CHEMICALLY CONTROLLED SINTERING AND MICROSTRUCTURAL DEVELOPMENT IN CERAMICS
KEMIJSKO NADZOROVANO SINTRANJE IN RAZVOJ MIKROSTRUKTURE V KERAMIKI
DRAGO KOLAR
Fakulteta za kemijo in kemijsko tehnologijo, Murnikova 5a, 1000 Ljubljana in Institut J. Štefan, Ljubljana Faculty of Chemistry & Chemical Engineering, University of Ljubljana, and J. Štefan Institute, Ljubljana
Prejem rokopisa - received: 1997-10-01; sprejem za objavo - accepted for publication: 1997-10-21
The sintering step in manufacturing most ceramic articles is dominated by the chemical reactions which occur during firing. Such chemical reactions may be caused by chemical heterogeneity of the constituents of the ceramic body, small amounts of additives or impurities. Even in chemically equilibrated multicomponent systems the capillary forces may cause dehomogenisation and influence the sintering mechanism. Reaction sintering in the BaTi03-CaZr03-Ti02 system is described as an example of a system where controlled chemical heterogeneity may optimise the electrical properties. Sintering of PbZri-xTix03 ceramic is described as an example of chemical dehomogenisation due to capillary forces.
Key words: reaction sintering, ceramic microstructure, barium titanate, lead zirconate-titanate, sintering mechanisms
Proces sintranja v proizvodnji večine keramičnih izdelkov določajo kemijske reakcije, ki potekajo med žganjem. Kemijske reakcije lahko povzročajo heterogenost sestavin keramike, majhne količine dodatkov ali nečistoče. Celo v kemijsko uravnoteženih večkomponentnih sistemih lahko kapilarne sile povzročijo dehomogenizacijo in vplivajo na mehanizem sintranja. Delo obravnava reakcijsko sintranje v sistemu BaTi03-Ti02-CaZrC>3 kot primer procesa, pri katerem lahko z nadzorovano kemijsko heterogenostjo optimiziramo električne lastnosti keramike. Sintranje PbZri-xTix03 keramike obravnavamo kot primer sistema, pri katerem se kemijsko homogena trdna raztopina med sintranjem prehodno dehomogenizira. Fenomen vpliva na mehanizem sintranja in električne lastnosti piezokeramike.
Ključne besede: reakcijsko sintranje. keramična mikrostruktura. barijev titanat, svinčev cirkonat-titanat, mehanizmi sintranja
1 INTRODUCTION
1.1 SWOT analysis of advanced technical ceramics
SWOT analysis (S = strength, W = weakness, O = opportunities, T = threat) is a frequently used method for evaluation of the potential of a particular activity. Such analysis applied to the manufacture of ceramic articles may give the following answers: (Many other arguments may be added).
S: Ceramics are the oldest artificial materials invented by mankind. In last 50 or 60 years the applicability of ceramics increased enormously. As a result, the mar-ket grovvth in many segments in recent decades main-tained a level of 8-10 % per year, vvhich is far above many other industrial activities. W: The weaknesses of ceramics are well knovvn: The producers and consumers complain of insufficient re-producibility, insufficient reliability and insufficient cost effectiveness of ceramic products. O: Experts agree on the great yet unexploited technical
potential of ceramics. T: There are several obstacles vvhich limit opportunities in the ceramic field. To be competitive and assure high performance products, high investments in knovvledge and manufacturing equipment are neces-sary.
In the following, we examine reasons for shortcom-ings of ceramics, listed in "W". In this article, we con-centrate on sintering.
Sintering is the final step in the ceramic fabrication process. It decisively influences the properties of products made from povvders by sintering and represents the last chance of optimising the product.
Sintering theories are based on physical arguments such as the tendency for reduction in free surface energy and on simplified assumptions such as simple particle shapes (spheres) and simple particle arrangements (two or a limited number of spheres).
The general validity of sintering theories developed in pioneertng works by Frenkel1, Ivensen2, Kucyznski3, Kingery4, Coble5, Johnson6 and others was proven by nu-merous experiments. Sintering theories are primarily concerned vvith material transport mechanisms and the kinetics of the sintering process. Knovvledge of both are of fundamental importance for designing a cost effective manufacturing process for ceramic and powder metal articles vvith optimal properties.
Hovvever, the value of sintering theories in the manufacturing practice of specific ceramic products is limited due to the fact that theoretical assumptions about sintering processes are not valid in practice7. Fundamental studies assume (I) a pure, a homogeneous, single compo-nent starting povvder, (II) uniform, small, spherical grains, (III) homogeneous body before sintering, and (IV) sintering at constant temperature. In contrast, prac-titioners in industrial produetion deal vvith (I) impure povvders, usually inhomogeneous mixtures, (II) povvders vvith a vvide distribution of particle sizes, aggiomerated to various degree, (III) vvith non-uniform density distri-
bution, and (IV) sintering at practically non-isothermal conditions.
Among serious inconsistencies between the assump-tions underlying theoretical and basic experimental re-seareh on sintering phenomenon, chemical reactions at sintering temperature among constituents in the sintered body play an important role.
Chemical reactions in many ceramic systems occur among the constituents which are not in equilibrium at high temperatures and form new compounds or solid so-lutions. When the sintering process involves purposely chemically heterogeneous mixtures vvhich are supposed to react during the sintering operation, one uses the term "reaction sintering". Examples are BaTi03 capacitor ce-ramics or ZnO based varistor ceramics. Most ceramic products are made of a basic compound vvith a small amount of additives which are intended to accelerate densification, suppress grain growth or modify the physi-cal properties of the fired ceramics. The role of additives was frequently analysed, for example in 8 and 9. In general, the additives may accelerate sintering by forming a liquid phase or control the mierostructure by forming solid second phases which pin grain boundaries and im-pede grain growth. More subtle effects, vvhich, however, profoundly influence the sintering process, are segrega-tion at the grain boundaries, change of lattice defect concentration and diffusion constants when additives form a solid solution vvith the major component and a change in the ratio of grain boundary energy to free surface energy.
It is clear that the design of the sintering process, based on theoretical physical arguments only, must be modified by taking into account the possible chemical reactions.
In the follovving section, we report the results of re-search conducted in our laboratory on the effects of various aspects of chemical reactions on sintering. Specifi-cally, we report effects of chemical heterogeneity on the sintering process in multicomponent systems. We dem-onstrate that even in single phase systems chemical de-homogenisation/homogenisation takes plače. The aim of the revievv is to stress the importance of knovvledge of chemical phenomena vvhich occur during sintering of ceramics in designing and control]ing the sintering process.
2 REACTION SINTERING IN THE BaTi03-CaZr03-Ti02 SYSTEM
The term "reaction sintering" or "reactive sintering" is used to describe a sintering process in vvhich a chemical reaction in the starting povvder mixture and the densification of the ceramic body occur in the same heating operation. Depending on the material system and proc-essing variables (particle size, temperature, pressure, etc.) the two processes, reaction and densification, can occur simultaneously, in sequence or as some mixture of these.
Temperature [°CJ
Figure 1: Dilatometric curves for BaTi03 - 2 mol % Ti02 (1), BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 (2), BaTi03 - 8 mol % CaZr03 (3) and CaZr03 (4).
Slika X: Krivulje krčenja za mešanice BaTi03 - 2 mol % Ti02 (1), BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 (2), BaTi03 - 8 mol % CaZr03 (3) in CaZr03 (4).
mol'/. CaZrOj
Figure 2: Average grain size vs. composition in BaTi03 - 2 mol % TiC>2 - CaZr03 ceramics sintered 1 hour at 1360°C. Slika 2: Poprečna velikost zrn v odvisnosti od sestave keramike BaTi03 - 2 mol % Ti02 - CaZr03, sintrane 1 uro pri 1360DC.
In ceramic fabrication practice, reaction sintering is usually avoided. The complexity of the processes caused by the chemical potential on one hand and the tendency to decrease the surface energy on the other renders the process difficult to control reproducibly. So, for example, in the manufacturing of soft Mn-Zn ferrites, the oxides are first calcined to form spinel solid solutions and then sintered. On the other hand, reactive sintering offers the possibility to optimise the properties of sintered products vvhich are intentionally not in chemical equilibrium. Such is the čase with several types of functional ceramics, for example varistors or high permittivity ceramic capacitors.
Frequently used compositions for ceramic capacitors are based on the solid solutions (Bai-xCax)(Tii.yZry)03. If Zr is incorporated into BaTiOi the maximum of the per-mittivity curve (at the Curie point) is shifted from 130°C to room temperature, vvhereas the non-ferroelectric Ca-TiOj broadens this maximum. CaTiOi has only limited solubility in BaTiOj and small undissolved particles of CaTiOj act as grain-growth inhibitors in the ferroelectric matrixl().
(Bai-xCax)(Tii-yZry)03 capacitor ceramic is usually prepared by reaction sintering of BaTi03 - CaZr03 mixed powders. BaTi03 and CaZrOs are not compatible at high temperatures. During sintering, chemical reaction takes plače resulting in formation of (Ba,Ca)(Ti,Zr)03 and (Ca,Ba)Ti03 solid solutions". The sintering kinetics and microstructural development of BaTi03-CaZr03 formu-lations are influenced by the kinetics of the chemical reaction, particularly at lower temperatures, i.e. in the in-itial sintering stage.
Research in our laboratory12-13 showed that the par-ticular batch of BaTi03 used started to shrink at around 1100°C and sintered considerably in the temperature re-gion 1300-1350°C (Fig. 1). CaZr03, being more refrac-
tory, started to shrink at around 1250°C and sintered to high density in the temperature region 1450-1500°C. A BaTi03 - 8 mol % CaZr03 mixture started to shrink at approximately the same temperature as BaTi03. Hovv-ever, after a few % shrinkage the kinetics slowed down and temperatures higher than 1500°C were needed for densification. The sinterability of the BaTi03-CaZr03 mixture was improved by avoiding solid state diffusion. A small addition of Ti02 (2 mol %), causing the formation of a low melting BaTiOs-BafiTinO^ eutectic at around 1310°C, strongly improved the sinterability of the mixture. Temperatures between 1300-1350°C were sufficient to achieve high density. The microstructures of dense BaTi03 - 2 mol % TiOj and BaTiOs - 10 mol % CaZr03-2 mol % Ti02 ceramics sintered at 1360°C are shown in Fig. 2. The microstructures also show that CaZr03 strongly reduces the exaggerated grain growth vvhich is common phenomenon in BaTi03 - Ti02 ceramics.
Comparison of the relative linear shrinkage of BaTiOs - 2 mol % Ti02 and BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 compacts as a function of time for isothermal sintering at 1260°C (Fig. 3) clearly indicated the differ-ences in the sintering mechanism. BaTi03 - 2 mol % Ti02 sintering kinetics in the initial stage can be de-scribed by an equation for first order kinetics Al/l = kt" in vvhich n is constant over the vvhole temperature range of sintering. In contrast, for BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 the graphical representation of ln(Al/l0)vs. In t shovvs a change in slope, indicating tvvo different densification regions vvith different n - values. The densification process is especially disturbed in the first stage of sintering at lovver temperatures (Fig. 4) due to the chemical reaction betvveen BaTi03 and CaZr03. When heated, BaTi03 and CaZr03 react to form a (Ba,Ca)(Ti,Zr)03 solid solution. The equilibrium solid solution limit is
ln t [mini
Figure 3: Comparison of the relative linear shrinkage of BaTi03 - 2 mol % Ti02 (a) and BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 (b) compacts as a function of time for the isothermal sintering at 1260°C. Slika 3: Primerjava med relativnim skrčkom oblikovancev iz BaTi03 -2 mol % Ti02 (a) in BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 (b) v odvisnosti od časa izotermnega sintranja pri 1260°C.
1300°C
ln t [min)
Figure 4: The relative linear shrinkage of BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 samples as a function of time for the isothermal sintering temperatures of 1260°C, 128CTC and 1300°C. Slika 4: Relativni linearni skrček oblikovancev iz BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 v odvisnosti od časa pri izotermnem sintranju pri 1260°C, 1280°C in 1300°C.
Time [min)
Figure 5: Degree of reaction (a) as a function of time for BaTi03 + CaZrOj —> BaZrOj + CaTi03 at various temperatures. Slika 5: Stopnja reakcije (a) v odvisnosti od časa za reakcijo BaTiOj + CaZrOj BaZrOj + CaTi03 pri različnih temperaturah.
around 16 mol % CaZr03. Larger amounts of CaZr03 cause formation of barium zirconate phase vvith some Ca and Ti in solid solution and a calcium titanate phase vvith some dissolved BaTi0310''4. In a 1:1 mole ratio BaTi03 -CaZr03 mixture the reaction vvas detectable by XRD analysis after only 1 hour's heating at 1100°C. The tvvo characteristic X-ray reflections of CaTi03 at 2 0 = 59.052 (d = 1.563) and 2 9 = 47.401 (d = 1.918) vvere used for identification of CaTi03 and to determine the amount of CaTi03 formed during the reaction.
The degree of reaction a as a function of time for the reaction BaTi03 + CaZr03 BaZr03 + CaTi03 at various temperatures is shovvn in Fig. 5. At lovver temperatures (1100°C) a remains under 20 % even after pro-longed heating. A considerably higher degree of reaction vvas achieved by heating the mixture of BaTi03 and CaZr03 at 1250°C and higher. The relationship a = f(t) at various temperatures has been evaluated using different mathematical expressions, theoretically derived for various models. The best fit vvas achieved vvith Jander's equation15:
[1 -(1 - a)"3]2 = kt,
vvhich describes three dimensional diffusion16. It may be concluded that the diffusion mechanism is rate con-trolling.
The reaction kinetics da/dt vvere compared vvith the sintering kinetics da'/dt. The densification parameter a' (a'= p-po/pih-po) was calculated from dilatometric meas-urements in the temperature region 1200-1300°C (Fig. 6).
Comparison of the reaction rates vvith the densification rates at various temperatures confirmed the priority of chemical reaction in the first stage of sintering, and densification in the second stage (at higher temperatures). Sintering conditions have a strong influence on the microstructural development. A slovver heating rate
_da' dt
Figure 6: Comparison of the reaction rate da/dt and densification rate da'/dt (a=p'-po/p,h-p0) for BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 samples at 1260°C.
Slika 6: Primerjava med reakcijsko hitrostjo doc/dt in hitrostjo zgoščevanja daVdt (a'=p'-po/pth-po) oblikovancev iz BaTi03 - 2 mol % TiOi - 8 mol % CaZr03 pri 1260°C.
(l°/min) leads to the development of a coarse microstructure vvith a broad grain size distribution, vvhereas more rapid heating (5° or 10°C/min.) results in a ftner microstructure vvith a narrovver grain size distribution. The results may be explained on the basis of the kinetic studies. A fast heating rate favours reaction sintering vvith hindered grain grovvth, vvhereas a slovv heating rate favours chemical reaction vvhich takes plače at lovver temperatures than densification. The heterogeneous structure vvith phase separations in the initial sintering stage favours discontinuous grain grovvth and a broad grain size distribution.
Microstructural development also depends on the solid state diffusion rate during heating. A small excess of Ti02, causing a liquid phase eutectic at sintering temperature, promotes grain grovvth, especially at slovv heating rates. When BaTi03 vvith a small excess of BaO vvas used, grain grovvth during sintering vvas considerably hindered, especially vvhen a slovv heating rate vvas used.
Coarsening of the microstructure as a result of pro-longed firing is accompanied by dielectric property changes. Different compositions of BaTi03 - CaZr03 based ceramics vvere fired at 1360°C for 15, 60, 480 minutes and the resulting dielectric properties vvere measured (Table 1). Whereas an e value belovv the Curie temperature decreases vvith increasing firing time, the peak at the Curie temperature increases in value and broadens. Changes in the intensity and broadness of the peak are accompanied by coarsening of the microstructure. The coarse microstructure shovvs a more pro-nounced Tc shift and higher permittivity, vvhile the room temperature permittivity is lovvered. Maximum values of permittivity vvere obtained vvith the composition 84 mol % BaTi03 - 16 mol % CaZr03 near to the limit of solid solubility in the system.
D. KOLAR: CHEMICALLY CONTROLLED SINTERING Table 1: Dielectric properties of BaZr03 based ceramics vvith 2 mol % TiOi addition vs. soaking rime at 1360°C (at 1 kHz)
Composition	90 m/o BaTi03-10 m/o CZ			84 m/o BaTi03-16 m/o CZ			80 m/o BaTi03-20 m/o CZ		
time (min)	15	60	480	15	60	480	15	60	480
£25°C	4753	4270	2030	4045	7200	10694	2780	3105	5210
tg5.104	107	105	200	79	92	130	45	74	105
Tc(°C)	60	60	70	20	25	35	-20	-15	-10
	6350	7820	8960	4600	7200	10767	3820	4650	6619
AC/C (%)	/	/	/	-31,0 + 14,0	-48,20	-71,2 +0,7	/	/	/
Chemically inhomogeneous ceramics in metastable ceramic equilibrium make it possible to make, by careful control of firing conditions, the temperature stable dielectric materials. This is ascribed to specific core-shell grain structure17. The relatively flat temperature charac-teristic of dielectric constant of BaTi03-CaZr03 based ceramics is determined by the superposition of the two permittivity/temperature maxima, those of BaTi03 core at - 120°C and (Ca,Zr) doped BaTi03 shell with maxi-mum at lower temperature. Core-shell structure is dis-cernible in TEM micrograph in Fig. 7. In the formation of the shell, reactive liquid phase plays important role, since the shell is formed by precipitation of dissolved matter on BaTiOj nuclei. Prolonged sintering provokes chemical homogenisation by solid state diffusion giving rise to inereased permittivity and inereased temperature dependence of permittivity.
3 SINTERING OF MONOPHASE COMPLEX CERAMICS: SINTERING OF Pb(Zro.5Tio.5)03 (PZT)
In analysing sintering phenomena, it is important to realise that chemical reactions may influence the densifi-cation mechanism even in chemically homogenous com-pounds.
In recent years, considerable progress in the processing of ceramics has been achieved by improving the quality of povvders. Attention is paid to vvet-chemical methods of povvder preparation, vvhich assure fine particle size, controlled morphology, high purity and high ho-mogeneity. Improved homogeneity is particularly desir-able in complex multicomponent ceramics.
It is frequently stressed that vvet chemical methods, such as coprecipitation or sol-gel methods, assure chemical homogeneity "on a molecular" level. Hovvever, this homogeneity may be temporarily lost during the sintering operation due to the very nature of the sintering process.
Kuczynski et alls pointed out that the vacancy gradient set up betvveen sintered particles by the sharp curva-ture of the neck betvveen particles can, under favourable conditions, produce considerable segregation in a com-pletely homogenised solid solution. When the diffusion coefficients of constituent atoms in solid solution are different (vvhich is frequently the čase), the neck area be-
comes enriched in the faster diffusing atoms, at least in the early stage of sintering vvhen the vacancy gradient due to the smallness of the radius of neck curvature is large. Such segregation must be a transient phenomenon, since segregation gives rise to a chemical potential gradient arising from the concentration gradient betvveen the neck and regions adjacent to the neck, and acts in a di-reetion opposite to the chemical potential gradient due to the neck curvature. When the radius of curvature of the neck inereases, the chemical potential vvhich causes the dehomogenization decreases. With accumulation of faster diffusing atoms in the neck, the concentration gradient also inereases and the maximum segregation is reached. After this, the chemical potential gradient due to the concentration gradient predominates and back diffusion from the neck to other regions occurs. As a result, homogeneity is re-established.
Kuczynski et al18 demonstrated segregation in Cu-In and Cu-Ag alloys. Mishra et al19 demonstrated dehomogenization of Au-Ag alloy and concluded that the initial neck grovvth as vvell as segregation take plače by surface
Figure 7: TEM photomicrograph of BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 ceramic, sintered 2 hours at 1260°C showing (A) solidified TiOi - rich phase at grain corners and along grain boundaries, (B) Ca-and Zr- modified domain free region and (C) ferroelectric BaTi03 grain core.
Slika 7: TEM posnetek keramike iz BaTi03 - 2 mol % Ti02 - 8 mol % CaZr03 keramike, sintrane 2 uri pri 1260°C. Posnetek kaže (A) strjeno talino, bogato na TiO? med zrni in vzdolž zrn, (B) s Ca in Zr bogat rob zrn brez domen in (C) ferroelektrično BaTi03 jedro z domenami.
IOOt
90-80-
D
H 70i
(Č
60-50-
40
200
400
600 T(oC)
800
1000
Figure 8: Sintering curve of Pb(Zro.5Ti».5)03 compact in air, pressed at 100 MPa. Heating rate: 10°C/minute.
Slika 8: Krivulja zgoščevanja oblikovanca iz PbfZro.sTio.OO,!, stisnjenega s tlakom 100 MPa. Hitrost segrevanja 10°C/min.
D H
300 t(min)
Figure 9: Density of Pb(Zro.5Tio.5)03 ceramics as a function of temperature and time of isothermal heating runs in an air atmosphere. Slika 9: Gostota Pb(Zro.5Tio.5)03 keramike v odvisnosti od temperature in časa pri izotermnem segrevanju v atmosferi zraka.
diffusion, whereas back diffusion occurs by a combina-tion of surface and volume diffusion.
It may be expected, due to capillary forces in the early sintering stage, that the dehomogenization effect occurs in complex ceramic systems as well. It should be particularly pronounced in sintering of nanosized pow-ders, where surface diffusion plays an important role. The segregation-homogenisation phenomenon should be reflected in densification curves.
To demonstrate the effect, we examined the sintering of fine sol-gel prepared powders of lead zirconate - lead titanate solid solution, Pb(Zr0.5Tio.5)03 (PZT)20.
Fig. 8 shows the densification curve of a Pb(Zr0 5Ti().5)03 compact, made of fine powder prepared by alkoxide sol-gel synthesis. The very rapid densification above ~950°C does not support the expected solid state sintering mechanism. Instead, the sudden and steep inerease in sintering rate and the well crystallised grains are indicative of liquid phase sintering. The liquid phase may be the PbO-PZT eutectic above ~840°C; however, the presence of PbO could not be detected by XRD or TEM in the starting powder.
The isothermal densification curves, presented in Fig. 9, show anomalous behaviour in the temperature region 750-800°C. The anomaly is an extended "induetion" period in the densification curves at 750 and 800°C. Anomalous densification of PZT in the initial sintering stage may be explained by preferential diffusion of PZT constituents.
Accumulation of faster diffusing species in the neeks between the particles, triggered by surface curvature, causes an inereased tendency for backward diffusion, sustained by the concentration gradient. Formation of the thermodynamically nonequilibrium phase and its sub-sequent annihilation interferes with the normal densification process, being reflected as an induetion period. Fur-ther densification commences only after neck curvature decreases and the material homogenises again.
At low sintering temperatures, material transport takes plače predominantly by surface diffusion and, when possible. by vapour transport. PZT is known for the high vapour pressure of PbO21. To the author's knovvledge, surface diffusivities of Pb, Zr and Ti ions in PZT have not been reported. Slinkina and Doncov22, using radioactive tracers, measured the effective self-diffu-sion coefficients in polycrystalline 99 % dense Pb(Zro.5Tio.5)03 ceramics. The effective diffusion coeffi-cient of Pb2+ was 5 - 40 times higher than that of O2" and almost two orders of magnitude higher than Dr, and DZr,
Figure 10: TEM micrograph of Pb(Zro.5Tio.5)03 ceramic, sintered at 900°C for I hour. Arrow points to Pb-rich inclusion among 3 perovskite grains (Courtesy G. Dražič).
Slika 10: TEM posnetek Pb(Zro.5Tio.5)03 keramike, sintrane pri 900°C 1 uro. Puščica kaže vključek, bogat s Pb, med 3 perovskitnimi zrni (posnetek G. Dražič).
vvhich were close to each other. Nakamura, Chandratreya and Fulrath21 and Kosec and Kolar24 reported that during the formation of PZT from PbTiO, and PhZrO,, titanium ions diffuse much faster than zirconium ions. The vapour pressure transport of PbO into the necks is also possible. To maintain electrical neutrality, diffusion of cations is accompanied by a simultaneous flow of oxygen ions, or through gas-phase transport. Faster diffusion of Pb and Ti, caused by high neck curvature in the initial sintering stage of fine-grained PZT compacts, causes accumula-tion of Pb and Ti or Zr in the necks, and corresponding depletion of both species in other regions of PZT grains. According to the phase diagram25, PZT may be Pb deft-cient up to 2 mol % PbO; however, PbO does not dis-solve in PZT. The simplified equation derived by Kuczynski et al18, makes it possible to estimate the maxi-mal excess concentration of the faster diffusing species in the neck area.
In simplified form, Kuczynski's equation reads:
2yQC Pc ~ AC kT
where
pc critical neck radius necessary to reverse the neck curvature driven atoms outvvard flow to concentration gradient driven inward flow y surface energy Q mean atomic volume
AC/C relative concentration gradient between neck area
and grain interior k Baltzman's constant T temperature
Neck radius is related to particle diameter (2a) and neck diameter (2x) by expression
Assuming surface energy =1 J/m2, mean atomic volume 1029 m3, and kT (at 1000 K) 1,4.10 20 J, one can estimate the maximal excess concentration of the faster diffusing specie in the neck area. With 200 nm size particles and neck diameter 140 nm, the calculated excess concentration is 8 mol %. In view of crudity of assump-tions, the result seems reasonable.
The proposed explanation of the PZT sintering anom-aly was supported by the following experimental obser-vations:
(1)XRD patterns of PZT after the early sintering stage conftrmed the presence of tetragonal and rhombohe-dral phases, whereas in the starting PZT powder only reflections of the tetragonal phase were present. Pb(ZrxTii-x)03 solid solution exhibits a phase trans-formation from the tetragonal to rhombohedral structure at x - 0.53. The appearance of rhombohedral phase indicates the shift in PZT composition. After prolonged sintering, the rhombohedral phase disap-
peared, confirming homogenisation to the initial composition with x = 0.5. (2) After the early sintering stage, PbO inclusions in the microstructure were observed by SEM and TEM ex-aminations (Fig. 10). Enrichment of PbO in the necks as compared with the grain interior was confirmed by quantitative EDX analysis in the transmission elec-tron microscope. On further heating, PbO inclusions redissolved.
The findings of the work presented disprove the cor-rectness of solid state sintering models assumed for PZT by several researchers. The transient presence of PbO, vvhich forms a eutectic in the early sintering stage, pro-vokes liquid phase sintering. This has several important consequences, for example that microstructural develop-ment in PZT must be sensitive to the heating schedule. Probably the most important result from this investiga-tion is that the dehomogenization phenomenon vvith the transient existence of metastable phases is likely to be a frequent occurrence in the initial stage of sintering of multicomponent ceramics. It is particularly to be ex-pected in sintering of fine povvders vvith a high driving force for sintering.
4	SUMMARY
The critical issues in mastering ceramic processing to assure the reliability and reproducibility of ceramic bod-ies are vvell knovvn. These are (a) appropriate ravv mate-rials (purity, finess, grain size distribution), (b) correct shaping of homogenous bodies vvithout macro defects and (c) suitable sintering parameters (T, t, dT/dt, etc.) But of particular importance for the reliable manufactur-ing process is a knovvledge and control of chemical reac-tions at high temperatures. This in turn demands knovvledge of high temperature phase relations, knovvledge of reaction kinetics and availability of data such as (1) sys-tematic trends in the periodic system, (2) the nature and strength of chemical bonds, (3) thermodynamics (AG,y) and (4) kinetic parameters (diffusion coefficients). In short, to achieve optimal properties of ceramic products and to assure competitiveness demands a high profes-sional knovvledge and appropriate equipment.
5	REFERENCES
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