Effect of Mechanical Milling on the Sintering Behaviour of Alumina

Transcription

1 J. Aust. Ceram. Soc. 44 [1] (28) Effect of Mechanical Milling on the Sintering Behaviour of Alumina J. S. FORRESTER *1, H. J. GOODSHAW 1, E. H. KISI 1, G. J. SUANING 1 and J. S. ZOBEC 2 1 School of Engineering, University of Newcastle, Callaghan, 238, NSW, Australia 2 Electron Microscope and X-ray Unit, University of Newcastle, Callaghan, 238, NSW, Australia * jenny.forrester@newcastle.edu.au ABSTRACT: Temperatures around 17ºC are normally required to sinter alumina powder into useful ceramic monoliths. Grain size reduction using high energy mechanical milling can be used to promote sintering at lower temperatures. Here, milling was conducted in alumina milling media at a low charge ratio to eliminate contamination which occurs when alumina is milled with milling media such as hardened steel. The milled powders were characterised by their apparent or agglomerate size and the true particle (or crystallite) size. X-ray diffraction peak broadening analysis showed considerable particle size reduction for long milling times. Milled powders were sintered and the progress of sintering monitored using in-situ high temperature XRD and contact dilatometry. Results show that some sintering occurs as low as 4ºC, and substantial sintering occurs at temperatures several hundred degrees below the normal sintering temperature. The low and high temperature regimes are discussed in terms of intra-agglomerate and inter-agglomerate consolidation. KEYWORDS: Alumina, mechanical milling, grain-size reduction, sintering temperature depression. INTRODUCTION Alumina (Al 2 O 3 -corundum structure) is an extremely useful ceramic in many existing and emerging technologies. These include bio-medical implants, the production of which often involves co-firing with metals. High density in alumina components is normally achieved by sintering compacted alumina powder above 17ºC [1,2]. However, many of the metals used in the biomedical industry have melting points close to or below the generally required sintering temperature of alumina (such as Ti (161ºC) and stainless steel ( 14ºC)). One list of biocompatible metals is given in Merrill et al. [3]. In addition, a reduced alumina sintering temperature would have other advantages. These include reduced processing costs and finer sintered grain sizes. Various methods have been employed to combat the generally high temperatures required to sinter alumina powder into useful ceramic monoliths. One such method is the introduction of dopants in the pre-fired alumina. Minor additions of oxides such as MnO 2 and/or TiO 2 are known to allow dense alumina to be synthesised at temperatures as low as 14ºC [1]. A combination of several dopants has seen the greatest decreases in densification temperatures [4,5]. The decrease in densification temperature may be due to the formation of a grain boundary film of liquid eutectic, or defects in the structure which increases bulk diffusion rates [6]. However, the use of dopants can have deleterious effects on the most useful properties of alumina such as high hardness and chemical inertness [6]. The question of biocompatibility of the dopants must also be considered. Although alumina is known to be biocompatible, the introduction of dopants can complicate the use of devices in the heavily regulated bio-medical research field. Lengthy trial periods are normally required prior to acceptance of a new material being classified as biocompatible. Therefore new or broad variations from existing dopant combinations are difficult to justify and implement. Another avenue for decreasing the densification temperature is to begin with a very fine grain size in the green state of the alumina [2]. This has been achieved using some specialised techniques, including sol-gel processing [7,8], the in-flight oxidation of nano-sized alumina using a plasma reactor [9], solution based combustion synthesis [8], and pulsed wire evaporation [1]. These processes, while interesting, may have limited applicability on an industrial scale due to the small quantities of powder produced, the availability and cost of such specialised equipment, and large operating costs. Grain size reduction can also be achieved using mechanical grinding processes. High energy mechanical milling has been used to produce nanometre sized particles in many materials including alumina [11]. Milled powders were sintered and 92% dense compacts with fine grains (.1-.2 μm) formed after sintering between ºC. Problems are often associated with mechanically milled powders, especially with an extremely hard material such as alumina. Firstly, contamination can be extensive from the hard alumina powder abrading the surfaces of the milling balls and vial. Karagedov and Lyakhov [11] found extensive contamination when milling alumina in hardened steel milling media, but they also demonstrated that the powder can be de-contaminated by boiling the milled powder in a 3-5% HCl solution. We have

2 48 J.S. Forrester, H.J. Goodshaw, E.H. Kisi, G.J. Suaning, J.S. Zobec also previously examined the mechanical milling of alumina using hardened steel milling media [12], and found that greater than 1% of the milled powder can be contamination from the milling media. Secondly, we compared the effects of milling alumina using several types of milling media, including hardened steel, tungsten carbide, zirconia and alumina milling media [13], and found that all of the milling media produced some degree of contamination. The lowest contamination was produced logically by the alumina milling media, however at a charge ratio of 1:1, the milling media were degraded rapidly. The second problem, agglomeration, promotes uneven sintering which sometimes results in mechanically weak and porous products. Alumina powder particles in agglomerates are known to increase the volume and size of pores in the green compact [14]. This causes a decrease in the densification rate in the sintered compact. To achieve the highest densities, the agglomeration needs to be controlled. In this work, the milling of alumina was studied to determine the suitability of this method for producing a sufficiently small grain size to promote sintering at lower temperatures. The milling was conducted in alumina milling media at low ball:charge mass ratio in order to eliminate contamination. Milled powders were sintered and the progress of sintering monitored using in-situ high temperature X-ray diffraction and contact dilatometry. EXPERIMENTAL Commercial alumina powder (99.9%, Aldrich, Australia) was milled in an alumina (Al 2 O 3 ) vial using Al 2 O 3 milling balls (for the purpose of differentiating between the powder and the milling media, alumina is used to describe the powder milled, and Al 2 O 3 refers to the milling media. The ball:powder ratio (charge ratio, CR) used was 3:1. This low charge ratio produced milled powders with little contamination from the alumina balls. Milling was carried out in a SPEX8 mixer/mill from 1 to 64 h. Powders were sieved following milling to break up loose agglomerates and to separate the powder from any fragments of the milling media. Following milling, XRD patterns were collected using a Philips 171 powder diffractometer using Cu K α radiation and fitted with a graphite monochromator. Patterns were collected in the range 1-9º 2θ, with a step size of.4º and a step counting time of 2 s. The milled powders were characterised in terms of their apparent or agglomerate size and the true particle (or crystallite) size. Crystallite size was determined by XRD peak broadening analysis. Sintering tests were conducted in-situ in a PANalytical X Pert X-ray diffractometer using Cu K α radiation. An Anton Parr heating attachment (HTK16) capable of temperatures up to 16ºC was used for sintering. Alumina powders were mixed with ethanol and pipetted onto a platinum heating strip used to support samples. A program of heating with intermittent XRD pattern collection was utilised. This program consisted of a heating rate of 5ºC/min to the selected temperature, followed by pattern collections of 3 mins each. For in-situ dilatometry tests, alumina powders were uniaxially pressed at 2 MPa into compacts approximately 6 mm thick. Samples were heated in air at 2.5ºC/min from 3 to 147ºC, soaked for 15 min then cooled at the same rate as heating. Samples of unmilled alumina, alumina milled for 64 h, and a sample milled for 64 h then sintered to 147ºC in the dilatometer were prepared for scanning electron microscopy (SEM) by setting in epoxy resin, followed by polishing to 1 μm finish using diamond paste. Secondary electron images were obtained using a Philips XL3 SEM, operated at 15 kv. As a check for minor levels of contamination, elemental fluorescent X-ray spectra were collected using an Oxford ISIS Si/(Li) energy dispersive spectroscopy (EDS) detector. Rietveld refinements using the XRD patterns were performed using the computer program LHPM [15]. The global parameters refined were the zero point, scale factor and four coefficients of a background polynomial. The following additional structural and instrumental parameters were refined: i. The lattice parameters (a and c). ii. The full width at half maximum of the Lorentzian component of the Voigt peak shape function (k) so that crystallite size could be estimated. iii. The Gaussian half-width parameter, U, as it is often associated with internal strains in a sample. The instrumental variables, i.e. the asymmetry parameter and the peak width parameters V and W, were fixed at.5, -.27 and.36 and respectively, following refinements using data recorded from a NIST alumina standard (#1976). RESULTS AND DISCUSSION Milled Alumina XRD patterns of the alumina powder milled from -64 h are shown in Fig. 1. There is a gradual decrease in the peak heights and a corresponding increase in the width of the peaks, indicating a decrease in crystallite size. This was further

3 J. Aust. Ceram. Soc. 44 [1] (28) hours 32 hours 16 hours 8 hours 4 hours 2 hours 1 hour Unmilled θ (degrees) Fig. 1: X-ray diffraction patterns from alumina milled for -64 h. The peak height decrease and broadening with increased milling time is evident. examined using Rietveld refinement and subsequent interpretation of the peak width parameters using the Scherrer equation, as shown in Fig. 2. broadening [16]. The initial crystallite diameter is ~448 nm, and after 1 h it has reduced to 69 nm. Further milling reduces the size to 29 nm after 64 h. Further milling may decrease the crystallite size slightly more. Other refined parameters for the milling series are shown in Table 1. Table 1: Refined lattice parameters, Gaussian FWHM, estimated crystallite size, and R B factors. Milling time (h) a(å) c(å) Gaussian FWHM (U) Crystallite size (nm) 4.759(2)* (5).32(2) (2) (7).14(3) (2) (9).177(9) (3) (9).14(1) (2) (1).26(1) (3) (1).29(2) (4) (1).35(2) (5) 13.(1).32(3) * The number in parentheses refers to the right-hand digit and represents the standard deviation estimated in the Rietveld analyses. Changes in the lattice parameters a and c are shown in Fig. 3. Both increase with milling time by approximately.1%. Lattice parameter c(a) R B Diameter (nm) Milling time (h) Fig. 2: Crystallite diameter as a function of milling time determined from XRD line broadening. Error bars are plotted at one combined standard deviation of the observed Lorentzian half-width parameter k obs and the instrumental Lorentzian half-width parameter k i. Lattice parameter a(a) Milling time (h) Milling time (h) Fig. 3: Lattice parameters a and c (top) of alumina as a function of milling time. Error bars are plotted at one standard deviation. We have relied here on the secθ dependence of particle size broadening and an assumed Lorentzian peak shape for the crystallite size component of the

4 5 J.S. Forrester, H.J. Goodshaw, E.H. Kisi, G.J. Suaning, J.S. Zobec powder. Fine and dense hard agglomerates (.2-4 μm) composed of many alumina crystallites are loosely packed into large (1-1 μm) soft agglomerates. The final micrograph shows the very dense surface of a sample milled for 64 h then sintered in the dilatometer. This will be discussed in the following section. Sintering All high temperature XRD patterns were collected from alumina milled for 64 h. Some of the results of these tests are shown in Figs. 5, 6 and 7. The first series comprised 1ºC steps from room temperature to 15ºC. A subset of these patterns is shown in Fig. 5. The peaks of the pattern recorded at room temperature show the broadness evident from milling for 64 h (Fig. 1). Some narrowing of the peaks is observed as low as 7ºC, especially at 66 and 68º 2θ (note the overlapping Pt peak from the heating element). Certainly by 11ºC, there is a clearly observable peak narrowing and peak height increase C C 2 7C 1 Room temperature Theta ( ) Fig. 5: High temperature XRD scans of milled alumina powder at increasing temperatures. The peaks at 4, 46 and 67º 2θ are from the Pt strip. Fig. 4: Secondary electron micrographs of (a) unmilled alumina powder (top figure), (b) the structure of the milled powder after milling for 64 h, (c) the surface a sample milled for 64 h, then sintered in the dilatometer sample. SEM micrographs of the microstructure at various stages of milling and sintering are shown in Fig. 4 and highlight differences in the particle size. Firstly, the large tightly packed particles of the unmilled alumina, approximately 5 μm in diameter. The micrograph in the centre of Fig. 4 shows an alumina sample that has been milled for 64 h. Two sizes of agglomerates are present in this The full width at half maximum (FWHM) of the 1, 11 and 2 peaks were calculated from this series of diffraction patterns at all the temperature steps, and these are shown in Fig. 6. The diffraction peaks from samples milled for 64 h are quite broad (FWHM of approximately.3 to.4. At 5ºC a reduction can be seen. This gradient is somewhat constant to around 9ºC. The slope then increases. Above 12ºC the slope again becomes less steep, an indication that the sintering rate decreases. Much of the sintering may already be complete at the lower temperature. By comparison, a fully sintered alumina doped with MnO and TiO 2 at room temperature had a measured FWHM of.15.

5 J. Aust. Ceram. Soc. 44 [1] (28) FWHM Final room temperature 15C 12C after 24 mins 12C after 12 mins C Temperature ( o C) Fig. 6: FWHM of selected diffraction peaks as a function of temperature. To further examine the sintering that occurs at low temperature, we conducted an isothermal series of sintering. A temperature of 11ºC was selected, as it was clear from Fig. 5 that substantial sintering occurs at this temperature. Figure 7 shows the XRD patterns of a sample heated to 11ºC, a pattern collected, soaked at this temperature for 3 min, another pattern collected, held another 3 min, another pattern collected, heated to 15ºC, a pattern collected, then decreased to room temperature where the last pattern was collected. This series showed that there was significant intensity increase and peak narrowing with time at this temperature. The increase to 15ºC produced little further change in the peaks, so it was concluded that most sintering had occurred prior to the increase to 15 ºC Final room temperature 15C 11C 6 mins 11C 3 mins 11C Room temperature Theta ( ) Fig. 7: High temperature XRD patterns of alumina soaked at 11ºC for various times then heated to 15ºC. Finally, a series was conducted under similar conditions to Fig. 7 apart from an increase in time between the recorded patterns and an increase in temperature from 11ºC to 12ºC. This series (Fig. 8) showed that there was no difference between sintering at 12ºC and 15ºC, and this result showed that the small grain size produced from mechanical milling does indeed promote sintering at lower temperatures. Room temperature Theta ( ) Fig. 8: Alumina after soaking at 12ºC for various times. Displacement (mm) Unmilled Milled 1h Milled 2h Milled 4h Milled 8h Milled 16h Milled 32h Milled 64h Temperature (ºC) Fig. 9: Dilatometer traces of unmilled alumina, and alumina milled from 1 h to 64 h during heating from 5ºC to ~147ºC and cooling to 5ºC. Dilatometry results are shown in Fig. 9. The unmilled sample shows little compaction until 11ºC and sintering has only partly occurred at 147ºC. It is evident that sintering is far from complete. Throughout the milling series, a steady increase in the displacement of the sample surface occurs. This is a strong indicator that densification is occurring. The notable points of the dilatometry results are firstly the increase in displacement with increased milling time, and secondly, the change in slope as a function of temperature. As milling time increases. The first part of each scan has a shallow slope, and at an elevated temperature, it increases in slope. It can be seen that with increased milling time, the temperature at which that slope increases reduces. Discussion This work shows that nanocrystalline alumina can be produced by mechanical milling using Al 2 O 3 milling media. In a previous study [13] it was shown that a charge ratio of 1:1 produced nanocrystalline alumina, however at the expense of degradation of the milling media. The loss of weight from the balls and milling vial were so significant (after 4 h of milling, the ball weight had decreased by 27%) that it is unlikely that this could

6 52 J.S. Forrester, H.J. Goodshaw, E.H. Kisi, G.J. Suaning, J.S. Zobec be a viable means to produce nanocrystalline powder. However in this study, the decrease in charge ratio to 3:1 has made a significant difference to the resultant powder. Damage to the milling media was minor, and although the time taken to produce a small crystallite size has increased, we have shown it can be done. XRD peak broadening is governed by the mean size of the diffracting domains and the effect of this is shown in Fig. 1. The crystallite size analysis (Fig. 2, and Table 1) performed during the Rietveld refinements shows clearly how the crystallite size has reduced to approximately 29 nm after 64 h of milling. These results can be compared to our earlier work using a CR of 1:1 [13]. In that work, the crystallite size of powder milled in Al 2 O 3 milled media decreased to around 28 nm. The advantage with the lower CR in this study is the reduction in contamination. During the milling and sintering process the alumina remains in the α-alumina form, but the lattice expands (Fig. 3). This may be due to the introduction of vacancies and other defects. Calculations from the data in Fig. 3 show that approximately.1% increase in the lattice parameters is produced by the milling process after 64 h. The SEM micrographs (Fig. 4) show the effect of milling on the alumina. Milling for 64 h breaks up the alumina particles into loosely packed agglomerates, within which nanocrystalline particles are randomly arranged. The sample sintered in the dilatometer appears dense, with little porosity. The high temperature XRD and dilatometry both show that some degree of sintering does occur at much lower than the 17ºC normally employed for this process. Indeed, indications are that some sintering occurs as low as 4ºC. It is suggested that the two different slopes that are seen in all of the dilatometry curves are the result of different stages in the sintering process. The first part of the curves is likely to be sintering within the agglomerates. The significant increase in slope at 9-12ºC is an indicator that there is significant sintering between agglomerates. 5. Conclusions 1. High energy mechanical milling in Al 2 O 3 media produces nanocrystalline alumina. 2. The alumina remains in the α-alumina form, but with lattice expansion due to vacancies and other structural defects. 3. A significant reduction in sintering temperature to below 12ºC was realised. 4. The powder can be sintered into dense ceramics much below the normally required sintering temperature of 17ºC. We also thank Mr David Phelan for assistance with the electron microscopy, and Mr Peter Garfoot for his assistance with sample preparation. References 1. I.V. Cutler, C. Bradshaw, C.J. Christensen and E.P. Hyatt, J. Amer. Ceram. Soc., 4 (1957) W.C. Johnson and R.L. Coble, J. Amer. Ceram. Soc. 61 (1978) D.R. Merrill, M. Bikson and J.G.R. Jeffreys, J. Neuroscience Methods, 141 (25) H. Erkalfa, Z. Misirli and T. Baykara, Ceramics International 21 (1995) A.R. Boccaccini and C. Kaya, Ceramics International 28 (22) M. Sathiyakumar and F.D. Gnanam, Ceramics International 28 (22) F.-S. Shiau and T.-T. Fang, Mater. Chem. and Phys. 6 (1999) L.C. Pathak, T.B. Singh, S. Das, A.K. Verma and P. Ramachandrarao, Mater. Lett. 57 (22) P.V. Ananthapadmanabhan, T.K. Thiyagarajan, K.P. Sreekumar and N. Venkatramani, Scripta Mater. 5 (24) J.H. Park, M.K. Lee, C.K. Rhee and W.W. Kim, Mater. Sci. and Eng. A, (24) G.R. Karagedov and N.Z. Lyakhov, Nanostructured Materials 11 (1999) H.J. Goodshaw, J.S. Forrester, G.J. Suaning and E.H. Kisi, J. Mater. Sci., 42 (27) C.B. Reid, J.S. Forrester, H.J.Goodshaw, E.H. Kisi and G.J. Suaning, Ceramics International, in press March S. Inada, T. Kimura and T. Yamaguchi, Ceramics International 16 (199) C.J. Howard and B.A. Hunter, A computer program for Rietveld analysis of X-ray and neutron diffraction patterns. (1997) ANSTO Lucas Heights Research Laboratories. 16. H.P. Klug and L.E. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, 2 nd Ed., J. Wiley and Sons, New York, 1974, p Acknowledgements The authors acknowledge a University of Newcastle RGC project grant for financial support.