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Near Net-Shaped Magnesium Aluminate Spinel by the Oxidation of Solid Magnesium-Bearing Precursors

CONTENTS

Pragati Kumar and Kenneth H. Sandhage
Department of Materials Science and Engineering
The Ohio State University, Columbus, OH 43210

ABSTRACT

The feasibility of fabricating near net-shaped spinel, MgAl2O4, by the oxidation of solid Mg- Al2O3-bearing precursors has been demonstrated. Dense disk-shaped and bar-shaped precursors were fabricated by the pressureless infiltration of molten Mg into porous (65-70% dense) Al2O3 preforms at 680-700°C. After solidification, some of the bar-shaped specimens were ground into rods of varied diameter. The shaped, solid Mg-Al2O3-bearing precursors were then oxidized in pure, flowing oxygen at 430°C/40 hours to 700°C/6 hours. The resulting mixtures of MgO and Al2O3 were converted into MgAl2O4 within 15 hours of annealing in oxygen at 1200°C. A subsequent 10 hour sintering treatment at 1700°C in flowing Ar yielded shaped spinel-bearing bodies with densities of 92.5% of theoretical. The dimensions of the final, sintered bodies were within 0.6% of those for the shaped Mg-Al2O3-bearing precursors. The phase content and microstructure at various stages of processing were evaluated with x-ray diffraction (XRD), scanning electron microscopy (SEM), and electron probe microanalysis (EPMA).

INTRODUCTION

Owing to its high melting point (2100°C), high hardness (16.1 GPa), resistance to chemical attack, and chemical compatibility with alumina, zirconia, and mullite, spinel (MgAl2O4) is an attractive matrix material for ceramic-matrix composites.1,2,3 Monolithic, transparent spinel is also used in of optical applications, such as windows for pressure vessels and bulletproof vehicles.

Spinel bodies are usually produced by ceramic powder-based processing. For example, spinel powder can be mixed with malleable organic material and then compacted and formed into a shaped green body (e.g., by injection molding, extrusion, etc.). After removal of the organic material by vaporization or pyrolysis at a modest temperature, the resulting porous spinel body is sintered at > 1600°C to obtain a dense, shaped component.4-6

Significant process optimization can be required during the organic burnout and sintering steps to avoid undesired defects within, and distortion of, the final ceramic body. Non-uniform or incomplete binder burnout can result in cracking or carbon contamination. Further, because the binder often occupies a significant volume fraction (20-40%) of the green body, the sintering shrinkage is often relatively large. As a result, dense ceramic parts produced by this conventional process will not retain the dimensions and, if densification is not uniform, the shape of the starting green body.

Recent work7-9,10,11,12 has demonstrated that near net-shaped ceramic bodies can be fabricated by an exciting new process: the oxidation of solid, alkaline-earth (AE) metal-bearing precursors. AE elements (Mg, Ca, Sr, Ba) are ductile and low melting (650 to 840°C).13

Hence, precursors containing AE metals can be formed into desired shapes by deformation processing or by melt infiltration into a porous ceramic preform. The molar volumes of AE metals tend to be larger than the molar volumes of the corresponding oxides (e.g., Vm[Mg] > Vm[MgO]). Such a volume reduction upon oxidation can be used to counter volume expansions associated with the oxidation of non-AE metals and/or with subsequent oxide-oxide reactions. Thus, by tailoring the precursor phase content, near net-shaped, all-ceramic bodies can be produced from AE-bearing precursors.

Consider the transformation of an equimolar Mg-Al2O3 precursor into MgAl2O4. The conversion of such a precursor can occur by: 1) the oxidation of Mg into MgO, and then 2) annealing of the resulting MgO + Al2O3 mixture to form MgAl2O4. The reduction in solid volume associated with magnesium oxidation can compensate for the increase in solid volume associated with the conversion of magnesia and alumina into spinel. As a result, the net volume change associated with spinel formation from an equimolar mixture of magnesium and alumina (reaction (1) below) is only 0.5%.

(1)

The theoretical dimensional change expected as a result of the transformation of such a Mg-Al2O3 precursor into spinel is only ~0.15% (assuming isotropic expansion, and equal pore fractions within the precursor and product bodies).

Shaped, dense Mg-Al2O3-bearing precursors to MgAl2O4 can be prepared by the infiltration of molten Mg into shaped, porous Al2O3 preforms. Porous oxide preforms of desired shape can be fabricated by extrusion, pressing, slip casting, or other traditional powder processes. After infiltration and solidification of the magnesium, the solid precursor can be machined or ground into a more complicated shape. The ductile magnesium can endow the precursor with the green strength required for such machining; that is, the magnesium can serve the role of a binder. However, unlike fugitive organic binders, magnesium is retained in the precursor (as magnesia) upon oxidation. In other words, the magnesium binder is "burned in" upon oxidation. The amount of porosity in the oxidized precursor should, therefore, be significantly less than the porosity generated by organic binder burnout for a conventionally-processed body (i.e., for equivalent amounts of binder in both cases). Hence, less sintering can be required with the AE-metal-bearing precursor process, so that better control of the shape and size of the fired body can be achieved. A variety of monolithic ceramics and ceramic-bearing composite materials for electronic14,15,16-17, magnetic,18 structural and biomedical19 applications have been synthesized from solid, AE-metal-bearing precursors.

Other oxidation-based techniques developed for the syntheses of ceramics include the self-propagating high temperature synthesis (SHS)20 and reaction-bonded aluminum oxide (RBAO)21-23,24 methods. The SHS process involves the propagation of a reaction-induced thermal wave through a specimen. Although relatively little applied energy is required to fabricate ceramics with the SHS process, components produced by this method tend to be porous and do not tend to retain the dimensions of the precursor. In RBAO processing, an intimate mixture of Al2O3 and Al powder is formed into a shaped, porous green body. The volume expansion associated with the oxidation of aluminum is used to compensate for the initial porosity present in the green compact, so that relatively low net shrinkage is observed after oxidation and sintering. In the present method, however, the as-infiltrated Mg-Al2O3-bearing precursor is dense and more mechanically robust than a porous Al-Al2O3 precursor, so that stresses incurred upon machining/grinding are less likely to lead to macrocracking.

The objective of this study was to demonstrate the feasibility of fabricating dense and near net-shaped spinel bodies by the oxidation and subsequent annealing of melt-infiltration-derived, Mg-Al2O3-bearing precursors. The phases and microstructures produced upon infiltration, oxidation, and post-oxidation annealing were investigated.

EXPERIMENTAL PROCEDURE

A. Precursor Preparation: A molten metal infiltration technique (Figure 1) was used to prepare dense, shaped Mg-Al2O3 precursors. This technique consisted of the following steps:

  1. Al2O3 preform synthesis: Porous Al2O3 preforms were fabricated by uniaxial pressing of Al2O3 powder. The Al2O3 powder, obtained from a commercial vendor (99.98% purity, Johnson Matthey, Inc., Ward Hill, MA), possessed a mean particle size of 7 microns, as determined with a particle size analyzer (Model SA-CP4, Shimadzu Corp., Kyoto, Japan). Uniaxial pressing was conducted in hardened-steel dies of circular (1 cm dia.) or rectangular (6 mm wide X 5 cm long) cross section using applied stresses of 800 or 100 MPa, respectively. The alumina disks and bars were then sintered in air at 1500°C for 5-12 hours so as to produce 65-70% dense preforms.
  2. Mg infiltration into Al2O3 preforms: The porous Al2O3 preforms were infiltrated with molten Mg. The porous preform was placed between Mg disks (15 mm dia. X 3 mm thick, 99.95% purity, Johnson Matthey, Inc.) within a MgO crucible. An excess amount of Mg (molar Mg/Al2O3 ratio = 20-25) was used to allow for complete filling of open pores within the porous alumina preform. The crucible containing the Al2O3 preform and the Mg disks was then placed in a silica tube purged with flowing Ar. The melting and infiltration of Mg into the preform was conducted by thrusting the silica tube into a furnace preheated to 600°C. The furnace temperature was then ramped to 680 or 700°C at a rate of 6°C/min. The quartz tube was held at this temperature for < 30 minutes, after which the power to the furnace was turned off, and the specimen was allowed to cool to room temperature. Most of the excess solid Mg surrounding the infiltrated disk was removed by cutting. The excess Mg close to the infiltrated disk surfaces was stripped from the disk by applying a modest shear stress. Light polishing with 1200 grit SiC-impregnated paper was then used to remove small, localized amounts of excess Mg adhering to the disk surface.
  3. Machining of Mg-Al2O3 bars: Infiltrated precursor bars were machined into more complex shapes (see Figure 9) with the use of a lubricated, cubic BN grinding wheel (18 cm dia.) operating at a rotation rate of 3450 rpm.

B. Heat Treatment: Isothermal oxidation of the Mg-Al2O3 precursor was conducted at 430 and 700°C in flowing oxygen for 40 and 6 hours, respectively. Post-oxidation annealing (for spinel formation) was conducted at 1200°C in flowing oxygen for 10-15 hours. Sintering was then conducted at 1700°C for 10 hours in a flowing Ar atmosphere. Figure 2 shows the time-temperature profiles associated with various heat treatments.

C. Characterization: X-ray diffraction (XRD), scanning electron microscopy (SEM), electron probe microanalysis (EPMA), and optical microscopy were used to characterize the specimens. XRD analyses were conducted with Cu-K radiation at room temperature on powder specimens produced by grinding with a stabilized-zirconia mortar and pestle. The ground powder was mixed with an x-ray transparent vacuum grease (Dow Corning, Inc., Midland, MI) and placed on a single crystal silicon substrate that had been ground so as to expose an irrational plane (i.e., so that silicon diffraction peaks would not be generated with the 2[theta] range examined). Al2O3 diffraction peaks were used as internal XRD standards. An external calibration method25,26 was adopted for quantitative XRD analyses of the amounts of MgO, Al2O3, and Mg2O4 within the samples. Specimens for optical and scanning electron microscopy were prepared by standard ceramographic techniques. The grain structure of sintered samples was revealed by thermal etching of polished surfaces for 4 hours at 1250°C in air. SEM and EDS analyses were conducted with a field emission gun microscope (Model XL-30, Philips Electronics N.V., Eindhoven, The Netherlands). The BSE and SE imaging were conducted on gold-coated specimens. EPMA/WDS analyses (Model SX-50, Cameca Instruments, Inc., Trumbull, CT) were conducted using a beam current of 20 nA and 10 kV with high-purity MgO and Al2O3 as standards.

Sample densities after various processing steps were determined using Archimede's method. The total porosity (open and closed) was evaluated with samples that were hermetically wrapped in a paraffin wax film of known density. Doubly-distilled water was used as the buoyant fluid. Archimede's measurements with and without the paraffin wax coating were used to evaluate the open porosity. Bulk density values were also obtained from dimensional and dry weight measurements. The flexural strengths of infiltrated Mg-Al2O3 precursor bars were evaluated with four-point bend tests conducted at room temperature as per ASTM standard C1161 (configuration A, inner span = 10 mm, outer span = 20 mm, cross-head speed = 0.2 mm/min.).27

RESULTS

The 65-70% dense Al2O3 preforms were infiltrated by molten Mg so as to obtain dense disk- shaped or bar-shaped Mg-Al2O3-bearing precursors. The closed porosity in the alumina preforms was only a few percent (e.g. 2.9% for the 70% dense preform). The molten Mg completely infiltrated the 1.5-1.8 mm thick disks of Al2O3 within 6 minutes at 680°C and the 5 mm thick bars within 30 minutes at 700°C under flowing Ar. No dimensional changes were detected after magnesium infiltration (Table I). The bulk densities of infiltrated precursors were measured to be 3.21 g/cm3 which was equivalent to 95.8% of theoretical density (TD = 3.35 g/cm3). The diffraction pattern obtained from an as-infiltrated precursor exhibited strong peaks for Mg and Al2O3, along with relatively weak peaks for MgO and Mg17Al12 (Figure 3). The largest peak labeled Mg in Figure 3 was shifted towards a higher 2 value than expected for pure Mg. The secondary electron images of polished cross-sections of as-infiltrated samples (Figure 4) revealed interconnected networks of Al2O3 and Mg-rich phases. EPMA conducted on the metallic phase yielded a composition of 8.7 at% Al and 91.3 at% Mg. Four-point bend tests on bar-shaped infiltrated specimens yielded an average (8 samples) flexural strength of 88.2 MPa (72-107 MPa).

Table I. Sample dimensions and cumulative variations (with respect to the porous alumina preform) after various processing steps.
Processing Step Diameter (d),
mm		 Thickness (t),
mm		 d/do
(%)		 t/to
(%)		 V/Vo
(%)
Preform 9.81 1.61 - - -
Infiltrated Preform 9.81 1.61 0 0 0
700°C, 6 h 9.83 1.62 0.20 0.62 1.03
1200°C, 15 h 10.05 1.66 2.45 3.11 8.21
1700°C, 10 h 9.86 1.62 0.51 0.62 1.65

After oxidation for 40 hours at 430°C, the specimen weight increased by 6.9%. Further oxidation at 700°C for 6 hours resulted in an additional weight gain of 1.04% (i.e., a total weight gain of 7.94%). No additional weight change was detected with further heat treatment. A small amount (3.5 wt%) of MgO, in the form of loose powder, was observed near the specimens after the 700°C heat treatment. The microstructure of an as-oxidized sample (after 700°C) is shown in Figure 5. The diffraction pattern obtained from the sample oxidized at 700°C (Figure 6) revealed large diffraction peaks corresponding to MgO and Al2O3, along with smaller peaks for MgAl2O4. Diffraction peaks for Mg-Al solid solutions or Mg-bearing intermetallic compounds (Al3Mg2, Al12Mg1713) were not detected after the 700°C treatment.

The transformation to spinel was completed between 10 and 15 hours of annealing at 1200°C, as revealed by the absence of MgO diffraction peaks in Figure 6. Weak Al2O3 diffraction peaks were also observed in Figure 6, however. The microstructure of a sample annealed for 15 hours at 1200°C (Figure 7) consisted of predominantly spinel with some isolated alumina particles and pores.

The 10 hour sintering treatment at 1700°C yielded samples with bulk densities of 3.33 g/cm3, which corresponded to 92.5% of TD (3.60 g/cm3). The amounts of the alumina and spinel phases in the sintered samples were 6.1 mol% (4.4 wt%) and 93.9 mol% (95.6 wt%), respectively (as determined by quantitative XRD analyses). Figure 8 reveals the microstructure of the sintered spinel after thermal etching. Small, isolated Al2O3 particles were detected within a spinel matrix.

The disks and machined bars retained the shapes of the precursors after complete transformation (Figure 9). Modest increases in dimensions were observed during transformation (Table I), with most of the increase occurring during spinel formation at 1200°C. After complete spinel formation at 1200°C and sintering at 1700°C, the dimensions of the disk and machined bar specimens returned to within 0.6% of the precursor values.

DISCUSSION

Precursor Preparation and Characterization

An equimolar mixture of Al2O3 and Mg corresponds to 64.6 vol% of Al2O3 and 35.4 vol% of Mg. Owing to such a relatively large ceramic content, deformation processing (extrusion, pressing, etc.) of an equimolar Mg-Al2O3 powder mixture was not thought to be a viable means of fabricating a high-density, shaped precursor body. However, because magnesium possesses a modest melting temperature (640°C) and a high affinity for oxygen28 (i.e., an affinity for wetting oxides), a melt-infiltration route (Figure 1) was considered.

In order to produce an equimolar Mg-Al2O3 precursor by the infiltration of liquid Mg into a porous, rigid Al2O3 preform, the open porosity within the preform should be close to 35.4%. Lower or higher values of porosity should yield alumina-rich or magnesia-rich compositions, respectively, upon complete transformation. Since stoichiometric or alumina-rich compositions were preferred over magnesia-rich compositions (owing to the tendency of magnesia to undergo hydration and cracking, and to the thermal expansion mismatch between magnesia and spinel29), the alumina preforms of the present work were sintered to achieve open porosity values of 27-33 vol%. Such sintering also yielded rigid preforms. That is, sufficient necking occurred between Al2O3 particles that the capillary pressure generated upon Mg infiltration did not result in particle rearrangement and contraction (dimensional changes were not detected after infiltration).

The x-ray diffraction pattern obtained from an infiltrated precursor exhibited diffraction peaks corresponding to Mg, MgO, Al2O3, and Mg17Al12 phases (Figure 3). The presence of MgO and Mg17Al12 peaks, and noticeable shifts in the positions of peaks labeled Mg in Figure 3 (relative to peak positions expected for pure Mg), indicated that the following displacement reaction occurred between liquid Mg and solid Al2O3 during infiltration.

(2)

where {} refers to a species in a liquid solution. The Al produced in this displacement reaction went into solution with liquid Mg at the infiltration temperature of 680-700°C. EPMA confirmed the presence of 8.7 at% Al in the metallic phase. As expected from the Mg-Al phase diagram, the Mg-Al liquid solidified into two phases when cooled to room temperature: a Mg-Al solid solution and a small amount of Mg17Al12. The presence of a Mg-Al solid solution within infiltrated specimens was substantiated by the shifts observed for the Mg-like diffraction peaks towards higher 2 values than expected for pure Mg (i.e., the lattice parameters of Mg-Al solid solutions are less than for pure Mg and decrease with increasing Al content30). The amount of Al in the Mg-Al solid solution was determined from the peak shift measurements to be 3.7 at%. Thus, the solidification of the Mg-Al melt can be expressed by the following reaction

(3)

where < > refers to a solid solution.

The ability to obtain complete pressureless infiltration of molten magnesium into the porous alumina preforms may have been enhanced by the observed displacement reaction (2). Indeed, other authors have shown that the pressureless penetration of molten aluminum into silica or silicate preforms was accompanied by a displacement reaction.31,32 Prior published work with the PRIMEX process33,34 has indicated that molten Al-Mg alloys can be pressureless infiltrated into porous Al2O3, if such infiltration is conducted in a N2-bearing atmosphere. Successful infiltration was obtained in the present work with the use of a flowing Ar atmosphere.

Quantitative XRD analyses, coupled with mass and volume balance calculations, indicated that the precursors were comprised of 47.9 mol% Al2O3, 18.5 mol% MgO, 29.2 mol% Mg-Al solid solution and 4.4 mol% Mg17Al12. The theoretical density of this phase mixture is 3.35 g/cm3. Since the measured precursor density was 3.21 g/cm3, the infiltrated precursors possessed a porosity of 4.2%. Much of this porosity could be ascribed to the shrinkage of the molten Mg-Al metal during solidification. The shrinkage associated with the solidification of a Mg-8.7at% Al liquid at 680°C (i.e., reaction (3)) is 9.0%. Since this liquid comprised 27.3 vol% of the precursor, the expected porosity in the precursor due to solidification shrinkage was 2.5%. The remainder of the 4.2% porosity was due to the small amount of closed, non-infiltrated pores present in the Al2O3 preform.

Infiltrated, bar-shaped specimens possessed an average flexural strength of 88.2 MPa (72 to 107 MPa). Such a value of green strength is an order of magnitude higher than is typically achieved with conventionally pressed ceramic green bodies and is also higher than the green strengths reported for porous Al-Al2O3 precursors (20-50 MPa). Such a high green strength value was a consequence of the continuous, percolative nature of the metallic Mg-Al alloy within, and the high density of, the infiltrated precursor (Figure 4).

Precursor Transformation

Owing to the occurrence of displacement reaction (2) during the infiltration of molten Mg into porous Al2O3, a metallic, Mg-rich Mg-Al liquid was produced at 680°C. The lowest solidus temperature for Mg-rich compositions in the Mg-Al system is 437°C.13 Because the initial intent of this work was to oxidize a solid metal-bearing precursor, the initial oxidation heat treatment was conducted at 430°C. However, after 40 hours at this temperature, the weight gain expected for complete oxidation was not achieved. Hence, a subsequent treatment for 6 hours at 700°C was used to complete the oxidation of the magnesium and aluminum within the infiltrated specimens (as confirmed by weight gain measurements and XRD analyses). It is likely that, with further optimization of the oxidation cycle, complete oxidation could be achieved in less time.

A mixture of MgO and Al2O3 produced by the oxidation of a mixture of Mg-Al solid solution and Mg17Al12 should possess 16.2% less volume than the latter metallic mixture. However, very small positive dimensional changes were detected after the 700°C oxidation treatment (Table I). Hence, the oxidized specimens must have contained more porosity than the as-infiltrated precursors. Comparison of Figure 4 and Figure 5 reveals that such an increase in internal porosity did indeed occur after the 700°C anneal. The pores generated upon oxidation were limited to dimensions less than or equal to those of the pores that had been present in the starting alumina precursor, however. Since little sintering of alumina occurred at 700°C, the alumina network in Figure 5 looks similar to that in Figure 4.

A small amount of spinel was detected after the 700°C anneal by XRD analyses. After further annealing for 10 hours at 1200°C, the solid-state reaction between magnesia and alumina yielded spinel as the major phase. After 15 hours at this temperature, spinel formation had been completed, as indicated by the absence of magnesia diffraction peaks (Figure 6). A small amount of residual Al2O3 was still detected, however. Since some loose magnesia powder had separated from the specimens during the oxidation at 430-700°C, the presence of a small amount of residual alumina upon complete spinel formation was not surprising.

After the 15 hour anneal at 1200°C, a volume increase of 7.2% was detected (relative to the specimen volume after the 700°C treatment). This value was not far from the 7.9% increase in volume expected upon the formation of spinel from magnesia and alumina:

(4)

Since a small amount of spinel was detected prior to the 1200°C anneal (Figure 6), the slightly lower measured value of 7.2% was in reasonable agreement with the calculated volume change.

After sintering for 10 hours at 1700°C in flowing Ar, 92.5% dense specimens were produced. This sintering treatment resulted in a volume decrease of only 6.4% (relative to the specimen volume after the 1200°C treatment). A shaped green body comprised of spinel and organic binder would have exhibited a volume decrease of more than 4 times this value, upon sintering to the same density (i.e., for the case where the amount of organic binder equals the amount of Mg-Al alloy in the Mg-Al2O3-bearing precursors of the present work).

After the 1700°C anneal, the dimensions of the sintered disks and bars were within 0.6% of the precursor dimensions. As seen in the macroscopic optical images in Figure 9, disk-shaped and machined bar specimens exhibited excellent shape/edge retention and were free of macrocracks after the 1700°C treatment. The before and after images in Figure 9 confirm the near net-shape capability of the AE-metal-bearing precursor process.

FIGURES
Figure 1: Schematic representation of precursor preparation.

Figure 2: Heat treatment schedule for Mg-Al2-O3-bearing precursor.

Figure 3: X-ray diffraction pattern obtained from an infiltrated precursor.

Figure 4: SE image of a polished cross-section of an as-infiltrated precursor.

Figure 5: SE image after oxidation at 700°C for 6 hours.

Figure 6: X-ray diffraction patterns after various heat treatments.

Figure 7: BSE image of a polished cross-section of a sample annealed at 1200°C for 15 hours.

Figure 8: BSE image of a spinel-righ sample sintered for 10 hours at 1700°C.

Figure 9: Optical micrographs of (a) an as-infiltrated, disk-shaped sample, (b) same as in (a), after sintering, (c) an infiltrated, machined preform and, (d) same as in (c), after sintering.


CONCLUSIONS

(1) Molten magnesium was found to completely infiltrate 1.5-1.8 mm thick, porous alumina preforms within 6 minutes at 680°C (pressureless infiltration). 5 mm thick preforms were completely infiltrated within 30 minutes at 700°C.

(2) Bar-shaped Mg-Al2O3-bearing precursors possessed flexural strengths of 72-107 MPa (average = 88.2 MPa).

(3) Oxidation of the shaped Mg-Al2O3-bearing precursors was completed after heat treatment at 430°C for 40 hours and 700°C for 6 hours in pure, flowing oxygen.

(4) Spinel formation was completed within 15 hours at 1200°C in oxygen.

(5) Sintering in flowing argon for 10 hours at 1700°C yielded 92.5% dense, spinel-bearing bodies.

(6) Near net-shaped disks and machined bars were produced (i.e., the dimensions of the Mg-Al2O3-bearing precursor and final MgAl2O4-bearing bodies agreed to within 0.6%).

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