosphere to avoid oxidation.

The x-ray diffraction (XRD) of the powders and thesintered samples was determined by an x-ray diffractometer(Philips PW 1700) using Cu Kα radiation at a scanning rateof 2◦ min−1. An internal standard substance (Si) was added.The β-sialon/(β-Si3N4 + β-sialon) ratio was calculated on thebasis of the peak heights of the (200) and (101) diffractions.The lattice parameters of β-sialon were found to depend on theAl content. Accordingly, to characterize the Al solid solutionin terms of z values, the lattice parameters were calculatedusing the following two equations [3] and considering a greatnumber of diffraction positions:

a = 0.7603 + 0.002 96 × z

c = 0.2907 + 0.002 55 × z.

(a)

(b)

Figure 2. Morphology of BP2 powder. (a) TEM, (b) HRTEM.Nanocrystalline β-Si3N4 is indicated by dashed lines.

Thermogravimetry and differential thermal analysis (STA409CD, simultaneous thermal analysis, NETZSCH, German) wereconducted in flowing high-purity nitrogen; the temperature wasincreased to 1500 ◦C (limited by apparatus) at a heating rate of10 ◦C min−1.

The bulk density of the sintered sample was measuredfrom the mass and dimensions. The milled powder wasobserved using field-emission HRTEM (JEM-3000F). Themicrostructure of a fracture surface of the sintered body wasobserved by scanning electron microscopy (Hitachi 5100).The grain size was estimated from the SEM image.

3. Results

3.1. High-energy mechanical milling

The XRD patterns of the BP0 and BP2 powders are shownin figure 1. After milling, the peaks of the starting powdersare significantly weakened and a highly diffuse backgroundstructure is observed. This indicates that the powders havebeen mostly transformed into an amorphous state. The residualβ-Si3N4 peaks are greatly broadened, suggesting a significantrefinement in the particle size of crystalline β-Si3N4.

Figure 2 shows TEM and HRTEM images of BP2powder. The powder features many particle agglomerates with

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Figure 3. Thermal analysis curves of as-received and milledpowders in flowing nitrogen.

Figure 4. Effect of sintering temperature on sintered density for thesame holding time of 5 min for BP0 and BP2.

diameters of 100–200 nm. Both the amorphization of startingpowders and the existence of nanosized Si3N4 particles areconfirmed by figure 2(b). Typical thermal analysis curves forthe as-received and milled powders heated in flowing nitrogenare shown in figure 3. Milling produces a new DTA exothermicpeak at about 1430 ◦C, which corresponds to the crystallizationof the amorphous fraction to β-sialon, as confirmed by XRDanalysis (not shown in this paper).

3.2. Spark plasma sintering

Both powders are spark plasma sintered, but at differenttemperatures for the same holding time of 5 min. Ascan be seen in figure 4, BP2 powder can be nearly fullydensified by treatment at 1550 ◦C for 5 min, while BP0 powderreaches nearly its theoretical density limit at 1650 ◦C. Thedensification temperature for BP2 is about 100 ◦C lower thanthat for as-received BP0 powder.

Figure 5 shows the β-sialon contents and the calculatedz values for sintered BP0 and BP2 samples for differentsintering temperatures, as calculated from XRD. Onlyβ-Si3N4 and β-sialon peaks can be detected for all samples.The β-sialon content for BP2 is 100%, while it is low forBP0, indicating that high-energy mechanical milling promotesthe phase transformation by amorphization of the startingpowders. The calculated z values of the β-sialon from BP2are always 1, as designed (the oxidation of starting powders

Figure 5. β-sialon content and z value of β-sialon after sintering atdifferent temperatures.

during milling has been avoided), whereas those for β-sialonfrom BP0 are greatly higher than 1, indicating that the β-sialongrains initially formed are richer in Al and O than the overallcomposition, which is in good agreement with the resultsreported in the literature [5, 6], although α-Si3N4 powder wasused as the starting powder there. As can also be seen fromfigure 5, the β-sialon content and calculated z value for BP0approaches equilibrium with increasing temperature.

Figures 6 and 7 show the fracture surfaces of denseceramics sintered at different temperatures starting fromBP2 and BP0 powders, respectively. The possibility ofachieving densification at lower temperature and the existenceof crystalline nano-Si3N4 after the milling treatment, as inthe BP2 specimens, can also influence the grain size of thefinal dense ceramics. Treatment of BP0 powder results incoarse equiaxed grains with an average grain size of about1 µm, while that of BP2 powder leads to homogeneouslydispersed fine equiaxed grains with average grain sizes of about50 nm (1550 ◦C) and 60 nm (1600 ◦C). It can be concludedthat the proposed method is effective for fabricating β-sialonnanoceramics.

4. Discussion

Many authors have studied the densification and phasetransformation of β-sialon ceramics from starting powders ofSi3N4, AlN and Al2O3 [5–8]. Transient liquid-phase sinteringis the term generally used to describe the process, which can bedescribed as proceeding in three (partly overlapping) stages:

Stage 1. Particle rearrangement, immediately followingthe formation of the liquid phase.

Stage 2. The solution–diffusion–reprecipitation process,which involves the solution of materials at the points ofcontact of particles and a centre-to-centre approach. This stageinvolves also solution of small grains and reprecipitation onlarge grains, so it is accompanied by grain growth. As thedensification proceeds, most of liquid phase is incorporatedinto the grains of newly formed sialon grains.

Stage 3. A coalescence stage, in which minordensification through grain coarsening is observed.

The formation of liquid phase in BP0 powder is basedon the reaction among Al2O3, SiO2, and small amounts ofAlN and Si3N4, where SiO2 is unavoidably present in thesurface of each Si3N4 particle. Due to the small amounts

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Figure 6. Fracture surface of a BP2 sample after sintering at different temperatures for 5 min. (a) 1550 ◦C, (b) 1600 ◦C, (c) 1650 ◦C.

Figure 7. Fracture surface of a BP0 sample after sintering at1700 ◦C for 5 min.

of sintering additives added and the low reaction rate, theamount of transient liquid phase is small. In BP2 powder,thanks to the high-energy mechanical milling, nearly all the

powders mix and form an amorphous Al–Si–O–N phase, whichfacilitates the formation of the transient liquid phase and furtheraccelerates the densification process.

Because the amorphous phase in BP2 is far fromthermodynamic equilibrium conditions, the first stage ofdensification (particle rearrangement) is accompanied by rapidprecipitation ofβ-sialon grains. Theβ-sialons precipitate fromthe liquid phase during densification of BP2 with a very highdriving force, which leads to isotropic grain growth [16]. Thelarge number of uniformly dispersed nano-β-Si3N4 particles inthe liquid phase can act as seeds, favour uniform precipitation,and limit the growth of each grain, leading to equiaxed nano-β-sialon particles.

The particle rearrangement is dependent on the particlesize and shape as well as the amount and viscosity of the liquidphase. Compared with the DTA curve, the precipitation ofβ-sialon from BP2 during spark plasma sintering is delayed tohigher temperature due to high heating rate, where a smallergrain size and greater amount of liquid phase are available forrapid rearrangement. This result is in good agreement with thesintering of α-sialon [17]. The shrinkage curve of BP2 powdersintered at 1600 ◦C for 5 min is plotted in figure 8. Nearly fulldensification is achieved through rearrangement within 2 min.

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Figure 8. Shrinkage curves of BP2 powder sintered at 1600 ◦C for5 min.

The large particle size and small amount of liquid phasein BP0 make it impossible for it to be fully densified atthe same densification temperature. Furthermore, a shapeaccommodating solution–diffusion–reprecipitation process isnecessary for complete densification of BP0, while it isnot necessary for the milled powder because of the grainprecipitation accompanying the densification [15].

After the formation of liquid phase during the sintering ofBP0 powder, with better wettability, AlN initially traps theoxide melt, preventing Si3N4 dissolution. Preferential dis-solution of AlN enriches the melt composition in Al, trig-gering transient precipitation of supersaturated β-sialon [5].With the dissolution of Si3N4 grains and the precipitation ofβ-sialon, the Al concentration is diluted. So the initiallyformed β-sialons are richer in Al and O than the overallcomposition [5, 6]. The solution–diffusion–reprecipitationprocess is time-consuming, so a longer time is requiredfor complete transformation. In contrast, the formation ofβ-sialons from BP2 is based on the rapid precipitation ofa homogeneous non-equilibrium amorphous phase, so theβ-sialons obtained always have the designed composition(z ∼ 1) and the phase transformation is rapid.

The full densification of BP2 could be achieved in therearrangement process, so the coarsening (solution–diffusion–reprecipitation) stage has been mostly avoided. For as-receivedBP0, the formation of β-sialons is based on the solution–diffusion–reprecipitation process; most β-sialons grow fromlarge β-Si3N4 grains [18], so a coarse microstructure hasbeen obtained. It should be noted that the ceramic from BP0

(1700 ◦C, 5 min) is composed of only about 40 wt% β-sialon;small grains in figure 6 seem to be unreacted β-Si3N4, so thegrain would be bigger if 100 wt% β-sialon was achieved.

5. Conclusion

(1) High-energy mechanical milling followed by sparkplasma sintering has proved to be an effective fabricationtechnique for obtaining single-phase β-sialon (z = 1)

nanoceramics.(2) The densification and phase transformation of as-received

powders are mainly based on the solution–diffusion–reprecipitation process.

(3) For milled powders, the main densification mechanismis particle rearrangement; and the phase transformationis based on precipitation from homogeneous amorphousAl–Si–O–N phase on nano-β-Si3N4 particles.

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