Journal of the European Ceramic Society - ANU College of ...users.cecs.anu.edu.au/~Qinghua.Qin/publications/pap248E-JECS.pdf · Journal of the European Ceramic Society 36 (2016) 671–678

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Journal of the European Ceramic Society 36 (2016) 671–678

Contents lists available at www.sciencedirect.com

Journal of the European Ceramic Society

jo ur nal home p ag e: www. elsev ier .com/ locate / jeurceramsoc

ast densification mechanism of bimodal powder during pressurelessintering of transparent AlON ceramics

ingchun Shana,∗, Zhaohui Zhanga, Xiannian Suna, Jiujun Xua,∗, Qinghua Qinb,iangtao Li c

Department of Materials Science and Engineering, Dalian Maritime University, Dalian 116026, ChinaResearch School of Engineering, Australian National University, Acton, ACT 2601, AustraliaTechnical Institute of Physics and Chemistry, Chinese Academy of Sciences, Beijing 100080, China

r t i c l e i n f o

rticle history:eceived 15 September 2015eceived in revised form 19 October 2015ccepted 20 October 2015vailable online 4 November 2015

a b s t r a c t

Both bimodal and unimodal AlON powders were pressureless sintered and their phase assemblages andmicrostructure evolutions at 1400–1900 ◦C were monitored to reveal the fast densification mechanismof bimodal powder. Phase assemblages analysis of samples sintered at 1400 ◦C suggests that fine AlONparticles were transformed into hexagonal �-Al2O3 and AlN, while coarse AlON particles were trans-formed into cubic �-Al2O3 and AlN. Correspondingly, only a small fraction of the unimodal mixture was

eywords:ensificationarticle size distributionluminum oxynitridehase transformation

�-Al2O3, while the relatively unstable �-Al2O3 kept dominating the bimodal samples and most of themwere directly reformed into AlON at a lower temperature of 1600–1700 ◦C for their similarity in crystalstructure. As a result, a bimodal grain size distribution was kept in the whole sintering process of thebimodal sample, which led to its fast densification. Using the bimodal AlON powder, highly transparentAlON ceramics were fully pressureless sintered at 1820 ◦C after holding for 2.5 h.

© 2015 Elsevier Ltd. All rights reserved.

. Introduction

In recent years, transparent aluminum oxynitride (�-AlON) withubic structure has attracted a growing interest due to its excellenttrength and high optical transparency, for that has many superiorpplications [1–5]. Transparent AlON can be fabricated by usingither pressing or pressureless sintering technique. Pressurelessintering technique is favored for its low cost and easy shapingn manufacturing transparent AlON compared to hot pressing orot isostatic pressing techniques [6,7]. Nevertheless, pressurelessintering transparent AlON by using traditional powder requires aigh sintering temperature about 1920–2000 ◦C and a long holdingime of ≥7 h, which is still high energy consumption.

Recently, we proposed a new starting powder having bimodalarticle size distribution (PSD) to further cut the pressureless sin-ering costs of transparent AlON. Our work has demonstrated that

his powder can be fast sintered for its fast heating of 40 ◦C /mincompared to ∼10 ◦C /min in other works) and a marginally lowerintering temperature of 1880 ◦C. More importantly, the holding

∗ Corresponding authors. Fax: +86 41184723376.E-mail addresses: [email protected] (Y. Shan), [email protected] (J. Xu).

ttp://dx.doi.org/10.1016/j.jeurceramsoc.2015.10.026955-2219/© 2015 Elsevier Ltd. All rights reserved.

time of pressureless sintering transparent AlON using this bimodalpowder can be remarkably shortened to 1.5–2.5 h [8]. The fast sin-tering of the transparent AlON ceramics using the new powder ismainly attributed to fast surface diffusion during heating, due tothe initial bimodal PSD of starting powder, and fast grain bound-ary diffusion during both heating and holding, due to the transientbimodal grain size distribution of AlON. It was noticed that fastheating and high relative density up to 99.11% once heating to1880 ◦C without holding indicated that the powder can be fast den-sified. However, the fast densification mechanism of this bimodalpowder was not very clear because no observation and measure-ment have been made before 1880 ◦C in that work. In this paper,a serial of experiments were performed at 1400–1900 ◦C for boththe bimodal and unimodal powder, and the densification processof both powders was monitored to reveal fast densification mech-anism of the bimodal powder.

Generally, contact fraction of solid particle, nucleation rate andsintering temperature are recognized as the key factors affectingthe solid state reaction rate. Contact fraction of solid particle iscontrolled by both particle size and particle size distribution. Mean-

while, experiment and theoretical analysis revealed that the phasetransformation kinetics is largely governed by the heterogeneousnucleation that is strongly sensitive to the driving force, dictated by

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72 Y. Shan et al. / Journal of the Europ

he majority phase of the starting powder, and that coherent nucle-tion is often favored when abundant isostructure nucleation sitesre available, i.e., nucleation will be easier if one or more reactantsave the similar structure with the product [9–11]. Higher temper-ture and long holding may be needed to eliminate the defects inhe product if grain boundary and volume diffusion during hold-ng is slow. Therefore, increasing contact fraction of powder orrain, controlling the crystal structure of intermediate reactantsan accelerate nucleation and phase transformation at low tem-erature. Fast grain boundary and volume diffusion during heatingnd holding is also favored to shorten holding time.

To obtain high contact fraction of powder, reducing particleize is a good choice in some cases. However, it is very costly tobtain the particle fine enough for most of ceramics, and the largehrinkage caused by fine particle sintering usually are harmful toample. Alternatively, one can manipulate the particle size distribu-ion by mixing the fine particle with a fraction of coarse particle tomprove contact fraction of raw powder. As a result, a mixture hav-ng bimodal PSD has been widely studied and applied in solid stateeaction. Research work on the effects of powder PSD on compactensity suggested that bimodal PSD mixtures provide higher greenensities due to its improved gap filling and particle rearrangementbility [12–17]. According to theoretical calculation, particle rear-angement is mainly carried out by surface diffusion in its sinteringrocess, which is associated with three types of particle contact,

.e., large–large, small–large and small–small [16]. Moreover, theearrangement parameters � (representation of rearrangement, thengle between the real motion vector of particle and the meaneld strain vector associated to particle) calculated by Martin [17]lso showed that the greatest value of � was reached for the par-icle assemblies characterized by a domination of large–large andmall–large contacts, which is the exact case for the bimodal pow-ers. Therefore, the proposed bimodal AlON powder can lead toast pressureless sintering benefited from both its increased contactraction and improved particle rearrangement ability associatedith particle contact type.

In the sintering process of transparent AlON ceramics from AlONowder, AlON particle is firstly transformed into Al2O3 and AlN at

ow temperature which then forms AlON again at high temper-ture (AlON is thermodynamically unstable below a temperatureetween 1600 ◦C and 1640 ◦C [18], which means that Al2O3, AlNr AlON may coexist in the mixtures during heating to sinteringemperature). In the homogeneous regions of transparent AlONharacterized by cubic crystal structure, the content of Al2O3 is7–80 mol% (depending on the temperature) [19]. It is evident thatl2O3 is the main reactant in the mixture to be sintered into AlON.s a matter of fact, Al2O3 can be present in many kinds of crystaltructures, such as cubic �-Al2O3, hexagonal �-Al2O3, and etc. It isorth noting that �-Al2O3 is more unstable than �-Al2O3 at high

emperature. In addition, it is well known that transparent AlONeramics has a cubic crystal structure for its high transparency,hich implies that �-Al2O3 has faster phase transformation kinet-

cs than �-Al2O3 because of the similarity in crystal structureetween �-Al2O3 and transparent AlON. Therefore, it appears that-Al2O3 is favored in pressureless sintering of transparent AlONeramics because it could be faster transformed into AlON at a loweremperature compared to �-Al2O3. However, the reduction mech-nism of Al2O3 from AlON powder is not well understood, whichakes it hard to control the crystal structures of transformed Al2O3.oreover, �-Al2O3 may be transformed into �-Al2O3 before it can

e transformed into AlON at high temperature.One possible way of controlling the phase of Al2O3 transformed

rom AlON powder is through adjusting the particle size. Study ofarticle size effect on crystal structure of Y2O3 particles synthesisevealed that there is a critical particle diameter of approximate.5 �m in the formation of phase of Y2O3 particles [20]. Particles

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significantly smaller than the critical diameter were all monoclinic,while those significantly larger were all cubic. The particle sizeeffect was also reported for heavy rare earth sesquioxides RE2O3(RE = Dy, Ho, Er, Tm and Yb) [21]. This particle size effect wasattributed to the interplay between surface energy and polymor-phism [22]. Therefore, it is highly possible for the proposed bimodalpowder to form unstable cubic �-Al2O3 during heating, which inturn contributes to fast sintering of transparent AlON and loweringthe sintering temperature. As a result, the phase assemblages ofall the sintered samples during heating were characterized in thisstudy. Indeed, unstable cubic �-Al2O3 was detected in the samplesduring heating.

Based on the above analyses and assumptions, in order to revealthe fast densification mechanism of the proposed bimodal powder,both a bimodal PSD AlON powder characterized by a domination oflarge–large and small–large contacts and a unimodal powder wereprepared to study the microstructure evolution and the densifica-tion mechanism of AlON ceramics between 1400 and 1900 ◦C bypressureless sintering. For the purpose of comparison, the parti-cle size of unimodal powder was chosen as same as that of smallparticle size of bimodal powder. At temperature <1700 ◦C, major-ity of hexagonal �-Al2O3 were found in the unimodal powder, buta large amount of cubic �-Al2O3 combined with a small fractionof �-Al2O3 were detected in the bimodal mixtures. By using thebimodal powder, AlON ceramics having high transparency werepressureless sintered at 1820 ◦C after holding for 2.5 h. The max-imum infrared transmittance for 3 mm thickness sample was upto 82.1% (94.0% of theoretical in-line transmittance). The fast den-sification mechanism of the bimodal AlON powder was presentedaccordingly.

2. Experiment procedure

2.1. Preparation of bimodal and unimodal PSD AlON powders

AlON powder was firstly synthesized by carbothermal reduc-tion and nitridation (CRN) method (detailed fabrication processwere given in [8]), then 0.5 wt% Y2O3 (Grade C, Starck, Germany)was added into the obtained AlON powder. Using Si3N4 ball (hav-ing 5, 8 and 10 mm diameter, weight ratio at 1.7:1:1.3) as millingmedia, keeping the ball-to-powder weight ratio at 7:1, 30 g of mix-ture powders of AlON and Y2O3 were ground in absolute ethylalcohol (weight ratio: absolute ethyl alcohol/powder = 3.93) by ballmill into both bimodal (at ∼1.1 �m and ∼2.2 �m) and unimodal (at∼1.1 �m) powder. It should be pointed out that 170 rpm for 24 h ofball mill was performed to obtain the bimodal powder, and 250 rpmfor 36 h of ball mill to gain the unimodal powder. The morphologyof the milled powders were observed by scanning electron micro-scope (FSEM; supra 55, Zeiss, Germany) as shown in Fig. 1. Theparticle size distribution of the obtained powders was measuredusing laser particle size analyzer (Mastersizer 2000, Malvern, UK)and their PSD are shown in Fig. 2.

2.2. Sintering of AlON powders

Two groups of sample, i.e., the bimodal and unimodal samples,were prepared using the obtained AlON powder. For each sample,1.6 g of the corresponding AlON powder was individually packedinto a pellet of 13 mm in diameter under 50 MPa. The pellet wasthen pressureless sintered within a graphite furnace in an atmo-sphere of 0.1 MPa N2. For the purpose of comparison, two identical

samples of the bimodal powder together with another two identi-cal samples of the unimodal powder were put into graphite furnaceeach time, and were heated at a heating rate of 40 ◦C /min to 1400 ◦C,1500 ◦C, 1600 ◦C, 1700 ◦C, 1800 ◦C and 1900 ◦C, respectively. Then

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Fig. 1. SEM images of the bimodal (a) and unimodal (b) AlON powders.

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throughout the whole sintering process (from 1400 ◦C to 1900 ◦C).The fraction of �-Al2O3 and �-Al2O3 in two groups of sample

ig. 2. Particle size distribution of the bimodal (a) and unimodal (b) AlON powders.

he furnace was cooled to room temperature to take out the sin-ered samples for testing and characteristic. In addition, one more

ellet packed with the bimodal powder was also sintered at 1820 ◦Cnd held for 2.5 h.

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2.3. Characterization of sintered samples

The phase assemblages of the samples heated to differenttemperature were characterized by X-ray diffractometry (XRD;D/Max-ULtima+, Rigaku, Japan) using Si as an internal standard. Thepolished samples of bimodal powder fabricated at 1800 ◦C, 1820 ◦Cand 1900 ◦C were hot etched at 1640 ◦C for 40 min in vacuum formicrographs observation. The SEM and optical microscope (GX51,Olympus, Japan) were employed to observe the morphology of allthe samples. Average grain sizes of AlON were statistically calcu-lated by using Image Pro plus software, where the average grainsize was the average value of diameters passing through objects’centroid.

The bulk density of all the samples sintered at different tem-perature was measured by Archimedes principle. The samples ofbimodal powder sintered at 1820 ◦C for 2.5 h were ground and pol-ished on both sides to a thickness of 3 mm for optical transmittancemeasurement. Optical transmittance in the wave of 2500–6000 nmwas recorded by a Fourier transform infrared spectroscopy (FTIR;Nexus 670, Thermo Nicolet, USA).

3. Results and discussion

3.1. Morphology of AlON powders

From Figs. 1 and 2, it can be seen that the AlON powder syn-thesized by CRN was successfully milled into both a bimodal (at∼1.1 �m and ∼2.2 �m) and a unimodal (at ∼1.1 �m) powder in thispaper. The unimodal powder has the equivalent size with that ofthe fine particle in the bimodal powder. As calculated in the previ-ous paper, the bimodal PSD AlON powder produced in this paper isin a domination of large–large and small–large contacts [8], whichcan improve its gap filling ability and is helpful to its rearrange-ment resulted from surface diffusion at the early and middle stageof sintering.

3.2. Microstructure and phase assemblages evolution

To better understand microstructure and phase assemblagesevolution of the sintered samples at different temperature, the SEMimages and the XRD patterns of both the bimodal and unimodalsamples after pressureless sintering at 1400–1900 ◦C without hold-ing are presented in Figs. 3 and 4, respectively.

At 1400 ◦C, grain morphology of both the bimodal and unimodalpowder developed toward sphere (Fig. 3a and b) compared to thestarting AlON powders shown in Fig. 1. However, the grain sizeand grain size distribution was found strongly depends on the PSDof the starting powder. A microstructure characterized by bimodalgrain size distribution and limited grain growth was observed inthe bimodal sample (Fig. 1a vs. Fig. 3a), on the contrary, a uniformgrain size and significant grain growth was observed in the uni-modal sample (Fig. 1b vs. Fig. 3b). As a result, the average grain sizeof the unimodal sample is larger than that of the bimodal sample,as shown in Fig. 5. Further detection of their phase assemblagesshowed that all the AlON powders were transformed into Al2O3and AlN at 1400 ◦C. Both cubic �-Al2O3 and hexagonal �-Al2O3 canbe found in both samples. However, it can be seen from Fig. 4athat the bimodal sample was dominated by cubic �-Al2O3, whileonly a small fraction of �-Al2O3 was detected in the unimodal sam-ple. Moreover, all the AlN detected in both group of sample have ahexagonal crystal structure and its crystal structure didn’t change

(as shown in Fig. 4a) and the similarity in PSD between the sin-tered sample at 1400 ◦C and its own starting powder suggests that

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Fig. 3. SEM images of the fracture surfaces of the samples heated to 1400 ◦C from bimodal (a) and unimodal powder (b), 1500 ◦C from bimodal (c) and unimodal powder(d), 1600 ◦C from bimodal (e) and unimodal powder (f), 1700 ◦C from bimodal (g) and unimodal powder (h), 1800 ◦C from bimodal (i) and unimodal powder (j), 1900 ◦C frombimodal (k) and unimodal powder (l), respectively.

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ne AlON particles were transformed into �-Al2O3 and AlN, whileoarse AlON particles were transformed into �-Al2O3 and AlN. Its worth noting that the reduction mechanism of Al2O3 from AlONowder has not been thoroughly studied so far. However, the pres-nce of �-Al2O3 will definitely be helpful to sinter the mixture intoransparent AlON ceramics because they all have a cubic crystaltructure.

Further heating up to 1600 ◦C, the observed grain morphologyf two powders was respectively similar to that at 1400 ◦C andhe grain size distribution kept its own original pattern of start-ng powder. It indicates that the main mass transport is surfaceiffusion for the bimodal powder at the early or middle stage ofintering. Correspondingly, surface diffusion combined with grainoundary diffusion contributed to the grain growth of the fine uni-odal powder at this stage due to its high specific surface energy.n the other hand, �-Al2O3 cannot be detected in the unimodal

ample from 1500 to 1600 ◦C, while �-Al2O3 still dominated inhe bimodal samples. It is reasonable to assume that the smallmount of �-Al2O3 found in the unimodal sample at 1400 ◦C wasransformed into �-Al2O3 because �-Al2O3 is unstable at this tem-

Fig. 4. XRD patterns of the bimodal and unimodal samples heated to 1400 ◦C (a), 1

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perature. Accordingly, only part of �-Al2O3 in the bimodal samplemay be transformed into �-Al2O3 because majority of Al2O3 in thebimodal sample was still �-Al2O3. However, it is not clear to us why�-Al2O3 in the bimodal samples remained in a dominated fraction.Further investigation is under way to clarify this observation.

At 1700 ◦C, only �-Al2O3 together with AlN can be found in theunimodal sample, which clearly demonstrates that �-Al2O3 cannotbe transformed into AlON at this temperature. Similarly, �-Al2O3and AlN can also be detected in the bimodal sample. However, �-Al2O3 can no longer be detected in the bimodal sample. Instead,a large amount of AlON was presented in the bimodal mixture. Itindicates that most of �-Al2O3 in the bimodal mixture was trans-formed into AlON at 1700 ◦C with a possibility of a small amountof �-Al2O3 being transformed into �-Al2O3 in this duration. It isevident that cubic �-Al2O3 is easier to be transformed into AlON ata lower temperature (between 1600 ◦C and 1700 ◦C) compared to

�-Al2O3 due to its instability at this temperature and similarity incrystal structure with transparent AlON ceramics. In addition, from1400 ◦C to 1700 ◦C, the pores in the bimodal mixture observed fromFig. 3 are much smaller and less than that in the unimodal sample

500 ◦C (b), 1600 ◦C (c), 1700 ◦C (d), 1800 ◦C (e) and 1900 ◦C (f), respectively.

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Fig. 5. Average grain size as a function of sintering temperature for the bimodal andunimodal samples.

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ig. 6. Effects of sintering temperature on relative density of the bimodal and uni-odal samples.

ue to its improved gap filling and rearrangement ability of bimodalSD.

Grain morphology at 1800 ◦C for both samples changed dra-atically due to the fast grain growth with all the �-Al2O3 being

ransformed into AlON because no �-Al2O3 can be detected inoth group of sample (Fig. 4e). It is evident that �-Al2O3 can beransformed into AlON at a temperature between 1700 ◦C and800 ◦C, which is higher than that for �-Al2O3 (between 1600 ◦Cnd 1700 ◦C). Moreover, only a few small pores can be observedn the bimodal sample at 1800 ◦C (Fig. 3i) and many vivid largeores can be seen in the unimodal sample at the same temperatureFig. 3j), which implies that the bimodal sample is more densifiedhan the unimodal one. As a matter of fact, the curve of relativeensity vs. temperature shown in Fig. 6 confirms that the rela-ive density of the bimodal sample (98.02%) is much higher thanhat of unimodal one (65.95%) at 1800 ◦C. From 1700 ◦C to 1800 ◦C,oth groups of sample had a quick increase in relative density.pparently, relative density of the bimodal sample (63.57–98.02%)

ncreases faster than that of the unimodal one (49.25–65.95%). Forhe unimodal sample, it is the transformation from �-Al2O3 to AlONhat contributes the most shrinkage in volume leading to densifi-ation. However, for the bimodal sample, only a small amount of-Al2O3 was left in the mixture at 1700 ◦C, which is much less than

hat in the unimodal sample (Fig. 4d). Therefore it is obvious that the

hrinkage in volume due to transformation from �-Al2O3 to AlONn the bimodal sample from 1700 ◦C to 1800 ◦C is very limited. Moreensification for the bimodal sample is due to grain diffusion andolume diffusion of AlON grains transformed by �-Al2O3 between

Fig. 7. Microstructures of the hot etched samples heated to 1800 ◦C (a) and 1900 ◦C(b) without holding.

1700 ◦C and 1800 ◦C, resulting in a rapid growth in AlON grain size,which can be confirmed by Fig. 5.

With all the Al2O3 being transformed into AlON for both groupsof sample at 1800 ◦C, only the diffusion of AlON grains can con-tribute to further densification of the product. The curve of averagegrain size vs. sintering temperature shown in Fig. 5 demonstratesthat the average AlON grain size of bimodal sample grew from6.76 �m to 15.32 �m which is much faster than that of unimodalsample (from 3.5 �m to 9.25 �m) in the same duration with tem-perature being raised from 1800 ◦C to 1900 ◦C, i.e., average graingrowth rate of 0.09 �m/ ◦C vs. 0.06 �m/ ◦C for bimodal vs. unimodalsample. It implies that AlON grain boundary diffusion and volumediffusion of the bimodal sample is much faster than that of the uni-modal sample at the final stage of sintering. The pores observedin the bimodal sample at 1800 ◦C (Fig. 3i) are much smaller andless than that in the unimodal sample at the same temperature(Fig. 3j), and even much smaller than that in the unimodal sampleat 100 ◦C higher temperature (Fig. 3l). Although a small amount ofAlN can be detected in the bimodal sample at 1800 ◦C (as shown inFig. 4e), these residual AlN were quickly dissolved into AlON beforethe sample was heated to 1900 ◦C (as shown in Fig. 4f). Microstruc-tures of the hot etched bimodal samples at 1800 ◦C and 1900 ◦Cgiven in Fig. 7 shows that the AlON grain at this stage has a transientbimodal grain size distribution. It illustrates that AlON grains in thebimodal sample can diffuse much faster (much more grain growth)and better (densified faster and had less and smaller defects, i.e.,

pores and residual AlN, which can be easier to be eliminated) thanthose in the unimodal sample in the final stage of sintering, whichcan only be attributed to its transient bimodal AlON grain size dis-

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(U) heated to 1400 ◦C, 1500 ◦C, 1600 ◦C, 1700 ◦C, 1800 ◦C and 1900 ◦C.

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Fig. 8. Optical images of the bimodal (B) and unimodal samples

ribution inherited from the bimodal starting powder. In turn, theressureless sintering of the bimodal sample into transparent AlONeramics needs a much less holding time.

.3. Fast densification mechanism of the bimodal AlON powder

From Fig. 6, it can be seen that the relative density of the bimodalample is consistently higher than that of the unimodal one athe same sintering temperature from 1400 ◦C to 1900 ◦C. This canlso be observed from the picture of all the samples sintered at400–1800 ◦C shown in Fig. 8, where the diameter of the bimodalample was smaller than that of the unimodal one at the same sin-ering temperature. It is evident that the bimodal AlON powder wasensified faster than the unimodal powder in the whole sinteringrocess.

At the early or middle stage of sintering (before the reforma-ion of AlON, <1700 ◦C), the bimodal sample has a better gap fillingnd rearrangement ability due to its bimodal PSD [17] of the Al2O3nd AlN mixture, which resulted in a higher relative density com-ared to that of the unimodal sample at the same temperature. Dueo the presence of both �-Al2O3(transformed from the fine AlONarticles) and �-Al2O3 (transformed from the coarse AlON parti-les) in the bimodal sample before the final stage of sintering, theransformation of AlON from Al2O3 didn’t accomplish simultane-usly. Instead, �-Al2O3 was transformed into coarse AlON grain at aemperature between 1600 ◦C and 1700 ◦C and �-Al2O3 was trans-ormed into fine AlON grain at a temperature between 1700 ◦C and800 ◦C. Different transform temperature of AlON from two typesf Al2O3 enables the mixture to keep its bimodal PSD. As a result,he reformed AlON grains have a bimodal grain size distributionhich contributes to a much faster (less holding) and better (less

nd smaller defects) densification at the final stage of sintering.Although �-Al2O3 can be transformed into AlON at a lower

emperature (between 1600 ◦C and 1700 ◦C), a high enough sin-ering temperature (between 1700 ◦C and 1800 ◦C in this paper)s still necessary to transform the �-Al2O3 in the bimodal mixturento AlON. Nevertheless, the sintering temperature of the proposedimodal powder is lower than that of reported powder due to its

mproved grain boundary diffusion and volume diffusion at thenal stage of sintering attributed to its transient bimodal AlON grainize distribution. Moreover, as shown in Fig. 9, the bimodal sampleintered at 1820 ◦C holding for 2.5 h shows too much light transmis-ion than that of the present unimodal sample sintered at the sameondition for 3 mm thickness sample (In fact, the unimodal sam-le is opaque). From Fig. 9, an up to 82.1% infrared transmittance

94.0% of theoretical in-line transmittance [8]) at a wavelength of3600 nm was achieved for the bimodal sample, which is evident

hat the proposed bimodal powder can be fully sintered into trans-arent AlON ceramics at 1820 ◦C with holding 2.5 h.

Fig. 9. Transmittance of the bimodal and unimodal AlON samples sintered at 1820 ◦Cheld for 2.5 h, and the optical images of bimodal and unimodal samples (a) for 3 mmthickness sample, the microstructure of bimodal sample (b).

4. Conclusions

1) Both a bimodal (at ∼1.1 �m and ∼2.2 �m) and a unimodal (at∼1.1 �m) AlON powder were successfully prepared by usingball mill. After these powders were heated to 1400 ◦C at a heat-ing rate of 40 ◦C /min, the AlON powders were transformed intoAl2O3 and AlN. Further phase assemblages analysis of both sam-ples sintered at 1400 ◦C suggests that fine AlON particles weretransformed into hexagonal �-Al2O3 and AlN, while coarse AlONparticles were transformed into cubic �-Al2O3 and AlN. It indi-cates that the phase of Al2O3 transformed from AlON powdercould be controlled by the AlON particle size.

2) Cubic �-Al2O3 was transformed into AlON at a temperaturebetween 1600 and 1700 ◦C and hexagonal �-Al2O3 was trans-formed into AlON between 1700 and 1800 ◦C, the lower phasetransformation temperature of �-Al2O3 to AlON is mainly ben-efited from the similarity in crystal structure between them.

3) The bimodal powder kept its bimodal PSD throughout the wholesintering process and therefore it can be densified faster andbetter than the unimodal powder. At the early or middle stageof sintering (before the reformation of AlON, <1700 ◦C), thebimodal sample has a better gap filling and rearrangement abil-ity due to its bimodal PSD. At the final stage of sintering, AlONgrains in the bimodal sample can diffuse much faster and betterthan those in the unimodal sample, which can only be attributedto the transient bimodal AlON grain size distribution inheritedfrom the bimodal starting powder. In other words, it was the

bimodal PSD that played a key positive role in the fast densifi-cation mechanism of bimodal powder.

4) The proposed bimodal powder can be fully sintered into trans-parent AlON ceramics at 1820 ◦C with holding 2.5 h.

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[Atmospheric pressure synthesis of heavy rare earth sesquioxidesnanoparticles of the uncommon monoclinic phase, J. Am. Ceram. Soc. 90 (11)

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cknowledgements

This work was supported by the Fundamental Research Fundsor the Central Universities (3132015097, 3132014323).

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