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Grain boundary chemistry of alumina by high-resolution imaging SIMS
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Acta mater. Vol. 47, Nos 15, pp. 4031±4039, 1999 # 1999 Published by Elsevier Science Ltd On behalf of Acta Metallurgica Inc. All rights reserved. Printed in Great Britain S1359-6454(99)00263-3 1359-6454/99 $20.00 + 0.00
GRAIN BOUNDARY CHEMISTRY OF ALUMINA BY HIGH-RESOLUTION IMAGING SIMS K. L. GAVRILOV 1{, S. J. BENNISON 2, K. R. MIKESKA 2 and R. LEVI-SETTI 1 Enrico Fermi Institute and Department of Physics, University of Chicago, Chicago, IL 60637, U.S.A. and 2E.I. DuPont de Nemours & Co. Inc., Central Research and Development, Experimental Station, Wilmington, DE 19880, U.S.A.
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AbstractÐThe unique capabilities of the high-resolution scanning ion microprobe developed at the University of Chicago (UC-SIM) are described and its utility is demonstrated in a study of grain boundary chemistry of alumina ceramics. When polycrystalline alumina is doped singly with either MgO or SiO2, strong segregation of the individual ions to grain boundaries is observed: (1) for Mg segregation Cgb =Cgrain 400; (2) for Si segregation Cgb =Cgrain 300. However, on codoping with both MgO and SiO2, grain boundary segregation is signi®cantly diminished by a factor of ®ve or more over single doping as both cations are redistributed into the bulk alumina lattice. A defect compensation mechanism is proposed to explain this mutual solid solubility of Mg and Si in alumina. One important consequence of this chemical redistribution is a change in abnormal grain growth morphology from facetted grains in SiO2 singly doped alumina, to non-facetted grains with curved boundaries in MgO and SiO2 codoped alumina. As the Mg/Si dopant ratio exceeds the equimolar concentration, abnormal grain growth development ceases. These ®ndings provide a physical mechanism to explain the role of MgO as a sintering aid to control microstructure evolution in alumina. Another signi®cant consequence of SiO2 redistribution on MgO doping is an observed improvement in the corrosion resistance of alumina to aqueous HF. Siliceous grain boundary ®lms readily corrode and compromise the intrinsically good corrosion resistance of bulk alumina. MgO doping in amounts greater than the SiO2 concentration prevents the formation of these corrodable silica-based phases leading to the development of aluminas for use in aqueous HF-containing ambients. # 1999 Published by Elsevier Science Ltd on behalf of Acta Metallurgica Inc. All rights reserved. Keywords: Ion microprobe; Secondary ion mass spectroscopy (SIMS); Ceramics; Microstructure; Corrosion
1. INTRODUCTION
High-resolution imaging secondary ion mass spectrometry (HRI-SIMS) has emerged as a valuable research technique to study interfacial chemistry of a variety of modern engineered materials such as metal alloys, composites and ceramics [1±6]. The application of HRI-SIMS has been driven by advances in the design of interface chemistry and structure to tailor electrical, magnetic and optical properties of materials. HRI-SIMS has been employed in the interfacial design of ceramics [7] where the in¯uence of trace impurities on sintering and microstructure development is exploited in the manufacture of materials with speci®ed physical performance [8]. A number of important innovations introduced in scanning ion microprobe (SIM) technology have created an opportunity to address the role of trace impurities in interfacial chemistry and associated properties. One innovation includes the development of a liquid-metal ion source (LMIS) that enables the theoretical lateral resolution limit for SIM instruments to be attained [3]. Another is the incorporation, in the University of Chicago scan{To whom all correspondence should be addressed.
ning ion microprobe (UC-SIM), of an ecient secondary ion transport system together with the addition of a Finnigan'90-based Magnetic Sector Mass Spectrometer that provides analytical imaging SIMS sensitivity in the range of 1 p.p.m. [3] (Fig. 1). In this paper we discuss the use of HRI-SIMS in combination with other analytical tools and experimental methodologies to understand the role of speci®c dopants and impurities in microstructure development and associated corrosion resistance of ceramics. Trace impurities and additives have long been known to strongly in¯uence the fabrication and properties of alumina-based ceramics. For example, MgO additions improve sintering characteristics and allow the manufacture of aluminas with controlled microstructures [9±11]. More commonly, impurities such as silica and glass-modifying ions tend to have a detrimental eect on processing since their presence often promotes abnormal grain growth [12±15]. These latter impurities often compromise properties that are sensitive to the presence of glassy grain boundary ®lms, such as electrical properties [16] and corrosion resistance [8, 17]. The detailed function of background impurities and trace additives has been dicult to ascertain due to problems in the detection of additives and their as-
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Fig. 1. Schematic of University of Chicago high-resolution imaging secondary ion mass spectrometer (UC-SIM).
sociated distribution in the microstructure. The HRI-SIMS is well suited to address the problem of trace element detection and distribution, particularly in interfacial regions such as grain boundaries. In this contribution we summarize the unique capabilities of the UC-SIM and discuss the advantages and limitations of this analytical tool. The use of HRI-SIMS is illustrated in three studies on alumina ceramics. The ®rst addresses segregation of dopants and impurities to grain boundaries. We show that codoping with magnesia and silica strongly aects their bulk mutual solid solubility and associated segregation characteristics to grain boundaries. The second study looks at microstructure evolution as a function of grain boundary chemistry. The segregation characteristics observed in the ®rst study play a determining role in grain morphology and growth kinetics. The third study considers the eect of segregated impurities on controlling corrosion resistance. We demonstrate that, with suitable control of grain boundary chemistry, polycrystalline aluminas may be engineered with corrosion resistance approaching that of single-crystal sapphire.
2. CAPABILITIES OF HIGH RESOLUTION IMAGING SIMS
In a scanning ion microprobe, a focused beam of primary ions is rastered over a specimen in order to
erode via sputtering the outermost surface layers. A fraction of the resulting ejected material is emitted as ionized atoms or molecules; this ionized fraction varies among elements and depends primarily on the electronic state of the sputtered material and on the chemical characteristics of the surface. The secondary ions and molecular clusters (either positively or negatively charged) are gathered and discriminated on the basis of their mass/charge ratio with a mass spectrometer. By recording the secondary ion signal as a function of the position of the scanning beam, two-dimensional compositional distribution maps are obtained. These SIMS maps can be digitally stored for subsequent image processing and analysis. HRI-SIMS is a powerful material characterization technique that simultaneously spatially locates and quanti®es dopant and impurity distribution in the microstructure. This simultaneous combination becomes especially important for advanced engineered materials where the detailed function of background impurities and additives has been debated for many years [11]. The UC-SIM instrument is ideally suited to address the questions concerning the role of impurities and additives, and has a number of advantages over other analytical techniques and ``conventional'' SIMS: (a) Sensitivity: HRI-SIMS can detect all elements and isotopes, many of them at part per million concentration levels. The signal-to-noise
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY 6
ratio can be as high as 10 Ðessentially, there is no background. (b) Spatial resolution: The spatial and depth resolutions of HRI-SIMS are excellent. A lateral resolution of 20±50 nm has been attained with the UC-SIM [3]. A depth resolution of 1±10 nm is common. However, there is a trade-o between sensitivity and resolution. (c) Speed of analysis: SIMS distribution maps can typically be acquired in 1±9 min time scales depending on the concentration of elements of interest. With this rate, it is possible to perform comprehensive analyses of multicomponent samples in practical times. (d) Simultaneous multigrain imaging: a large number of grain boundaries are readily analyzed during construction of an elemental map. For 2 example a map of a 2020 mm region will contain several hundred grain boundaries for a material with an average grain size of 2 mm. (e) Specimen preparation: bulk materials polished with standard ceramographic procedures can be studied. Additionally, the specimen surface can be sputter-cleaned in situ with the primary ion beam or with an auxiliary high-current ion beam. In practice this means that surface artifacts introduced during preparation can frequently be identi®ed and sputter-erased. Spatial resolution approaching the probe size (typically 035 nm) in scanned SIMS maps can be attained when the corresponding secondary ion signal is suciently large. Thus, the ``analytical resolution'' in elemental maps is conceptually dierent from ``structural resolution'' in the context of light or electron micrographs. Even for dopants or impurities present at the trace level in bulk, it is generally possible to detect grain boundary segregation with the UC-SIM. We have demonstrated this capability routinely in the analysis of grain boundaries in polycrystalline ceramics. Further details of this instrument and its capability in analyzing ceramic and composite microstructures can be found elsewhere [1, 18, 19]. Quanti®cation of the secondary ion signal is often problematic and is the primary limitation of the SIMS method. This problem arises from variations in the secondary particle sputtering yield and ion fraction as a function of the matrix structure and composition, and by the emission of neutral species. This limitation can be overcome using either rigorous calibration standards [20] or by combining SIMS with other characterization techniques. In many cases internal standards may be used such as stoichiometric precipitates in the microstructure. Also, relative concentration measurements may be employed to quantify the extent of grain boundary segregation relative to the bulk.
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A further complication in SIMS analysis concerns discrimination between sputtered complexes of similar mass/charge ratio. For example, in the case of 28 + alumina with SiO2 impurities, sputtered Si ions + have a similar mass/charge ratio to AlH and resolving these complexes is dicult. In principle, these two species could be resolved by the use of the Finnigan'90-based Magnetic Sector Mass Spectrometer, but at the expense of SIMS sensitivity. A de®nitive solution is obtained for the study of dopants by using isotopes with appropriate mass where possible. In the case of Si dopants in 29 alumina, Si-containing SiO2 may be used to 29 + + dierentiate sputtered Si from AlH complexes [7].
3. EXPERIMENTAL STUDIES
3.1. Sample preparation Sumitomo AKP-30 alumina was used as the base powder for sample preparation. This powder, according to Sumitomo, is 99.995% pure, with a 0.5 mm mean crystallite size and 50% of agglomerates below 1 mm. Powder handling was carried in ``semi-clean'' conditions using a laminar ¯ow hood 1 and clean Te¯on laboratory ware. Doping was achieved by adding aliquots of various metal salt complexes dissolved in water or ethanol followed by drying and/or hydrolyzing and deagglomeration. Additions of MgO were made using Mg(NO3)24H2O (99.99% purity). Additions of SiO2 were made using tetraethyl-orthosilicate (TEOS) (99.99% purity). Green pellets were prepared by dry pressing at 75 MPa followed by cold isostatic pressing at 224 MPa. Sample pellets were ®red in clean covered alumina crucibles containing protective powder identical in composition to the pellets. A resistance heated air furnace was employed for ®ring. The ®ring schedule consisted of calcining at 10008C for 2 h, ramping to a peak soak temperature of 16508C, and annealing for 60 min at 16508C. Specimens were cooled in a furnace powero condition, which results in a cooling rate on the order of 6008C/h at temperatures near the sintering temperature. 3.2. Microstructural characterization The sintered aluminas were polished to ¯at mirror surface ®nishes using successively ®ner grades of diamond abrasives down to 0.5 mm. Charging was not a problem with these materials providing that a thin coating of gold was applied to the polished section prior to imaging with the SIMS microprobe. Images were collected after sputtering o the gold layer and any surface contamination resulting from specimen preparation (i.e. polishing). Estimates of grain boundary segregation were made by comparing ion yields from grain boundary regions, Ygb, to ion yields from grain centers, Ygrain:
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Enrichment factor
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY
Ygb Ion probe diameter : 1 Ygrain Grain boundary width
It is assumed that matrix eects are uniform across grain boundaries. A 40 nm diameter ion probe and a boundary width of 1 nm were used to calculate grain boundary enrichments for both Mg [4] and Si segregation [21]. We acknowledge that the segregation characteristics of Mg and Si are likely to be dierent. Magnesium probably segregates in a classical solid±solution sense and Si most likely forms a distinct grain boundary ®lm. However, the enrichment width scale for both these distinct segregation mechanisms is similar and on the order of 1 nm. The uncertainty in the estimated enrichment values is 220% of the mean. Scanning electron microscopy (SEM) and optical microscopy were used to characterize the microstructural morphology of the sintered alumina samples. Specimens for microscopy were polished as described above. The polished sections were subsequently thermally etched at 15508C for 10 min in air to reveal the microstructure morphology. All specimens were sputter-coated with a thin (03 nm) layer of gold/palladium to mitigate sample charging in the SEM and enhance re¯ectivity for optical microscopy. The sequence of sample characterizations was: (1) examine microstructure by SEM or optical microscopy to identify areas of interest; (2) mark areas of interest using Vickers indentations; (3) repolish to remove features created by thermal etching; (4) sputter coat with thin gold layer; (5) locate regions of interest in HRI-SIMS from residual hardness impressions; (6) sputter o gold layer and surface contamination; and (7) analyze region of interest by HRI-SIMS. 3.3. HF corrosion test protocol The HF corrosion test protocol consisted of subjecting specimens to an azeotropic 38.26 HF/61.74 H2O (w/o) solution at 9020.58C. The details of the protocol are discussed elsewhere [7]. The HF conÿ centration was 20 M and the F ion concentration ÿ2 was 8:9 10 M at the azeotropic composition. Ceramic specimens of approximately cubic geome3 try with a nominal volume of 1 cm were subjected to corrosion testing by immersing individually in 150 ml of the HF/H2O azeotrope contained in a Te¯on reaction ¯ask. An eective linear corrosion rate, D_ , was determined for each specimen from weight-loss measurements. D_ is typically reported in ml/y or mm/y and is given by @W D_ @ tsr
2
where @ W W0 ÿ Wi (W0 is the initial sample weight and Wi the instantaneous or ®nal sample weight), s is the sample surface area, r is the material density, and t is time. D_ is a measure of the rate of material thinning or the rate at which the
Fig. 2. SIMS maps of singly doped polycrystalline alumina sintered at 16508C: (a) Si+ map for alumina doped with Si/Al=1000 p.p.m.; (b) Mg+ map for alumina with Mg/ Al=500 p.p.m.
material uniformly corrodes away at its exposed surfaces. 4. RESULTS AND DISCUSSION
4.1. Segregation of Si and Mg to grain boundaries The distribution of SiO2 dopant in a 1000 p.p.m. Si/Al specimen sintered for 480 min at 16508C is shown in Fig. 2(a). The grain morphology displays a small number of elongated (abnormal) grains containing patches of ®ner grains, which is typical of aluminas contaminated or doped with small amounts of glass-forming compounds [12±15]. The long facets of these elongated grains are typically of {0001} orientation and contain thin glassy ®lms
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Fig. 3. Grain boundary enrichment ratio (grain boundary ion-yield signal/grain center ion-yield signal) plotted for polycrystalline alumina doped singly with MgO or SiO2 and codoped with MgO and SiO2. Note the strong drop in both Mg and Si segregation to grain boundaries on codoping with both elements.
[14]. Using an integrated secondary ion image (ISI) topography map of this region (not shown) there is 28 + no indication of a Si signal corresponding to the location of surface polishing scratches in this region. Strong segregation of Si to grain boundaries and pore surfaces can be seen in Fig. 2(a). The ratio of Si in the boundary region to grain center is Cgb =Cgrain 0300 as shown in Fig. 3. It should be 28 + noted that the Si signal used to form these images contains a component made up of sputtered + AlH molecular ions that form due to a reaction between sputtered Al and residual hydrogen in the vacuum chamber and/or alumina itself. This hidden 28 + contribution to the apparent Si signal means that the grain boundary enrichment ratio is most likely to be an underestimate of the true Si grain boundary segregation. We have solved this problem 29 in a separate study using an Si-enriched SiO2 source and shown that Si indeed segregates strongly to alumina grain boundaries and is slightly soluble in bulk alumina [7]. Figure 2(b) presents a SIMS image of polycrystalline alumina singly doped with 500 p.p.m. Mg/Al sintered for 60 min at 16508C. The contrast in this + image scales with the Mg signal and shows clear segregation of Mg to grain boundaries. Also, there is evidence for bulk solid solubility of Mg in the grain centers. The ratio of grain boundary Mg to bulk Mg is estimated to be Cgb =Cgrain 0400 as shown in Fig. 3. The bright particle in the ®eld of view is believed to be a precipitate of MgAl2O4 spinel. No clear evidence for glass forming impurities, speci®cally Si, was seen in this specimen. Results of grain boundary segregation on codoping alumina with both Mg (Mg/Al=500 p.p.m.) and Si (Si/Al=1000 p.p.m.) sintered for 60 min at 16508C are also shown in Fig. 3. The measurement of the 28 + 28 + ratio of Si in the grain boundary region to Si in the bulk central grain region reveals a signi®cant
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drop in the enrichment ratio to Cgb =Cgrain 080, down from a ratio of 300 in singly doped 28 Si/Al=1000 p.p.m. alumina. The measurement of the ratio of Mg in the grain boundary region to Mg in the bulk is estimated to be Cgb =Cgrain 070, down from 0400 for singly doped Mg=Al 500 p:p:m: alumina. Clearly codoping has reduced the segregation of both Mg and Si to grain boundaries and has probably redistributed these dopants into the bulk alumina. In order to check whether the dopants are drawn into the bulk alumina structure, + we compare the ratio of the Mg signal from the grain centers in alumina codoped with MgO and SiO2 to alumina singly doped with MgO. This ratio is increased by a factor of 17. This estimate is accurate only to a factor of two due to variability in machine tuning and associated collection eciency of secondary. Note that this is not an issue in estimating boundary enrichment since we are comparing ion signals from one image at constant eciency. Despite this uncertainty in the estimate, we conclude that magnesia and silica are indeed brought into bulk solid solution in alumina. Our observations that MgO additions to alumina redistribute SiO2 provide direct support for the hypothesis by Handwerker et al. that MgO enhances the solid solubilities of glass-forming ions, such as Si, in alumina [13]. Further support for such increases in mutual solid solubility can be found in the work by Roy and Coble [22] who discovered that the solubility of TiO2 in alumina is signi®cantly enhanced by equimolar additions of MgO. Roy and Coble proposed a defect compensation reaction to account for this enhanced mutual solid solubility. We propose a similar defect reaction to explain the increased mutual solid solubility of both SiO2 and MgO in alumina on codoping. Using Kroger±Vink notation: Al2 O3
MgO SiO2 ÿ ÿ4 ÿ Mg 0 Al SiAl 3OxO x Mg 0 Al SiAl 3OxO 4 Mg 0 Al ±SiAl 3OxO :
3a 3b
The resulting Mg±Si defect complex is charge neutral and its solubility is primarily determined by the associated lattice strain. Also, charge neutralization by formation of this defect complex will reduce the interaction potential with grain boundaries. The new interaction potential will also be determined only by strain considerations and is consistent with the observed reduction in grain boundary segregation on codoping. We acknowledge that this is a simpli®ed view of the defect chemistry since the role of other unknown background impurities will in¯uence the exact defect equilibrium. In the next two sections we show the eect of changes in segregation of Si and Mg and solid solubility on alumina microstructure evolution and corrosion resistance.
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Fig. 4. Images of thermally etched alumina showing: (a) equiaxed grain structure for MgO singly doped alumina (Mg=Al 500 p:p:m:); (b) facetted abnormal grain structure for SiO2 singly doped alumina (Si=Al 1000 p:p:m:); (c) abnormal grain structure in MgO/SiO2 codoped alumina showing less facetted growth morphology.
4.2. Microstructure evolution In this section we compare microstructure morphology evolution for single doped (Mg or Si) and codoped (Mg plus Si) alumina and correlate observations of growth morphology with dopant distributions determined by HRI-SIMS. Figure 4(a) shows an example of the grain structure for MgO singly doped alumina (Mg/ Al=500 p.p.m.) sintered at 16508C. The morphology shows the classic equiaxed structure that is characteristic of MgO-doped aluminas [9, 10]. The underlying Mg distribution in this material is that shown in the HRI-SIMS map presented as Fig.
2(b), namely strong segregation of Mg to grain boundaries and some bulk solid solubility. Our observations are consistent with those made by Thompson and co-workers [4, 5] who used HRISIMS to identify segregation of MgO to grain boundaries in alumina. Coupled with kinetic observations of grain growth suppression on doping with MgO [23, 24], Thompson et al. [6] have proposed that MgO functions primarily by retarding grain growth through a solute drag mechanism. Figure 4(b) shows an example of abnormal grain growth for SiO2 singly doped alumina (Si/ Al=1000 p.p.m.) sintered at 16508C. The elongated abnormally growing grains have a distinct facetted characteristic. The long facets of these elongated grains are typically of {0001} orientation and contain thin glassy ®lms [14]. Strong segregation of Si to grain boundaries and pore surfaces in these aluminas can be seen in the HRI-SIMS map presented as Fig. 2(a). These observations are consistent with several studies that show the presence of small amounts of a glassy grain boundary phase leads to the initiation of abnormal grain growth. It has been proposed that this strong facetted morphology results from non-uniform wetting of the liquid ®lm and associated mobility dierences between wetted and non-wetted boundaries [14]. However, the HRI-SIMS image presented is direct evidence for complete boundary wetting. Figure 4(c) is an example of abnormal grain growth for alumina codoped with both MgO (Mg/ Al=500 p.p.m.) and SiO2 (Si/Al=1000 p.p.m.) sintered at 16508C. The abnormally growing grains display signi®cantly dierent growth morphology compared to alumina singly doped with SiO2 [Fig. 4(b)]. The grains are non-facetted and display curved boundaries in the codoped material as compared to SiO2 singly doped alumina. The codoping level is less than the equimolar Mg/Si concentration; however, measurements of dopant segregation to grain boundaries (Fig. 3) reveal that segregation of each cation is reduced signi®cantly. We believe that this reduction in Si segregation is sucient to prevent the formation of glassy grain boundary ®lms. The absence of such ®lms will signi®cantly aect the growth morphology by either: (1) changing the equilibrium grain shape; (2) adjusting mobility dierences between dierent boundary misorientations; or (3) aecting the local energy balance between intersecting boundaries and the local dihedral angle. Unequivocal identi®cation of the mechanism awaits further study. Microstructure evolution in codoped aluminas where the concentration of Mg exceeds that of Si leads to the suppression of abnormal grain growth and a grain morphology that is similar to MgO singly doped alumina shown in Fig. 4(a). In this case silica grain boundary ®lms are believed to be suppressed totally. Figure 5 plots critical Si and Mg concentrations for the initiation of abnormal grain
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Fig. 5. Critical Mg and Si doping concentrations for abnormal grain growth. The solid line represents the boundary for onset of the abnormal grain growth for samples sintered at 16508C for 1 h. Open circles represent specimens in which copious abnormal grain growth was observed. Filled circles represent specimens in which no abnormal grain growth was observed. The ®lled square symbols represent specimens in which only limited abnormal grain growth was observed.
growth during alumina sintering. The ®gure shows that microstructures will contain normal grains for Si/Mg concentration ratios less than unity, and abnormal grains for Si/Mg ratios greater than unity. Points on the plot are experimental observations. Samples with equimolar Si and Mg additions (i.e. 500 p.p.m. Mg/500 p.p.m. Al and 1000 p.p.m. Mg/1000 p.p.m. Al) are at the boundary between abnormal and normal grain morphologies. These transition samples primarily have uniaxial, normal grains with occasional inclusions of abnormal grains. The 458 line on the plot between normal and abnormal morphology in Fig. 6 re¯ects the underlying defect-compensation mechanism that determines segregation and mutual solid solubility as shown in equations (3a) and (3b). Our observations that MgO additions to alumina redistribute SiO2 provide direct support for the hypothesis that the important function of MgO as a sintering aid is to act as a ``scavenger'' that mitigates the deleterious eect of glass-forming ions on microstructure evolution. This hypothesis was originally made by Handwerker et al. [13] from studies of grain growth and microstructure evolution in aluminas of controlled purities. Once glassy ®lms are eliminated, the resulting segregation behavior acts to control grain growth through solute redistribution. This redistribution also stabilizes the microstructure against initiation of further abnormal grains [25]. 4.3. Tailoring grain boundary chemistry for corrosion resistance to aqueous HF The need to identify corrosion resistant ceramics for use under extreme environmental conditions has become increasingly important as ceramics are considered for materials of construction in increasingly
Fig. 6. Dependence of corrosion rate of polycrystalline alumina as a function of codoping with MgO and SiO2: (a) constant background of 500 p.p.m. Mg/Al, increasing Si additions; (b) constant background of 1000 p.p.m. Si/ Al, increasing Mg additions. In both cases, corrosion rate is signi®cantly diminished when the Mg concentration is increased above the equimolar codoping concentration.
aggressive chemical contacting applications at high temperature and pressure. Recently, it has been shown that the corrosion performance of polycrystalline ceramics in aqueous HF at elevated temperature is often controlled by the preferential dissolution of grain boundary phases [7]. Although the bulk crystal may be inherently resistant to corrosion, impurities such as silica that segregates to grain boundaries, are readily attacked resulting in the disintegration of the polycrystalline structure. The tendency for preferential attack by HF at siliceous grain boundaries occurs in most polycrystalline oxide ceramics. Additions of MgO to alumina have been shown to improve HF corrosion resistance by controlling grain boundary chemistry and microstructure evolution during sintering [8]. High purity alumina samples with MgO additions can be fabricated with corrosion resistances approaching single crystal sapphire.
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In this study alumina samples have been prepared with a constant background level of MgO (Mg/Al=500 p.p.m.) and monotonically increasing SiO2 additions. Figure 6(a) shows that the HF corrosion is mitigated up to SiO2 additions of Si/ Al=500 p.p.m., i.e. the equimolar codoping concentration. At SiO2 concentrations greater than the equimolar codoping concentration (Si/Mg concentration ratios greater than unity), the corrosion increases. Figure 6(b) shows HF corrosion for alumina samples codoped with increasing MgO concentrations and a constant background concentration of 1000 p.p.m. SiO2. The ®gure shows that corrosion is mitigated for samples with Mg/Al concentrations greater than the Si/Al level, i.e. Mg=Si > 1. Samples containing MgO concentrations less than the equimolar codoping level demonstrate corrosion rates that are four orders of magnitude greater than for samples with Mg concentrations above the equimolar codoping level. These observations suggest that MgO additions modify the grain boundary chemistry by removing SiO2 from the grain boundaries. Quantitative SIMS analysis of grain boundary distributions of MgO and SiO2 shown in Fig. 3 indicate that MgO indeed diminishes SiO2 segregation to grain boundaries on codoping. The removal of silica-based glassy grain boundary ®lms results in an increase in corrosion resistance as shown in the corrosion data. The implication is that silica ®lms are removed from the grain boundaries by redistribution into the bulk by a mutual defect compensation mechanism, as discussed above and shown in equations (3a) and (3b). To mitigate corrosion, the Mg/Si concentration ratio must be equal to or greater than unity in order to remove silica-based, glassy grain boundary ®lms. Samples with Mg/Si concentration ratios less then unity behave similar to silica (i.e. rapid corrosion of siliceous grain boundaries) and sample ratios greater than unity behave similar to sapphire (i.e. corrosion of bulk alumina). The manufacture of corrosion-resistant aluminas by grain boundary engineering has been exploited in HF-containing reaction engineering [8]. 5. CONCLUSIONS
HRI-SIMS is a powerful technique for the study of dopant and impurity eects in the processing and properties of ceramics. It provides key insights into compositional characteristics and may be readily integrated into an experimental program and combined with other characterization techniques. HRI-SIMS provides direct evidence for the segregation of either SiO2 or MgO dopants to grain boundaries in polycrystalline alumina. On codoping with both MgO and SiO2, grain boundary segregation is signi®cantly diminished over single doping as both cations are redistributed into the bulk alumina lattice. A defect compensation mechanism
is proposed for this reaction. Dopant redistribution changes abnormal grain growth morphology from facetted grains in SiO2 singly doped alumina, to non-faceted grains with curved boundaries in MgO and SiO2 codoped alumina. As the Mg/Si dopant ratio exceeds the equimolar concentration, abnormal grain growth development ceases. These ®ndings provide a physical mechanism to explain the role of MgO as a sintering aid to control microstructure evolution in alumina. MgO doping is observed to improve the corrosion resistance of alumina to aqueous HF through the removal of corrodable silica-based grain boundary ®lm doping. AcknowledgementsÐWe thank Dan Schaeer (DuPont) for experimental assistance. Rowland Cannon (LBL) and Carol Handwerker (NIST) are acknowledged for stimulating discussions. This work was supported by the NSF under award DMR-9625354 and through the MRSEC program under NSF award DMR-9400379.
REFERENCES 1. Levi-Setti, R., Chabala, J. M., Gavrilov, K. L., Mogilevsky, R. and Soni, K. K., Scanning Microscopy, 1993, 7(4), 1161. 2. Soni, K. K., Kang, H. G., Grant, P. S., Cantor, B., Adriaens, A. G., Gavrilov, K. L., Mogilevsky, R., Levi-Setti, R., Tseng, M. W. and Williams, D. B., J. Microsc., 1994, 177(3), 414. 3. Chabala, J. M., Soni, K. K., Li, J., Gavrilov, K. L. and Levi-Setti, R., Int. J. Mass Spectrometry Ion Processes, 1995, 143, 191. 4. Soni, K. K., Thompson, A. M., Harmer, M. P., Williams, D. B., Chabala, J. M. and Levi-Setti, R., J. appl. Phys., 1995, 66(21), 2795. 5. German, R. M., Messing, G. L. and Cornwall, R. G., in Sintering Technology. Marcel Dekker, New York, 1996, p. 309. 6. Thompson, A. M., Soni, K. K., Chan, H. M., Harmer, M. P., Williams, D. B., Chabala, J. M. and Levi-Setti, R., J. Am. Ceram. Soc., 1997, 80(2), 373. 7. Gavrilov, K. L., Bennison, S. J., Mikeska, K. R., Chabala, J. M. and Levi-Setti, R., J. Am. Ceram. Soc., 1999, 82(4), 1001 8. Bennison, S. J. and Mikeska, K. R., HF-resistant ceramics and use thereof. U.S. Patent 5,411,583, May, 1995. 9. Coble, R. L., J. appl. Phys., 1961, 32(5), 793. 10. Coble, R. L., Transparent alumina and method of preparation. U.S. Patent 026,210, March, 1962. 11. Handwerker, C. A., Blendell, J. E. and Kaisser, W. A., in Ceramic Transactions, Vol. 7. The American Ceramic Society, Westerville, OH, 1990, p. 13. 12. Kaisser, W. A., Sprissler, M., Handwerker, C. A. and Blendell, J. E., J. Am. Ceram. Soc., 1987, 70(5), 339. 13. Handwerker, C. A., Morris, P. A. and Coble, R. L., J. Am. Ceram. Soc., 1989, 72(1), 130. 14. Bateman, C. A., Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1989, 72(7), 1241. 15. Bae, S. I. and Baik, S., J. Am. Ceram. Soc., 1993, 76(4), 1065. 16. Yan, M. F., Mater. Sci. Engng, 1981, 48, 53. 17. Mikeska, K. R., Bennison, S. J. and Gavrilov, K. L., J. Am. Ceram. Soc., submitted. 18. Chabala, J. M., Gavrilov, K. L., Mikeska, K. R., Bennison, S. J. and Levi-Setti, R., Proceedings of the
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY Microscopy and Microanalysis `97 Meeting, Cleveland, OH, 10±14 August. Springer, Berlin, 1997, p. 525. 19. Gavrilov, K. L., Chabala, J. M., Levi-Setti, R., Mikeska, K. R. and Bennison, S. J., Proceedings of the XI International Conference on Secondary Ion Mass Spectrometry (SIMS-XI), Orlando, FL, 1997. 20. Benninghoven, A., Rudenauer, F. G. and Werner, H. W., in Secondary Ion Mass Spectrometry. John Wiley, New York, 1987, p. 277. 21. Brydson, R., Chen, S.-C., Riley, F. L., Milne, S. J.,
22. 23. 24. 25.
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Pan, X. and Ruhle, M., J. Am. Ceram. Soc., 1998, 81(2), 369. Roy, S. K. and Coble, R. L., J. Am. Ceram. Soc., 1968, 5(1), 1. Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1983, 66(5), C90. Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1985, 68(1), C22. Gavrilov, K., Bennison, S. J., Mikeska, K. R., and Levi-Setti, R., J. Am. Ceram. Soc., submitted.
Acta mater. Vol. 47, Nos 15, pp. 4031±4039, 1999 # 1999 Published by Elsevier Science Ltd On behalf of Acta Metallurgica Inc. All rights reserved. Printed in Great Britain S1359-6454(99)00263-3 1359-6454/99 $20.00 + 0.00
GRAIN BOUNDARY CHEMISTRY OF ALUMINA BY HIGH-RESOLUTION IMAGING SIMS K. L. GAVRILOV 1{, S. J. BENNISON 2, K. R. MIKESKA 2 and R. LEVI-SETTI 1 Enrico Fermi Institute and Department of Physics, University of Chicago, Chicago, IL 60637, U.S.A. and 2E.I. DuPont de Nemours & Co. Inc., Central Research and Development, Experimental Station, Wilmington, DE 19880, U.S.A.
1
AbstractÐThe unique capabilities of the high-resolution scanning ion microprobe developed at the University of Chicago (UC-SIM) are described and its utility is demonstrated in a study of grain boundary chemistry of alumina ceramics. When polycrystalline alumina is doped singly with either MgO or SiO2, strong segregation of the individual ions to grain boundaries is observed: (1) for Mg segregation Cgb =Cgrain 400; (2) for Si segregation Cgb =Cgrain 300. However, on codoping with both MgO and SiO2, grain boundary segregation is signi®cantly diminished by a factor of ®ve or more over single doping as both cations are redistributed into the bulk alumina lattice. A defect compensation mechanism is proposed to explain this mutual solid solubility of Mg and Si in alumina. One important consequence of this chemical redistribution is a change in abnormal grain growth morphology from facetted grains in SiO2 singly doped alumina, to non-facetted grains with curved boundaries in MgO and SiO2 codoped alumina. As the Mg/Si dopant ratio exceeds the equimolar concentration, abnormal grain growth development ceases. These ®ndings provide a physical mechanism to explain the role of MgO as a sintering aid to control microstructure evolution in alumina. Another signi®cant consequence of SiO2 redistribution on MgO doping is an observed improvement in the corrosion resistance of alumina to aqueous HF. Siliceous grain boundary ®lms readily corrode and compromise the intrinsically good corrosion resistance of bulk alumina. MgO doping in amounts greater than the SiO2 concentration prevents the formation of these corrodable silica-based phases leading to the development of aluminas for use in aqueous HF-containing ambients. # 1999 Published by Elsevier Science Ltd on behalf of Acta Metallurgica Inc. All rights reserved. Keywords: Ion microprobe; Secondary ion mass spectroscopy (SIMS); Ceramics; Microstructure; Corrosion
1. INTRODUCTION
High-resolution imaging secondary ion mass spectrometry (HRI-SIMS) has emerged as a valuable research technique to study interfacial chemistry of a variety of modern engineered materials such as metal alloys, composites and ceramics [1±6]. The application of HRI-SIMS has been driven by advances in the design of interface chemistry and structure to tailor electrical, magnetic and optical properties of materials. HRI-SIMS has been employed in the interfacial design of ceramics [7] where the in¯uence of trace impurities on sintering and microstructure development is exploited in the manufacture of materials with speci®ed physical performance [8]. A number of important innovations introduced in scanning ion microprobe (SIM) technology have created an opportunity to address the role of trace impurities in interfacial chemistry and associated properties. One innovation includes the development of a liquid-metal ion source (LMIS) that enables the theoretical lateral resolution limit for SIM instruments to be attained [3]. Another is the incorporation, in the University of Chicago scan{To whom all correspondence should be addressed.
ning ion microprobe (UC-SIM), of an ecient secondary ion transport system together with the addition of a Finnigan'90-based Magnetic Sector Mass Spectrometer that provides analytical imaging SIMS sensitivity in the range of 1 p.p.m. [3] (Fig. 1). In this paper we discuss the use of HRI-SIMS in combination with other analytical tools and experimental methodologies to understand the role of speci®c dopants and impurities in microstructure development and associated corrosion resistance of ceramics. Trace impurities and additives have long been known to strongly in¯uence the fabrication and properties of alumina-based ceramics. For example, MgO additions improve sintering characteristics and allow the manufacture of aluminas with controlled microstructures [9±11]. More commonly, impurities such as silica and glass-modifying ions tend to have a detrimental eect on processing since their presence often promotes abnormal grain growth [12±15]. These latter impurities often compromise properties that are sensitive to the presence of glassy grain boundary ®lms, such as electrical properties [16] and corrosion resistance [8, 17]. The detailed function of background impurities and trace additives has been dicult to ascertain due to problems in the detection of additives and their as-
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Fig. 1. Schematic of University of Chicago high-resolution imaging secondary ion mass spectrometer (UC-SIM).
sociated distribution in the microstructure. The HRI-SIMS is well suited to address the problem of trace element detection and distribution, particularly in interfacial regions such as grain boundaries. In this contribution we summarize the unique capabilities of the UC-SIM and discuss the advantages and limitations of this analytical tool. The use of HRI-SIMS is illustrated in three studies on alumina ceramics. The ®rst addresses segregation of dopants and impurities to grain boundaries. We show that codoping with magnesia and silica strongly aects their bulk mutual solid solubility and associated segregation characteristics to grain boundaries. The second study looks at microstructure evolution as a function of grain boundary chemistry. The segregation characteristics observed in the ®rst study play a determining role in grain morphology and growth kinetics. The third study considers the eect of segregated impurities on controlling corrosion resistance. We demonstrate that, with suitable control of grain boundary chemistry, polycrystalline aluminas may be engineered with corrosion resistance approaching that of single-crystal sapphire.
2. CAPABILITIES OF HIGH RESOLUTION IMAGING SIMS
In a scanning ion microprobe, a focused beam of primary ions is rastered over a specimen in order to
erode via sputtering the outermost surface layers. A fraction of the resulting ejected material is emitted as ionized atoms or molecules; this ionized fraction varies among elements and depends primarily on the electronic state of the sputtered material and on the chemical characteristics of the surface. The secondary ions and molecular clusters (either positively or negatively charged) are gathered and discriminated on the basis of their mass/charge ratio with a mass spectrometer. By recording the secondary ion signal as a function of the position of the scanning beam, two-dimensional compositional distribution maps are obtained. These SIMS maps can be digitally stored for subsequent image processing and analysis. HRI-SIMS is a powerful material characterization technique that simultaneously spatially locates and quanti®es dopant and impurity distribution in the microstructure. This simultaneous combination becomes especially important for advanced engineered materials where the detailed function of background impurities and additives has been debated for many years [11]. The UC-SIM instrument is ideally suited to address the questions concerning the role of impurities and additives, and has a number of advantages over other analytical techniques and ``conventional'' SIMS: (a) Sensitivity: HRI-SIMS can detect all elements and isotopes, many of them at part per million concentration levels. The signal-to-noise
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY 6
ratio can be as high as 10 Ðessentially, there is no background. (b) Spatial resolution: The spatial and depth resolutions of HRI-SIMS are excellent. A lateral resolution of 20±50 nm has been attained with the UC-SIM [3]. A depth resolution of 1±10 nm is common. However, there is a trade-o between sensitivity and resolution. (c) Speed of analysis: SIMS distribution maps can typically be acquired in 1±9 min time scales depending on the concentration of elements of interest. With this rate, it is possible to perform comprehensive analyses of multicomponent samples in practical times. (d) Simultaneous multigrain imaging: a large number of grain boundaries are readily analyzed during construction of an elemental map. For 2 example a map of a 2020 mm region will contain several hundred grain boundaries for a material with an average grain size of 2 mm. (e) Specimen preparation: bulk materials polished with standard ceramographic procedures can be studied. Additionally, the specimen surface can be sputter-cleaned in situ with the primary ion beam or with an auxiliary high-current ion beam. In practice this means that surface artifacts introduced during preparation can frequently be identi®ed and sputter-erased. Spatial resolution approaching the probe size (typically 035 nm) in scanned SIMS maps can be attained when the corresponding secondary ion signal is suciently large. Thus, the ``analytical resolution'' in elemental maps is conceptually dierent from ``structural resolution'' in the context of light or electron micrographs. Even for dopants or impurities present at the trace level in bulk, it is generally possible to detect grain boundary segregation with the UC-SIM. We have demonstrated this capability routinely in the analysis of grain boundaries in polycrystalline ceramics. Further details of this instrument and its capability in analyzing ceramic and composite microstructures can be found elsewhere [1, 18, 19]. Quanti®cation of the secondary ion signal is often problematic and is the primary limitation of the SIMS method. This problem arises from variations in the secondary particle sputtering yield and ion fraction as a function of the matrix structure and composition, and by the emission of neutral species. This limitation can be overcome using either rigorous calibration standards [20] or by combining SIMS with other characterization techniques. In many cases internal standards may be used such as stoichiometric precipitates in the microstructure. Also, relative concentration measurements may be employed to quantify the extent of grain boundary segregation relative to the bulk.
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A further complication in SIMS analysis concerns discrimination between sputtered complexes of similar mass/charge ratio. For example, in the case of 28 + alumina with SiO2 impurities, sputtered Si ions + have a similar mass/charge ratio to AlH and resolving these complexes is dicult. In principle, these two species could be resolved by the use of the Finnigan'90-based Magnetic Sector Mass Spectrometer, but at the expense of SIMS sensitivity. A de®nitive solution is obtained for the study of dopants by using isotopes with appropriate mass where possible. In the case of Si dopants in 29 alumina, Si-containing SiO2 may be used to 29 + + dierentiate sputtered Si from AlH complexes [7].
3. EXPERIMENTAL STUDIES
3.1. Sample preparation Sumitomo AKP-30 alumina was used as the base powder for sample preparation. This powder, according to Sumitomo, is 99.995% pure, with a 0.5 mm mean crystallite size and 50% of agglomerates below 1 mm. Powder handling was carried in ``semi-clean'' conditions using a laminar ¯ow hood 1 and clean Te¯on laboratory ware. Doping was achieved by adding aliquots of various metal salt complexes dissolved in water or ethanol followed by drying and/or hydrolyzing and deagglomeration. Additions of MgO were made using Mg(NO3)24H2O (99.99% purity). Additions of SiO2 were made using tetraethyl-orthosilicate (TEOS) (99.99% purity). Green pellets were prepared by dry pressing at 75 MPa followed by cold isostatic pressing at 224 MPa. Sample pellets were ®red in clean covered alumina crucibles containing protective powder identical in composition to the pellets. A resistance heated air furnace was employed for ®ring. The ®ring schedule consisted of calcining at 10008C for 2 h, ramping to a peak soak temperature of 16508C, and annealing for 60 min at 16508C. Specimens were cooled in a furnace powero condition, which results in a cooling rate on the order of 6008C/h at temperatures near the sintering temperature. 3.2. Microstructural characterization The sintered aluminas were polished to ¯at mirror surface ®nishes using successively ®ner grades of diamond abrasives down to 0.5 mm. Charging was not a problem with these materials providing that a thin coating of gold was applied to the polished section prior to imaging with the SIMS microprobe. Images were collected after sputtering o the gold layer and any surface contamination resulting from specimen preparation (i.e. polishing). Estimates of grain boundary segregation were made by comparing ion yields from grain boundary regions, Ygb, to ion yields from grain centers, Ygrain:
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Enrichment factor
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY
Ygb Ion probe diameter : 1 Ygrain Grain boundary width
It is assumed that matrix eects are uniform across grain boundaries. A 40 nm diameter ion probe and a boundary width of 1 nm were used to calculate grain boundary enrichments for both Mg [4] and Si segregation [21]. We acknowledge that the segregation characteristics of Mg and Si are likely to be dierent. Magnesium probably segregates in a classical solid±solution sense and Si most likely forms a distinct grain boundary ®lm. However, the enrichment width scale for both these distinct segregation mechanisms is similar and on the order of 1 nm. The uncertainty in the estimated enrichment values is 220% of the mean. Scanning electron microscopy (SEM) and optical microscopy were used to characterize the microstructural morphology of the sintered alumina samples. Specimens for microscopy were polished as described above. The polished sections were subsequently thermally etched at 15508C for 10 min in air to reveal the microstructure morphology. All specimens were sputter-coated with a thin (03 nm) layer of gold/palladium to mitigate sample charging in the SEM and enhance re¯ectivity for optical microscopy. The sequence of sample characterizations was: (1) examine microstructure by SEM or optical microscopy to identify areas of interest; (2) mark areas of interest using Vickers indentations; (3) repolish to remove features created by thermal etching; (4) sputter coat with thin gold layer; (5) locate regions of interest in HRI-SIMS from residual hardness impressions; (6) sputter o gold layer and surface contamination; and (7) analyze region of interest by HRI-SIMS. 3.3. HF corrosion test protocol The HF corrosion test protocol consisted of subjecting specimens to an azeotropic 38.26 HF/61.74 H2O (w/o) solution at 9020.58C. The details of the protocol are discussed elsewhere [7]. The HF conÿ centration was 20 M and the F ion concentration ÿ2 was 8:9 10 M at the azeotropic composition. Ceramic specimens of approximately cubic geome3 try with a nominal volume of 1 cm were subjected to corrosion testing by immersing individually in 150 ml of the HF/H2O azeotrope contained in a Te¯on reaction ¯ask. An eective linear corrosion rate, D_ , was determined for each specimen from weight-loss measurements. D_ is typically reported in ml/y or mm/y and is given by @W D_ @ tsr
2
where @ W W0 ÿ Wi (W0 is the initial sample weight and Wi the instantaneous or ®nal sample weight), s is the sample surface area, r is the material density, and t is time. D_ is a measure of the rate of material thinning or the rate at which the
Fig. 2. SIMS maps of singly doped polycrystalline alumina sintered at 16508C: (a) Si+ map for alumina doped with Si/Al=1000 p.p.m.; (b) Mg+ map for alumina with Mg/ Al=500 p.p.m.
material uniformly corrodes away at its exposed surfaces. 4. RESULTS AND DISCUSSION
4.1. Segregation of Si and Mg to grain boundaries The distribution of SiO2 dopant in a 1000 p.p.m. Si/Al specimen sintered for 480 min at 16508C is shown in Fig. 2(a). The grain morphology displays a small number of elongated (abnormal) grains containing patches of ®ner grains, which is typical of aluminas contaminated or doped with small amounts of glass-forming compounds [12±15]. The long facets of these elongated grains are typically of {0001} orientation and contain thin glassy ®lms
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY
Fig. 3. Grain boundary enrichment ratio (grain boundary ion-yield signal/grain center ion-yield signal) plotted for polycrystalline alumina doped singly with MgO or SiO2 and codoped with MgO and SiO2. Note the strong drop in both Mg and Si segregation to grain boundaries on codoping with both elements.
[14]. Using an integrated secondary ion image (ISI) topography map of this region (not shown) there is 28 + no indication of a Si signal corresponding to the location of surface polishing scratches in this region. Strong segregation of Si to grain boundaries and pore surfaces can be seen in Fig. 2(a). The ratio of Si in the boundary region to grain center is Cgb =Cgrain 0300 as shown in Fig. 3. It should be 28 + noted that the Si signal used to form these images contains a component made up of sputtered + AlH molecular ions that form due to a reaction between sputtered Al and residual hydrogen in the vacuum chamber and/or alumina itself. This hidden 28 + contribution to the apparent Si signal means that the grain boundary enrichment ratio is most likely to be an underestimate of the true Si grain boundary segregation. We have solved this problem 29 in a separate study using an Si-enriched SiO2 source and shown that Si indeed segregates strongly to alumina grain boundaries and is slightly soluble in bulk alumina [7]. Figure 2(b) presents a SIMS image of polycrystalline alumina singly doped with 500 p.p.m. Mg/Al sintered for 60 min at 16508C. The contrast in this + image scales with the Mg signal and shows clear segregation of Mg to grain boundaries. Also, there is evidence for bulk solid solubility of Mg in the grain centers. The ratio of grain boundary Mg to bulk Mg is estimated to be Cgb =Cgrain 0400 as shown in Fig. 3. The bright particle in the ®eld of view is believed to be a precipitate of MgAl2O4 spinel. No clear evidence for glass forming impurities, speci®cally Si, was seen in this specimen. Results of grain boundary segregation on codoping alumina with both Mg (Mg/Al=500 p.p.m.) and Si (Si/Al=1000 p.p.m.) sintered for 60 min at 16508C are also shown in Fig. 3. The measurement of the 28 + 28 + ratio of Si in the grain boundary region to Si in the bulk central grain region reveals a signi®cant
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drop in the enrichment ratio to Cgb =Cgrain 080, down from a ratio of 300 in singly doped 28 Si/Al=1000 p.p.m. alumina. The measurement of the ratio of Mg in the grain boundary region to Mg in the bulk is estimated to be Cgb =Cgrain 070, down from 0400 for singly doped Mg=Al 500 p:p:m: alumina. Clearly codoping has reduced the segregation of both Mg and Si to grain boundaries and has probably redistributed these dopants into the bulk alumina. In order to check whether the dopants are drawn into the bulk alumina structure, + we compare the ratio of the Mg signal from the grain centers in alumina codoped with MgO and SiO2 to alumina singly doped with MgO. This ratio is increased by a factor of 17. This estimate is accurate only to a factor of two due to variability in machine tuning and associated collection eciency of secondary. Note that this is not an issue in estimating boundary enrichment since we are comparing ion signals from one image at constant eciency. Despite this uncertainty in the estimate, we conclude that magnesia and silica are indeed brought into bulk solid solution in alumina. Our observations that MgO additions to alumina redistribute SiO2 provide direct support for the hypothesis by Handwerker et al. that MgO enhances the solid solubilities of glass-forming ions, such as Si, in alumina [13]. Further support for such increases in mutual solid solubility can be found in the work by Roy and Coble [22] who discovered that the solubility of TiO2 in alumina is signi®cantly enhanced by equimolar additions of MgO. Roy and Coble proposed a defect compensation reaction to account for this enhanced mutual solid solubility. We propose a similar defect reaction to explain the increased mutual solid solubility of both SiO2 and MgO in alumina on codoping. Using Kroger±Vink notation: Al2 O3
MgO SiO2 ÿ ÿ4 ÿ Mg 0 Al SiAl 3OxO x Mg 0 Al SiAl 3OxO 4 Mg 0 Al ±SiAl 3OxO :
3a 3b
The resulting Mg±Si defect complex is charge neutral and its solubility is primarily determined by the associated lattice strain. Also, charge neutralization by formation of this defect complex will reduce the interaction potential with grain boundaries. The new interaction potential will also be determined only by strain considerations and is consistent with the observed reduction in grain boundary segregation on codoping. We acknowledge that this is a simpli®ed view of the defect chemistry since the role of other unknown background impurities will in¯uence the exact defect equilibrium. In the next two sections we show the eect of changes in segregation of Si and Mg and solid solubility on alumina microstructure evolution and corrosion resistance.
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Fig. 4. Images of thermally etched alumina showing: (a) equiaxed grain structure for MgO singly doped alumina (Mg=Al 500 p:p:m:); (b) facetted abnormal grain structure for SiO2 singly doped alumina (Si=Al 1000 p:p:m:); (c) abnormal grain structure in MgO/SiO2 codoped alumina showing less facetted growth morphology.
4.2. Microstructure evolution In this section we compare microstructure morphology evolution for single doped (Mg or Si) and codoped (Mg plus Si) alumina and correlate observations of growth morphology with dopant distributions determined by HRI-SIMS. Figure 4(a) shows an example of the grain structure for MgO singly doped alumina (Mg/ Al=500 p.p.m.) sintered at 16508C. The morphology shows the classic equiaxed structure that is characteristic of MgO-doped aluminas [9, 10]. The underlying Mg distribution in this material is that shown in the HRI-SIMS map presented as Fig.
2(b), namely strong segregation of Mg to grain boundaries and some bulk solid solubility. Our observations are consistent with those made by Thompson and co-workers [4, 5] who used HRISIMS to identify segregation of MgO to grain boundaries in alumina. Coupled with kinetic observations of grain growth suppression on doping with MgO [23, 24], Thompson et al. [6] have proposed that MgO functions primarily by retarding grain growth through a solute drag mechanism. Figure 4(b) shows an example of abnormal grain growth for SiO2 singly doped alumina (Si/ Al=1000 p.p.m.) sintered at 16508C. The elongated abnormally growing grains have a distinct facetted characteristic. The long facets of these elongated grains are typically of {0001} orientation and contain thin glassy ®lms [14]. Strong segregation of Si to grain boundaries and pore surfaces in these aluminas can be seen in the HRI-SIMS map presented as Fig. 2(a). These observations are consistent with several studies that show the presence of small amounts of a glassy grain boundary phase leads to the initiation of abnormal grain growth. It has been proposed that this strong facetted morphology results from non-uniform wetting of the liquid ®lm and associated mobility dierences between wetted and non-wetted boundaries [14]. However, the HRI-SIMS image presented is direct evidence for complete boundary wetting. Figure 4(c) is an example of abnormal grain growth for alumina codoped with both MgO (Mg/ Al=500 p.p.m.) and SiO2 (Si/Al=1000 p.p.m.) sintered at 16508C. The abnormally growing grains display signi®cantly dierent growth morphology compared to alumina singly doped with SiO2 [Fig. 4(b)]. The grains are non-facetted and display curved boundaries in the codoped material as compared to SiO2 singly doped alumina. The codoping level is less than the equimolar Mg/Si concentration; however, measurements of dopant segregation to grain boundaries (Fig. 3) reveal that segregation of each cation is reduced signi®cantly. We believe that this reduction in Si segregation is sucient to prevent the formation of glassy grain boundary ®lms. The absence of such ®lms will signi®cantly aect the growth morphology by either: (1) changing the equilibrium grain shape; (2) adjusting mobility dierences between dierent boundary misorientations; or (3) aecting the local energy balance between intersecting boundaries and the local dihedral angle. Unequivocal identi®cation of the mechanism awaits further study. Microstructure evolution in codoped aluminas where the concentration of Mg exceeds that of Si leads to the suppression of abnormal grain growth and a grain morphology that is similar to MgO singly doped alumina shown in Fig. 4(a). In this case silica grain boundary ®lms are believed to be suppressed totally. Figure 5 plots critical Si and Mg concentrations for the initiation of abnormal grain
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY
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Fig. 5. Critical Mg and Si doping concentrations for abnormal grain growth. The solid line represents the boundary for onset of the abnormal grain growth for samples sintered at 16508C for 1 h. Open circles represent specimens in which copious abnormal grain growth was observed. Filled circles represent specimens in which no abnormal grain growth was observed. The ®lled square symbols represent specimens in which only limited abnormal grain growth was observed.
growth during alumina sintering. The ®gure shows that microstructures will contain normal grains for Si/Mg concentration ratios less than unity, and abnormal grains for Si/Mg ratios greater than unity. Points on the plot are experimental observations. Samples with equimolar Si and Mg additions (i.e. 500 p.p.m. Mg/500 p.p.m. Al and 1000 p.p.m. Mg/1000 p.p.m. Al) are at the boundary between abnormal and normal grain morphologies. These transition samples primarily have uniaxial, normal grains with occasional inclusions of abnormal grains. The 458 line on the plot between normal and abnormal morphology in Fig. 6 re¯ects the underlying defect-compensation mechanism that determines segregation and mutual solid solubility as shown in equations (3a) and (3b). Our observations that MgO additions to alumina redistribute SiO2 provide direct support for the hypothesis that the important function of MgO as a sintering aid is to act as a ``scavenger'' that mitigates the deleterious eect of glass-forming ions on microstructure evolution. This hypothesis was originally made by Handwerker et al. [13] from studies of grain growth and microstructure evolution in aluminas of controlled purities. Once glassy ®lms are eliminated, the resulting segregation behavior acts to control grain growth through solute redistribution. This redistribution also stabilizes the microstructure against initiation of further abnormal grains [25]. 4.3. Tailoring grain boundary chemistry for corrosion resistance to aqueous HF The need to identify corrosion resistant ceramics for use under extreme environmental conditions has become increasingly important as ceramics are considered for materials of construction in increasingly
Fig. 6. Dependence of corrosion rate of polycrystalline alumina as a function of codoping with MgO and SiO2: (a) constant background of 500 p.p.m. Mg/Al, increasing Si additions; (b) constant background of 1000 p.p.m. Si/ Al, increasing Mg additions. In both cases, corrosion rate is signi®cantly diminished when the Mg concentration is increased above the equimolar codoping concentration.
aggressive chemical contacting applications at high temperature and pressure. Recently, it has been shown that the corrosion performance of polycrystalline ceramics in aqueous HF at elevated temperature is often controlled by the preferential dissolution of grain boundary phases [7]. Although the bulk crystal may be inherently resistant to corrosion, impurities such as silica that segregates to grain boundaries, are readily attacked resulting in the disintegration of the polycrystalline structure. The tendency for preferential attack by HF at siliceous grain boundaries occurs in most polycrystalline oxide ceramics. Additions of MgO to alumina have been shown to improve HF corrosion resistance by controlling grain boundary chemistry and microstructure evolution during sintering [8]. High purity alumina samples with MgO additions can be fabricated with corrosion resistances approaching single crystal sapphire.
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In this study alumina samples have been prepared with a constant background level of MgO (Mg/Al=500 p.p.m.) and monotonically increasing SiO2 additions. Figure 6(a) shows that the HF corrosion is mitigated up to SiO2 additions of Si/ Al=500 p.p.m., i.e. the equimolar codoping concentration. At SiO2 concentrations greater than the equimolar codoping concentration (Si/Mg concentration ratios greater than unity), the corrosion increases. Figure 6(b) shows HF corrosion for alumina samples codoped with increasing MgO concentrations and a constant background concentration of 1000 p.p.m. SiO2. The ®gure shows that corrosion is mitigated for samples with Mg/Al concentrations greater than the Si/Al level, i.e. Mg=Si > 1. Samples containing MgO concentrations less than the equimolar codoping level demonstrate corrosion rates that are four orders of magnitude greater than for samples with Mg concentrations above the equimolar codoping level. These observations suggest that MgO additions modify the grain boundary chemistry by removing SiO2 from the grain boundaries. Quantitative SIMS analysis of grain boundary distributions of MgO and SiO2 shown in Fig. 3 indicate that MgO indeed diminishes SiO2 segregation to grain boundaries on codoping. The removal of silica-based glassy grain boundary ®lms results in an increase in corrosion resistance as shown in the corrosion data. The implication is that silica ®lms are removed from the grain boundaries by redistribution into the bulk by a mutual defect compensation mechanism, as discussed above and shown in equations (3a) and (3b). To mitigate corrosion, the Mg/Si concentration ratio must be equal to or greater than unity in order to remove silica-based, glassy grain boundary ®lms. Samples with Mg/Si concentration ratios less then unity behave similar to silica (i.e. rapid corrosion of siliceous grain boundaries) and sample ratios greater than unity behave similar to sapphire (i.e. corrosion of bulk alumina). The manufacture of corrosion-resistant aluminas by grain boundary engineering has been exploited in HF-containing reaction engineering [8]. 5. CONCLUSIONS
HRI-SIMS is a powerful technique for the study of dopant and impurity eects in the processing and properties of ceramics. It provides key insights into compositional characteristics and may be readily integrated into an experimental program and combined with other characterization techniques. HRI-SIMS provides direct evidence for the segregation of either SiO2 or MgO dopants to grain boundaries in polycrystalline alumina. On codoping with both MgO and SiO2, grain boundary segregation is signi®cantly diminished over single doping as both cations are redistributed into the bulk alumina lattice. A defect compensation mechanism
is proposed for this reaction. Dopant redistribution changes abnormal grain growth morphology from facetted grains in SiO2 singly doped alumina, to non-faceted grains with curved boundaries in MgO and SiO2 codoped alumina. As the Mg/Si dopant ratio exceeds the equimolar concentration, abnormal grain growth development ceases. These ®ndings provide a physical mechanism to explain the role of MgO as a sintering aid to control microstructure evolution in alumina. MgO doping is observed to improve the corrosion resistance of alumina to aqueous HF through the removal of corrodable silica-based grain boundary ®lm doping. AcknowledgementsÐWe thank Dan Schaeer (DuPont) for experimental assistance. Rowland Cannon (LBL) and Carol Handwerker (NIST) are acknowledged for stimulating discussions. This work was supported by the NSF under award DMR-9625354 and through the MRSEC program under NSF award DMR-9400379.
REFERENCES 1. Levi-Setti, R., Chabala, J. M., Gavrilov, K. L., Mogilevsky, R. and Soni, K. K., Scanning Microscopy, 1993, 7(4), 1161. 2. Soni, K. K., Kang, H. G., Grant, P. S., Cantor, B., Adriaens, A. G., Gavrilov, K. L., Mogilevsky, R., Levi-Setti, R., Tseng, M. W. and Williams, D. B., J. Microsc., 1994, 177(3), 414. 3. Chabala, J. M., Soni, K. K., Li, J., Gavrilov, K. L. and Levi-Setti, R., Int. J. Mass Spectrometry Ion Processes, 1995, 143, 191. 4. Soni, K. K., Thompson, A. M., Harmer, M. P., Williams, D. B., Chabala, J. M. and Levi-Setti, R., J. appl. Phys., 1995, 66(21), 2795. 5. German, R. M., Messing, G. L. and Cornwall, R. G., in Sintering Technology. Marcel Dekker, New York, 1996, p. 309. 6. Thompson, A. M., Soni, K. K., Chan, H. M., Harmer, M. P., Williams, D. B., Chabala, J. M. and Levi-Setti, R., J. Am. Ceram. Soc., 1997, 80(2), 373. 7. Gavrilov, K. L., Bennison, S. J., Mikeska, K. R., Chabala, J. M. and Levi-Setti, R., J. Am. Ceram. Soc., 1999, 82(4), 1001 8. Bennison, S. J. and Mikeska, K. R., HF-resistant ceramics and use thereof. U.S. Patent 5,411,583, May, 1995. 9. Coble, R. L., J. appl. Phys., 1961, 32(5), 793. 10. Coble, R. L., Transparent alumina and method of preparation. U.S. Patent 026,210, March, 1962. 11. Handwerker, C. A., Blendell, J. E. and Kaisser, W. A., in Ceramic Transactions, Vol. 7. The American Ceramic Society, Westerville, OH, 1990, p. 13. 12. Kaisser, W. A., Sprissler, M., Handwerker, C. A. and Blendell, J. E., J. Am. Ceram. Soc., 1987, 70(5), 339. 13. Handwerker, C. A., Morris, P. A. and Coble, R. L., J. Am. Ceram. Soc., 1989, 72(1), 130. 14. Bateman, C. A., Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1989, 72(7), 1241. 15. Bae, S. I. and Baik, S., J. Am. Ceram. Soc., 1993, 76(4), 1065. 16. Yan, M. F., Mater. Sci. Engng, 1981, 48, 53. 17. Mikeska, K. R., Bennison, S. J. and Gavrilov, K. L., J. Am. Ceram. Soc., submitted. 18. Chabala, J. M., Gavrilov, K. L., Mikeska, K. R., Bennison, S. J. and Levi-Setti, R., Proceedings of the
GAVRILOV et al.: GRAIN BOUNDARY CHEMISTRY Microscopy and Microanalysis `97 Meeting, Cleveland, OH, 10±14 August. Springer, Berlin, 1997, p. 525. 19. Gavrilov, K. L., Chabala, J. M., Levi-Setti, R., Mikeska, K. R. and Bennison, S. J., Proceedings of the XI International Conference on Secondary Ion Mass Spectrometry (SIMS-XI), Orlando, FL, 1997. 20. Benninghoven, A., Rudenauer, F. G. and Werner, H. W., in Secondary Ion Mass Spectrometry. John Wiley, New York, 1987, p. 277. 21. Brydson, R., Chen, S.-C., Riley, F. L., Milne, S. J.,
22. 23. 24. 25.
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Pan, X. and Ruhle, M., J. Am. Ceram. Soc., 1998, 81(2), 369. Roy, S. K. and Coble, R. L., J. Am. Ceram. Soc., 1968, 5(1), 1. Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1983, 66(5), C90. Bennison, S. J. and Harmer, M. P., J. Am. Ceram. Soc., 1985, 68(1), C22. Gavrilov, K., Bennison, S. J., Mikeska, K. R., and Levi-Setti, R., J. Am. Ceram. Soc., submitted.