Control of the γ-alumina to α-alumina phase transformation for an optimized alumina densification

Abstract

In this work, we studied the aptitude to sintering green bodies using γ-Al₂O₃ transition alumina as raw powder. We focused on the influence of the heating rate on densification and microstructural evolution. Phase transformations from transition alumina γ → δ → θ → α-Al₂O₃ were studied by in situ X-rays diffraction from the ambient to 1200 °C. XRD patterns revealed coexistence of various phase transformations during the heating cycle. DTA and dilatometry results showed that low heating rate leads to a significant reduction of the temperature of the α-Al₂O₃ alumina formation. Around 1190, 1217 and 1240 °C were found when using 5, 10 and 20 °C/min of heating rate, respectively. The activation energy for θ-Al₂O₃ → α-Al₂O₃ transformation calculated by Kissinger and JMA equations using dilatometry method were 464.29 and 488.79 kJ/mol, respectively and by DTA method were 450.72 and 475.49 kJ/mol, respectively. In addition, the sintering of the green bodies with low heating rate promotes the rearrangement of the grains during θ-Al₂O₃ → α-Al₂O₃ transformation, enhancing the relative density to 95% and preventing the development of a vermicular structure.

Resumen

En este trabajo, se ha estudiado la capacidad de sinterización de muestras en verde a partir de γ-Al₂O₃ de transición en forma de polvo. El trabajo se ha focalizado en la influencia de la velocidad de calentamiento sobre la densificación y la evolución microestructural. Las transformaciones de fase de alúminas de transición γ→δ→θ→α-Al₂O₃ se han estudiado in situ mediante Difracción de Rayos X (DRX) desde temperatura ambiente hasta 1.200 °C. Los diagramas de XRD han revelado la coexistencia de diversas transformaciones de fase durante el ciclo de calentamiento. Los Análisis Térmicos Diferenciales (ATD) realizados y los datos de dilatometría han mostrado que velocidades de calentamiento bajas conducen a una reducción significativa de la temperatura de formación de α-Al₂O₃. Detectándose alrededor de 1.190, 1.217 y 1.240 °C cuando se utilizan 5, 10 y 20 °C/min como velocidad de calentamiento, respectivamente. La Energía de Activación para la transformación θ-Al₂O₃→α-Al₂O₃ calculada mediante las ecuaciones de Kissinger y JMA usando métodos dilatometricos han sido 464,29 y 488,79 kJ/mol, respectivamente, y mediante ATD 450,72 y 475,49 kJ/mol, respectivamente. Además, la sinterización con velocidad de calentamiento baja promueve la reorganización de los granos durante la transformación θ-Al₂O₃ → α-Al₂O₃, el aumento de la densidad relativa al 95% y la prevención del desarrollo de una estructura vermicular.

Introduction

Due to its interesting physico-chemical and mechanical properties, alumina is one of the most used technical ceramics as structure parts, applied in medical use (orthopaedic implants) and electronics [1–5]. To obtain dense alumina products, a perfect processing control from powder synthesis to sintering [6,7] is required. Gibbsite (Al(OH)₃) and boehmite (AlOOH) are the most used precursors for the preparation of α-Al₂O₃. During the heat treatment, aluminium hydroxides transform into transition alumina in the form of metastable structure before ending as α-Al₂O₃ thermodynamic stable phase (Fig. 1) [8,9]. Boehmite transforms into γ-Al₂O₃ transition alumina under a temperature range of 500–550 °C with a departure of structural water [10]. The transformation of γ-Al₂O₃ monoclinic phase (d = 3.56 g cm⁻³) to α-Al₂O₃ hexagonal phase (d = 3.98 g cm⁻³) is accompanied by a volume reduction of about 10% causing a large density increase [11–14]. The transformation of θ-Al₂O₃ transition alumina to α-Al₂O₃ takes place in the temperature range 1050–1200 °C. This transformation occurs through a nucleation and grains growth mechanism [15–18] and is influenced by several parameters such as grain size [19], chemical composition [20] and heating rate [21]. The θ-Al₂O₃ → α-Al₂O₃ transformation is very significant for the sintering process and consequently for the control of microstructure [22]. According to the sintering conditions, the formation of crystalline α-Al₂O₃ can produce a vermicular microstructure form consisting of a wide pores network, which develops during the α-Al₂O₃ transformation [23,24]. In addition, it is difficult to study with precision the transformation of transition alumina phases due to their instability and high reactivity [21]. Different studies showed that a high energy milling influences the phase transformation from transition alumina to α-Al₂O₃ stable phase [25–29]. Specific diagrams defining the pathway of phase transformations were established according to the heating temperature and particle grain size [10,30]. Some studies were conducted using coarse and fine grained gibbsite in order to examine the transformation paths using in situ X-rays diffraction [19]. Others concern the gibbsite and boehmite to α-Al₂O₃ phase transformation cases. They noticed that small α-Al₂O₃ crystallites of spherical form are obtained by boehmite transformation and larger ones in plate form from gibbsite [31]. The raw powder and the milling mode are therefore fundamentally important for the sintering process and the densification. Many techniques are used to improve the densification process like α-Al₂O₃ seeds [32,33] and the alumina doping agents [34].

Fig. 1.

Transformation sequences of gibbsite (Al(OH)₃) to α-A1₂O₃[10].

In this work we used γ-Al₂O₃ transition alumina as raw powder to study the densification process and the microstructure of sintered samples using various heating rates. The γ-Al₂O₃ to α-Al₂O₃ transformation was characterized in situ by XRD, and the activation energy accompanying the θ-Al₂O₃ to α-Al₂O₃ transformation was evaluated.

Experimental procedures

The used raw powder is γ-Al₂O₃ transition alumina obtained by boehmite calcinations at 600 °C (boehmite being obtained by partial dehydration of gibbsite at 450 °C). The chemical composition expressed in weight % is given as follows: Al₂O₃(88.34%), SiO₂(3.86%), F(0.89%), Na₂O(0.34%), CaO₂(0.18%), Fe₂O₃(0.096%) and structural water(6%). The powder was attrition milled during 3 h using alumina balls of 2 mm diameter. A 0.25% mass ratio of DARVAN C as a dispersant and 1% of duramax as a bonder were added. During the milling process, the pH was adjusted at 10.4 [35]. After milling, the suspensions were completely dried at 110 °C. The powder was sieved to obtain aggregates smaller than 45 μm. In order to highlight the phase transformations, the thermal behaviour of the powder was studied using an analyser SETERAM TGA 92 allowing a differential thermal analysis (DTA). Thermal cycles were recorded from ambient up to 1600 °C with three heating rates (5, 10 and 20 °C/min). A powder mass of nearly 40 mg was used in all cases. The analysis by X-rays diffraction was carried out using Bruker D8 Advance diffractometer working at high temperatures. Green samples thermal shrinking was also studied between room temperature and 1300 °C. Dilatometry tests were carried out using a NETZSCHDIL402C dilatometer. The green compacts were initially shaped by uniaxial pressing then by cold isostatic pressing at 300 MPa. The samples were sintered at 1700 °C for 2 h. The final densities were determined by Archimedes's method using distilled water. The SEM micrographs were obtained using a JOEL 840A scanning electronic microscope. A 20 kV voltage and a working distance between 10 and 15 mm were used. The gold metalized polished surfaces were observed after a thermal treatment at 1600 °C during 1 h.

Results and discussions

Raw powder

The morphology of the used raw boehmite is illustrated in Fig. 2. It consists of agglomerates of around 75 μm. The raw powder was attrition milled reducing its mean diameter (D₅₀) down to 0.29 μm (Fig. 3). The absolute density of the milled powder measured with helium pycnometer was 3.08 g/cm³. To obtain the γ-Al₂O₃ transition alumina phase, the milled boehmite was heated at 600 °C in order to completely eliminate the structural water according to established methods [30].

Fig. 2.

SEM micrographs of boehmite powder before milling.

Fig. 3.

Particle size distribution for boehmite after 3 h attrition milling.

Differential thermal analysis

The thermal behaviour (DTA) of the γ-Al₂O₃ transition alumina powder, for different heating rates, is illustrated in Fig. 4. During the heating cycle between ambient and 1600 °C, the curves revealed that:

• -

  For a 5 °C/min heating rate, the points A, B and C, three exothermic peaks of transition alumina transformation phases are observed. The first peak at around 800 °C is related to the transformation of γ-Al₂O₃ to the δ-Al₂O₃, the second at around 1080 °C is due to the δ-Al₂O₃ → θ-Al₂O₃ transformation and finally the third peak at 1150 °C to the θ-Al₂O₃ → α-Al₂O₃ phase transformation.

• -

  For 10 and 20 °C/min heating rates, only the θ-Al₂O₃ → α-Al₂O₃ phase transformation appeared. They respectively start at 1217 °C (Point E) and 1240 °C (point F). Our results correspond to Weffer and Misra diagram [10]. A shift in the θ-Al₂O₃ → α-Al₂O₃ transformation peak for the three heating rates is also observed. The lowest heating rate reduces the temperature of α-Al₂O₃ formation. The transformation seems more progressive, with a larger and less intense peak than for higher heating rates.

Fig. 4.

DTA curves of the γ-Al₂O₃ raw powder up to 1600 °C for different heating rates.

Dilatometry

The shrinkage curves and their derivatives of the pressed samples are shown on (Fig. 5). They indicate two notable variations along the heating cycle. The samples shrinkage was disrupted by a small expansion indicating the θ-Al₂O₃ → α-Al₂O₃ phase transformation at 1056, 1082 and 1140 °C for the heating rate of 5, 10 and 20 °C/min respectively. The lowest heating rate promotes a maximum transformation rate at 1150 °C. This type of expansion was observed by different authors under the influence of alpha alumina doping seeds, sample compaction pressure rates and doping agents [8,34]. From the crystallographic point of view, the θ-Al₂O₃ → α-Al₂O₃ phase transformation occurs by crystallographic structure change (oxygen anion moving positions) from cubic face centred towards a compact hexagonal structure.

Fig. 5.

Dilatometry curves (a) and their derivatives (b) of the γ-Al₂O₃ transition alumina samples heated up to 1300 °C for different heating rates.

This transformation takes place by nucleation and growth mechanisms. It occurs with high activation energy depending on the heating rate [16–36]. In such transformation, the activation energy is mostly used in the nucleation process. This induces the temperatures shifts at the beginning of the θ-Al₂O₃ → α-Al₂O₃ phase transformation caused by the crystallographic reorganizations mechanisms [37]. In order to complete and evaluate the sintering process, a dilatometric cycle of γ-Al₂O₃ transition alumina of pressed green sample from ambient to 1700 °C with a 5 °C/min heating rate was made (Fig. 6). This test enables us to make a shrinkage process study. No volume variation was observed at the beginning of sintering process. This can be explained by the fact that the transition alumina phase transformations are characterized by structural changes without shrinkage. The θ-Al₂O₃ → α-Al₂O₃ phase transformation initiates at 1050 °C and reaches its maximum at 1192 °C.

Fig. 6.

Dilatometry curve of the γ-Al₂O₃ transition alumina sample heated up to 1700 °C at 5 °C/min heating rate.

DRX in situ

Fig. 7 shows the XRD patterns carried in situ from ambient to 1200 °C on boehmite attrition milled raw powder. The choice of this monohydrate comes from the fact that γ-Al₂O₃ transition alumina is produced from the dehydration of fine boehmite [10]. The XRD data were collected on a Bruker D8 Advance diffractometer with angular sweeping from 2θ = 42° until 2θ = 50° at 0.1°/min. The phases were identified by comparison with the JCPDS tables (γ-Al₂O₃: ICDD 50-741, δ-Al₂O₃: ICDD 16-394, θ-Al₂O₃: ICDD 23-1009, α-Al₂O₃: ICDD 46-1212). From the ambient to 450 °C, the diffractograms shows that the observed peaks correspond to crystalline boehmite. Boehmite is completely transformed to γ-Al₂O₃ in the temperature range 550 to 850 °C. The γ-Al₂O₃ and δ-Al₂O₃ phases coexist up to 950 °C. Beyond this temperature, the θ-Al₂O₃ starts to form progressively. The spectra of those “pseudo-amorphous” δ-Al₂O₃ and θ-Al₂O₃ phases decrease gradually with formation of the α-Al₂O₃ stable phase. Beyond 1200 °C, only the α-Al₂O₃ is observed.

Fig. 7.

XRD in situ patterns of the raw boehmite attrition milled showing transformation from ambient to 1200 °C.

Activation energy

Dilatometry and DTA tests were undertaken to evaluate the activation energy of θ-Al₂O₃ → α-Al₂O₃ transformation. The activation energy (E_a) values were obtained from transformation temperature peak measurements [20,38] using the three heating rates (Fig. 8). The activation energy values for θ-Al₂O₃ → α-Al₂O₃ transformation were estimated by Kissinger and (JMA) methods [39] using the following Eqs. (1) and (2), according to the peak temperature T_p at the different heating rates:

where ‘ϕ’ is the heating rate (K/s), T_p is the temperature corresponding to the maximum of θ → α-Al₂O₃ transformation peak (K), E_a is the activation energy of the θ → α-Al₂O₃ transformation (kJ/mol) and R is the ideal gas constant (8.314 J/mol K).

Fig. 8.

θ to α-Al₂O₃ maximum transformation peak of dilatometry derivative curves (a) and DTA curves (b) at various heating rates.

The relation characterizing the variable x:

where A_T is the area of the exothermic peak in the DTA curve at temperature T and A₀ is the area under the peak.

The plots obtained in Figs. 9 and 10, (ln(ϕ/T_p²) vs 1/T_p) and (ln (ϕ) vs 1/T_p) respectively, are characterized by straight lines with different slopes E_a at different heating rates. The obtained activation energy value by dilatometry was about (464 kJ/mol). The one obtained by DTA is about 450 kJ/mol. These close values remain lower than those reported by Ma. Krell and Wang (517 kJ/mol) [38,39]. According to these results, the γ-Al₂O₃ transition alumina raw powder attrition milled contributes to lowering activation energy for the formation of the stable α alumina phase. On the other hand, the value of the exhibitor Avrami parameter [40,41], which reflects the transformation rate by germination and growth, is defined by (Eq. (3)) from the endothermic shape of crystallization obtained by DTA. The morphology of the crystal growth can be obtained at the transformation ratios (75% and 25%) [42].

where ΔT is the width of the crystallization peak at half maximum value. The values of the Avrami constants are listed in Table 1. They increase from 1.83 to 2.73 with increasing heating rate. The average values of t_0.75/t_0.25 for each heating rate are also listed in Table 1. These values are very close suggesting that the growth morphology is the same at the different heating rates.

Fig. 9.

Plots of ln (Φ/T_p²) versus 1/T_p (a) by dilatometry curves, (b) by DTA curves.

Fig. 10.

Plots of lnΦ versus 1/T_p (a) by dilatometry curves, (b) by DTA curves.

Table 1.

Avrami parameter (n) values at different heating rates.

Heating rate (°C/min)	n	t0.75/t0.25
5	1.83	1.1.
10	2.25	1.08
20	2.73	1.05

Microstructure

The SEM micrographs of green bodies’ sintered samples at 1700 °C are shown in Fig. 10. The γ-Al₂O₃ samples heated with a heating rate 20 °C/min had a porous structure with 85% of relative density and 1.6 μm grain mean size (Fig. 10a). For the sample fired with a 10 °C/min heating rate, a microstructure improvement was noticed. The grain boundaries are quite apparent with important intergranular porosity. The relative density is 90% and the mean grains size is 1.54 μm (Fig. 10b). On the other hand, the samples heated at 5 °C/min had a dense and homogeneous microstructure with low residual porosity and 95% relative density and 3.2 μm mean grain size (Fig. 10c). The observed high density and homogeneous crystal structure with clearly apparent grain boundaries are a caused by a well established particles rearrangement during the θ-Al₂O₃ → α-Al₂O₃ transformation process (Fig. 11). The residual porosity in intra-granular parts was difficult to eliminate [19]. Otherwise, the presence of impurities in the starting alumina powder could be at the origin of liquid silicates formation. Generally, this glassy phase is located in grain boundaries and triple points. The formation of the amorphous phase facilitates alumina densification during sintering (Fig. 12).

Fig. 11.

SEM images of γ-Al₂O₃ green bodies fired at 1700 °C with various heating rates: 20 °C/min (a), 10 °C/min (b), and 5 °C/min (c).

Fig. 12.

SEM images showing grains boundaries of γ-Al₂O₃ samples sintered at 1700 °C with 5 °C/min heating rate.

Conclusion

The influence of the heating rate on the densification and microstructures of green bodies using γ-Al₂O₃ raw powder was studied during sintering. With a low heating rate (5 °C/min), the DTA curves showed a γ-Al₂O₃ → α-Al₂O₃ phase transformations. For higher heating rates (10 and 20 °C/min), only the θ-Al₂O₃ → α-Al₂O₃ phase transformation was detected. The dilatometry curves showed that the temperature at the beginning of the θ to α alumina transformation decreases with decreasing heating rates. The in situ XRD results for the boehmite raw powder used showed coexistence of different transition alumina during the phase transformations. The diffraction patterns in the temperature range (550–850 °C) correspond to the γ-Al₂O₃ phase. At 950 °C, this phase is still present and coexists with the δ-Al₂O₃ phase. Beyond this temperature, we observed the θ-Al₂O₃ → α-Al₂O₃ phase transformation. The activation energy for this transformation was evaluated using Kissinger equation from the results obtained by both DTA and dilatometry techniques. The SEM micrographs showed that samples heated at a 5 °C/min rate had a high relative density (95%) and homogeneous microstructure with low residual porosity.