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Microstructural observations on sintered alpha alumina
Materials Chemistry 4 (1979) 721 - 730 © CENFOR S.R.L.- Printed in Italy
SHORT COMMUNICATIONS
MICROSTRUCTURAL OBSERVATIONS ON SINTERED ALPHA ALUMINA This paper reports a study of microstructural features observed on tape casted, high surface finish high purity (~ 99%) alumina substrates prepared on a laboratory scale. Some of the features observed may indirectly affect the material behaviour as regards both electronic circuitry deposition and adhesion, and in-service performance. The processing technology of the substrates, described in detail elsewhere 1, consists essentially in (i) the preparation of a mix of MgO-doped (~ 4500 ppm) alumina powder with the appropriate amounts of organic solvents, binders, deflocculents and plasticizers; (ii) tape casting the above slip on a moving cellulose triacetate bed by using a specially designed "doctor blading" apparatus I ; (iii) drying and punching the tape and finally (iv) sintering in air. Firing temperatures ranging from 1350°C to 1600"C and annealing times from 0 to 25 hrs were used to obtain materials having intermediate microstructures. Microstructural analyses were performed by qualitative and quantitative X-ray diffractometry, SEM and WDS. Microstructural changes related to the degree of material densification could be interpreted2 on the basis o'f existing theories of highly porous 3 and dense 4 materials. The additional microstructural features reported in this study consist in the appearance of magnesium aluminate spinel precipitates at the grain boundaries of the alumina crystals and in the establishment of a pronounced crystallographic texture of the ~-A1203 grains, both dependent on sintering temperature and time.
722
b)
a)
I
c)
d)
Fig. 1 - Mg mapping of: a) as-fired substrate surface sintered at 1350°C for 24 hrs; b) as-fired substrate surface sintered at 1400°C for 1 hr; c) as-fired substrate surface sintered at 1450°C for 1 hr; d) as-fired substrate surface sintered at 1500°C for 1 hr (1900 x).
723 Magnesium aluminate spinel precipitation was revealed by X-ray diffraction analysis although the overall MgO content in the material did not exceed a few thousand ppm. The suggestion can be made that, even after sintering, most of the MgO concentrates at the grain boundaries of the alumina crystals. This view agrees with the solubility limit of MgO in AI 2 0 3 reported as being a few hundred ppm at the higher temperatures used in the present studyS and decreasing at lower temperatures. Mg mapping at various sintering temperatures and times allows the sequence of magnesium aluminate spinel precipitation to be determined (Figs. 1a-d). Fig. 2 represents a combined SEM and Mg mapping picture of a Substrate fired in an industrial kiln at 1470°C with an overall firing cycle of 72 hrs (cold-to-cold). At the lowest annealing temperature (1350°C) magnesium aluminate spinel precipitation occurs only after prolunged annealing time (see Fig. la). Smaller and more frequent magnesium aluminate spinel
a) 1470°C 72 h (industrial cycle)
b) 1470°C 72 h (industrial cycle~
Fig. 2 - a) Polished and etched surface o f a substrate sintered at 1470°C in an industrial kiln with a firing cycle, o f 72 hrs, and b) Mg mapping o f the same zone as in Fig. 2a in which magnesium aluminate spinel grains are clearly evidenced (1900x).
724 precipitates appear at increasing sintering temperatures (see Figs. 1 b-c-
-d). The general theory of statistical fluctuations gives the following equation 6, 7 for the nucleation rate I of a new phase, i.e. the number of nuclei of critical size capable of further growth formed per unit volume and unit time at temperature T: A~* (1)
I=N*N~:(Ne
- kT ) ( A e
A~ kT )
AS* where N* = N e kX is the number of nuclei of critical size per unit volume; N ; = A e kX is the number of elementary events per unit time converting critical nuclei into stable nuclei; N is a quantity proportional to the number of sites where nuclei can appear; A is a quantity which depends on the activation mechanism of the growth of the stable nuclei. The nucleation rate shows a maximum which occurs approximately at:
(2)
Wm=Ws/3(1+ A~b*)/[1+ 1/3 A~b* \)
in which Ts is the melting point of the nucleated crystal. Tm increases with increasing A ~ , i.e. the overall thermodynamic potential barrier which must be overcome to convert a critical nucleus into a stable one. When A ~ = A~*, i.e. it equals the overall increment of thermodynamic potential of the system in the formation of a nucleus of the new phase of critical size, then T m = 1/2 Ts. The above findings have been successfully applied in the development of nucleation theories of simple one-component viscous liquids (glasses) where a specification of the terms A~b~ and A~b* is possibile. No explicit relations can be developed in the present case for describing the thermodynamic functions appearing in Eqn. (1) on the basis of suitable models and physico-chemical material parameters. Further difficulty is encountered because of microstructural changes, which occur as sintering proceeds, both in pore substructure and grain size that
725 directly reflect on the heterogeneous nucleation process involved: the rate determining species must diffuse to the reaction sites via volume, surface or grain boundary transport mechanisms. This of course leads to microstructure-dependent apparent activation energies for nucleation and grain growth. Nevertheless, if the system is considered as a diluted dispersion of one of the reactants (MgO) in a homogeneous alumina-pore matrix, an attempt can be made to search for some possible relation between the experimental behaviour observed for the nucleation (precipitation) of magnesium aluminate spinel and the simple representation suggested by Tamman8 for glasses, including the stages of both nucleation and crystal growth. Although only largely qualitative deductions may be expected due to the strongly oversimplified model assumed, when the nucleation (precipitation) rate of magnesium aluminate spinel, determined by the combined analysis of SEM micrographs and Mg maps, is plotted as a function of the sintering temperature, the resuiting curve strictly resembles the behaviour expected on the basis of Tamman's concepts (see Fig. 3). The decrease of the average size of the magnesium aluminate spinel precipitates at increased sintering temperatures is also consistent with the above model for which increasing probability for stable nuclei formation occurs in the rising left hand side of the nucleation curve: because of the limited amount of MgO in the material, the more the nuclei formed, the less the magnesia available for growth of the nuclei. A detailed analysis of the preferred orientation observed on the a-A12 03 grains was performed by quantitative X-ray diffraction analysis. Since the intensities of the (0006) and (000.12) basal reflections were very weak, the stronger reflections (10]'.10) and (1159) were used. These reflections are very close to (0001) and nearly indistinguishable from it, being the angle between (10]'.10) and (0001) only 15 ° and the angle between (1129) and (0001) only 18". Comparative analysis performed on the random powder, on the dried tape and on the specimens fired at various temperatures and/or annealing times (see Table 1) showed the aligning effect of the tape casting operation to be of very minor importance, whereas the increasing crystallographic texture with increasing firing temperatures and longer annealing times suggests
726
I,G
J
'
!
1 / I
•
a)
i
) /
I
¢ /
20°C
011
-,-Supercooting
,,
W~ToC Tm.p.
,!I
!
1600 1BOO Tin. P. Temperature °C
Fig. 3 - a) Schematic for the crystallization of a viscous liquid: T, heating temperature; I, nucleation rate; G, rate o f crystal growth. Curve I represents the nucleation rate; curve 2 represents the rate o f crystal growth. (g) Metastable supercooling zone; (~) metastable zone o f hig[~ viscosity. b) Experimental curve o f nucleation (precipitation) rate (arbitrary units) o f magnesium aluminate spinel (Tm = 2135°C).
that a reconstructive phenomenon takes place during sintering and grain growth. A moderate texture perpendicular to the plane of the substrate for tape-cast A1203 was previously observed by Pentecost et al. 9 Di Marcello et al.] 0 reported a m u c h stronger basal-plane fibre texture in both bulk and surface suggesting a connection between texture formation
727 Table 1 - Crystallographic texture in the o~-A1203 grains on the substrate surface (intensity values normalized to those of the random powder). specimen
I/I R
random powder dried tape fired at 1500°C (0 hrs) fired at 1550°C (25 hrs) fired at 1600°C (3 hrs) fired at 1550°C (1 hrs) + 1600°C (3 hrs) fired at 1470°C (72 hrs) (industrial cycle)
1.0 1.2 1.6 1.8 1.8 3.4 3.3
and grain growth during sintering. Sundhal et al. 11 proposed that surface-controlled selective grain growth could act as a possible mechanisln of texture formation. Taking into account the strong sensitivity of texture appearance to sintering atmosphere and considering that the driving force for grain growth during sintering is the minimization of the surface free energy of the system, it seems probable (Nakada et al. I 2) that the preferred orientation might form during grain growth when growth rates vary for grains with different planes exposed at the surface. Also, the texture in the bulk may be somewhat less than at the surface since grain boundaries between successive iso-oriented grains may not remain a perfect twist boundary, i.e. a configuration of a lower free energy than that of a high-angle mixed boundary. The experimentally observed basal plane fibre texture, in fact, corresponds to the close packing plane in the hexagonal structure of 0~-A120 3 , i.e. the plane with the lowest surface free energy. Taking as "/b the grain-boundary free energy per unit grain boundary area and A'ys as the average difference in surface free energy per unit surface area between surface grains with exposed planes of lowest free energy to grains with exposed planes of different orientation, Foster et al. 1 3 suggested a model to predict the growth rate G of surface-ener-
728 gy controlled secondary grains in 3% Si-Fe sheets. They obtained the following relation: (3)
G=M
1.25 % ATs + + K r 2r
where M is the grain boundary mobility, r the average grain radius and K a term accounting for the contribution of grain boundary inclusions to the total free energy. The additional driving force associated with surface grains whose exposed surface is that of the lowest free energy is represented by the term A'Ys/2r. Although no exact calculation can be attempted for evaluating the relative importance of the two first terms on the right hand side of Eqn. (1), due to the uncertainty in the definition of the material parameters M, ~'b and ATs, the above model also might apply in the present case. The contribution of the term ATs to Eqn. (1) would probably have considerably importance, as can be inferred from the faceting of o~-A12O a grains on the substrate surface. A
Fig. 4 - TEM micrograph o f an as-fired substrate surface sintered at 1500°C f o r 4 hrs showing faceting o f the alumina grains.
729 TEM micrograph
of hexagonally-shaped
basal faceted planes taken on a
s u b s t r a t e s u r f a c e f i r e d a t 1 5 0 0 ° C f o r 4 h r s is s h o w n in Fig. 4.
A . B E L L O S I , P. V I N C E N Z I N I C.N.R., Research Laboratory f o r Ceramics Technologies - F A E N Z A - Italy.
REFERENCES
1. 2. 3. 4. 5. 6.
7. 8. 9.
10. 11.
12. 13.
C. FIORI -- Ceramurgia, 9, 1979. A. BELLOSI, P. VINCENZINI -- Ceramurgia International, 5, 89, 1979. C. GRESKOVICH, K.W. LAY - - J . Amer. Ceram. Soc., 55, 142, 1972. A. MOCELLIN, W.D. KINGERY -- ]. Amer. Ceram. Soc., 56, 309, 1973. S.K. ROY, R.L. COBLE - - J . Amer. Ceram. Soc., 51, 1, 1968. R. BECKER - A n n . Physik, 32, 128, 1938. V.N. FILIPOVICH - The Glassy State, Vol. 1, Izd. Akad. Nauk. SSSR, Moscow-Leningrad, 1963, p. 9. G. TAMMAN -- The States o f Aggregation, D. Van Nostrand Co., New York, 1925. J.L. PENTECOST, C.H. WRIGTH -- in Advances in X-Ray Analysis, Vol. 7, G.R. Mallett, Marie Fay and W.M. Mueller Eds., Plenum Press, New York, p. 174, 1964. F.V. DI MARCELLO, P.L. KEY, J.C. WILLIAMS -- J. Am. Cerarn. Soc., 55, 1964, p. 174. R.C. SUNDHAL, L. BERRIN -- in The Science o f Ceramics Machining and Surface Finishing, S.J. Schneider, Jr. and R.W. Rice Eds., Nat. Bur. Stand. (US) Spec. Publ. n. 348, 1972. Y. NAKADA, T.L. SCHOCK - ] . Amer. Ceram. Soc., 5 8 , 4 0 9 , 1975. K. FOSTER, J.J. KRAMER, G.W. WIENER -- Trans. Metall. Soc. AIME, 227, 185, 1963.
SHORT COMMUNICATIONS
MICROSTRUCTURAL OBSERVATIONS ON SINTERED ALPHA ALUMINA This paper reports a study of microstructural features observed on tape casted, high surface finish high purity (~ 99%) alumina substrates prepared on a laboratory scale. Some of the features observed may indirectly affect the material behaviour as regards both electronic circuitry deposition and adhesion, and in-service performance. The processing technology of the substrates, described in detail elsewhere 1, consists essentially in (i) the preparation of a mix of MgO-doped (~ 4500 ppm) alumina powder with the appropriate amounts of organic solvents, binders, deflocculents and plasticizers; (ii) tape casting the above slip on a moving cellulose triacetate bed by using a specially designed "doctor blading" apparatus I ; (iii) drying and punching the tape and finally (iv) sintering in air. Firing temperatures ranging from 1350°C to 1600"C and annealing times from 0 to 25 hrs were used to obtain materials having intermediate microstructures. Microstructural analyses were performed by qualitative and quantitative X-ray diffractometry, SEM and WDS. Microstructural changes related to the degree of material densification could be interpreted2 on the basis o'f existing theories of highly porous 3 and dense 4 materials. The additional microstructural features reported in this study consist in the appearance of magnesium aluminate spinel precipitates at the grain boundaries of the alumina crystals and in the establishment of a pronounced crystallographic texture of the ~-A1203 grains, both dependent on sintering temperature and time.
722
b)
a)
I
c)
d)
Fig. 1 - Mg mapping of: a) as-fired substrate surface sintered at 1350°C for 24 hrs; b) as-fired substrate surface sintered at 1400°C for 1 hr; c) as-fired substrate surface sintered at 1450°C for 1 hr; d) as-fired substrate surface sintered at 1500°C for 1 hr (1900 x).
723 Magnesium aluminate spinel precipitation was revealed by X-ray diffraction analysis although the overall MgO content in the material did not exceed a few thousand ppm. The suggestion can be made that, even after sintering, most of the MgO concentrates at the grain boundaries of the alumina crystals. This view agrees with the solubility limit of MgO in AI 2 0 3 reported as being a few hundred ppm at the higher temperatures used in the present studyS and decreasing at lower temperatures. Mg mapping at various sintering temperatures and times allows the sequence of magnesium aluminate spinel precipitation to be determined (Figs. 1a-d). Fig. 2 represents a combined SEM and Mg mapping picture of a Substrate fired in an industrial kiln at 1470°C with an overall firing cycle of 72 hrs (cold-to-cold). At the lowest annealing temperature (1350°C) magnesium aluminate spinel precipitation occurs only after prolunged annealing time (see Fig. la). Smaller and more frequent magnesium aluminate spinel
a) 1470°C 72 h (industrial cycle)
b) 1470°C 72 h (industrial cycle~
Fig. 2 - a) Polished and etched surface o f a substrate sintered at 1470°C in an industrial kiln with a firing cycle, o f 72 hrs, and b) Mg mapping o f the same zone as in Fig. 2a in which magnesium aluminate spinel grains are clearly evidenced (1900x).
724 precipitates appear at increasing sintering temperatures (see Figs. 1 b-c-
-d). The general theory of statistical fluctuations gives the following equation 6, 7 for the nucleation rate I of a new phase, i.e. the number of nuclei of critical size capable of further growth formed per unit volume and unit time at temperature T: A~* (1)
I=N*N~:(Ne
- kT ) ( A e
A~ kT )
AS* where N* = N e kX is the number of nuclei of critical size per unit volume; N ; = A e kX is the number of elementary events per unit time converting critical nuclei into stable nuclei; N is a quantity proportional to the number of sites where nuclei can appear; A is a quantity which depends on the activation mechanism of the growth of the stable nuclei. The nucleation rate shows a maximum which occurs approximately at:
(2)
Wm=Ws/3(1+ A~b*)/[1+ 1/3 A~b* \)
in which Ts is the melting point of the nucleated crystal. Tm increases with increasing A ~ , i.e. the overall thermodynamic potential barrier which must be overcome to convert a critical nucleus into a stable one. When A ~ = A~*, i.e. it equals the overall increment of thermodynamic potential of the system in the formation of a nucleus of the new phase of critical size, then T m = 1/2 Ts. The above findings have been successfully applied in the development of nucleation theories of simple one-component viscous liquids (glasses) where a specification of the terms A~b~ and A~b* is possibile. No explicit relations can be developed in the present case for describing the thermodynamic functions appearing in Eqn. (1) on the basis of suitable models and physico-chemical material parameters. Further difficulty is encountered because of microstructural changes, which occur as sintering proceeds, both in pore substructure and grain size that
725 directly reflect on the heterogeneous nucleation process involved: the rate determining species must diffuse to the reaction sites via volume, surface or grain boundary transport mechanisms. This of course leads to microstructure-dependent apparent activation energies for nucleation and grain growth. Nevertheless, if the system is considered as a diluted dispersion of one of the reactants (MgO) in a homogeneous alumina-pore matrix, an attempt can be made to search for some possible relation between the experimental behaviour observed for the nucleation (precipitation) of magnesium aluminate spinel and the simple representation suggested by Tamman8 for glasses, including the stages of both nucleation and crystal growth. Although only largely qualitative deductions may be expected due to the strongly oversimplified model assumed, when the nucleation (precipitation) rate of magnesium aluminate spinel, determined by the combined analysis of SEM micrographs and Mg maps, is plotted as a function of the sintering temperature, the resuiting curve strictly resembles the behaviour expected on the basis of Tamman's concepts (see Fig. 3). The decrease of the average size of the magnesium aluminate spinel precipitates at increased sintering temperatures is also consistent with the above model for which increasing probability for stable nuclei formation occurs in the rising left hand side of the nucleation curve: because of the limited amount of MgO in the material, the more the nuclei formed, the less the magnesia available for growth of the nuclei. A detailed analysis of the preferred orientation observed on the a-A12 03 grains was performed by quantitative X-ray diffraction analysis. Since the intensities of the (0006) and (000.12) basal reflections were very weak, the stronger reflections (10]'.10) and (1159) were used. These reflections are very close to (0001) and nearly indistinguishable from it, being the angle between (10]'.10) and (0001) only 15 ° and the angle between (1129) and (0001) only 18". Comparative analysis performed on the random powder, on the dried tape and on the specimens fired at various temperatures and/or annealing times (see Table 1) showed the aligning effect of the tape casting operation to be of very minor importance, whereas the increasing crystallographic texture with increasing firing temperatures and longer annealing times suggests
726
I,G
J
'
!
1 / I
•
a)
i
) /
I
¢ /
20°C
011
-,-Supercooting
,,
W~ToC Tm.p.
,!I
!
1600 1BOO Tin. P. Temperature °C
Fig. 3 - a) Schematic for the crystallization of a viscous liquid: T, heating temperature; I, nucleation rate; G, rate o f crystal growth. Curve I represents the nucleation rate; curve 2 represents the rate o f crystal growth. (g) Metastable supercooling zone; (~) metastable zone o f hig[~ viscosity. b) Experimental curve o f nucleation (precipitation) rate (arbitrary units) o f magnesium aluminate spinel (Tm = 2135°C).
that a reconstructive phenomenon takes place during sintering and grain growth. A moderate texture perpendicular to the plane of the substrate for tape-cast A1203 was previously observed by Pentecost et al. 9 Di Marcello et al.] 0 reported a m u c h stronger basal-plane fibre texture in both bulk and surface suggesting a connection between texture formation
727 Table 1 - Crystallographic texture in the o~-A1203 grains on the substrate surface (intensity values normalized to those of the random powder). specimen
I/I R
random powder dried tape fired at 1500°C (0 hrs) fired at 1550°C (25 hrs) fired at 1600°C (3 hrs) fired at 1550°C (1 hrs) + 1600°C (3 hrs) fired at 1470°C (72 hrs) (industrial cycle)
1.0 1.2 1.6 1.8 1.8 3.4 3.3
and grain growth during sintering. Sundhal et al. 11 proposed that surface-controlled selective grain growth could act as a possible mechanisln of texture formation. Taking into account the strong sensitivity of texture appearance to sintering atmosphere and considering that the driving force for grain growth during sintering is the minimization of the surface free energy of the system, it seems probable (Nakada et al. I 2) that the preferred orientation might form during grain growth when growth rates vary for grains with different planes exposed at the surface. Also, the texture in the bulk may be somewhat less than at the surface since grain boundaries between successive iso-oriented grains may not remain a perfect twist boundary, i.e. a configuration of a lower free energy than that of a high-angle mixed boundary. The experimentally observed basal plane fibre texture, in fact, corresponds to the close packing plane in the hexagonal structure of 0~-A120 3 , i.e. the plane with the lowest surface free energy. Taking as "/b the grain-boundary free energy per unit grain boundary area and A'ys as the average difference in surface free energy per unit surface area between surface grains with exposed planes of lowest free energy to grains with exposed planes of different orientation, Foster et al. 1 3 suggested a model to predict the growth rate G of surface-ener-
728 gy controlled secondary grains in 3% Si-Fe sheets. They obtained the following relation: (3)
G=M
1.25 % ATs + + K r 2r
where M is the grain boundary mobility, r the average grain radius and K a term accounting for the contribution of grain boundary inclusions to the total free energy. The additional driving force associated with surface grains whose exposed surface is that of the lowest free energy is represented by the term A'Ys/2r. Although no exact calculation can be attempted for evaluating the relative importance of the two first terms on the right hand side of Eqn. (1), due to the uncertainty in the definition of the material parameters M, ~'b and ATs, the above model also might apply in the present case. The contribution of the term ATs to Eqn. (1) would probably have considerably importance, as can be inferred from the faceting of o~-A12O a grains on the substrate surface. A
Fig. 4 - TEM micrograph o f an as-fired substrate surface sintered at 1500°C f o r 4 hrs showing faceting o f the alumina grains.
729 TEM micrograph
of hexagonally-shaped
basal faceted planes taken on a
s u b s t r a t e s u r f a c e f i r e d a t 1 5 0 0 ° C f o r 4 h r s is s h o w n in Fig. 4.
A . B E L L O S I , P. V I N C E N Z I N I C.N.R., Research Laboratory f o r Ceramics Technologies - F A E N Z A - Italy.
REFERENCES
1. 2. 3. 4. 5. 6.
7. 8. 9.
10. 11.
12. 13.
C. FIORI -- Ceramurgia, 9, 1979. A. BELLOSI, P. VINCENZINI -- Ceramurgia International, 5, 89, 1979. C. GRESKOVICH, K.W. LAY - - J . Amer. Ceram. Soc., 55, 142, 1972. A. MOCELLIN, W.D. KINGERY -- ]. Amer. Ceram. Soc., 56, 309, 1973. S.K. ROY, R.L. COBLE - - J . Amer. Ceram. Soc., 51, 1, 1968. R. BECKER - A n n . Physik, 32, 128, 1938. V.N. FILIPOVICH - The Glassy State, Vol. 1, Izd. Akad. Nauk. SSSR, Moscow-Leningrad, 1963, p. 9. G. TAMMAN -- The States o f Aggregation, D. Van Nostrand Co., New York, 1925. J.L. PENTECOST, C.H. WRIGTH -- in Advances in X-Ray Analysis, Vol. 7, G.R. Mallett, Marie Fay and W.M. Mueller Eds., Plenum Press, New York, p. 174, 1964. F.V. DI MARCELLO, P.L. KEY, J.C. WILLIAMS -- J. Am. Cerarn. Soc., 55, 1964, p. 174. R.C. SUNDHAL, L. BERRIN -- in The Science o f Ceramics Machining and Surface Finishing, S.J. Schneider, Jr. and R.W. Rice Eds., Nat. Bur. Stand. (US) Spec. Publ. n. 348, 1972. Y. NAKADA, T.L. SCHOCK - ] . Amer. Ceram. Soc., 5 8 , 4 0 9 , 1975. K. FOSTER, J.J. KRAMER, G.W. WIENER -- Trans. Metall. Soc. AIME, 227, 185, 1963.