Alumina : sintering and optical properties
Citation for published version (APA):
Peelen, J. G. J. (1977). Alumina : sintering and optical properties. Technische Hogeschool Eindhoven.
https://doi.org/10.6100/IR4212
DOI:
10.6100/IR4212
Document status and date:
Published: 01/01/1977
Document Version:
Publisher’s PDF, also known as Version of Record (includes final page, issue and volume numbers)
Please check the document version of this publication:
• A submitted manuscript is the version of the article upon submission and before peer-review. There can be
important differences between the submitted version and the official published version of record. People
interested in the research are advised to contact the author for the final version of the publication, or visit the
DOI to the publisher's website.
• The final author version and the galley proof are versions of the publication after peer review.
• The final published version features the final layout of the paper including the volume, issue and page
numbers.
Link to publication
General rights
Copyright and moral rights for the publications made accessible in the public portal are retained by the authors and/or other copyright owners and it is a condition of accessing publications that users recognise and abide by the legal requirements associated with these rights.
• Users may download and print one copy of any publication from the public portal for the purpose of private study or research. • You may not further distribute the material or use it for any profit-making activity or commercial gain
• You may freely distribute the URL identifying the publication in the public portal.
If the publication is distributed under the terms of Article 25fa of the Dutch Copyright Act, indicated by the “Taverne” license above, please follow below link for the End User Agreement:
www.tue.nl/taverne
Take down policy
If you believe that this document breaches copyright please contact us at:
openaccess@tue.nl
providing details and we will investigate your claim.
OPTICAL PROPERTIES
PROEFSCHRIFT
TER VERKRlJGING VAN DE GRAAD V AN DOCTOR IN DE TECHNISCHE WETENSCHAPPEN AAN DE TECHNISCHE HOGESCHOOL EINDHOVEN, OP GEZAG VAN DE RECTOR MAGNIFICUS, PROF. DR. P. VAN DER LEEDEN, VQOR EEN COMMISSIE AANGEWEZEN DOOR HET COLLEGE VAN DEKANEN IN HET OPEN· BAAR TE VERDEDIGEN OP DINSDAG 17 MEl 1977
TE 16.00 UUR
DOOR
JAN GERARD JACOB PEELEN GEBOREN TE RENKUM
Aan Evelien, Marc en Janine Aan mijn oudas
lIe! onderzoek beschreven in dit procfschrift is uitgevoerd op hct Natuurkundig Laboratorium van de N.Y. Philips' Gioeibmpenfabrieken te Eindhoven. lk ben de dirccti.;; van dit iabocatoriul1l erkentelijk voor de mij geboden gelegcnhcid om het onder7.oek in dele vorm te publiceren,
Gaarne wil ik mijn dank betuigen aan allen, die bij de uitvocring van de experimcn-ten betrokken zijn geweest en aan de lotstandkoming van dit procfschrift hebben bijgcdragcn, Ecn belangrijk deel van het ondcrzock is geda~n in samenwcrking met Dr. R. Metselaar. Yoor dCl.e samenwe!'king ben ik hem vcc1 dank verschuldigd. MijI\ dank gaat verder uit naar mijn coHcga's, die het manuscript van dlt proefschrift krilisch hcbbcn doorgelezen.
I. INTRODUCTION
1.1 Properties of alumina . . . .
1.2 Application of alumina as a lamp envelope 2
1.3 The present investigation 2
REFERENCES . . .
4
2. SINrERJNG OF ALUMlNA AND THE lNFLUENCE OF DOPES 5 2.1 Short introduction to the sintering process . 5 .2.2 Sintering of alumina, Review of the literature 8
2.2.1 lnfluence of impurities. . . 8
2.2.2 Influence of grain boundaries. . 'I
2.2.3 Influence of additives on sintering. Defect structure. 11 2.2.4 Influence of additives on the microstructure. , . 12 2.2.5 Conclusion . . . 14 2,3 Influence of MgO on the evolution of the microstructure 14
2.3.1 Introduction . . . . 14
2.3.2 Experimental prooedures . . . 16 2.3.3 ReSults . . . , . . . . , . . , 18 2.3.3.1 Influence of MgO on the dens.ity 18 2.3.3.2 lnfIuence
of
MgO on the grain size 19 2.3.3.3 lnl1uence of extra addition ofeaO and Y103 2423.3.4 Auger spectroscopy 26
2.3.4 Discussion . 29
2.3,5 Conclusion. 31
REFERENCES . . . 32
3. HOT PRESSING OF ALUMINA 37
3.1 Introduction to hot pressing 37
3.2 Continuous hot pressing 38
3.3 Results and discussion 40
3.3.1 Influence of the atmosphere 41
3.3.2 Influence of the hot-pressing parameters 43 3.3.3 Influence of powder properties . . . 45 3.3.4 The mic\"ostructure of hot-pressed alumina 46 3.3.5 Crystallographic texture of the grains 49
3.4 Conclusion. 52
REFERENCES . . . 53
4. OPTICAL PROPERTIES OF ALUMINA 55
4.1 General introduction . . . . 55
4.1.1 Review of the literature 56
4.2.1 Scattering theory. 59 4.2.2 Calculation of the in-line tf<lmmission 63
4.2.2.1 Pores with a fixed radius 63
4.2.2.2 Porc siLC diSlribution 68
4.2.2.3 Inl1ucIlcc of the pore size distribution on the scattering
coefficient . 69
4.2.3 Measurements of the in-line transmission 71 4.2.3.1 IntlL!cnce ,)f the pore siLc . 73
4.2.3.2 Influence of the MgO content 75
4.2.3.3 Int1uence of extra addition of CaO and Y103 77 4.2.4 Dekrminalion of microstructural paramotcrs [rom the measured
transmission spectra. 78
4.2.5 Conclusion. 83
4.3 Transparent and lranslucent alumina 83
4.}.1 Introduction , 8}
4.3.2 Theoretical c()nsiderations 85
433 Experimental part 86
4.3.4 Rc~ults and disC\lssion 89
4.3.S Conclusion, 92
REFERENCES 94
SUMMARY 97
1.1 Properties of alumina
In ceramic literature the term "alumina" is used rather loosely to denote L aluminous material of all types taken collectively;
2. the anhydrous and hydrous aluminium oxides taken indiscriminately;
3. the calcined or substantially water-free aluminium oxides, without distinguishing the phases present, and
4. corundum Dr alpha alumina, specifically.
Gitzen I) in his standard review on alumina uses the term in the sense of the second definition. In the present work we wHl use the term in the SenSe of the fourth defini· tion. Many other phases like gamma, delta, cIa, kappa, chi, rho, theta alumina, all corresponding to the molecular formula Al2 03 , are described in the literature. They are all transition phases and some of them are of doubtful existence or the differences are mainly baSed on somewhat subtle X-ray diffraction differences.
Beta alumina is not a monotropic form but it is a mixed oxide containing alkali or alkaline earth atoms. More information on these compounds can be found in the literature 1,2).
Alpha alumina, the only thermally stable oxide of aluminium, has the corundum structure. This structure may be described as a slightly distorted hexagonal close-packing of oxygen ions. The aluminium ions occupy two-thirds of the octahedral inlerstices, while one third is empty. This structure is extremely stable. The formation energy of alumina is about 400 kcal/mol, only exceeded by the oxides of some rarc earth metals, like La.
Consequences are an extremely low vapour pressure and a high melting point (2045 °C). The deviation from the stoichiometric composition of alumina at normal atmospheric conditi(ms is
<
5 x 10'$ even at 1600°c
~).Alumina has a unique combination of useful electrical, mechanical and chemical jJroperties. Electrically it possesses high resistivity (around 1017 nem at room temperature), good dielectric strength and a low dielectriC loSS factor at high frequencies. These electrical pr<Jpertjes were already applied early in the technical development of sintered alumina as insulators for spark plugs. The electrical con-ductivity is very sensitive to the purity of the oxide.
Mechanically, alumina possesses great hardness, resistan~e to abrasive wear and dimensional stability. The strength properties of ~i!1tered alumina are strongly influenced by the microstructure: porosity, grain size, pore size, second phases. In the application of alumina as cutting tools use is made of its special mechanical properties.
environments. Even sodium does not provoke any appreciable reaction. Alumina maintains these characteristics to high temperatures: the working temperature can be \lP to 1500
°c.
I.! Application of alumina as a lamp envelope
The combination of the above·mentioned properties makes sintered alumina a very suitable material to be used as a lamp envelope in the high·pressure sodium lamp. This lamp is based on the phenomenon that the n{)[mally narrow spectrum of sodium light can be broadened to cover most of the visible spectrum, if the sodium pressure and temperature inside the lamp are increased 4).The lamp produces a golden.white light with a very high efficiency. The material of the discharge tube has to resist the ,mack from ~odium vapour at 1250 QC, and of course the light loss in the visible part of the spectrum should be as low as possible.
To meet the latter requirement a material of very low porosity has to be used. In the fifties t.he problem in sintcring alumina waS the effect. of discontinuous grain growth. The somewhat larger grains with more strongly curved boundaries tend to grow very fast, so that pores become enclosed in the grains and are unable to diffuse out. In 1960 Coble and Burke ., ,6) slIcceeded in suppressing discontinuous grain growth by adding 0.25 wt % MgO to the alumina. This drastically reduces the numbet of pores that Me trapped inside the grains and prolonged Sinterlng then produces a material that transmits light. reasonably wdl, thoush very diffusely. The material is oalled LucaloJ( (derived from translucent aluminium oxide) or DGA (derived from "door-schijncnd gasdicht aluminiumoxidc") and has found wide application. The lamp envelopes made with it proved to be capable of withstanding attack from the hot sodium vapour for more than 20000 hours.
1.3 The present investigation
After this important. practical result a number of fundamental questions remained. How MgO could suppress the discontinuous grain growth in alumina W<lS not under-stood. In numerous publications many ideas about the- role of a sintering additive were developed. A complication in the study of the sintering of alumina is the important role th~t impurilies i.e. not intentionally added ~dditives rMy play. It is very diftkult to prepare alumina on a largo soale with impurities in the ppm range, and many sintering studies have been carried Ollt with impure powders, making the results questionable to say the Iea~l.
Another unanswered question waS how the transmission of light through sintcrod alumina is related \.0 tbe mi;;;rostrllcture. The material Lucalox was introdl\Ced as an essentially pore-free material due to the action of MgO 7). A high density and a large grain size were thought to be necessary for a high light transmission, while the influence of the pore sile was not noticed. The lack of transparency of sintered
alumina was ascribed to guin boundary scattering and the intrinsic birefringence of alumina. Some recent developments point in another direction. It proved to be possible by continuous hot pressing at a relatively low temperature to produce a transparent material with a 25 times lower grain size than the normally sintered alumina 8). This indicate~ that not the size of the grains but the size of the often neglected pores determines the optical properties of dense alumina. Moreover, application of the theory of light scattering showed that the transmission of poly crystalline alumina could be described very well with the parameters pore size, number of pOres and spread in the pore size distribu tion 9).
This thesis is the result of an investigation to answer the above-mentioned funda-mental q\lestions. Dense alumina samples with micrustructures widely varying in grain size and pore size have been prep;ned by normal sinteting and hot pressing.
nle~e samples have been used to verify the statement that pores mainly determine the transmission of light through polycrystalline alumina 8,9). The in t1ucn~e of MgO on the sintering behaviour of alumina was investigated by studying the evolution of the microstructure for samples with increasing amounts of additive 10 ).
Chapter 2 of this thesis gives a survey of the intluence of impurities and additives on the sintering behaviour of alumina. It also describes the influence of MgO on the evolution of the microstructure of alumina.
Chapter 3 deals with the hot pressing of alumina powder. The microstructures of normally Sinlercd and hot-pressed alumina are compared.
Chaptet 4 deals with the uptical properties of polycrystalline, dense alumina. Jt gives relations between the light transmission and the microstructure of materials sintered in the normal way as described in chapter 2, Or hot-pressed as described in chapter 3. The questi(Jn whether a material should be called transparent Or trans· lucent is also treated in chapter 4.
REFERENCES
I) W.H. Gitzen, AhLrnina as a Ceramic Material, Am. Ceram- Soc., 1970. 2) B.C. Lippens, Structure and Texture of Aluminas, Thesis, Delft, 1961. 3) R,W. Vest, citeda~privatecommunlcaticmin JJ. Mills, J.Phys.Chem_
Solids 31, 2577 (1970).
4) L.B. Beijer, H _1.1 _ van Boort and M. Ko~dam, Ughting Design & Application, 4,15 (July 1974).
5) R.L Coble and I.E. Burke, Proc_4th Int. Symp. on the Rw;tivity of Solids, Eds 1.11. de Boer et aI., Elsevier Pub!. Co., Amsterdam, 1961, p_38. 6) RL- Coble, 1. App!. Phys. 32, 793 (1961).
7) C.L brochure on Lucalox, Lamp Glass Department, General Flectri(; Co .. Cleveland, Ohio, J 970_
8) 1 _G _1. Peden, Science of Ceramics 6, XVH (1973).
9) 1 _G.J. Peden and R- Metselaar, 1. App!. Phys_ 45,216 (1974). 10) J.C.J. Peelen, Mat. Sci. Res. to,443
(1975)-2. SINTERING OF ALUMINA AND THE INFLUENCE OF DOPES
2.1 Short introduction to the sintering process
The process to convert powder compacts into dense products is called ~intering. It can be described as the thermally activated proce~s of densification of compacts at temperatures below the melting point. During sintering many changes ()CCUL The essential characteristics are tho increase in strength of the powder compact caused by the formation of bonds between the particles and the shrinkage of the compact as the void spaces between the particles decrease in size and are ultimately almost completely eliminatcd.·Another characteristic is the increase in average grain size. Sometimes very large grains can develop during sintering. This complex process of sintcring is influenced by a variety of factors: the nature of the powder, impurities, intentional additions, pressing conditions, sintering temperature, time and atmo· sphere. Because of the complexities involved, it is impossible to fit the sintering process in onc model.
Looking closer at the process of sintering, we can distinguish three different stages of sintering:
1. DUring the initial stage particles begin to adhere together and necks grow between the particles. At the end of this stage grain boundaries are established and grain growth begin~ to occur. Shrinkage is only a few percent.
2. During the intermediate stage grain growth continues. Continuous pore channels are formed along the grain edge~. The cross·section of these channels dCcreascs gradually until at the end of this stage, at a relative density of about 95%, the channels are pinched.
3. In the final stage dosed pores are formed at the grain boundaries and grain COrnerS. This final stage m~y lead to an almost completely dense material by
removal of these pores, or, alternatively, the pores may be trapped inside the grains when the boundaries break aWilY from them: the problem of pore·grain boundary interaction becomes dominant. P(lre growth will occur as well. Figure 2.1 illustrates the initial stage of the sintering of alumina. It shows the original powder particles and the necks growing between them.
Sintcring requires displacement of atoms, of cations and anions, and this process is determined by diffusion. Several mass transport mechanisms have been proposed as contributing to sintering, some processes producing shrinkage and others producing no shrinkage, In all cases matter flows into the neck between the powder particles. The six mechanisms for this flow of matter, as ccmsidered by Ashby I), have been Collected in table 2.1. In figure 2.2, which is a plane section through an assembly of three spheres, the possible diffUSion paths are shown.
Fig. 2.1. Scanning electronmicrograph of the initial stage of sintering process, showing neck growth between the particles. White bar is 1 lim.
GRAIN
\'BOUNDARY
Fig. 2.2. Three-sphere model showing the six transport paths from table 2.1 (page 7). In all
Table 2.1
Mechanism No. Transport path Source of matter Sink of matter
I Surface diffusion Surface Neck
2 Lattice diffusion Surface Neck
3 Vapour transport Surface Neck
4 Boundary diffusion Grain boundary Neck
5 Lattice diffusion Grain boundary Neck
6 Lattice diffusion Dislocations Neck
Mechanisms only leading to rounding off of the pores without shrinkage include evaporation of matter from the convex surface of the particles and condensation in the neck between the particles. It is doubtful whether this method of mass trans-port is very imtrans-portant in the sintering of an oxide with a low vapour preSSure such as aiumina. Another mechanism by which matter can be transported without shrinkage is surface diffusion. This can be a competitive process with other processes that produce Shrinkage, especially at low sintering temperature.
The shrinkage of alumina can best be explained by assuming atomic diffusion along two different paths of diffusion. Volume diffusion or lattice diffusion involves the transport of material from the grain boundary tluough the lattice int(1 the neck between the particles, while in grain boundary diffusion the grain boundary itself might be the path along which the atoms move. Both transport mechanisms will contribute simultaneollsly to the sintering process.
Johnson ~) has proposed sintering rate equations for conCUHent volume, grain boundary and even surface diffusion.
The driving force for sintering is the excess free energy of a powder compared with that of a dense material. The decrease in free energy occurring on sintering a powder of 1 pm particle size is about 1 cal/g. Part of the initial surface energy is used to form grain boundaries. This grain boundary energy delivers the driving force for grain growth. The sintering process cannot be treated in all details here. Suffice it to remark that it is a great step from a two or three particles model to the kinetics of densification of real powders having a particle size distribution, no spherical shape, etc, Therefore, discrepancies between theory and practice are often found. In ionic compounds effects from a difference in diffusivity of cations and anions have to be considered. This has often been underestimated in the ceramic literature. Ready 3) and Reynen ~) have pointed to the effect of non "stoichiometry on sintering, which sometimes can be very large. Review papers on sintering are included in the list of references S-s).
Some aspects of importance to an understanding of the sintering behaviour of alumina are treated in the next section, These aspects include the influence of impurities and grain boundaries on sintering, the defect structure of alumina and the
inl1ucncc of additives on the Il1icro~tructurc of alumina_ Section 23 is devoted mainly to the ini1uence of magnesia on the ev()luti()n of the microstructl1[e of alurnina.
2_2 Sintering of alumina. Review of the literature
2.2.1 In/luencro oIimpuritfes
Since sintcdng is dependent on diffusion processes, it is important to know what kind of point defects are present in alumina and in what concentration, In the caSe of intrinsic diffu~ion the fraction of point defects is determined only by the temperature and the energy, hj; it costs to create the defects. Compared with NaCl, for instance, alumina is a rather unknown material. In the case of alumina there is still doubt as to whether a Schottky or a Frenkel type defect structure is present (see sec- 2_2_3) and the energy for defect formation is unknown. Oishi and Kingery 9) measured
jS 0 diffusion both in poly crystalline arid in monocrystalline Alz 03 , It can be deduced from their results that it costs an energy of 20.5 eV to create 2 Al vacancies
~nd 3 0 vacancies.
Fryer 10,1 t ) analyses his results on pressure sinlcring of undoped alumina in a different way _ He believes that the activation energy he observes can be ascribed to intrinsic diffusioIl arid he arrives at a Schottky formation energy of 10 eV_ His analysis taken over by others 12), however, is very doubtfuL It is unlikely that his starting powder, Linde A, permits intrinsic behaviour at temperatures as low as
1200 - 1400
or-
In a recent publication Dienes et aL 13) give a calculated value of 28.5 eV necessary to create a Schottky quintuplet or 20 eV to Create a Frenkel pair in the Allatlice_ B()(h values arc very high compared withii/=
2.3 eV for NaC!. As a conScqllcrlcc, tho concentration of intrinsic vacancies will be extremely low, even at high temperatures. With the value of 20.5 eV the fraction of vacant calion sites at 1600°c
can be estimated as 10-10, 1'his is much lower than the impurity level of the purest alumina!It is dear from this that only extrinsic diffusion behaviour carl b~ expected: the vacancy conccntJ'ation will be controlled by the number and nature of impurities present. ~nd will be independent of tempel'ature. Add to this that many impurities have a very limited wlubilily in alumina Blld thet'efore, they will be present as a dispersed second phase. The presence of these impurities in not well characterized starting materials may explain the many contradictory reslllts, which typify the investigations of diffusion in alumina, such as sintcring behaviour, electrical conductivity etc, (see figure 23).
As a consequence, alumina is by no meanS a suitable model material for sintcring research 8,14).
Anion impurities can be as effective as cation impurities 15), although they have usually been neglected.
4 5 10 15 <0 10'{7'rK)
Fig. 2 .3. R~po,td lit¢,at"," data fOr tho electrical cunciuctivity of alumina (f,om ref. 16).
2.2.2 Influence oigrain boundaries
Sintering requires displacement of both cations and anions by diffusion. The main diffusion paths lie in the volume of the grains and in the grain boundaries. The relative effectiveness of both will be affected by the microstructure. It is very difficult to assign an exact grain boundary width. Mistler and Coble 17), analysing grain growth and sintering data, give a value of about 100 A.
At a grain boundary the atoms are less densely packed and this may result in a lower activation energy for boundary diffusion and a greater atomic mobility.
Oishi and Kingery 9) demonstrated that the oxygen ion self-diffusion coefficient is enhanced by the presence of grain boundaries. Paladino and Kingery 18) measured the diffusion of 26 AI onJy in polycrystalline alumina and they state that diffusion of AI is not influenced by the presence of grain boundaries. The combined results are given in figure 2.4. Based on these results and on an analysis of the apparent diffu~ion coefficient in sintering and creep experiments, Paladino and Coble 19) conclude that volume diffusion of Al is the rate-controlling process in the sintcring of fine.grained alumina, while 0 diffuses more rapidly along grain bound-aries. For grain sizes larger than 20 p,m they expect the diffusion of oxygen to be rate·controlling.
In a later analysis Mistler and ~oble 20) expect the change-over from Al to 0 rate control to take place at a grain size of 5 .urn. It is very difficult to deduce the
Ternll"!"t'11.'r'$II.';'
19~O le.SCl liM 1650 I~~O 1450 13!:£1 1~'5O
Fig. 2.4. Combined ,osul" of the oxygen diffusion measurements of Oishi and Kin~ry·) and
[he aluminium diffusion measurements of Pa.Lldino and Kingt;;:t'y 11').
kinetics from calculated values of the diffUSion coefficient or activation energies for a process such as sintering, because the values of Do and DAI are S() similar. This may explain why the sintering mcchanism has been ascribed to volume diffusion and grain-boundary diffusion 21). Enhanced oxygen diffusion along grain boundaries has also been reported for other oxides, like MgO 2~), Fe203 23) and CoO 24). Enhanced cation diffusion has been reported for U02 2S).
Grain boundaries may int1ucnce the diffusion ofions in another way. It has been shown 26 ,l7) lhat the concentration of defects near vacancy sources and sinks like grain boundaries differs from the bulk concentration.
The reasOn is that the energies needed to creat~ anion and cation vacancies or, in
the case of a Frenkel defect strllcture, vacancies and interstitials, are different. The consequence is that the grain boundary may carry an electric charge resulting from the presence of excess ions of one sign. This charge is compcnsated by a spacc-charge cloud of the opposite sign adjacent to the boundary. The distance the space charge region extends in to the crystal is assumed to be 20 to 100
A.
The sign and the magnitude of the electrostatic potential at the boundary are determincd by both the solute concentration, if present, and the temperature 2~). This will be true especially for alumina with its low concentration of thermally indl1ccd lattice defccts.lGngery 28) ha~ demon~trated the existence of a boundary charge in a number of oxides by observing the bowing of grain boundaries in an electric field at 1650 °C. He observed a positive sign of the boundary charge in an alumina sample containing about 1000 ppm MgO. The charge of the grain boundaries may possibly explain the preference for grain boundary diffusion of only oxygen 29) . However, as the con· cent ration and valence of impurities affects the charge of the boundaries, one has to be very careful with conclusions. There is hardly any knowledge of these effects on grain-boundary diffusion.
2.2.1 Influence of additives on sintering. Defect strncture
Many investigators have studied the initial stage of sintering of deliberately and un· intentionally doped alumina to see, what defect model might explain the sintering data. Not only the sintering behaviour has been studied, also the steady state creep behaviour, thermal grooving and electrical conductivity, which are all diffusion-controlled processes. Failure to recogni~e the importance of purity, concentration of additives, tho presence of second phases, the intemlation b"tween ~int"ring and grain growth, and different sintering mechanisms like grain boundary and volume diffusion, led to great confusion in the literature (see e.g. fig. 23).
Until 1971 the basis of every explanation was a Schottky defect structure, i.e. with AI and 0 vacancies as the main native defects. Addition of 4+ or 5+ cations would create AI vacandes, addition of 2+ cations 0 va~ancies. In 1970 Bagleyet a!. 30) found an increase of the sintering rate of alumina, when doped with Ti02 • In 1972 McAllister and Cutler 31) reported that addition of both MnO and Ti01 to alumina increased the rate of thermal grooving. In 1973 Rao and Cutler 32) also found that addition of FeO increased the sintering rate of alumina.
These observations are difficult to explain with a Schottky type defect structure, since the assumption of Schottky pairs as the predominant defects leads to the conclusion that the sintering rate of Fe.dopcd alumina is controlled by the diffusion of 0 vacancies, whereas the sintering rate of Ti-doped alumina is controlled by the diffusion of A1 vacancies. It is diffucult to understand such a cnange in the rate· controlling species. If a rapid diffusion path for the oxygen ions is assumed to explain the sintcring data for Ti"doped alumina, it is difficult to explain why the sintering rate of alumina is enhanced by the presence of divalent cations.
In 1971 Brook, Vee and Kroger Ja), in an investigation of the electrical conductivity of doped and undored alumina, were the first to suggest the existence of a Frenkel· type defect structure, that is to say AI interstitials and Al vacancies as the main native defects. An A1 interstitial is a possible defect in view of the octahedral holes in the oxygen sublattice, one third of which are empty. Rao and Cutler 32) tried to verify this by measuring the power by which the volume diffusion coefficient depends on the total Fe content. With the same intention Hollenberg and Gordon "") studied to what power the creep rate of Fe-doped alumina depends on the oxygen
partial pressure. They all found that the scatter in their data did not enable them to determine the rate-controlling species Al
j ..
orVo _
The calculations of Dienes et a1. I') of the formation energy of defects in alumina favour the idea of a Schottky defect structure. In another calculation thc formatioll energyor
Schottky defects is estimated t() be greater than that of Fre)l kel defects 35) The rapid ratc of reduction of Mg-doped alumina compared with that of Ti-dopcd alumina led Cox 36) to ~uggest the presence of the very mobile interstitial cation in Mg-doped samples. Recently Dutt and Kroger }1) concluded fr<)m the oxygen pressure depen-dence of the electrical condl1ctivity of Fc-doped alumina that AJi - -must be the
majodty defect ~pecies_The assumption of Frenkel defects is attractive, because it makes the sintcring and creep data consistent. The sintering rate of Fe-doped alumina can then be explained
as being due to diffusion of Al by an interstitialcy medlanism. whereas the sintering rate of Ti-dop~d alumina is controlled by the diffusion of Al through a vacancy mechanism. ThllS in either case the sintering rate is contrulled by the djffu~ion of Al ions with no chang~ in tbe rate-controlling species.
For MO the incorporation reaction becomes:
(2.1) and hr MOz :
3(MOz )s ,., 3M
A
i + VAi'
+ 600 (2_2)2.2.4 Injluem:f! o/addiliv&s all the microstmcture
Apart from the theoretical considerations of the previous section, another goal of additives is to lower the sintering temperature or to produce a desired microstructure by controlling the relative rates of the competitive reactions, which occur during the h~ating of a powder compact. The critical stage during Sinlering is the final stage. Sint~ril1g may thm lcadto an almost completely deme material, if the pores remain on the grnin boundaries and are able to migrate with the grain boundaries_
Alternatively, grain boundary migration may proceed too rapidly, trapping the pores inside the grains. This will never re~ult in a very dense materiaL
This so called discontinuous grain growth (fig. 2.5) occurs when undopcd alllInina is sintered at high temperatures.
The migration rate of grain boundaries is strongly influenced by the number, size, geometry, and mobility of the pores, which exert a drag on the boundaries. A gimilar drag can be exerted by impurities or additives, and the presence of a sotid or liquid phase ',t the boundaries has also a strong inf1llence ()O the grain boundary mobility. A basis for an understanding of the conditions, which determine pore migration and grain boundary impurity drag was bid by Brook 38) and extended recently by Carpay 39)_
. ' "
,
..
.
.
.
,.
'..
," -.; , . !.
:,', , .... " :.
.
: '...
..
:~~
- : '.. ,I ., ,'. .. .. '. . ~-
.
.
.
,: " " .''
.
,.
.
.':
.~. ': . : ~".' -',., ,.
',.
.
:' . .... ., . ..
..
Fig. 2.5. Undoped alumina sintered for 10 h at 1850 °C, showing discontinuous grain growth.
The effect of many additives, such as oxides, fluorides, metals, has been investigated to overcome the discontinuous grain growth in alumina. A summing up can be found in ref.
40).
In a very recent paper Kovatschev 41) describes the influence of16
dif-ferent additives. Cahoon and Christensen 42) and Coble and Burke 43,44) were the first to demonstrate that discontinuous grain growth could be inhibited by adding0.25
wt%
MgO to the alumina powder, reSUlting in a densely sintered product. Warman and Budworth 4S) even published criteria for the selection of additives to enable the sintering of alumina to proceed to full density. Their two criteria are: volatility of the additive to get a uniform distribution on the grain boundaries and insolubility of the additive in alumina for it to remain on the grain boundaries. They found that ZnO, NiO, CoO and Sn02 gave the same result as MgO.Rossi 46) and Rossi and Burke 47), reporting their work on finding additives other than MgO that would permit sintering of alumina to full density, were unable to confirm earlier work, except for NiO. They observed that many additives like CaO, SrO, BaO, Y203 and Zr02 when sintered at high temperatures, yielded a micro-structure similar to that shown in figure
2
.
6
.
This microstructure is characterized by clusters of fine pores at the centre of about
150 11m large grains. Surrounding each fine pore cluster is a region which is free of
fine pores, but it may contain very coarse pores, 10 -
20
11m in diameter. The struc-ture also shows an intergranular phase which is presumed to have been liquid at the sintering temperature. All these additives form eutectic compositions melting below the sintering temperature. The authors make this structure plausible, proposing a sequence of events leading to three generations of grains.Fig. 2.6. Alumina doped with 0.3 wt % Y 20, sintered for 3 hat 1900
°e,
showing poreclusters in grain centres and pore-free regions with occasional large pores (from ref. 47).
The above illustrates the strong influence of the presence of a solid or liquid phase at the grain boundaries on the evolution of the microstructure. Less work has been done on the influence which powder characteristics have on the occurrence of dis-continuous grain growth in alumina. Agglomerates in the powder and density fluctuations in the powder compact can strongly influence the microstructure 48 -51).
2.2
.
5 Conclusion
Densification and grain growth during sintering of alumina are influenced by a
complexity of factors. Contradictory results reported in the literature are often
caused by using poorly characterized starting powders. The diffusion during sintering is impurity-controlled. The kinetics of the sintering process are strongly influenced by the presence and the charge of the grain boundaries. Sintering experiments with deliberately doped alumina can be explained consistently, if a Frenkel type defect structure is assumed. Impurities or additives play the leading part in the evolution of the microstructure.
2.3 Influence of MgO on the evolution of the microstructure
2.3.1 Introduction
Without doubt MgO is the most intensively studied of the additives that promote
the sintering of alumina to high densities. The great improvement in the optical and
mechanical properties of the sintered material justifies all these investigations. Ryshkewitch
52)
was the first to suggest MgO as the ideal additive for the sinteringof alumina, this system having a eutectic point at 1925 °C, well above the practical sintering temperature of alumina.
His explanation is that a dispersed phase will form a system of barriers between the oxide particles,inhibiting their mutual contact and coalescence_
Coble 4~,44) demonstrated that alumina could be sintered with the aid of 0.25 wt % MgO to a nearly pore-free condition. Other necessary conditions are the Use of a carefully selected powder and sintering in an atmosphere of a gas that can readily diffuse out of closed pores in the partially sintered compact 53). Coble's result marked thc first achievement of a nearly pore· free mlltedaL Since that time other materials have been developed with extremely low porosities, e.g. MgO 54),
La-doped Ph (Zr .Ti)03 for opto-eiectronic devices '5), (Ni,Z)) )Fe~ 0 4 for recording heads 56), garnets for magneto-optical applications 57), Th-doped YZ03 for
poly-crystalline laser rods ,8), MgAl2 0 4 ,9).
MgO is often said to play the role of grain growth inhibilor. However. it does not inhibit normal grain growth, as will be shown in sec. 2.3-4, but it only inhibits the discontinuous grain growth. The result is that pores are nOt trapped inside the grains, but can disappear by diffusion of vacancies_ Several theories have been put forward to explain the role of MgO, Ryshkewitch 52) mentioned the physical separation of the alumina grains. The action of second-phase particles of spinel located preferen-tially at grain boundaries in pinning these boundaries and decreasing the mobility has generally been realized 60-63). Jorgensen and Westbrook <S4) assumed that MgO dissolves preferentiaUy in the grain boundary region, called solute segregation_ This would cause a decrease in the grain-boundary mobility. Although they could not give convincing experimental evidence, this assumption has found the most wide-spread belief. Haroun and Budworth 63,65) conclude that MgO distributes itself throughout the specimen by vapour-phase transport and then remains at the grain boundaries as a tUm around the grains. However, an additive in solid solution can also suppress discontinuous grain growth by reduction of the grain-boundary mobility or increase of the pore mobility_ This wM earlier pointed out by Coble 44), and later by Rossi and Burke 4'), although they were unable to give the precise mechanism. The goal of the investigation to be described in this chapter was to study the evolution of the microstructure of alumina with increasing amounts of MgO, ~tarting far below the solubility limit. This WOrk is only possible if the available alumina powder has an impurity content in the ppm range. Up to now no work with such low impurities has been published_ However, only then it is possible and significant to use MgO as an additive in concentrations as small as 50 ppm. The density and grain size of the sintered alumina will be given as a function of the MgO content. Auger electwn spectroscopy (AES) is a very sensitive technique to see whether different concentrations of additive are present in the boundary and bulk regions. Marcus and Fine 66) in 1972, who were the tlrst to apply this technique to look for an enrichment of MgO at the grain boundarie:j>, could not detect any MgO at all, but to their surprise they found a strong Ca cnrichmenL Before them Tong and Williams 67)
reported an increased Mg concentration at the grain boundaries_ They used the technique of spark source mass spectrometry, which makes craters about 0.3 J.lm
deep at the (;;rain boundaries. Analyses of the data of ref. 66 yielded a rather high detection limit for Mg, so it seemed worth while to repeat their experiments. During this work Tayl(Jr et al. 68) reported that they found an enrichment of Mg
at the grain boundaries by a factor of Z-They used the technique of X-ray photo-electron spectroscopy (XPS), which has a somewhat greater sensitivity than
AES-23.2 Experimental procedures
All experiments described in this chapter were carried out with an alumina powder, obtained from Rubis SyntMtique des Alpes, code A 15 RZ. This powder, made by calcination of alum, consists of 85% Ct-AI20" with a mean particle si7.e of 0_3 pm and 15% r-Ab 03 with a mean particle sir.e of 0_02 /lill. The decon1p()~iti()n reaction during calcination is accompanied by sintering, in which process luge
agglomerates can grow_ Since these agglomerate~ can have an unfavourable influence on the sintering process, we used a p()wder deagglomerated by the manu factureI', which reduces the mean sizo of the ag(;;lomerates from 15 pm to 4 ).Im-Figure 2.7 is an eiect.ronphot.omicrograph of this powder, showing the particles of Ct-and
r-AI2 03 -The specific surface of the powder determined by the BET adsorption method is 15 m2 (g. The impurity content was analysed spectrochemically except for the impurities Na, K and Ca, for which the spectrochemical method is wther insensitive_ Their cuncentration was determined by atomic absorption analysis. The results can be summarized as follows:Fe 5, Ga 4, Mg 0.8, Si 30, Na 10, K 30, Ca .;;; IO ppm.
The additive MgO was added as a solution of Mg acetate-4 H10 (Merck, reagent grade) in absolute alcohol to a suspension of the alumina powder in the S'lme alcohoL This suspension was dried while continuously stirring. After further drying the powder was sioved and prepressed in a plexiglass die to avoid CQntamirl3tion into pellet.s 20
mm
in diameter and 10 mm thick. These pellets were then isost.atically pressed at 100 MNhn2 and preheated in oxygen at700°C to decompose the acetate to oxide. The sinr.ering took place in a high-purity ahlmina tube heated in a molybdenum resistance furnace.
Great care was taken to measure the apparent densities of the sintered spc(;imcns because of the small difference with the theoretical density of alumina. The X-ray density was calculated frum our experiment;llly determinod lattice parameters a'" 4_7585 ± 0.0002 A and c = 12_9942;1; 0.0005 A to be 3_9859 j; 0.0005 g/cm3• The density was determined by the method of Prokic 69), which takes into aCl;()Unt the counterbalancing force of the air. The density of the sarnple at temperature r,
D1,
can be calculated fromMt Di - Do
"t
; ,
0'
...
.
O.llJm
Fig, 2,7. Elec[ronpho[omicrograph of alumina powder used in [he sinrering experimenrs, showing
where MI is the weighl of the sample in air, M2 is the wm of the weight of the sample in air and the weight of .a thin nylon thread fixed to a hook of the balance and submerged in distilled and boiled·out water of density
DJ .
M,ds the weight of the sample tied to the nylon thread and submerged in the water, Do is the density of the air, dependent on temperature, barometric pressure and relative air humidity-The density of water was checked using the density of a highly perfect silicon single crystal as a reference 10). An accuracy in the apparent density of ± 0_0005 gfcm' can be achieved. This means an inaccuracy in the pOf()sity of ± 2.5 x 10.4. The microstructure of the sintered samples was studied by microscopic observation of polished and etched samples. Polishing was done on a vibratory poliShing machine with diamond paste_ Etching was carried out by heating the polished samples at 1450°c
for 1.5 h (thermal etching). The mean grain size was calculated using the relation 11):G"'
1.5T,
whereT
is the mean grain intercept of random lines Dn photo-graphs of the polished and etched cfoss-sections- The value obtained in this way is a good approximation if the size distribution of the grains is not too wide, i_e_ if no discontinuous grain growth has occurred_For ttie Auger experiments a commercial Auger spCdrometer was used, consisting of a Physical Electronics Cylindrical Mirror Analy:.:er and a standard Auger ultra-high vacuum system (VItek-Perkin Elmer). To prevent charging effects ()n alumina we used the grazing inddcncc electron gun, the diameter of the electron beam being about 50 Mm- In-depth profiles could be measured by means ofaXe·ion gun. Rods 2 mm in diameter were prepared from the sample and fractured in a specially constructed break apparatus under ultra-high vacuum (10'10 Iorr). The fresh sllrfacc was then analysed with the Auger spectron1eler.
2.3.3 Results
2.3,3.1 Influence of MgO on the density
Sintering experiments were carried out on alumina powder compacts with increasing amounts of MgO dope, starting with the undoped alumina fOf comparison_
Figure 2.8, Cllrve A, gives the density of the specimens after sintering at 1630"C for 1.5 h in a humid H2 atmosphere (dew point 20 ·C) to prevent yolatilization of MgO. It is seen that doping with 50 ppm MgO is already sufficient to increase the density of alumina substantially_ The density reaches a maximum at 300 ppm dope level; at larger dope concentrations a deCrease in density is found, Second.phase particles can be detected in the specimens with 300 ppm MgO and more with the optical microscope_ This is in agreement with the solubility data of Roy and Coble 72), They found for the solubility of MgO in alumina in vacuum:
where Xis the atomic fraction Mg/Al and Tthe absolute temperature. According to this equation the solubility of MgO in alumina at 1630
°c
is 250 ppm.The highest density is thu~ obtained when the amount of MgO corresponds to the solubility limit in alumina. With electron microprobe analysis Mg and Al could be detected in the second-phase particles. The MgfAl ratio in the spinel phase was about 20 % lower than in stoichiometric spinel. This corresponds to the high solubility of Al103 in MgAl204 at high temperatures 73).
,...
....
';j; L.OO c'"
-.::J 3.96 196 3.94 lS2 3.90o
50 100 300 A: 1.5h 1630'C B: 1.S h 1630'C +'0 h ,850'C 1000 3000 ppm MgOfig. 2.8 lnflu~no. of MgO content on the density of .inteNd alumina.
All specimens underwent a second heat treatment at 1850
°c
for 10 h, This resulted in an increase in the density of all samples; see figure 2.8, curve B. There is no distinct peak now because all samples with a MgO content between )00 and 1000 ppm reach nearly full density. Higher MgO contents result in a decrease of the density, This decrease is greater than can be ex:plained by the slight decrease in theoretical density due to the presence of the second phase. According to eq. (2.4) the solubility limit at 1850°c
is 1350 ppm MgO,2,3.3,2 Influence of MgO on the grain size
The microstructure of the speciJ1len~ was evaluated and the mean grain size deter-mined for the various amounts of MgO dope. Figure 2.9 gives the result after the sintering treatment at 1630
°c
for 1,5 h, It appears that MgO promotes grain growth until the solubility limit is reached. The mean grain siu decreases when more MgO is present, probably because of the dragging effect of the second phase particles on the grain boundaries. In the single-phase region we observe a distinct increase in grain size with the MgO content together with an increase in density.E
2,8/
Il!) QJ7
N ';jj c/.
nI6
I-. /
l!)5
t.
1.5 h 1630·C
~o
,r' --~~~--~---50 100 300 1000 3000 ppm MgOFig. 2.9. Influence of MgO content on the mean grain size of alumina sintered for 1.S h at
1630°C.
Pores lying on a
grain
boundary obviously influence grain growth more or less like
a
second phase
.
The fewer the pores, the more freely the grain boundaries can move.
Fig. 2.10. Undoped alumina sintered for 1.S h at 1630 °c.
Figure 2.10 is a photomicrograph of the microstructure
of
undoped alumina
specimen. No discontinuous grain
growth
has occurred at the temperature
of
Fig. 2.11. Alumina doped with 300 ppm MgO sintered for 1.5 h at 1630 °e.
Figure 2.11 shows the microstructure of the sample doped with 300 ppm MgO, while figure 2.12 gives the microstructure of the sample doped with 3000 ppm MgO. This last picture shows the many second-phase particles present, mostly at the corners of the grains.
Fig. 2.12. Alumina doped with 3000 ppm MgO sintered for 1.5 hat 1630
0c.
During the second sintering period at 1850
°c
for 10 h discontinuous grain growth..
'.
.
',..
", "..
/.
.
•
.
/
. ... .:
..
.
~,;.:
;
'.
~
(
..
' . . : ' = /'•
.
.
:', '.
.
.
..
.
.
,.
.
'.
..
.
.
.
'
.
,.
, / • '0 ", o • ': ~.'
.
"0 • ~. O • • ', ', '0 '.
...
".
.'
.
.
".
.
.
Fig. 2.13. Undoped alumina sintered for 10 h at 1850 cc.
The microstructure contains many large grains of 100,um and more, while most
of
the pores are trapped inside the grains. The
specimen
doped with 50 ppm MgO
contains cracks
and
shows a non·uniform microstructure
:
many very large
grains,
even
larger than 100 ,urn, are visible within a matrix of much smaller grains. Grain growth
occurs in a
stage at
which most pores have disappeared
.
The large grains now are
essentially pore-free, in contrast to the
case of
undoped alumina
.
Large pores are
found between the large
and
the
small
grains
.
The same holds, but to
a
much lesser
extent,
for the sample with 100 ppm MgO. This sample contains no cracks, however.
The non-uniform microstructure is illustrated in figure
2
.14
.
-'
50pm
~ ) /,
.•
,"
"
I
,.
'\.. • . .J.
:
r-~'.
~. \ 'j..
t
:
".
\ >-.... t•
.
,
-".
r,
':.
". .n 1 ~~~ -"~' ',1 •. / '( .( " -: . ~ .:...
~ \ { .~,
. . ~.
....
"
: - /\ ( ' .. .. ".1; .... ... /...-: .. /1 r".J . • r~Fig. 2.14. Alumina doped with 100 ppm MgO sintered for 10 h at 1850 cc, showing non-uniform grain growth.
All samples with more MgO show a very regular distribution of grain sizes. Illustra-tions are given in figure 2.15, which shows the microstructure of the sample with 300 ppm MgO, containing occasionally a second-phase particle, and in figure 2.16, which shows the microstructure of the sample with 3000 ppm MgO, containing large coarsened spinel particles.
Fig. 2.15. Alumina doped with 300 ppm MgO sintered for 10 h at 1850 °C.
The mean grain size of these samples is given in figure 2.17. Essentially the Same dependence on the MgO content is found as in the previous case after sintering at 1630 °C; increasing grain size in the singlo-phase region, decreasing grain size in the second-phase region . .---. E 2:-I~ <I> 25 N III c:
.§
20
(!) 15 10 0 50 100 300 15 h 1630·C -10 h 18S0 'C1000
3000 ppm MgOfig. 2.17. Influence of MgO come", on th~ mo<ln grain size of alumina aft~r a seeond 5intering period at 1850 °c for 10 h.
Although parI of the second phase can originate from precipitation during cooling, the solubility rate of the spinel particles, present at lho grain corners, is apparently very low.
2.3.3.3 Influence of extra addition of CaO and Y 2 O~
Many combinations of MgO and some other oxide as additive have been tried to ~ee
whdher the pwpertics of sintered alumina could be improved. The patenlliterature in particular gives many examples of such investigations. Most attention has been paid to the combinations MgO + CaO 14) and MgO + Y 203 75), To see whether the extra addition of CaO or Y20] could enhance the transmission of alumina sintered at relatively low temporatures, we carried out $inlcring experiments starting with the alumina powder already doped with 300 ppm MgO, To this powder 0.1 wt % Y20, was added as an emulsion of very pure Y 1 O~ in absQlute alcohol. The SO ppm
extra CaO was added by immersion of a compact in a solution of Ca acetate. These pellets were already isostatically pressed and presintercd at 1250
°c
in oxygen. This treatment increases the relative density tll 56 % and gives the necessary mechanical strength, Knowing the constant amount of water these compactsabsorb, one can calculate the required concentration of the Ca acetate solution to add 50 ppm CaO.
Comparing the microstructure of the samples with and without
Y2 0
3 , it can beconcluded that Y 203 has an inhibiting effect on grain growth. But also the density
is lower after the same sintering treatment. The spread in grain size distribution is
larger without the occurrence of discontinuous grain growth. Electron microprobe
analysis reveals solid phases of two types: a phase containing only Mg and Al, often
in the form of needles, and a phase containing only Y and Al, probably Y 3 Als
0
12particles with a diameter of 1· to 2 pm, lying on the grain boundaries. This phase
must be responsible for the decrease in grain size compared with the samples without
the extra Y2
0
3 dope.Addition of extra CaO causes a much more irregular microstructure (Figure 2.18).
Although the mean grain size is not much larger, the grain size distribution is very wide. The same is true of the pore size distribution. The density is equal or stays a
little behind. The phase diagram of the system Al2
0
3 --CaO predicts low meltingeutectics. The straight grain boundaries in the microstructure point indeed to the
presence of a liquid phase
76).
With electron microprobe analysis two phases canagain be distinguished: the well-known spinel phase with extra Alz
0
3 in solidsolution, and a more irregularly occurring phase containing Mg, Ca and AI,
presumably CaO . 6 AI203 with part of the Ca atoms substituted for Mg atoms.
Both phases are clearly distinguishable by the different colour of the luminescence,
'
.
\.
/,
/"."
~- \Fig. 2.18. Alumina doped with 300 ppm MgO and 50 ppm CaO sintered for 10 h at 1850
°c,
which occurs when the sample is irradiated with a broad defocussed electron beam (figure 2.19). The colour of the spinel phase is green, that of the Ca-containing phase is black, while the main phase is violet-blue.
The results of the transmission measurements on these samples and on the samples
containing only MgO are given in section 4.2.3.
Fig. 2.19. Luminescence of the same sample as in fig. 2.18 in the electron microprobe. The
colour of the spinel phase is green, that of the Ca-containing phase is black.
2.3.3
.4
Auger spectroscopyAuger spectroscopy was used to investigate the possible grain boundary segregation of MgO and CaO in alumina.
Samples containing 1000 ppm MgO and 45 ppm CaO were used for these experi-ments. The bulk concentrations were verified by atomic absorption analysis. The
powder compacts were sintered at 1630
°c for 1.5 hand 1850 °c for 10 h in a
humid
H2
atmosphere. Rods 2 mm in diameter were prepared from the sample.Figure 2.20 is a scanning electronmicrograph of a fracture surface, which shows that the fracture is mainly intergranular.
Other rods were fractured under high vacuum. The Auger spectrum of the fresh
fracture surface is shown in fig. 2.21. Besides the peaks of Al and
0,
the Ca peak at292 eV is very clear in the spectrum. The Ca concentration at the surface could
be fixed
77)
at about 6 at%.
This means an enrichtment of the surface by afactor of 1000.
Signal averaging techniques were used to decrease the detection limit for Mg. A part of the resulting spectrum is shown in figure 2.22.
Fig. 2.20. Scanning electronmicrograph of a fracture surface of an alumina sample with 1000 ppm MgO and 45 ppm CaO, showing mainly intergranular fracture. White bar is 10 /-1m.
Auger signal
i
- 3 x Cal2921 AI AIo
_ Energy I eV I ! I I I ,o
200 400 600 800 1000 1200 1400 1600Fig. 2.21. Auger electron spectrum of a fresh fracture surface of the sample in fig. 2.20.
Quantization of the Mg peak yields a Mg concentration of about 0.1 at
%.
Thisagrees with the analysis of the original powder. No grain boundary enrichment
could thus be established. It should be remarked that the accuracy of the Auger
analysis is a factor of 2. A further remark should be that part of the MgO may have evaporated during the sintering treatment. In fact, analysis of the sintered
piece showed that the MgO content had decreased by 20
%.
A minorII
~
~
Port Cit tne group of hnliJs, of AbO)at ~ 1350 .v
_E~ergy I.vl
1040 1200 13S0
Fig. 2.22. Part of the A1)g~r el~c[rOIl spectrum of the ~<1.lYIplt!" in figure 2.20, obtl:iined by means
of .ign.l .v<rag;ng techni<jues.
By simultaneous sputtering of the fracture ~urfacc with Xe iom and measuring the Auger spectra, the profile normal to the surface was investigated. The remlt, of
~ 5 sec sputter line ..".,.., 30
1
-Iv-
1 min ~ ~ ~ 10 ..~
40 Cthese experiments are shown in figure 2.23. The figure indicates that Ca is concen-trated in a region nea, the intergranular f,acture surface. One has to be very careful in translating the sputter time into a depth measured in Angstroms. The intergranular surface is very rough and the Xe ion beam hit the surface at grazing incidence. The sputtering of the surface can therefore be expected to be non-uniform.
2.3.4 Discussion
The results of sees 2.3.3 and 2.3.4 demonstrate that MgO influences both the densification kinetics and grain growth of alumina in the same way. This suggests that densification and grain growth are governed by the same mechanism. This could also be concluded from intermediate stage sintering data of CoO ll). Very small amounts of MgO promote the sintering of alumina. A second phase of spinel is obviously not necessary. The experiments with Auger spectroscopy do not point to an important grain boundary segregatiun uf MgO or tu its presence as a
mOl
around the grains, which would decrease the grain boundary mobility. This is contrary to the conclusion arrived at by Jorgensen and Westbrook (>l) andBudworth 65). Our conclusions are supported by the work of Johnson and Stein 78). With broad beam AES they observed Mg peaks representing not more than twice the bulk level. With scanning Auger microprobe they were able to determine that most of the Mg detected on the fracture surface is confined to discrete particles and that the overall coverage between the particles corresponds to the bulk level in the alumina. The overall increased concentration of Mg at the grain boundaries found by Taylor et aI. 68, n) must be the result of precipitation rather than segregation. It should be realized that the results of grain boundary analysis will be influenced by the cooling rate after sintering. During slow cooling segregation and precipitation may strongly increase.
Our eJ\periments and the eady eJ\perilnents of Marcus and Fine 66) show that ell is a pronounced segregant. Since the size misfit for the Ca ion in alumina is large, it may be that elastic strain energy makes an important contribution to its tendency to segregate. The ionic radius of Ca is about 0.99 A, significantly larger than that of Al, which is about 0.50 A and Mg, which is about 0.65 A. 1n spite of its strong tendency to segregate, CaO as the only additive to alumina is not effective in pre-venting discontinuous grain growth in alumina 45). Nor did the combination of MgO and CaO in our experiments improve the density of the samples or their transmission properties (section 4.2.3). The same holds for the combination MgO and Y,O).-.
111ere is still another reasun for rejecting the idea of solute segregation of Mg. Figure 2.9 (page 20) shows that in the single-phase region the grain size increases with the MgO content. In the case of segregation of Mg to the grain boundaries one would expect the reverse to occur. This reverse effect has been found in Y 203 containing various amounts of Th02 , and was ascribed to segregation of III 30).
Coble 44) already observed that addition of MgO to alumina did not irlhibit the rate of normal grain growth, and Coble and Burke $) listed possiblo reasons for this. However, they added 0.25 wt % MgO, which far e)(cecds the solubility limit. Our results indicate that MgO even promotes normal grain growth as long as the solubility limit has not bem reached.
II sCCms that the average grain growth rate is completely controlled by the volume fraction and si<l:e of the pores. The more MgO is present (below the solubility limit) the mOrC the sintering !"ate and thus the rate of pore removal is enhanced. This means that there are fewer restraints on the grain boundaries and the grain boundary velocity can more closely approach the velocity of a free boundary. The essential action of MgO is then the enhancement of the densificatiOl\ rate of alumina and nO! th", reduction of grain boundary mobility. Rao and Cutler 32) found that the sintering rate of alumina doped with fe2T ions increas~s with the total Fe (;On tent , and it seems that the same is tmc of alumina doped with Mg2 + ions. This makes Ollr results consistent with t.he defect model for alumina diSC\lssed in sec. 2.2.3. We must conclude that the inhibition of discontinous grain growth by MgO is not caused by a considerable red\lction in grain boundary mobility due to ci ther a soltltc segregation mechanism or to the presence of spinel a~ second-phase particle~.
More important is the enhancement of the sint.ering rate or rale of pON removal, which averts the condilion for the occurrence of discontinu()l[s grain growth. A second effect of MgO may be an increase of the pore mobility, so that the pore can f·oUow the grain boundary more easily. Possibly both mechanisms contribllte to the inhibition of discontinuous grain growth by MgO. The additive might also change the poro geometry, i.e. the grain boundary angle of a pore situated on a grain boundalY. This dihedral angle is determintd by the fa tio of the grain boundary energy and the SIIrface energy. Jorgensen 51 ), trying to detect pos~ible
solute segregation, states that addition of 0.1 wt % MgO inGreases the angle by a faGtor of 2. However, we were unable to observe this large increase in our samples. Above the solubility limit scc()nd'phase particles are present and this changes the situalion completely. The second·phase inclusions lie on grain edges and at grain corners and influence the grain growth process,
An increase irl the MgO content corresponds to an increase in the average number of second·phase part.icles per grain boundary and/or iln increase in their average size. This results in an increasing drag on the grain bOllndaI'ies and e)(plains the decreasing mean grain size found experimentally. MocelUn and Kingery ~2), who have annealed samples of MgO-doped al\lmina, als(l find that second-phase draggjng probably con· troIs the grain growth kinetics.
Another aspect of an increasing amount of second-phase particles is the decrease ill sintering rate, which is expressed in lower densities (figure 2.S, page 19). The same phenomenon has beon found in other systems. Rcijncn 4) has put forward the
hypothesis that grain boundary sliding is essential for fast sintering, and suggests that a dispersed second phase would suppress this grain boundary sliding for purely geo-metrical reasons. If this hypothesis holds in this case, longer sintering times should give higher densities. This has indeed been confirmed. Another explanation involves an overall decrease of mobility in the boundary when the boundary is 'blocked' by a dispersed phase.
Finally a remark on the non-uniform microstructure in the samples with 100 ppm MgO and less, as illustrated in fig. 2.14 (page 22). It is possible that this is caused by a non-uniform distribution of MgO. Locally mare MgO may be present and this would then result in a higher rate of pore removal and more rapid grain growth. These large grains oan then grow at the expense of others. This might explain why the large grains are nearly pore-free (in contrast with the large grains in the case of undoped alumina) and why large pores are often found between large and small grains. A similar phenomenon has been found by Groscovkh 83) in the ThO~ "doped YZ03 system.
2.3.5 Conclusion
The addition of MgO promotes both the densification and grain growth of alumina as long as the solubility limit has not been reached. The most effectIve dope level corresponds to the amount that can be brought into solid solution. Auger spectro-scopy shows no important enrichment of the grain boundary area with MgO. These results do not support the theory of grain boundary segregation. The essential action of MgO seems to be enhancement of the pore removal rate. At higher dope levels, when seccmd·phase particles are present, grain growth is slowed down and the densification is also negatively influenced.
REFERENCES
I) M.F. Ashby, Acta Met. 22,275 (J974). 2) D.L. Johnson, J. App!. Phys. 40,192 (1969). 3) D.W.Re~dey, J.Am.Cerarn.Soc.49,366(1966).
4) P.J.L. Reynen, Problems of Nonstoichi(lmElry, Ed. A. Rabenau, North-Holland PlIbL Co., Amsterdam, 1970, p. 219,
5) R.L. Coble and 1.E. Burke, Progr. Cmm. Sci. 3,197 (1963). 6) LB. Cutler, High Temperature OXides, part
nr,
Ed. A.M. Alper, AcademicPress, New York, )970, p. 129.
7) G.c. Kupynski, Adv.Coll.lrlterf.Sci,3,275(l972). 8) A.L. Stuijts, Ann. Rev. Mal. Sci. 3, 363 (1973).
9) Y. Oishi and W.D. Kingery, J. Chern. Phys. 33,480 (1960). 10) C.M. Fryer, Trans. Brit. Ccram. Soc. 68,191 (1969). II) C.M. Fryer, Trans. BriL Ceram. Soc. 71,231 (1972).
12) H.P.R. Frederibe and W.R. I-[osler, Mat. Sci. Res. 9,233 (1975).
13) G.1. Dienes, D.O. Welch, C.R. Fisher, R.D. Hatcher, 0, Lazareth and M. Sam berg, Phys. Rev. B. 11,3060 (1975).
J4) R.1. Brook, J. Am. Ccrarn. Soc. 55, 114 (1972).
15) M.H. Leipold and C.M. Kapadia, j. Am. Ceram. Soc. 56,200 (1973). 16) A.A. Bauer and LL. Bates, Battelle Mem. Ins!. Rep!., 1930, July 31
(1974).
17) R.E, Mistler and R.L. Coble, 1. App!. Phys. 45,1507 (1974). J8) A.E. Paladino and W .D, Kingery, J. Chern. Phys. 37,957 (1962). 19) A.E. Paladino and R.L. COble, J. Am. Ceram. Soc. 46,133 (1963). 20) R.E. Mistler and R.L. Coble, J. Am. Ce,am. Soc. 54, 60 (1971).
