chapter 2 powder elaboration and...

Chapter II: Powder elaboration and characterization

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Chapter 2 Powder elaboration and characterization

2.1 Composite powder elaboration by surface modification of a commercial alumina powder

Ceramic composites can be prepared by using various processing routes, starting from powder mixture from composite powders obtained by wet chemical methods, as described in the previous chapter. Here, the composite powder, having the final composition 95 vol.% Al2O3-5 vol.% ZrO2, named AZ5, has been obtained by modifying a commercial α-alumina powder. Precisely, the adopted procedure involves the surface modification of the α-alumina particles, which result coated by the zirconia precursor.

Figure 1 shows the typical flow chart of the adopted procedure: first, the alumina powder is dispersed in distilled water, in order to reach a fine and narrow particle size distribution. Then, an aqueous solution of zirconium chloride (0.38 M) is drop-wise added to the alumina slurry, in a proper amount to obtain the final desired composition. After homogenisation by stirring for 1 h, the suspension, containing the zirconia precursor, can be dried, inducing the precipitation of the zirconia precursor onto the alumina particles surface, by liquid evaporation. After drying, the obtained product contains alumina particles coated by the amorphous zirconia precursor. The thermal treatment of the dried powder promotes the zirconia crystallisation, yielding the alumina-zirconia composite powder.

Figure 1 Flow chart of the adopted procedure for the alumina-zirconia powder elaboration


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Some elaboration steps are crucial to control the powder features:

� The alumina powder dispersion: even if a starting sub-micrometer or nano powder is selected, often a certain level of agglomeration can be observed in the as-received powder. Therefore, the dispersion of the starting powder is needed in order to break the particle agglomerates and to obtain a uniform and stable suspension.

� The dilution and the pH of the modified solution: as depicted in the previous chapter, the stability of a ceramic powder slurry depends on several factors, such as the suspension pH, the concentration, the ionic strength 1,2. Since the modified solution contains zirconium chloride dissolved in distilled water and its pH is very acid (less than 1), the effect of its addition on the alumina suspension stability must be controlled, to avoid flocculation and heterogeneity. Moreover, the alumina solubility into water depends on the solution pH, as well as on the temperature and other ionic species. It is well known that the solubility of corundum increases dramatically as pH decrease, while the minimum solubility is at a pH of about 6-7. Finally, such an acid working pH can induce other drawbacks, for instance the corrosion of equipment steel parts. Consequently, for all these reasons, the doping solution pH must be controlled and suitably modified, to achieve the best compromise in terms of operative conditions and final results.

� The drying method: during this stage, the parameters must be set up in order to keep the suspension homogeneity in the dried product. The main feature of the dried powder is the homogeneous distribution of the zirconia precursor onto the alumina particles. Hence, the influence of different drying methods, for instance oven- and spray-drying, on this feature will be discussed.

� The calcination conditions: during thermal treatments, several phenomena can occur, such as the by-products burn out, the zirconia crystallisation, and moreover, depending on the treatment temperature, the formation of necks between powder particles can be promoted. The powder calcination before forming is necessary to avoid a large weight loss during the sintering step, which can affect densification and induce microcracking or undesired deformation. Generally, the powder thermal treatments can be set up on the ground of thermogravimetric analyses, identifying the lowest temperature that provides a complete by-product burn out and negligible or limited particle agglomeration. Moreover, when phase transitions are involved, also the thermal treatment influence on the phase evolution must be taken into account, playing a role on the final phases development, as well as on their distribution and size. In our system, i.e. the surface modified α-alumina powder, the zirconia crystallisation is promoted during thermal treatments, and the influence of the temperature and the time on the final zirconia features will be discussed.

The following sections deal with the set up of several processing parameters.


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2.1.1 Set up of the modification procedure

2.1.1.1 Alumina dispersion

A commercial α-alumina powder (TM-DAR TAIMICRON, supplied by Taimei Chemicals Co., Japan) has been selected to develop the Al2O3-ZrO2 composites. The main features of this powder, as provided by the supplier, are reported in Table I 3. Figure 2 shows a TEM micrograph (a) and the XRD pattern (b) of the as-received powder. The XRD analysis has further demonstrated that only α-Al 2O3 is present in the starting powder (JCPDF n°10-0173).

Table I: Technical data of TM-DAR TAIMICRON alumina powder

Crystalline phase Alpha

B.E.T. Specific surface area (m2g-1) 14.5

Primary particle size * (µm) 0.10

Purity ** 99.99%

Impurities ** (ppm) Si (10) Fe (8) Na (8)

* measured by SEM **measured by ICP-AES

(a) 10 20 30 40 50 60 70

αααα

ααααInte

nsity

(a.

u)

2 Theta (degrees)

αααα

αααα

αααα

αααα

αααα

αααα

αααα

αααα

αααα

αααα

(b)

Figure 2 TEM image (a) and XRD pattern (b) of the as received pure alumina powder (α= alpha-alumina)

The importance of a suitable alumina powder dispersion has been explained in the previous paragraph. For this purpose, a preliminary study on the dispersability of this powder has been carried out. By laser granulometric measurements, the as-received alumina powder was highly agglomerated, the agglomerate size being about 30 µm (distribution by volume).


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In order to improve the powder dispersion, the influence of three pH values on the dispersability and rheology of the alumina suspension was investigated, precisely working at pH 3, 4.5 and 6.5, the latter being the natural pH of the starting powder. The influence of basic pHs was not investigated, in view of the following steps in the procedure. In fact, in order to obtain the Al2O3-ZrO2 composite powder, a solution containing Zr4+ ions, is added to the dispersed alumina slurry. As it will be better explained in the following section, the investigated alumina suspension pHs are limited to acidic or quasi-neutral values, for avoiding uncontrolled zirconium hydroxide precipitation when the zirconium salt solution is added to the alumina slurry.

Thus, three suspensions were prepared containing 50 wt% alumina in distilled water, and the pHs were modified by HCl addition to reach the desired values. The alumina suspensions were ball milled by using α-Al 2O3 spheres (diameter = 2.0-2.5 mm), the powder/spheres weight ratio was 1/5.

The particle size distributions reported in Figure 3 were reached by ball milling the alumina suspensions for 3 h. The results of laser granulometric analyses performed on the three alumina suspensions show that similar particle size distributions were reached under the investigated pH conditions: d50 was equal to 0.16 µm, 0.17 µm and 0.18 µm, respectively for pH 6.5, 4.5 and 3. Being these pHs far from the isoelectric point of alumina, that is about 8.5 4, such results were reasonably expected.

For all the pH conditions, a good deflocculation was achieved after ball milling, the mean particle size being comparable to that reported by the powder supplier.

Since the alumina powder exhibits a good dispersion in a wide pH range, up to pH 6.5, the alumina suspension pH was set to 4.5, for the following processing. Moreover, as it will be discussed in the next paragraph, the doping solution pH will be modified, up to 4, so that the suspension pH is kept almost constant during the procedure.

0.1 10

5

10

15

20

25

30

35

40

pH 3pH 4.5pH 6.5

Vol

ume

perc

enta

ge %

Particle size [µm]

Figure 3 Particle size distributions, by volume percentage, for the alumina suspensions at different pHs


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2.1.1.2 Zirconium salt solution

The solution of zirconia precursor contains the inorganic zirconium salt, precisely ZrCl4, dissolved in distilled water. Zr4+ ions show a strong hydrolysis, so that the solution pH is less than 1, and the precipitation of Zr(OH)4 in absence of a complexing agent begins at a considerably low pH of about 3 5. However, the Zr(OH)4 precipitation can be retarded adding a complexing agent: for example Choy et al.5 reported that when citric acid is used as chelating agent, the hydroxide precipitation occurs above pH 7, being the [Zr4+] concentration equal to 0.1 M. Therefore, the solution pH was modified to a value of about 4, by adding a proper amount of tribasic ammonium citrate (molar ratio ammonium citrate:ZrCl4 equal to 2:1). The tribasic ammonium citrate acts as chelating agent, so that the solution pH can be increased up to 4, without inducing Zr(OH)4 precipitation. The pH increase also allows to avoid the metal parts corrosion: for example, as it will be descried later, the modified suspension is fast dried in a spray dryer, and when highly acid slips were processed, initial corrosion was observed on some steel components in the drying apparatus, such as the screws and the metallic flanges.

Thus, during the mixing procedure, the solution containing zirconium tetrachloride and tribasic ammonium citrate, having a pH of 4-4.5, is drop-wise added to the alumina slip, in a proper amount to obtain the desired final composition. This modified suspension is kept under magnetic stirring for 1 h before drying, to assure homogeneity.

2.1.1.3 Drying step

As before stated, the zirconia precursor precipitation occurs during drying, induced by the solvent evaporation. The aim is to obtain a zirconia precursor layer homogeneously distributed onto the alumina particle surface. Some preliminary drying tests were carried out in oven, at 105°C. In this case, in the slowly evaporated solution and precursor segregation was observed. Moreover, the samples obtained by using the oven-dried powder showed an uneven second phase distribution, just due to the uncontrolled zirconia precursor precipitation. An example of microstructural dishomogeneity observed in the sintered body obtained by pressing the oven-dried and calcined composite powder is reported in Figure 4.

Figure 4 SEM micrograph of AZ5 sample, showing an uneven zirconia distribution (white particles) into the alumina matrix (dark grains)


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In order to improve the homogeneity of the zirconia precursor distribution, a fast drying was preferred, in order to avoid the salt segregation. Therefore, in the set up procedure, the suspension was spray dried. For this purpose, the suspension must be diluted down to 4 wt% of solid content and sprayed into the drying chamber, being the inlet temperature 140°C. The spray drying equipment is described in detail in the Appendix devoted to the experimental techniques.

Drying of the atomized suspension occurs in a very short time, thanks to the small volume of the droplets and the high surface area when evaporation takes place. In this condition, the suspension is brought to supersaturation very quickly, and finally a more homogeneous distribution of zirconia precursor onto the alumina particle surface is achieved.

In Figure 5 a, a SEM image of the granules obtained by spray drying the suspension is reported. The granules are spherical and the mean diameter is about 3.3 µm. The size distribution obtained by Image Analysis Technique is reported in Figure 5 b.

(a)0.5 0.8 1.5 2.6 4.5 7.8 13.5 23.6

0

5

10

15

Per

cent

age

(%)

Granule size (µm) (b)

Figure 5 SEM image of the spray dried granules (a), and related size distribution, by number (b)

Figure 6 TEM image of spray-dried AZ5 granules


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In Figure 6, TEM images of the as-dried AZ5 powder are reported. The spray dried granules are aggregates of the primary alumina particles, which are coated by the zirconia precursor.

The as-dried powder was submitted to XRD analyses: the detected phases were α-alumina and ammonium chloride (JCPDS n° 07-0007), as shown in Figure 7. The maximum drying temperature, i.e. 140°C, is not sufficient to promote zirconia crystallization; consequently, as expected, a broad peak at low angles, characteristic of the presence of an amorphous phase, is visible, which does not appear in the XRD pattern of pure alumina (see Figure 2 b). These results are in agreement with the data reported by several Authors 6, 7, 8, 9 on the preparation of oxide powder by using citric acid as chelating agent, which refer the obtainment of an amorphous precursor, which crystallize after calcination.

10 20 30 40 50 60 70

***

*

αααα

ααααInte

nsity

(a.

u)

2 Theta (degrees)

αααα

αααα

αααααααα

αααα

αααα

αααα

αααα

αααα

αααα

*

Figure 7 XRD pattern of the as-dried AZ5 powder (α=alpha alumina,*=NH4Cl)

2.2 Thermal behaviour

From the previous results, a thermal treatment to promote by-products burn-out and zirconia crystallization was investigated.

The choice of the treatment temperature and time is of primary importance, since it influences the composite powder features, for instance the degree of particle agglomeration as well as the second phase size and distribution.

First, the thermal behaviour of the powder has been followed by TG-DTA analyses, carried out on 115 mg of powder, heated in 20% N2-80%O2 atmosphere at 5°C min-1 up to 1000°C. The TG-DTA curves are reported in Figure 8. The total mass loss of about 24% was completely recorded in the 100-600°C temperature range, imputable to by-products (such as ammonium citrate and chloride) thermal decomposition. The exothermic peaks, between 300°C and 600°C, are given by several overlapped phenomena, such as decarboxilation, organic skeleton destruction, carbon burning, chloride release10, and also zirconia


56

crystallisation into tetragonal phase, as supported by XRD analyses that will be extensively presented in the next sections.

The studies performed by Petrova et al. 10 about the thermal decomposition of zirconium and yttrium citric complex in both ethylene glycol and water media can help to understand the involved phenomena in the thermal treatment of the dried alumina-zirconia powder. To synthesize Zr-complexes, they prepared an aqueous solution containing the metal salt (ZrCl4) and citric acid, and then acetone was added to isolate the complexes. The isolated precipitates were dried, and the thermochemical behaviour was studied by TG-DTA analyses, reported in Figure 8, and IR spectroscopy. They proposed a possible schema of the process taking place during the thermal decomposition of the Zr-complexes, as follows:

Stage I Dehydration: The dehydration proceeds in two steps, namely the first one up to 120°C and the second one up to 145°C, suggesting that two type of water molecules are bonded to the complexes, by adsorption or coordination.

Stage II Intramolecular dehydration, chlorine partial release: the dehydration process continues up to 270°C and a loss of intramolecular water is supposed, with the formation of C=C bond, i.e. with transformation of the citrate into aconitate.

Stage III-VII Decarboxilation, organic skeleton destruction, carbon burning, chlorine complete release: the decarboxilation takes place transforming the aconitates into itaconates (stage III) and continues up to 430°C (stage IV), followed by the complete destruction of the organic skeleton up to 500°C (stage V) and burning of the residual carbon (stage VI) and final chlorine release (stage VII).

TG

DTG

TGTG

DTG

(a)

DTADTA

(b)

Figure 8 TG-DTG (a) and DTA (b) curves of zirconium citric complexes, as reported by Petrova et al. 10


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With respect to the results obtained by Petrova et al., the TG-DTA curves of this study (Figure 9) show some discrepancies, even if differences are expected being different the composition of the analyzed systems. In fact, Petrova et al. 10 studied a system containing isolated Zr-complexes, obtained by using citric acid as chelating agent, whereas in this work, the thermal analyses were performed on alumina particles coated with a zirconia precursor, obtained by following the already described procedure, and using tribasic ammonium citrate, as complexing agent. Consequently, the effective composition of the complexes could likely vary, leading to some differences in the thermal behaviour. In Table II, the peak temperatures of DTG and DTA curves are compared to the data obtained by Petrova et al.10 and related to the decomposition stages. It is worth noting that in the case of alumina-zirconia powder, a marked endothermic peak appears at 286°C, corresponding to a peak in the DTG at 290°C, that could be related to the decarboxilation phenomenon. Moreover, an exothermic peak in DTA curve is observed at 380°C, which was not reported by Petrova et al.

100 200 300 400 500 600 700 800 900 1000-25

-20

-15

-10

-5

0

TG

(%

)

Temperature (°C)

490°C

360°C

290°C

DT

G

115°C

(a)

0 200 400 600 800 1000-5

0

5

10

15

20

220°C115°C

380°C 483°C

Hea

t flo

w (

mV

)

Temperature (°C)

286°C

exo

(b)

Figure 9 Figure 9 TG (solid line) and DTG (dashed line) curves (a), and DTA curve of the as AZ5 powder (b)


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Table II Peak temperatures of DTG and DTA analysed carried out on the AZ5 powder compared to the data in Reference [10]

AZ5 powder Zr-Citrate complex [10]

DTG peak (°C) DTA peak (°C) DTG peak (°C) DTA peak (°C)

Stage I 115 115 (-) 120 140 (-)

Stage II - 220 (-) 205 220 (-)

Stage III 290 286 (-) - -

Stage IV 360 380 (+) 380 430 (+)

Stage V 490 483 (+) 480 -

Stage VI - - - 550 (+), 720 (+)

Stage VII - - - 910 (+)

(+) and (-) indicate the exothermic and endothermic signal, respectively

2.3 Investigation of the crystallization mechanism

In this section, the study of the zirconia amorphous-to-crystalline phase transformation for AZ5 powder is described. The kinetics and the mechanism of the crystallisation were investigated to provide a deeper knowledge of the intermediate steps between the starting amorphous phase on the alumina surface and the final developed microstructure, whereas a strict control of the nucleation and growth of zirconia crystals is needed to achieve well-tailored microstructures.

This aim was pursued by a precise follow up of the zirconia crystallisation at the surface of α-alumina grains during thermal treatments via XRD analyses and TEM observations.

Three different aspects about the zirconia crystallisation were analysed:

� The zirconia phase evolution, investigated by conventional XRD analyses and Raman Spectroscopy.

� The zirconia crystallisation kinetics, investigated by both conventional XRD and in-situ High Temperature X-Ray Diffraction (HTXRD), and concerning the zirconia crystalline fraction and the crystal size.

� The zirconia crystallisation mechanism, studied by coupling the TEM and HRTEM observations and the results concerning the crystallisation kinetics.

These analyses were performed on the thermally treated AZ5 powder under different conditions; for a sake of simplicity, three samples series can be distinguished:


59

- Set A: The AZ5 powder samples were treated in the furnace, up to 300°C, 400°C, 500°C, 600°C, 800°C and 1000°C, respectively, at a heating rate of 10°C min-1, for soaking times of 1 h and 10 h at the maximum temperature. The following characterization analyses were carried out on selected samples:

● Conventional XRD analyses (2θ range: 5°-70°) allowed to investigate the phase evolution as a function of treatment temperature.

● TEM observations of the powders allowed studying the nucleation and growth of zirconia crystallites on alumina grains surface.

● HRTEM was also performed on selected thermally treated powders to observe in detail the zirconia crystallite features.

- Set B: Small samples of AZ5 powder were also submitted to isothermal treatments, by quickly plunging them into a furnace already at a steady temperature (i.e., 500°C, 600°C, 800°C, 1000°C), and then holding them in temperature for 1, 2, 5, 15, 30 min, 1, 2, 3, 6, 9, 12, 24 h, respectively. Only in the case of the isotherm at 500°C, a longer duration of 65 h was exploited. These samples were submitted to conventional XRD analyses (2θ range: 24°-33°).

- Set C: The AZ5 powder was in-situ treated, by using a High Temperature X-Ray diffractometer equipped with a furnace allowing the samples to be heated up to 1200°C. Two kinds of experiments were carried out:

● (Set C-I) The AZ5 samples were heated from room temperature up to 1000°C (temperature ramps of 1.5°C min-1, 4.2°C min-1, 5.6°C min-1, 30°C min-1 respectively). The XRD analyses were in-situ collected, during heating (2θ range: 24.2°-38.3°).

● (Set C-II) The AZ5 samples were heated up to 500°C, 600°C, 800°C, 1000°C or 1200°C, respectively, followed by an isotherm soaking at the maximum temperature for 12 h (temperature ramp of 30°C/min). The XRD analyses were in-situ collected during the isotherm treatment (2θ range: 24.2°-38.3°). Moreover this sample set was submitted to the Raman Spectroscopy analyses to study the phase evolution as a function of treatment temperature.

From the collected XRD patterns, the main information about zirconia polymorphism, crystallized fraction and crystallite size were obtained. Particularly, the area of the main zirconia peak (101) was measured and used to calculate a crystallinity degree index. Therefore, the crystalline fraction f of each sample was calculated as the ratio between the area of its main zirconia peak and that presented by a sample in which zirconia was fully crystallized.

The average crystallite size was calculated by the line-broadening method, using the Scherrer’s equation:

θβλ

cos⋅⋅= k

D (1)

where D is the crystalline size, λ is the wavelength of the CuKα line (1.54060 Å), k is the Scherrer constant equal to 0.9, and β is the full width at half maximum of the main zirconia peak (2θ≅30.25°), assuming a Gaussian profile.


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2.3.1 Phase evolution

A first investigation of the phase evolution has been performed by treating the as-dried AZ5 powder in a furnace at different temperature (heating and cooling rate = 10°C min-1) for a soaking time of 1 h (samples belonging to the Set A). Precisely, XRD analyses were carried out on AZ5 powder calcined at 300°C, 400°C, 500°C, 600°C, 800°C and 1000°C.

From XRD results, tetragonal zirconia starts to crystallize at 400°C (see Figure 10), and this phase remains the only one detected, near α- alumina, up to 1500°C, although no phase stabilizer was added (JCPDF n°80-0965). As often mentioned in literature, the first phase to crystallize from an amorphous zirconium oxide precursor11,12,13,14,15 is usually tetragonal zirconia. Here, the tetragonal zirconia is obtained after all the thermal treatments, even if monoclinic phase is the thermodynamically stable one at room temperature. Bhagwat et al.7

synthesized nanocrystalline zirconia powder by the amorphous citrate route, and performed in-situ HTXRD, revealing that the starting amorphous zirconia mainly crystallized (92 vol.%) into tetragonal symmetry while the remaining 8 vol.% was monoclinic zirconia; this tetragonal-monoclinic phase ratio was stable up to 750°C, thanks to the small crystal size. Garvie 13 suggested that, for very small crystals, the tetragonal phase could be stable thanks to the smaller surface energy, and he found that when zirconium oxide samples, obtained by precipitation, were heated, the tetragonal crystallites grew to a maximum of 30 nm, before a complete transformation to the monoclinic phase. As a consequence, a critical size for transformation, Dc, can be calculated, on the basis of the surface free energy for the two phases. The surface free energy of tetragonal nanoparticles depends on several factors, for example on the state of aggregation. Razaei et al. 14 calculated a Dc of about 9 nm, for isolated spherical zirconia particles, and of 33 nm, when zirconia is within aggregates. In the AZ5 system, zirconia crystals remained tetragonal in all the calcined samples even at high temperature (1500°C for 1 hrs), and they reached a maximum size of about 52 nm. It is therefore probable that the interface energy between α-alumina and tetragonal zirconia contributes to lower the total free energy of tetragonal phase respect to the monoclinic one, hindering the related transformation.

20 30 40 50 60 70

1500°C

800°C

1000°C

600°C

500°C

400°C

Inte

nsity

(a.

u)

2 Theta (degrees)

300°C

ααααT

αααα

αααα

αααα

Tαααα

αααα

ααααααααTαααα αααα

Figure 10 XRD patterns of AZ5 powder calcined at different temperatures for 1h (α=alpha alumina, T= tetragonal zirconia)


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The phase evolution was also followed by Raman Spectroscopy since this technique is preferable to investigate the primary crystallization and possible phase transformation, being more sensitive to intermediate-range order, while the phase detectable by XRD must have a periodicity over a length of 5 nm 16.The analyses were performed on both pure alumina samples and AZ5 powders after treatments at 500°C, 600°C, 800°C, 1000°C and 1200°C for 12 h (samples belonging to the Set C II). The obtained spectra are reported in Figure 11. The Raman spectrum of the alumina powder presents several peaks at 378, 416, 429, 451, 576, 644 and 750 cm-1, in good agreement with the literature 17. In the spectra of the composite powder others peaks at 144, 268, 314, 458 and 599 cm-1 appear, which are imputable to zirconia tetragonal structure 18. No characteristic peaks of the monoclinic 19 or the cubic structure 20 are observed even after treatment at 1200°C.

In both XRD patterns and Raman spectra of the composite samples, also calcined at high temperature, only tetragonal zirconia is detectable, confirming that, thanks to its very small crystallite size, tetragonal structure is preserved even after prolonged high-temperature treatments.

100 200 300 400 500 600 700 800

TTT

T

α

α

ααα

α

1200°C

1000°C

800°C600°C500°C

Inte

nsity

(a.

u)

Wave number (cm-1)

Pure alumina

αT

Figure 11 Raman spectra of pure alumina and AZ5 powders (Set C II) after treatments at 500°C, 600°C, 800°C, 1000°C and 1200°C (α=alpha alumina, T= tetragonal zirconia)

2.3.2 Crystallization kinetics

As explained in the Section 2.3, the zirconia crystallization kinetics was studied on Set B and C samples by evaluating zirconia crystalline fraction and zirconia crystal size, as a function of the main thermal treatment parameters, i.e. the treatment temperature, the duration or the heating rate.


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The evaluation of zirconia crystalline fraction and crystal size in the differently treated samples allowed to obtain a large experimental data amount, which was analysed on the ground of Johnson-Mehl-Avrami (JMA) theory, which is briefly presented in the following section.

2.3.2.1 Theoretical background: the JMA theory

The zirconia crystallization from amorphous precursor belongs to the nucleation and growth transformations. The rate of transformation for nucleation and growth process depends both on the rate at which stable nuclei form and on their growth rate, the slowest mechanism limiting the overall transformation rate. Two important characteristics of the nucleation and growth transformation are reported, as stated by Christian 21 in his general introduction to the transformation of metals and alloy:

� Dependence on Time. At any temperature, the amount of transformation increases with time, until a state of minimum energy is reached. In practice, Christian 21

specified that the transformation at some temperature may be so slow that it cannot be detected in any observable period of time.

� Dependence on Temperature. In principle, the transformation amount does not depend on temperature, except in the sense that the equilibrium state is itself a function of temperature. At any temperature, if a sufficient time is given, the transformation would continue until equilibrium is reached. The transformation rate varies with temperature, and for any transformation a temperature can be found below which the change proceeds at a negligible rate.

In the generic nucleation and growth transformation from an initial phase i to a final phase e, the number of new e regions nucleated in the time interval between τ and τ+dτ is equal to

τdVI ie ⋅⋅ , being Ie the nucleation rate per unit volume, iV the untransformed volume at the time t=τ. During the initial transformation stages, the untransformed volume is comparable to the total volume, VV i ≅ . The growth rate in any direction is Y, assumed isotropic, so that the transformed region are spherical. The volume of an e region originated at the time t=τ is:

( )33

3

4 τπ −

= tYve (t >τ) (2)

and the total transformed volume eV at time t is:

( )∫=

−⋅

=t

t

ee dtYIVVτ

ττπ 33

3

4 (3)

Assuming Ie constant, the equation (3) cab be integrated, obtaining the transformed fraction:

43

3tYI

V

Vf e

e

⋅⋅⋅

== π (4)

this implies that in the early stage of transformation, the total transformed volume is proportional to the fourth power of time.


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Avrami 22,23,24 extended this model taking into account the mutual interference of regions growing from separate nuclei. In fact in the previous equations, the possible impingement of two region has not be considered, while this event becomes more probable as the transformation goes on. Avrami assumed that when one grain impinges upon another one, the growth ceases, and defined an “extended” volume of the transformed phase e, e

exV , that is the

volume obtained without impingement phenomena. As in the case of the actual transformed volume expressed in the equation (2) and (3), the “extended” volume is:

( )∫ −⋅

=t

eeex dtYIVV

0

33

3

4 ττπ (5)

The equation (5) differs from the equation (3) because the eexV includes also the region

nucleated in the already transformed material, at t<τ. The relation between the “extended” and the real volume is:

eex

ee dV

V

VdV

−= 1 (6)

or:

−−=

V

VVV

ee

ex 1ln (7)

and replacing the equation (7) into the equation (5):

( ) ( )∫ −⋅⋅

=−t

e dtYIf0

33

3

41ln ττπ (8)

This equation may be integrated if Ie and Y are assumed constant with time, obtaining:

⋅⋅⋅−−= 43

3exp1 tIYf eπ

(9)

In general, Ie is not constant with time. For example, Avrami referred that the nucleation occurs only at certain preferred sites, the “germ nuclei”, which already exist in the starting phase and may consist in “tiny blocks” or “crystal molecules” of the new phase. Consequently the nucleation rate depends on the number of these “germ nuclei”, which are gradually exhausted because some of them become active growth nuclei, or because they are swallowed by growing grains of new phase. Their number for unit volume decreases from the initial

quantity N to ( )tNN = . Thus, the nucleation rate is not constant with time:

dtdNtII ee =≡ )( , and, in the general case, the equation (8) must be used.

Avrami demonstrated that the following general relation can be obtained by integrating the equation (8):

( )( )nKtf −−= exp1 (10)


64

where K is a rate constant, and n is the Avrami exponent. The Avrami exponent can be determined from the logarithmic form of the previous equation:

( ) ( )tnKnf

lnln1

1lnln ⋅+⋅=

− (11)

The rate constant K can be expressed in the form of an Arrhenius equation:

−⋅=RT

QKK exp0 (12)

Thus the relation between lnK and 1/T should be linear and the slope of the straight line is the apparent Activation Energy, Q.

The Avrami exponent can be calculated also by non-isothermal methods, when the heating rate is kept constant, just modifying the equation (11) as follows 25:

( ) ( )vnKnf

lnln1

1lnln ⋅−⋅=

− (13)

where v is the heating rate, and f is the crystallized fraction as determined by the area of the main zirconia XRD peak.

From the Avrami treatment, the exponent n is related to the crystallization mechanism, particularly to the assumptions regarding the growth, and it results:

3≤ n ≤ 4 for a three-dimensional nucleation and growth process;

2≤ n ≤ 3 for a two-dimensional growth (plate-like growth);

1≤ n ≤ 2 for a one-dimensional growth (linear).

However, this Avrami exponent interpretation is based on the previously described hypotheses about Ie and Y . In the next sections, it will be shown that the here analysed transformation is not exhaustively explained only by applying Avrami’s law, because of the unusual low exponent values characterizing the zirconia crystallization on the alumina surface. Gremillard 26 demonstrated that the Avrami exponent was not only dependent on the physical mechanism governing the nucleation and growth, but also on the kinetic characteristics, particularly on the ratio between nucleation and growth rates. These considerations could help to understand the results reported in the following sections.

2.3.2.2 Crystalline fraction evolution during isothermal treatments

Figure 12 reports a detail of the XRD patterns collected on the Set B samples calcined at 500°C for increasing times, showing the evolution of the main peak of tetragonal zirconia, which was already detectable after 5 min soaking at this temperature. The peak area, proportional to the crystallized zirconia amount, increased as a function of time, reflecting the


65

evolution of the second phase during the isothermal treatments, as detailed below. Similar results have been obtained for the other Set B samples, treated at higher temperatures.

2728

2930

3132

3310

100

1000

Inte

nsity

(a.

u)

Time (m

in)2 Theta (degrees)

Figure 12 Evolution of the main diffraction peak of tetragonal zirconia in AZ5 powders calcined at 500°C for increasing times (Set B)

Figure 13 shows the crystalline fraction vs. time, measured on Set B samples, isothermally treated in the temperature range 500-1000°C. The crystallisation rate increases with the treatment temperature increase. The crystallized fraction after 1 min is about zero for the powder treated at 500°C, and about 26 %, 51 % and 80 % for the samples treated at 600°C, 800°C and 1000°C, respectively. It was not possible to evaluate an induction time for crystallisation, except for the isothermal treatment at 500°C, since in this case crystalline zirconia is not detectable after 1 min, but a crystallinity degree of about 20 % is reached just after 2 min.

The maximum zirconia crystalline fraction reaches the value of 1 only at 1000°C, whereas maximum value of 0.72, 0.73 and 0.90 vol.% are obtained during the treatment at 500, 600 and 800°C respectively, and a further increase of f was not appreciated after the maximum treatment time exploited in this study. We cannot exclude that a longer treatment time leads to complete zirconia crystallization, even at temperature lower then 1000°C.


66

1 10 100 10000,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1000°C 800°C 600°C 500°C

Cry

stal

line

frac

tion

Time (min)

Figure 13 Zirconia crystalline fraction vs. time, in AZ5 powder, submitted to isothermal treatments in furnace at different temperatures (Set B)

The crystallization kinetics were analysed using the Avrami equation (9), for isothermal heating. The Avrami exponent n and the rate constant K were determined from equation (11), and the apparent Activation Energy, Q was calculated by the slope of the Arrhenius plot, obtained applying equation (12).

The Avrami plots are reported in Figure 14.

0 2 4 6 8

-1,5

-1,0

-0,5

0,0

0,5

1,0

1,5

2,0

2,5

1000°C 800°C 600°C 500°C

ln(ln

(1/(

1-f)

))

ln(t)

Figure 14 Avrami plots for isothermally treated AZ5 powder (Set B)


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The calculated Avrami exponents for isothermal treatments are in the range of 0.12-0.26, and they decrease as the treatment temperature increases, while the apparent Activation Energy Q is 132 KJ⋅mol-1 (see Figure 15).

0,8 1,0 1,2 1,4

-5

0

5

ln(K

)

1/T x 103 (K -1)

Figure 15 Arrhenius plot for isothermally treated AZ5 powder (Set B)

2.3.2.3 Crystalline fraction evolution during isochronal treatments

In view of a more precise follow-up of zirconia crystallization, based on a large data amount collected by carrying out faster experiments, in-situ HT-XRD analyses were performed, obtaining the Set C I samples. Since a real isothermal treatment cannot be achieved in a diffractometer, zirconia crystallisation was monitored under a constant heating rate.

In these experiments the phase evolution was followed on samples treated at several heating rates, kept constant during each experiment, as explained in Section 2.3. Figure 16 plots the trend of the crystalline fraction vs. temperature, for the different heating rates. The crystallization rate increases as the heating rate decreases, and the calculated Avrami exponents are in the range 0.09-0.49, while the apparent Activation Energy Q is 111 KJ mol-1 (see Figure 17 and 18).


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400 600 800 1000 12000,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

30°C min -1

1.5°C min -1

4.2°C min -1

5.6°C min -1

Cry

stal

line

frac

tion

Temperature (°C)

Figure 16 Zirconia crystalline fraction vs. temperature, in in-situ treated AZ5 powder at different heating rates (Set C I)

0 2 4-4

-2

0

ln(ln

(1-/

(1-f

)))

ln(ν)

450°C 950°C 750°C 550°C 1150°C

Figure 17 Avrami plots for isothermally treated AZ5 powder (Set C I)


69

0,8 1,0 1,2 1,4-8

-6

-4

-2

0

2

4

1/T x 103(K -1)

ln(K

)

Figure18 Arrhenius plot for isochronal treated AZ5 powder (Set C I)

The Avrami exponents, calculated by exploiting both the isothermal and isochronal methods, are very low, and they are difficult to be explained, if only correlated to the nucleation and growth mechanism, as explained in the Section 2.3.2.1. Instead, further investigations need to achieve a complete understanding of the transformation mechanism at an issue, such as the study of crystal size evolution, and consequently of the growth rate, and the microstructural analysis by TEM, which allows following the zirconia crystallization.

2.3.2.4 Crystallite size evolution

The mean crystallite size evolution was first followed as a function of temperature during isochronal treatments, performing in-situ XRD analyses on the sample Set C I, and evaluating the crystallite size by the Scherrer’s equation. These results are reported in Figure 19.

400 600 800 1000 1200

6

8

10

12

14

16

18

20

22

24

26

Cry

stal

lite

size

(nm

)

Temperature (°C)

30°C min -1

1.5°C min -1

4.2°C min -1

5.6°C min -1

Figure 19 Mean zirconia crystallite size vs. temperature, in in-situ treated AZ5 powder at different heating rates (Set C I)


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The growth rate is roughly independent by the heating rate, in the investigated range (1.5 - 30°C min-1), but strongly affected by the temperature. Below 800°C, crystallite size remains almost constant (roughly 8 - 10 nm), whereas it varies from 10 to more than 21 nm in the temperature range 800-1200°C.

Similar results were reported by Upadhyaya et al. 27, 28, who studied the crystallization behaviour of a 3Y-TZP powder doped with different amounts of Al2O3. They investigated the progress of crystallization in the 3Y-TZP-Al2O3 powders by hot-stage XRD technique, and they described a initial slow increase in the primary particle size followed by a faster raise, starting at about 900°C for the 3Y-ZTP powder, and at higher temperature for the powder containing Al2O3. They suggested that the growth at low temperature is an interface-controlled reaction, involving a short-range ions diffusion, while at higher temperature a long range diffusion governs the rapid crystallite growth 27, 28. However, for the system here investigated, the curves presented in Figure 19 can be fitted by a growth exponential function, reaching a good correlation factor (R2 between 0.956 and 0.998), therefore it is difficult to distinguish two different mechanisms governing the growth at low and high temperature, respectively, and the abrupt raise in the growth rate can be explained by the diffusion coefficient increase with temperature.

The observed trend for crystal growth during isochronal treatment is also confirmed by the evolution of the crystallite size vs. time during the isothermal treatments, reported in Figure 20 a and b. The mean crystallite size grows as the isotherm temperature increases, and it is almost constant with time at the lower temperatures, between 500°C and 800°C. The treatment time starts to play a role in crystallite growth at 1000°C. The differences in size, observed by comparing the data from the two methods (i.e. isothermal treatments in the furnace or in the HTXRD), are reasonably imputable to the different real temperature of the samples. Anyway, the observed trends are in good agreement.

For the powders isothermally treated in HTXRD, the zirconia crystallites are about 8 nm in size at 500°C, and they grow up to 32 nm after 12 h at 1200°C; however, the phase transformation from tetragonal to monoclinic zirconia was not observed, even if no stabilisers were present, as already noticed for Set A end Set C samples, in the Section 2.3.1. The activation energy for crystallite growth of 43 KJ mol-1 was calculated from the large number of available data collected by HTXRD (Figure 21). The calculated value is in agreement with literature data, related to the growth activation energy of nanocrystalline grains 27-29. In those papers, a lower activation energy for nanocrystalline growth, with respect to the value for bulk materials, was imputed to the grain-rotation-induced grain coalescence mechanism, possible in free-standing powder 29.


71

1 10 100 1000

0

2

4

6

8

10

12

14

16

18

Cry

stal

lite

size

(nm

)

Time (min)

1000°C 800°C 600°C 500°C

(a)

10 100 10005

10

15

20

25

30

35

500°C 600°C 800°C 1000°C 1200°C

Cry

stal

lite

size

(nm

)

Time (min) (b)

Figure 20 Mean zirconia crystallite size vs. time, for the AZ5 powder samples isothermally treated in furnace, Set B (a) and in HT-XR diffractometer, Set C II (b)


72

0,6 0,7 0,8 0,9 1,0 1,1 1,2 1,3 1,4-8

-6

-4

-2

0

2

4

6

8

10

ln K

1/T X 103 (K -1)

Figure 21 Arrhenius plot for Activation Energy of crystal growth, for AZ5 powder samples, Set C II

2.3.3 TEM and HRTEM investigations

Several powder samples have been observed by TEM after different thermal treatments to investigate the nucleation and growth mechanism from an additional point of view, complementary to the ones reported in the previous section and mainly dealing with transformation kinetics. In fact, from the previous results, only partial information about the nucleation and growth mechanism of zirconia crystallization are obtained. For example, a short or undetectable induction time has been observed by XRD performed on thermal treated AZ5 powder, consequently we can deduce that a very fast nucleation occurs, but no hypothesis on the nature of this mechanism, i.e. homogeneous or heterogeneous nucleation, can be drawn out from XRD analyses. Otherwise, the beginning of crystallization was observed by TEM, on powder treated at low temperature (Set A), revealing that zirconia crystallization occurs by homogeneous nucleation. Figure 22 is an example of the observed homogeneous nucleation of zirconia into AZ5 powder treated at 600°C.


73

Figure 22 TEM image of AZ5 treated at 600°C (Set A)

A survey of the zirconia crystallite formation in AZ5 powder (Set A) are shown in Figure 23, over a wide range of treatment temperature and time. After calcination at 500°C for 1 h, alumina grains are surrounded by an amorphous layer (Figure 23 a). Few, small zirconia crystallites are observable inside the amorphous phase, almost completely detached from the alumina surfaces. After 10 h at 500°C, zirconia crystallites are slightly grown and partially aggregated (Figure 23 b). After treatment at 600°C, a lower amount of amorphous phase was detected all around the alumina grains (Figure 23 c), in agreement with the results previously shown in Figure 13, and, prolonging the treatment, the trend was similar to that observed at 500°C: discrete amorphous phase pockets appeared in which crystallites tend to agglomerate (Figure 23 d). Since in this temperature range, i.e. 500-600°C, almost negligible crystal growth has been detected by XRD analyses, as discussed in the previous section, the decrease in zirconia crystallized fraction can be likely imputed to nucleation rather than growth. Consequently, at low temperature, at the beginning of transformation, the mechanism limiting the overall transformation rate is growth.

Treatments at 800°C and 1000°C led to the progressive crystallite growth (Figures 23 e and 23 g), as expected by the results already reported in Figure 20. These observations strengthened the previous results: the treatment duration seems to have a poor influence on crystallite growth at the lower treatment temperatures, while at higher temperatures relevant zirconia growth is observed during isothermal treatment. Anyway, even if the treatment duration does not influence zirconia grains size at low temperature, it is misleading concluding that time does not affect the final composite microstructure. In fact, another important feature was pointed out only by TEM observation: due to their low affinity to the alumina grains surface, the small zirconia crystallites are drained by the amorphous phase flow into discrete pockets in which they tend to aggregate. This is of primary importance since such segregation could lead to coalescence and growth of the zirconia grains during sintering, hindering the microstructural control and the preservation of the nanoscale.


74

Figure 23 BF TEM images of AZ5 powder (Set A) after treatments at different temperatures and times: a) 500°C-1 h, b) 500°C-10 h, c) 600°C-1 h, d) 600°C-10 h, e) 800°C-1 h, f) 800°C-

10 h, g) 1000°C-1 h, h) 1000°C-10 h. The arrows indicate the zirconia nuclei drained into discrete pockets


75

HR-TEM was also performed for a deeper understanding of zirconia nucleation process. In Figure 24 a a detail of the amorphous phase surrounding an alumina grain after calcination at 500°C is presented: nanosized zirconia crystallites are dispersed in the amorphous phase and only occasionally they are in contact with the alumina grains surface. These observations clearly stated that zirconia starts to crystallize from the amorphous phase, promoted by the precursor decomposition, undergoing a homogeneous nucleation, maybe in correspondence of preferential sites, characterized by a local atomic order similar to tetragonal symmetry. Figures 24 b and 24 c give an experimental evidence of the evolution of the contact angle between zirconia nano-crystallites and alumina grains as the amorphous layer progressively disappears, after calcination at 600°C (Figure 24 b) and 1000°C (Figure 24 c). It can be qualitatively observed that the contact angle decreases with the increase of the calcination temperature, supporting the hypothesis of a homogeneous nucleation of zirconia. At 500°C, a large number of very small crystallites are formed. By increasing the temperature, zirconia crystallites undergo coalescence and improve their surface adhesion to alumina grains.


76

Figure 24: HRTEM images of the AZ5 powder (Set A) calcined at a) 500°C for 1 h (black circles point out zirconia crystallites), (b) 600°C for 1 h, (c) 1000°C for 1 h


77

Further information about the evolution of zirconia crystallites after the different heat treatments (Set A) were determined from Dark Field (DF) and Annular Dark Field (ADF) TEM technique, which allow to precisely determine the zirconia grain size and distribution (more details about this technique are reported in the Appendix). The DF TEM images of AZ5 powder treated at different temperatures are compared in Figure 25. The mean zirconia crystallite sizes are 9, 11, 17 and 25 nm for the samples treated at 500, 600, 800 and 1000°C respectively, in agreement with the above reported data.

Figure 25 DF TEM images of AZ5 powder after treatments at different temperatures: a) 500°C-1 h, b) 600°C-1 h, c) 800°C-1 h, d) 1000°C-1 h (Set A)

In conclusion, some important information have been obtained by TEM observations on AZ5:

1) zirconia homogeneously nucleates into the amorphous precursor layer;


78

2) due to the low affinity between alumina and zirconia, when the amorphous phase is present, zirconia crystallites only occasionally approach the alumina grain surface, but yielding a contact angle larger than 90° (see Figure 26);

3) due to the low affinity, during prolonged thermal treatment at low temperature, the Zr-rich amorphous phase can flow onto the alumina grain surface, in order to reduce the extension of their interface, so that also the zirconia nuclei are drained into discrete pockets among alumina particles. As a consequence, the starting homogeneity of zirconia distribution can be lost;

4) at high temperature, zirconia crystallization is faster, growth prevails on nucleation, the amorphous phase disappears without draining phenomena, so that the final zirconia grains are larger but more homogeneously distributed onto the alumina particles, if compared to the samples obtained by prolonged low temperature treatments (see Figure 27 a and b).

Figure 26 Detail of zirconia grain in AZ5 powder treated at 500°C for 1 h (Set A) by HRTEM

(a) (b)

Figure 27 HRTEM image of AZ5 powder treated at 1000°C for 1 h (Set A)


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2.3.4 An overall discussion about zirconia crystallization in AZ5 powder

Combining the results from XRD and TEM, several general comments on the mechanism of zirconia crystallization onto alumina grain surface can be pointed out.

Tetragonal zirconia starts to crystallize at very low temperature (about 400°C), as shown by XRD in Section 2.3.1. By TEM, it was also shown that zirconia crystallites start to nucleate into the amorphous layer on the alumina grains, yielded by the thermal decomposition of the zirconia precursor, following a mechanism of homogeneous nucleation. Only improving the treatment temperatures, due to the progressive disappearance of the residual amorphous phase, the zirconia crystals are forced to approach the alumina grains surface and a related evolution of the contact angle between the two phases was observed.

A very short induction time for crystallization was observed only in the case of an isothermal treatment at 500°C, supporting the hypothesis of a very fast nucleation. Instead, considering the evolution of crystal size as a function of temperature and time, a very slow crystal growth was determined in the medium-low temperature regimes even for long treatment times. In contrast, an appreciable increase of the mean zirconia dimension during isothermal treatments was observed during treatments at a temperature of 1000°C or higher.

The calculated Avrami exponents for isothermal and isocronal treatments are lower than 0.5, and they decrease with increasing temperature. These values are in agreement with the data of Ghosh et al. 28, who explained their results by the dominance of interfacial control mechanism over the bulk diffusion in the growth behaviour of zirconia. The crystallization from amorphous zirconia as a diffusionless transformation has been already reported in previous studies 11,12, based on the investigation of the structure of amorphous hydrous ZrO2. In the above studies, they described the structure of the amorphous zirconia as made of alternated layers of zirconium and oxygen atoms, with interatomic distances similar to that of tetragonal zirconia. Moreover, by electron diffraction patterns, they claimed that amorphous zirconia seemed to be a bulk sample of crystals nuclei with a structure very close to that of tetragonal phase, but with very short periodicity 11. The hypothesis of these Authors 11, 12 can therefore account for a very fast nucleation, as also resulted by the present investigations, and it also in agreement with the homogeneous nucleation observed by TEM in this study.

Therefore, the low value of the Avrami exponent and its dependence on temperature could be explained, by considering the contribution of both nucleation and growth. The nucleation rate cannot be supposed constant with temperature: in fact, it is very high at the beginning of the transformation, at temperatures lower than 600°C, and it becomes moderate at high temperature. On the contrary, the growth rate is supposed to be low at the temperatures below 1000°C, due to a poor diffusion rate. The Avrami exponent results to be influenced not only by the nucleation and growth mechanism but also by the kinetics characteristics, depending on the ratio between the growth and nucleation rates. When nucleation prevails over growth, no incubation time is observed, the crystallized fraction increase very fast at the beginning of transformation, but, as the transformation proceeds, it increases very slowly, because growth is almost negligible.


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2.3.5 A possible scenario for zirconia crystallization

A possible scenario for the zirconia crystallization can be proposed starting from the collected results. After drying the modified powders are made of alumina grains coated with a layer of amorphous zirconia precursor (Figure 28 a). Crystallization is then promoted by thermal treatments; by calcination at low temperature, a homogeneous nucleation of zirconia nano-crystals in the amorphous layer is already observed in the initial stage of the thermal treatment (Figure 28 b), while their growth rate is very slow, so that, at the intermediate stage, a slight decrease of amorphous phase is observed, mainly imputable to continuous nucleation (Figure 28 c). However, prolonging the treatment at low temperature, the amorphous phase, even if still present, is preferentially drained into discrete pockets among the alumina particles, to reduce the respective contact surface (Figure 28 d). Consequently, the calcination time at low temperature could have a negative effect on the final microstructure after sintering. In a first case, the powder calcined for short time contains uniformly distributed zirconia nuclei (Figure 28 c), so that, the resulting microstructure is homogeneous and a narrow size distribution of zirconia grains is obtained (Figure. 28 e). Otherwise, the powder calcined at low temperature for longer time contains zirconia nuclei preferentially located into pockets among the alumina particles where the amorphous phase is placed, and few isolated zirconia nano-grains (Figure 28 d). This non-uniform morphology of the composite powder could reasonably give rise to an inhomogeneous microstructure in the sintered materials, where a bimodal distribution size of zirconia grains could be expected, being the larger grains originated by the coalescence of the close nuclei into pockets and the smaller ones by the isolated grains, and also abnormal grain growth could occur into the alumina matrix (Figure 28 f).

Instead, treating at high temperature, a faster crystallization takes place (Figure 28 g), the amorphous phase disappears at shorter times, comparing with the low temperature treatments, and crystal growth prevails over nucleation, so that larger zirconia crystals are yielded, homogeneously distributed onto the surface of the alumina particles (Figure 28 h). In this case, the hypothesized microstructure of sintered bodies could be homogeneous (Figure 28 i), because the starting composite powder contains uniformly distributed zirconia on the alumina particles surface, but the mean grain size, both for the matrix and the second phase, is expected to be greater than in the case of materials obtained from the powder treated at low temperature for short time. Furthermore, the powder treatment at high temperature, as 1000°C, could induce an incipient sintering between particles, leading to the formation of hard agglomerates, difficult to break, which could be the cause inhomogeneity in the final microstructure, such as residual porosity.


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Figure 28 Possible zirconia crystallization mechanism


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2.3.5.1 Investigation of the powder thermal treatment influence on the sintered microstructures

A preliminary investigation about the influence of powder thermal treatment on sintered microstructures was performed, in order to confirm the hypotheses presented in the Section 2.3.5.

With this aim, Al2O3 - 10 vol.% ZrO2 (AZ10) powder, prepared following the post-doping procedure, was processed, assuming the AZ10 behaviour similar to the AZ5 one. AZ10 powder was divided into two batches, respectively treated at 600 °C for 1 h and for 20 h, then uniaxially pressed at 300 MPa, and finally sintered at 1500°C for 3 h. The final microstructures were observed by BSE-SEM (Figure 29). Homogeneous zirconia distribution can be observed in the sample obtained by the powder submitted to the shorter treatment (Figure 29 a and b). Otherwise, some zirconia aggregates are present in the sample obtained by powder submitted to a prolonged treatment. The difference in zirconia size histograms, obtained by the image analyses on the BSE-SEM micrographs, further demonstrates the uneven zirconia distribution in the latter sample, due to the dishomogeneous arrangement of zirconia nano-grains into the starting powder, as reported in the previous sections.

2 µµµµm2 µµµµm(a)

400 800 1200 1600 20000,00

0,05

0,10

0,15

0,20

0,25

0,30

Vol

umic

frac

tion

(%)

ZrO2 Particle Size (nm) (b)

2 µµµµm2 µµµµm2 µµµµm(c)

400 800 1200 1600 20000,00

0,05

0,10

0,15

0,20

0,25

0,30

Vol

umic

frac

tion

(%)

ZrO2 Particle Size (nm) (d)

Figure 29 BSE-SEM micrographs (a and c) and zirconia size distribution (b and d), for AZ10 sintered bodies, obtained by differentially thermal treated powder


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Thus, it possible to conclude that powder thermal treatment actually influences the final composite microstructure, and the uneven zirconia distribution into the powder leads to an uncontrolled final microstructure in the composite. As a consequence, a detrimental effect on mechanical properties and reliability can be due to unsuitable powder thermal history.

2.4 Conclusions

This chapter described the elaboration of an alumina-zirconia composite powder. The selected procedure to obtain AZ powder involved the surface modification of a commercial alumina by adding a zirconium inorganic salt as zirconia precursor.

The following elaboration parameters have been controlled and set up:

1) alumina dispersion,

2) pH and dilution of the zirconium salt solution,

3) modified suspension drying,

4) composite powder thermal treatment,

these parameters play a crucial role in the final microstructure of sintered bodies. Particularly, this latter issue was deeply investigated because the heat treatment leads to the crystallization of zirconia nano-crystals onto the alumina surfaces, whose size and distribution are influenced by the treatment temperatures and times. Consequently, the crystallization kinetics was mainly studied by XRD analyses, applying the JMA theory, and the nucleation and growth mechanisms were thoroughly investigated by TEM and HRTEM observations. The synergy among different investigation techniques allowed to complete the survey on the powder transformation, identifying the proper treatment in view of the sintered bodies production.

As a conclusion, the procedure set up in this study, is a promising method to obtained composite powders with the desired composition. Moreover, the importance of the nano-powder engineering, based on a deep knowledge of the involved transformation mechanisms, was recognized, when the study aims to a full control of the powder features and final microstructures.


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chapter 2 powder elaboration and...

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