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Oxidation of φ′-aluminium oxynitride
Corrosion Science 53 (2011) 939–945
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Oxidation of u0 -aluminium oxynitride E. Xolin, Y. Jorand ⇑, C. Olagnon, L. Gremillard Université de Lyon, INSA-Lyon, MATEIS, UMR CNRS 5510, 20, av Albert Einstein, F-69621 Villeurbanne cedex, France
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Article history: Received 16 June 2010 Accepted 17 November 2010 Available online 23 November 2010 Keywords: A. Ceramic B. TEM B. Weight loss C. High temperature corrosion C. Oxidation
a b s t r a c t The oxidation in air of single crystal u0 -aluminium oxynitride (AlON) grains has been characterized by thermogravimetry and X-ray diffraction in the 1273–1673 K range. Two oxidation stages have been observed, suggesting the formation of a transitional phase. Below 1473 K, oxidation results in the apparition of platelets and noodle-like crystals on the surface of the initially faceted single crystals. Above 1473 K, low density a-alumina polycrystals start forming on the grain surface and grow towards the grain core with increasing temperature or time. Their low density is mainly due to the presence of a network of nano-porosities. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Aluminium oxynitride, also called AlON is a compound of AlN and Al2O3. It has been discovered about 50 years ago, and was described as a c-alumina phase stabilized by nitrogen ions. It has been studied at first in order to improve the properties of silicon nitride, aluminium nitride and alumina [1]. Nowadays it is used for its mechanical, optical and chemical properties [2–4]. Several authors have successively worked on the AlN–Al2O3 phase diagram in which several phases and polytypes were reported. The first attempt was proposed by Lejus [5] who found two oxynitride phases: at high temperature, at increasing alumina content one observes the c phase of spinel structure and 1.86Al2O3–AlN composition, then the tetragonal d phase of composition 13Al2O3–AlN. Different polytypes have also been observed by Adams [6], Sakai [7] and Bartman [8]. In 1972, Michel [9] observed an additional u0 phase of tetragonal structure and 5Al2O3–AlN composition. He identified the diffraction peaks which are still used as reference for X-ray diffraction indexation. This same phase was later called e from an extension of other phase diagram [10]. MacCauley [11] then published an extended version of the phase diagram in which he restored the first name of the phase (u0 ) as originally proposed by Michel. Recently the existence of the three AlON phases (c, d and u0 ) was confirmed by an extended study of the AlN–Al2O3 binary diagram by Tabary [1], and their structure was analyzed in details. In particular Tabary verified the structure of the u0 phase and established its stability domain from 2203 to 2323 K. He proposed a calculated phase ⇑ Corresponding author. Tel.: +33 (0)4 72 43 83 51; fax: +33 (0)4 72 43 85 28. E-mail address: [email protected] (Y. Jorand). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.11.023
diagram in agreement with his experimental results where the composition of the u0 phase ranges between 4Al2O3–AlN and 8Al2O3–AlN. Although the oxidation behaviour and thermal stability of the c-AlON have been extensively studied [12–17], to the authors knowledge no such study has been conducted yet for the u0 phase. Since the structure and compositions of c and u0 phases are relatively close, it is expected that the behaviour of the former may give some insight into the behaviour of the latter. Some aspects of the different studies are not consistent, such as oxidation kinetics or temperature onsets. In addition the mechanisms of oxidation are difficult to compare, probably because of the different testing atmospheres and c-AlON compositions. There is however agreement in the formation of an intermediate metastable phase during the process, leading to a-alumina formation. The first studies, conducted by Goursat et al. [12,13] on two different powders in oxygen, showed that the oxidation starts at about 923 K and that the weight gain reaches a maximum at about 1373 K, attributed to the formation of a metastable phase (c0 ). This phase was described as a c-alumina like structure that retains nitrogen and contains vacancies. At higher temperature, this phase is transformed to a low density a-alumina (75% of the theoretical density). This very low density was attributed to the creation of additional vacancies due to nitrogen departure, associated to the observed weight loss. Further heat treatments at higher temperature increased the density of the formed alumina to the theoretical value. Lefort [15] also conducted tests on powders and agreed with the structure evolution, but he found the oxidation to start at 1073 K, to accelerate at 1273 K to finally lead to a weight loss at 1473 K. He studied the kinetics at high temperature and showed that oxidation was controlled by interface reaction. Zheng [16] studied the evolution
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of the cell parameters and the weight loss of c-AlON powders during oxidation but only recorded a weight increase. Wang [17] studied the oxidation of powders and plates in air. He also observed either a peak or a step in the weight gain during non-isothermal oxidation tests, suggesting two distinct stages of oxidation. The first mechanism starts at 1273 K and the second one at 1653 K; they show different activation energy, related to a control of the oxidation kinetics by the reaction rate initially and then by diffusion. In the same study the formation of macro-porosities and cracks on the specimens was reported. Finally c-AlON oxidation appears to be a complex phenomenon involving several reactions, and very sensitive to structural and microstructural parameters that are difficult to control. The aim of this work was to study the oxidation behaviour of a u0 -AlON in the 1273–1673 K temperature range. Isothermal and non-isothermal oxidation kinetics of powders were measured in air and related to structure and microstructure analyzed by XRD, SEM and TEM. The oxidation reactions are described. 2. Experimental procedures The present aluminium oxynitride single-crystals were obtained by melting a blend of alumina and aluminium nitride as described elsewhere [18,19]. Milling and sieving followed in order to obtain 200 mesh granulometry. The grain size distribution was determined by laser diffraction under dry conditions (Malvern, Master Sizer 2000). The experimental data were fitted to a normal distribution, giving a median grain diameter of 72 lm and a standard deviation of 24 lm. As shown in Fig. 1, the grains present faceted shapes with rough surfaces due to the presence of both very small particles and cleavage steps. The main crystallographic phase of the initial powder is the u0 phase, as determined by X-ray diffraction (XRD, Rigaku, JCPDS file 28-0028), with 5 wt.% aluminium nitride as the main impurity. An overall nitrogen weight ratio of 1.98 wt.% was measured using a chemical analysis by combustion technology with a thermal conductivity detector (HORIBA, EMGA620W/C). Oxidation kinetics were determined by isothermal and non-isothermal thermogravimetric analyses (TGA) on a system having a detection limit of 10 lg (Setaram TGA92), using about 0.25 g of AlON powder in alumina crucibles. For non-isothermal conditions, after the sample was introduced in the furnace, dry air was flown in the reaction tube at a fixed flow of 3 l/h (Air Liquide Alphagaz 1, 20 mol.% ±1 oxygen, H2O <3 mol ppm). The furnace was heated from room temperature to 1773 K at 5 K/min. For isothermal TGA experiments, the sample was maintained in an argon flux (Alphagaz 1, O2 <3 mol ppm, H2O <3 mol ppm) during heating in order to prevent any oxidation
Fig. 1. SEM micrographs of as-received u0 -AlON grains. The enlargement shows small particles on the grain surface.
below the dwell temperature. For the same purpose a 20 K/min heating rate was used to reach the prescribed temperature. To ensure temperature stabilization, a 10 min delay was applied before argon was replaced by dry air. Then the powder was oxidized during 4 h under a constant air flow of 3 l/h. Finally, oxidized powder was quenched by a fast cooling at 20 K/min. For both isothermal and non-isothermal TGA experiments, a blank obtained on calcined alumina with the same thermal cycle was employed to correct the data for system errors and buoyancy effects due to gas flow as the temperature increases. The correction was ascertained by weighting the powder before and after oxidation on an alternate balance. The higher amount of oxidized product necessary for physicochemical characterization was obtained in a furnace under static ambient air with a heating rate of 3 K/min and 2 h at dwell temperature. The powder was observed by scanning electron microscopy before and after oxidation. XRD was used to identify the structural evolutions. The structure of powders heat treated at 1573 K was observed by conventional transmission electron microscopy (MET, JEOL, 200CX). For this purpose, thin foils were prepared by embedding the powders in epoxy resin and cutting 200 lm thick slices. The thickness was further decreased to electron transparency by diamond polishing and subsequent ion-milling with argon at 4.5 keV (PIPS, Gatan). The density of the powders was measured by helium pycnometry (Micromeritics AccuPyc 1330) on samples previously dried at 473 K for 24 h. To evaluate the volume of the cracks formed during oxidation, mercury intrusion porosimetry measurements were performed on powders oxidized at 1673 K and dried at 473 K for 24 h just before analysis (Micromeritics AutoPore III, 0.35–400 MPa). 3. Results and discussion Fig. 2 shows the non-isothermal thermogravimetric curve. Weight gain initiates at 1150 K and accelerates at 1200 K. At 1440 K the weight gain rate decreases and an inflexion is observed, before it increases again as the temperature reaches 1550 K. Oxidation approaches completion at 1770 K. These observations suggest that u0 -AlON oxidation proceeds in two stages. This will be further discussed with the results of the XRD analyses of the intermediate oxidation products. As reaction is complete, a 2.7% net weight gain is noticed. Tabary [1] shown that the composition range of stable u0 -AlON is All1O15.45N0.55 to All1O14.75N1.25.
Fig. 2. Nonisothermal thermogravimetric curves.
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Therefore the net weight gain during total oxidation should lays within the range 1.65–1.91 wt.%, which is much less than the experimental value. This mismatch could be attributed to the AlN impurity. Under this hypothesis one obtains a calculated AlN content ranging from 3.53 to 4.63 wt.%, consistent with the rough estimation given by the XRD analysis of the raw material (5 wt.%). Fig. 3a presents the isothermal oxidation curves performed between 1373 and 1673 K, temperatures chosen on the basis of anisothermal experiments. The conversion ratio (a) is the ratio of the weight gain at a given time divided by the experimental net weight gain for the oxidation to be complete, i.e., 2.7 wt.%. Oxidation seems to be complete after 4 h at 1623 K and less than 2 h at 1673 K. This indicates there is no residual nitride present in the material after such treatments. In order to gain insight in the mechanisms, the reaction curves have been fitted to various kinetic models. Kinetic equations can all be represented by the general form expression:
FðaÞ ¼ k t
ð1Þ
where t is the reaction time and k the reaction rate constant, independent of reaction time. k depends on the temperature T by the
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classical relation, k = k0 exp (Ea/RT) with k0 the Arrhenius preexponential factor, R the ideal gas law constant and Ea the activation energy. Various forms of the expression F(a), summarized in Table 1, have been proposed in order to account for the effects of the numerous other parameters of the reacting system such as composition, controlling reaction mechanisms, geometrical considerations. The fluid–solid reaction considered here proceeds in two steps. First, the mobile reactant must reach the solid reactant, by diffusion through the product layer, and then the reaction itself can proceed at the phase boundary. When the reaction rate is controlled by bulk diffusion through the product layer, models developed by Ginstling and Brounshtein [21] or Carter [22] may be preferentially considered. If the chemical reaction at the interface is slow compared to diffusion, the model proposed by Sharp [23] should be the most pertinent one in our case. A third group of F(a) expressions must be considered if nucleation and growth are the rate-controlling mechanisms. The expressions proposed by Erofeyev [25] are usually adequate to describe such kinetics. None of these different models allows a fine description of our results for all temperatures. Many reasons may explain this discrepancy. First, these models apply to idealized particle shapes (spherical or planar) which is not the case of angular and irregular
Fig. 3. Isothermal oxidation curves as function of time for temperature between 1323 and 1673 K. (a) Raw data. (b) Ginstling–Brounshtein model. (c) One-dimensional rodlike growth Erofeyev model.
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Table 1 Summary of fluid–solid reaction models. Ref.
Geometry
Expression
Product layer diffusion control Jander [20] Ginstling and Brounshtein [21] Carter [22]
Pseudo-spherical Spherical Spherical
kt = (1 (1 a)1/3)2 kt = 1 2/3 a (1 a)2/3 kt = [z (1 + (z 1) a)2/3 (a 1)(1 a)2/3]/(z 1)
Phase boundary reaction control Sharp [23] Sharp [23] Jach [24]
Spherical Cylindrical Cubic
kt = 1 (1 a)1/3 kt = 1 (1 a)1/2 a = 8 k3 t3 12 k2 t2 + 6 kt
Three-dimensional isotropic Two-dimensional planar One-dimensional rod-like One-dimensional homogeneous parabolic
(kt)4 = ln (1 a) (kt)3 = ln (1 a) (kt)2 = ln (1 a) (kt)1 = ln (1 a) (kt)0.5 = ln (1 a)
Nuclei-growth control Erofeyev [25]
Nucleus activation Nucleus growth
particles examined here. Second, one can suppose that the limiting phenomenon may not be the same during the whole oxidation process; however it has not been possible to establish clear combination of mechanisms: although it was expected that Sharp’s model should provide a correct fit of the experimental data for short times – before the formation of a consequent product layer – this was never observed. Third, as pointed out by Shimizu [26], the particle size distribution is not taken into account in the models; however, in the present case a narrow granulometric distribution has been selected thus this should have only little effect. Last, evolutions of the structure of the particles, in this case the occurrence of extensive cracking during the oxidation phenomenon as will be described later in this article, are not taken into account in the models. However some fair fits are obtained with Ginstling–Brounshtein and Erofeyev (rod-like case) models for limited temperature ranges. The transformed curves of these two models are respectively given in Fig. 3b and c. Such transformed plot, F(a) vs. reaction time, should be linear (the slope of the line would give k). This is observed at intermediate temperatures (1473–1573 K) for Ginstling–Brounshtein model, and at high temperatures (above 1623 K) for Erofeyev’s. This tends to show that oxidation reaction is diffusion limited at intermediate temperatures, then limited by the growth of a new phase at high temperatures. A more consistent analysis needs further information on the particle structure evolution, which is discussed below. The evolutions of the phases, measured by XRD as a function of the oxidation temperature, are presented in Table 2. Evolutions of the diffraction patterns are given Fig. 4. The first detected structure modification starts at 1273 K with the apparition of a small fraction of c0 -alumina in addition to the main u0 -AlON phase. At 1473 K, the fractions of these phases match and an additional small content of a-alumina is present. At 1573 K, the a-alumina becomes the main phase while the c0 -alumina content strongly decreases and the u0 -AlON disappears. Above this temperature, only a-alumina is detected. The formation of the intermediary phase during oxidation is in agreement with the two step behaviour of the non-isothermal thermogravimetric results (Fig. 2). Oxidation of
Fig. 4. Evolution of the XRD patterns with oxidation temperature.
Table 2 Qualitative evaluation of the phase content as measured by XRD. Temperature (K)
u0 -AlON
c0 -Al2O3
1273 1373 1473 1573 1673
++++ ++
+ ++ +
a-Al2O3 + ++++ +++++ +++++
Fig. 5. Grain density variation and phases detected by XRD with oxidation temperature.
u0 -AlON is therefore characterized by the formation of the same intermediary phase (c0 -alumina) as that of c-AlON reported in
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Fig. 6. SEM micrographs of grains oxidized at different temperatures: (a) 1273 K, (b) 1373 K, (c) 1473 K, (d) 1573 K and (e and f) 1673 K.
Fig. 8. SEM micrograph showing a grain broken by oxidation. The arrows show intergranular porosities that are also observed by TEM.
Fig. 7. Mercury intrusion porosimetry on as-received and fully oxidized grains.
the literature. This suggests that the different oxidation reaction stages are similar for both u0 - and c-AlON and occur for the same temperature range.
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Fig. 9. TEM micrograph of oxidized grain showing both the large intergranular porosities and a network of nanometric intragranular porosities.
The density of the reaction products was measured by helium pycnometry. This technique takes into account all the different open pores and therefore provides values close to the absolute density, except in the presence of closed porosity. The results are presented in Fig. 5 as a function of oxidation temperature. Density decreases with increasing temperature down to a minimum value for an oxidation temperature of 1673 K, above which a slight increase is observed. This density evolution between 1573 and 1773 K is surprising since the only phase detected by XRD was a-alumina. X-rays depth penetration in that case is about 20 lm, which is shorter than the grain radius. The grain core phase is therefore not detected by XRD. Thus at 1573 K c0 -alumina may well be still present in the core. This hypothesis can not explain the density increase between 1673 and 1773 K. In any cases the density measured on a-alumina obtained from u0 -AlON oxidation is lower than the theoretical one (3.4 instead of 3.98 g/cm3). As pointed out in the introduction, Goursat et al. [12] have observed on c-AlON the same phenomenon; it was attributed to the presence of vacancies in high concentration. In the present case, it should be mentioned that such a low density could be also explained by the presence of closed pores; this will be discussed later. SEM observations (Fig. 6) show that grain surface and structure are strongly modified by oxidation. At low temperature 1273 K (Fig. 6a), the grain surface is no longer smooth but mainly covered by hexagonal platelets. X-ray data suggest that these platelets are c0 -alumina crystals. At 1373 K (Fig. 6b), a noodle-like surface structure also appears while a-alumina is detected. At 1473 K (Fig. 6c), platelets have disappeared, and the surface is covered of a polycrystalline layer with microcracks at grain boundaries. At 1573 K
(Fig. 6d), cracks are more open and reveal the presence of a columnar grain structure. This significant cracking of the grains is probably caused by the important volume increase (9%) associated to the density decrease and the recrystallisation. Both induce stresses released by cracking. At the final temperature of 1673 K (Fig. 6e and f), a-alumina columnar grains grow toward grain core. This suggests that at high temperature the rod-like Erofeyev model is adequate to describe the u0 -AlON oxidation, especially since diffusion cannot be a limiting process due to the extensive cracking that allows fast transport of gases to the reaction site. This is coherent with the above analysis of oxidation kinetics. Crack volume can be estimated by mercury intrusion porosimetry. Fig. 7 shows mercury intrusion volume per solid volume unit vs. pore access diameter calculated by the Washburn equation. In both cases the curves can be split into two parts. The first one corresponds to a large intrusion volume increase for a diameter of 30 and 50 lm for the as-received and oxidized grains respectively. These diameters and volume increases can be attributed to the intergranular packing porosity. Both the intrusion volume and the mean pore diameter are higher for the oxidized material, suggesting that the oxidized powders exhibit a lower bulk packing efficiency (however, since bulk packing is very sensitive to vibrations and pouring, this variation is probably meaningless). The second part of the curve is more relevant. For the as-received grains it displays a plateau without further mercury intrusion. In the same diameter range of 10–0.01 lm, the oxidized grains undergo a constantly increasing mercury intrusion. It is attributed to mercury penetration into the cracks whose volume is estimated to 6 vol.% of the total grain volume. The SEM observations showing cracks of approximately triangular shape and maximum opening of about 5–10 lm support the interpretation (Fig. 6d and e). Additional SEM observations (Fig. 8) have been conducted on a broken oxidized grain and confirmed that the columnar growth reaches the grain center upon oxidation completion. A detailed analysis of Fig. 8 also reveals the presence of 100–200 nm diameter intergranular porosities. TEM observations (Fig. 9) let out a more surprising result: the aalumina formed after oxidation is nanoporous, i.e., it contains a quite regular network of 30 nm diameter porosities. The low density of this phase can be explained by these nano-porosities, as shown by the simple geometric calculation below. Assuming a cubic arrangement of these porosities of a given radius (r = 15 nm, measured on TEM images) separated by a distance d of 45 nm (also measured on the TEM images), one can calculate the resulting density of the overall material. This calculated density q is given by:
q ¼ qth 1
4pr 3
3
3d
Fig. 10. Schematic drawing of the grain microstructure evolution with oxidation temperature.
ð2Þ
E. Xolin et al. / Corrosion Science 53 (2011) 939–945
with qth the theoretical density of a-alumina (3.98 g/cm3). The calculated density is about 3.4 g/cm3, which is in agreement with the one measured by helium pycnometry. This proves that, with our experimental conditions, the low density of the alumina formed after oxidation of u0 -AlON is mainly due to the presence of closed, nanosized porosities rather than vacancies as previously proposed by Goursat [12]. Broadly speaking, one can assume that the balance between these two phenomena depends on the structure of raw materials and the thermal treatment conditions as theses nanoporosities may result from vacancies coalescence. As a summary, the overall evolution of the grain structure induced by oxidation is schematized in Fig. 10: (a) 1273 K. In a first step oxidation begins at the surface of the initially faceted u0 -AlON single crystal with the apparition of platelets and noodle-like crystals, which seem to be c0 -alumina. (b) 1373 K. The initial grain morphology is retained, equivalent quantities of u0 -AlON and c0 -alumina are detected by XRD. aAlumina polycrystals just start forming on the grain surface. Oxidation kinetics seems to be controlled by a diffusion process. (c) 1473 K. The main detected phase is low density a-alumina with traces of c0 -alumina (probably located at the depth detection limit of XRD). Surface microcracking starts due to the volume expansion associated to the oxidation reactions. Nevertheless, oxidation kinetics still appears to be controlled by a diffusion process. (d) 1573 K. a-Alumina is the only phase detected. The surface aalumina polycrystals undergo a columnar growth towards the grain core, then associated to an extensive cracking of the grains. Oxidation kinetics is controlled by the growth of the columnar grains described by Erofeyev rod-like model. (e) 1673 K. Oxidation is complete, grains structure is a loose assembly of columnar low density a-alumina crystals. Nitrides species are completely oxidized, since conversion ratio reaches 100% after 2 h. 4. Conclusion The oxidation of u0 -AlON has been characterized and described as a function of time and temperature. It results in the formation of highly cracked, low density a-alumina polycrystals. These polycrystals arise at about 1473 K from an intermediate phase of c0 alumina, itself formed above 1273 K. The large volume expansion occurring during the transformations explains the cracking. Contrary to what was supposed in the literature, we show here that the low density of the a-alumina resulting from the oxidation process is due to the presence of a regular – if not periodical – lattice of nanometric porosities throughout the whole material, and not to the presence of oxygen or nitrogen vacancies. How these porosities are formed and organized remains to be clarified. References [1] P. Tabary, Etude du diagramme de phase Al2O3–AlN (A study of the Al2O3–AlN phase diagram), Ph.D. thesis, Paris XI, Orsay, 1997, pp. 120. [2] K. Takeda, T. Hosaka, Characteristics of new raw material AlON for refractories, Interceram 38 (1989) 18–22.
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[3] G.W. Kollenberg, H. Schneider, Microhardness of mullite at temperatures to 1000 °C, J. Am. Ceram. Soc. 72 (1989) 1739–1740. [4] J.W. MacCauley, Structure and Properties of Aluminum Nitride and AlON Ceramics, US Army Research Lab Aberdeen Proving Ground MD Weapons and Materials Research Directorate, 2002, pp. 20. [5] A. Lejus, Sur la formation à hautes températures de spinelles non stochiométriques et de phases dérivées (On the high temperature formation of non stoechiometric spinels and derived phases), Rev. Int. Hautes Temp. Refract. 1 (1964) 53–95. [6] I. Adams, T.R. Aucoin, G.A. Wolf, Luminescence in the system Al2O3–AlN, J. Electrochem. Soc. 109 (1962) 1050–1054. [7] T. Sakai, Hot-pressed oxynitrides in the system AlN–Al2O3, in: D. Kolar, S. Pejovnik, M.M. Ristic (Eds.), Sintering – Theory and Practice, Materials Science Monographs, Elsevier, Amsterdam, 1982, pp. 591. [8] S.F. Bartman, G.A. Slack, Al10N8O3 and Al9N7O3, two new repeated-layer structures in he Al2O3 system, Acta Crystallogr. B 35 (1979) 2281–2283. [9] D. Michel, Contribution à l’étude de phénomènes d’ordonnancement de défauts dans des monocristaux de matériaux réfractaires à base d’alumine et de zircone (A contribution to the study of defects organisation in monocristals of alumina and zirconia based refractories), Rev. Int. Hautes Temp. Refract. 9 (1972) 225–242. [10] P. Bassoul, A. Lefevre, J.C. Gilles, Phases epsilon obtenues par décomposition des spinelles non stoechiométriques dans les systèmes Ga2O3–MgO, Al2O3– NiO, Al2O3–AlN et Al2O3–NiO (Epsilon phases obtained by decomposition of non stoechiometric spinels in Ga2O3–MgO, Al2O3–NiO, Al2O3–AlN and Al2O3– NiO systems), Mater. Res. Bull. 11 (1976) 11–14. [11] J.W. MacCauley, N.D. Corbin, High temperature reactions and microstructures in the Al2O3–AlN system, in: F.L. Riley (Ed.), Progress in Nitrogen Ceramics, Proc. NATO Adv. Study Institute on Nitrogen Ceramics, held at Univ. of Sussex, 27 July to 7 August, 1981, Martinus Nijhoff, Boston, 1983, pp. 111–118. [12] P. Goursat, P. Goeuriot, M. Billy, Contribution à l’étude du système Al/O/N. IRéactivité de l’oxynitrure d’aluminium’’(A contribution to the study of the Al/ O/N system, I – aluminum oxynitride reactivity), Mater. Chem. 1 (1976) 131– 149. [13] P. Goursat, M. Billy, P. Goeuriot, Contribution à l’étude du système Al/O/N. II: Rétention d’azote dans les produits d’oxydation de l’oxynitrure d’aluminium (A contribution to the study of the Al/O/N system, II – nitrogen retention in the oxydation products of aluminum oxynitride), Mater. Chem. 6 (1981) 81–93. [14] D. Goeuriot, Réactivité, frittage et caractérisation des céramiques dans les systèmes alumine-oxynitrure d’aluminium gamma et nitrure de bore (Reactivity, sintering and characterisation of the gamma aluminum oxynitride–boron nitride ceramic systems), Ph.D. thesis, Lyon, Université Claude Bernard – Lyon I, 1987, pp. 319. [15] P. Lefort, G. Ado, M. Billy, Oxydation de l’oxynitrure d’aluminium pulvérulent dans l’oxygène au-dessus de 1100 °C (Oxydation of aluminum oxynitride powders in oxygen above 1100 °C), J. Phys. 47 (1989) 229–237. [16] J. Zheng, B. Forslund, On the thermal stability and oxidation behaviour of aluminum oxynitride (AlON) powder, Chin. J. Mater. Res. (1998) 497–506. [17] X. Wang, S. Du, L. Wenchao, Kinetic studies of the oxidation of c-aluminum oxynitride, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 33 (2002) 201–207. [18] D. Dubots, P. Faure, Direct nitridation process of metals having a low melting point, European Patent EP 0494 129G, 1992. [19] E. Xolin, Etude de propriétés physiques et mécaniques de matériaux granulaires cohésifs: application aux meules abrasives à liants vitreux, Ph.D. thesis, INSA-Lyon, 2005, pp. 198. [20] W. Jander, Reaktionen im festen Zustande bei hoheren Temperaturen (Reactions in solid state at high temperature), Z. Anorg. Allg. Chem. 163 (1927) 1. [21] A.M. Ginstling, B.I. Brounshtein, Concerning the diffusion kinetics of reactions in spherical particles, Zh. Prikl. Khim. 23 (1950) 1327–1338. [22] R.E. Carter, Kinetic model for solid-state reactions, J. Chem. Phys. 34 (1961) 2010–2015. [23] J.H. Sharp, G.W. Brindley, B.N. Narahari Achar, Numerical data for some commonly used solid state reaction equations, J. Am. Ceram. Soc. 49 (1966) 379–382. [24] J. Jach, The thermal decomposition of NaBrO3. Part I unirradiated material, J. Phys. Chem. Solids 24 (1963) 63–73. [25] B.V. Erofeyev, A generalized equation of chemical kinetics and its application in reactions involving solids, C.R. (Dokl.) Acad. Sci. URSS 52 (1946) 511–514. [26] A. Shimizu, Suitability of the kinetic model for estimation of powder reaction rate, Powder Technol. 100 (1998) 24–31.
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Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Oxidation of u0 -aluminium oxynitride E. Xolin, Y. Jorand ⇑, C. Olagnon, L. Gremillard Université de Lyon, INSA-Lyon, MATEIS, UMR CNRS 5510, 20, av Albert Einstein, F-69621 Villeurbanne cedex, France
a r t i c l e
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Article history: Received 16 June 2010 Accepted 17 November 2010 Available online 23 November 2010 Keywords: A. Ceramic B. TEM B. Weight loss C. High temperature corrosion C. Oxidation
a b s t r a c t The oxidation in air of single crystal u0 -aluminium oxynitride (AlON) grains has been characterized by thermogravimetry and X-ray diffraction in the 1273–1673 K range. Two oxidation stages have been observed, suggesting the formation of a transitional phase. Below 1473 K, oxidation results in the apparition of platelets and noodle-like crystals on the surface of the initially faceted single crystals. Above 1473 K, low density a-alumina polycrystals start forming on the grain surface and grow towards the grain core with increasing temperature or time. Their low density is mainly due to the presence of a network of nano-porosities. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Aluminium oxynitride, also called AlON is a compound of AlN and Al2O3. It has been discovered about 50 years ago, and was described as a c-alumina phase stabilized by nitrogen ions. It has been studied at first in order to improve the properties of silicon nitride, aluminium nitride and alumina [1]. Nowadays it is used for its mechanical, optical and chemical properties [2–4]. Several authors have successively worked on the AlN–Al2O3 phase diagram in which several phases and polytypes were reported. The first attempt was proposed by Lejus [5] who found two oxynitride phases: at high temperature, at increasing alumina content one observes the c phase of spinel structure and 1.86Al2O3–AlN composition, then the tetragonal d phase of composition 13Al2O3–AlN. Different polytypes have also been observed by Adams [6], Sakai [7] and Bartman [8]. In 1972, Michel [9] observed an additional u0 phase of tetragonal structure and 5Al2O3–AlN composition. He identified the diffraction peaks which are still used as reference for X-ray diffraction indexation. This same phase was later called e from an extension of other phase diagram [10]. MacCauley [11] then published an extended version of the phase diagram in which he restored the first name of the phase (u0 ) as originally proposed by Michel. Recently the existence of the three AlON phases (c, d and u0 ) was confirmed by an extended study of the AlN–Al2O3 binary diagram by Tabary [1], and their structure was analyzed in details. In particular Tabary verified the structure of the u0 phase and established its stability domain from 2203 to 2323 K. He proposed a calculated phase ⇑ Corresponding author. Tel.: +33 (0)4 72 43 83 51; fax: +33 (0)4 72 43 85 28. E-mail address: [email protected] (Y. Jorand). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.11.023
diagram in agreement with his experimental results where the composition of the u0 phase ranges between 4Al2O3–AlN and 8Al2O3–AlN. Although the oxidation behaviour and thermal stability of the c-AlON have been extensively studied [12–17], to the authors knowledge no such study has been conducted yet for the u0 phase. Since the structure and compositions of c and u0 phases are relatively close, it is expected that the behaviour of the former may give some insight into the behaviour of the latter. Some aspects of the different studies are not consistent, such as oxidation kinetics or temperature onsets. In addition the mechanisms of oxidation are difficult to compare, probably because of the different testing atmospheres and c-AlON compositions. There is however agreement in the formation of an intermediate metastable phase during the process, leading to a-alumina formation. The first studies, conducted by Goursat et al. [12,13] on two different powders in oxygen, showed that the oxidation starts at about 923 K and that the weight gain reaches a maximum at about 1373 K, attributed to the formation of a metastable phase (c0 ). This phase was described as a c-alumina like structure that retains nitrogen and contains vacancies. At higher temperature, this phase is transformed to a low density a-alumina (75% of the theoretical density). This very low density was attributed to the creation of additional vacancies due to nitrogen departure, associated to the observed weight loss. Further heat treatments at higher temperature increased the density of the formed alumina to the theoretical value. Lefort [15] also conducted tests on powders and agreed with the structure evolution, but he found the oxidation to start at 1073 K, to accelerate at 1273 K to finally lead to a weight loss at 1473 K. He studied the kinetics at high temperature and showed that oxidation was controlled by interface reaction. Zheng [16] studied the evolution
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of the cell parameters and the weight loss of c-AlON powders during oxidation but only recorded a weight increase. Wang [17] studied the oxidation of powders and plates in air. He also observed either a peak or a step in the weight gain during non-isothermal oxidation tests, suggesting two distinct stages of oxidation. The first mechanism starts at 1273 K and the second one at 1653 K; they show different activation energy, related to a control of the oxidation kinetics by the reaction rate initially and then by diffusion. In the same study the formation of macro-porosities and cracks on the specimens was reported. Finally c-AlON oxidation appears to be a complex phenomenon involving several reactions, and very sensitive to structural and microstructural parameters that are difficult to control. The aim of this work was to study the oxidation behaviour of a u0 -AlON in the 1273–1673 K temperature range. Isothermal and non-isothermal oxidation kinetics of powders were measured in air and related to structure and microstructure analyzed by XRD, SEM and TEM. The oxidation reactions are described. 2. Experimental procedures The present aluminium oxynitride single-crystals were obtained by melting a blend of alumina and aluminium nitride as described elsewhere [18,19]. Milling and sieving followed in order to obtain 200 mesh granulometry. The grain size distribution was determined by laser diffraction under dry conditions (Malvern, Master Sizer 2000). The experimental data were fitted to a normal distribution, giving a median grain diameter of 72 lm and a standard deviation of 24 lm. As shown in Fig. 1, the grains present faceted shapes with rough surfaces due to the presence of both very small particles and cleavage steps. The main crystallographic phase of the initial powder is the u0 phase, as determined by X-ray diffraction (XRD, Rigaku, JCPDS file 28-0028), with 5 wt.% aluminium nitride as the main impurity. An overall nitrogen weight ratio of 1.98 wt.% was measured using a chemical analysis by combustion technology with a thermal conductivity detector (HORIBA, EMGA620W/C). Oxidation kinetics were determined by isothermal and non-isothermal thermogravimetric analyses (TGA) on a system having a detection limit of 10 lg (Setaram TGA92), using about 0.25 g of AlON powder in alumina crucibles. For non-isothermal conditions, after the sample was introduced in the furnace, dry air was flown in the reaction tube at a fixed flow of 3 l/h (Air Liquide Alphagaz 1, 20 mol.% ±1 oxygen, H2O <3 mol ppm). The furnace was heated from room temperature to 1773 K at 5 K/min. For isothermal TGA experiments, the sample was maintained in an argon flux (Alphagaz 1, O2 <3 mol ppm, H2O <3 mol ppm) during heating in order to prevent any oxidation
Fig. 1. SEM micrographs of as-received u0 -AlON grains. The enlargement shows small particles on the grain surface.
below the dwell temperature. For the same purpose a 20 K/min heating rate was used to reach the prescribed temperature. To ensure temperature stabilization, a 10 min delay was applied before argon was replaced by dry air. Then the powder was oxidized during 4 h under a constant air flow of 3 l/h. Finally, oxidized powder was quenched by a fast cooling at 20 K/min. For both isothermal and non-isothermal TGA experiments, a blank obtained on calcined alumina with the same thermal cycle was employed to correct the data for system errors and buoyancy effects due to gas flow as the temperature increases. The correction was ascertained by weighting the powder before and after oxidation on an alternate balance. The higher amount of oxidized product necessary for physicochemical characterization was obtained in a furnace under static ambient air with a heating rate of 3 K/min and 2 h at dwell temperature. The powder was observed by scanning electron microscopy before and after oxidation. XRD was used to identify the structural evolutions. The structure of powders heat treated at 1573 K was observed by conventional transmission electron microscopy (MET, JEOL, 200CX). For this purpose, thin foils were prepared by embedding the powders in epoxy resin and cutting 200 lm thick slices. The thickness was further decreased to electron transparency by diamond polishing and subsequent ion-milling with argon at 4.5 keV (PIPS, Gatan). The density of the powders was measured by helium pycnometry (Micromeritics AccuPyc 1330) on samples previously dried at 473 K for 24 h. To evaluate the volume of the cracks formed during oxidation, mercury intrusion porosimetry measurements were performed on powders oxidized at 1673 K and dried at 473 K for 24 h just before analysis (Micromeritics AutoPore III, 0.35–400 MPa). 3. Results and discussion Fig. 2 shows the non-isothermal thermogravimetric curve. Weight gain initiates at 1150 K and accelerates at 1200 K. At 1440 K the weight gain rate decreases and an inflexion is observed, before it increases again as the temperature reaches 1550 K. Oxidation approaches completion at 1770 K. These observations suggest that u0 -AlON oxidation proceeds in two stages. This will be further discussed with the results of the XRD analyses of the intermediate oxidation products. As reaction is complete, a 2.7% net weight gain is noticed. Tabary [1] shown that the composition range of stable u0 -AlON is All1O15.45N0.55 to All1O14.75N1.25.
Fig. 2. Nonisothermal thermogravimetric curves.
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Therefore the net weight gain during total oxidation should lays within the range 1.65–1.91 wt.%, which is much less than the experimental value. This mismatch could be attributed to the AlN impurity. Under this hypothesis one obtains a calculated AlN content ranging from 3.53 to 4.63 wt.%, consistent with the rough estimation given by the XRD analysis of the raw material (5 wt.%). Fig. 3a presents the isothermal oxidation curves performed between 1373 and 1673 K, temperatures chosen on the basis of anisothermal experiments. The conversion ratio (a) is the ratio of the weight gain at a given time divided by the experimental net weight gain for the oxidation to be complete, i.e., 2.7 wt.%. Oxidation seems to be complete after 4 h at 1623 K and less than 2 h at 1673 K. This indicates there is no residual nitride present in the material after such treatments. In order to gain insight in the mechanisms, the reaction curves have been fitted to various kinetic models. Kinetic equations can all be represented by the general form expression:
FðaÞ ¼ k t
ð1Þ
where t is the reaction time and k the reaction rate constant, independent of reaction time. k depends on the temperature T by the
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classical relation, k = k0 exp (Ea/RT) with k0 the Arrhenius preexponential factor, R the ideal gas law constant and Ea the activation energy. Various forms of the expression F(a), summarized in Table 1, have been proposed in order to account for the effects of the numerous other parameters of the reacting system such as composition, controlling reaction mechanisms, geometrical considerations. The fluid–solid reaction considered here proceeds in two steps. First, the mobile reactant must reach the solid reactant, by diffusion through the product layer, and then the reaction itself can proceed at the phase boundary. When the reaction rate is controlled by bulk diffusion through the product layer, models developed by Ginstling and Brounshtein [21] or Carter [22] may be preferentially considered. If the chemical reaction at the interface is slow compared to diffusion, the model proposed by Sharp [23] should be the most pertinent one in our case. A third group of F(a) expressions must be considered if nucleation and growth are the rate-controlling mechanisms. The expressions proposed by Erofeyev [25] are usually adequate to describe such kinetics. None of these different models allows a fine description of our results for all temperatures. Many reasons may explain this discrepancy. First, these models apply to idealized particle shapes (spherical or planar) which is not the case of angular and irregular
Fig. 3. Isothermal oxidation curves as function of time for temperature between 1323 and 1673 K. (a) Raw data. (b) Ginstling–Brounshtein model. (c) One-dimensional rodlike growth Erofeyev model.
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Table 1 Summary of fluid–solid reaction models. Ref.
Geometry
Expression
Product layer diffusion control Jander [20] Ginstling and Brounshtein [21] Carter [22]
Pseudo-spherical Spherical Spherical
kt = (1 (1 a)1/3)2 kt = 1 2/3 a (1 a)2/3 kt = [z (1 + (z 1) a)2/3 (a 1)(1 a)2/3]/(z 1)
Phase boundary reaction control Sharp [23] Sharp [23] Jach [24]
Spherical Cylindrical Cubic
kt = 1 (1 a)1/3 kt = 1 (1 a)1/2 a = 8 k3 t3 12 k2 t2 + 6 kt
Three-dimensional isotropic Two-dimensional planar One-dimensional rod-like One-dimensional homogeneous parabolic
(kt)4 = ln (1 a) (kt)3 = ln (1 a) (kt)2 = ln (1 a) (kt)1 = ln (1 a) (kt)0.5 = ln (1 a)
Nuclei-growth control Erofeyev [25]
Nucleus activation Nucleus growth
particles examined here. Second, one can suppose that the limiting phenomenon may not be the same during the whole oxidation process; however it has not been possible to establish clear combination of mechanisms: although it was expected that Sharp’s model should provide a correct fit of the experimental data for short times – before the formation of a consequent product layer – this was never observed. Third, as pointed out by Shimizu [26], the particle size distribution is not taken into account in the models; however, in the present case a narrow granulometric distribution has been selected thus this should have only little effect. Last, evolutions of the structure of the particles, in this case the occurrence of extensive cracking during the oxidation phenomenon as will be described later in this article, are not taken into account in the models. However some fair fits are obtained with Ginstling–Brounshtein and Erofeyev (rod-like case) models for limited temperature ranges. The transformed curves of these two models are respectively given in Fig. 3b and c. Such transformed plot, F(a) vs. reaction time, should be linear (the slope of the line would give k). This is observed at intermediate temperatures (1473–1573 K) for Ginstling–Brounshtein model, and at high temperatures (above 1623 K) for Erofeyev’s. This tends to show that oxidation reaction is diffusion limited at intermediate temperatures, then limited by the growth of a new phase at high temperatures. A more consistent analysis needs further information on the particle structure evolution, which is discussed below. The evolutions of the phases, measured by XRD as a function of the oxidation temperature, are presented in Table 2. Evolutions of the diffraction patterns are given Fig. 4. The first detected structure modification starts at 1273 K with the apparition of a small fraction of c0 -alumina in addition to the main u0 -AlON phase. At 1473 K, the fractions of these phases match and an additional small content of a-alumina is present. At 1573 K, the a-alumina becomes the main phase while the c0 -alumina content strongly decreases and the u0 -AlON disappears. Above this temperature, only a-alumina is detected. The formation of the intermediary phase during oxidation is in agreement with the two step behaviour of the non-isothermal thermogravimetric results (Fig. 2). Oxidation of
Fig. 4. Evolution of the XRD patterns with oxidation temperature.
Table 2 Qualitative evaluation of the phase content as measured by XRD. Temperature (K)
u0 -AlON
c0 -Al2O3
1273 1373 1473 1573 1673
++++ ++
+ ++ +
a-Al2O3 + ++++ +++++ +++++
Fig. 5. Grain density variation and phases detected by XRD with oxidation temperature.
u0 -AlON is therefore characterized by the formation of the same intermediary phase (c0 -alumina) as that of c-AlON reported in
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Fig. 6. SEM micrographs of grains oxidized at different temperatures: (a) 1273 K, (b) 1373 K, (c) 1473 K, (d) 1573 K and (e and f) 1673 K.
Fig. 8. SEM micrograph showing a grain broken by oxidation. The arrows show intergranular porosities that are also observed by TEM.
Fig. 7. Mercury intrusion porosimetry on as-received and fully oxidized grains.
the literature. This suggests that the different oxidation reaction stages are similar for both u0 - and c-AlON and occur for the same temperature range.
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Fig. 9. TEM micrograph of oxidized grain showing both the large intergranular porosities and a network of nanometric intragranular porosities.
The density of the reaction products was measured by helium pycnometry. This technique takes into account all the different open pores and therefore provides values close to the absolute density, except in the presence of closed porosity. The results are presented in Fig. 5 as a function of oxidation temperature. Density decreases with increasing temperature down to a minimum value for an oxidation temperature of 1673 K, above which a slight increase is observed. This density evolution between 1573 and 1773 K is surprising since the only phase detected by XRD was a-alumina. X-rays depth penetration in that case is about 20 lm, which is shorter than the grain radius. The grain core phase is therefore not detected by XRD. Thus at 1573 K c0 -alumina may well be still present in the core. This hypothesis can not explain the density increase between 1673 and 1773 K. In any cases the density measured on a-alumina obtained from u0 -AlON oxidation is lower than the theoretical one (3.4 instead of 3.98 g/cm3). As pointed out in the introduction, Goursat et al. [12] have observed on c-AlON the same phenomenon; it was attributed to the presence of vacancies in high concentration. In the present case, it should be mentioned that such a low density could be also explained by the presence of closed pores; this will be discussed later. SEM observations (Fig. 6) show that grain surface and structure are strongly modified by oxidation. At low temperature 1273 K (Fig. 6a), the grain surface is no longer smooth but mainly covered by hexagonal platelets. X-ray data suggest that these platelets are c0 -alumina crystals. At 1373 K (Fig. 6b), a noodle-like surface structure also appears while a-alumina is detected. At 1473 K (Fig. 6c), platelets have disappeared, and the surface is covered of a polycrystalline layer with microcracks at grain boundaries. At 1573 K
(Fig. 6d), cracks are more open and reveal the presence of a columnar grain structure. This significant cracking of the grains is probably caused by the important volume increase (9%) associated to the density decrease and the recrystallisation. Both induce stresses released by cracking. At the final temperature of 1673 K (Fig. 6e and f), a-alumina columnar grains grow toward grain core. This suggests that at high temperature the rod-like Erofeyev model is adequate to describe the u0 -AlON oxidation, especially since diffusion cannot be a limiting process due to the extensive cracking that allows fast transport of gases to the reaction site. This is coherent with the above analysis of oxidation kinetics. Crack volume can be estimated by mercury intrusion porosimetry. Fig. 7 shows mercury intrusion volume per solid volume unit vs. pore access diameter calculated by the Washburn equation. In both cases the curves can be split into two parts. The first one corresponds to a large intrusion volume increase for a diameter of 30 and 50 lm for the as-received and oxidized grains respectively. These diameters and volume increases can be attributed to the intergranular packing porosity. Both the intrusion volume and the mean pore diameter are higher for the oxidized material, suggesting that the oxidized powders exhibit a lower bulk packing efficiency (however, since bulk packing is very sensitive to vibrations and pouring, this variation is probably meaningless). The second part of the curve is more relevant. For the as-received grains it displays a plateau without further mercury intrusion. In the same diameter range of 10–0.01 lm, the oxidized grains undergo a constantly increasing mercury intrusion. It is attributed to mercury penetration into the cracks whose volume is estimated to 6 vol.% of the total grain volume. The SEM observations showing cracks of approximately triangular shape and maximum opening of about 5–10 lm support the interpretation (Fig. 6d and e). Additional SEM observations (Fig. 8) have been conducted on a broken oxidized grain and confirmed that the columnar growth reaches the grain center upon oxidation completion. A detailed analysis of Fig. 8 also reveals the presence of 100–200 nm diameter intergranular porosities. TEM observations (Fig. 9) let out a more surprising result: the aalumina formed after oxidation is nanoporous, i.e., it contains a quite regular network of 30 nm diameter porosities. The low density of this phase can be explained by these nano-porosities, as shown by the simple geometric calculation below. Assuming a cubic arrangement of these porosities of a given radius (r = 15 nm, measured on TEM images) separated by a distance d of 45 nm (also measured on the TEM images), one can calculate the resulting density of the overall material. This calculated density q is given by:
q ¼ qth 1
4pr 3
3
3d
Fig. 10. Schematic drawing of the grain microstructure evolution with oxidation temperature.
ð2Þ
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with qth the theoretical density of a-alumina (3.98 g/cm3). The calculated density is about 3.4 g/cm3, which is in agreement with the one measured by helium pycnometry. This proves that, with our experimental conditions, the low density of the alumina formed after oxidation of u0 -AlON is mainly due to the presence of closed, nanosized porosities rather than vacancies as previously proposed by Goursat [12]. Broadly speaking, one can assume that the balance between these two phenomena depends on the structure of raw materials and the thermal treatment conditions as theses nanoporosities may result from vacancies coalescence. As a summary, the overall evolution of the grain structure induced by oxidation is schematized in Fig. 10: (a) 1273 K. In a first step oxidation begins at the surface of the initially faceted u0 -AlON single crystal with the apparition of platelets and noodle-like crystals, which seem to be c0 -alumina. (b) 1373 K. The initial grain morphology is retained, equivalent quantities of u0 -AlON and c0 -alumina are detected by XRD. aAlumina polycrystals just start forming on the grain surface. Oxidation kinetics seems to be controlled by a diffusion process. (c) 1473 K. The main detected phase is low density a-alumina with traces of c0 -alumina (probably located at the depth detection limit of XRD). Surface microcracking starts due to the volume expansion associated to the oxidation reactions. Nevertheless, oxidation kinetics still appears to be controlled by a diffusion process. (d) 1573 K. a-Alumina is the only phase detected. The surface aalumina polycrystals undergo a columnar growth towards the grain core, then associated to an extensive cracking of the grains. Oxidation kinetics is controlled by the growth of the columnar grains described by Erofeyev rod-like model. (e) 1673 K. Oxidation is complete, grains structure is a loose assembly of columnar low density a-alumina crystals. Nitrides species are completely oxidized, since conversion ratio reaches 100% after 2 h. 4. Conclusion The oxidation of u0 -AlON has been characterized and described as a function of time and temperature. It results in the formation of highly cracked, low density a-alumina polycrystals. These polycrystals arise at about 1473 K from an intermediate phase of c0 alumina, itself formed above 1273 K. The large volume expansion occurring during the transformations explains the cracking. Contrary to what was supposed in the literature, we show here that the low density of the a-alumina resulting from the oxidation process is due to the presence of a regular – if not periodical – lattice of nanometric porosities throughout the whole material, and not to the presence of oxygen or nitrogen vacancies. How these porosities are formed and organized remains to be clarified. References [1] P. Tabary, Etude du diagramme de phase Al2O3–AlN (A study of the Al2O3–AlN phase diagram), Ph.D. thesis, Paris XI, Orsay, 1997, pp. 120. [2] K. Takeda, T. Hosaka, Characteristics of new raw material AlON for refractories, Interceram 38 (1989) 18–22.
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[3] G.W. Kollenberg, H. Schneider, Microhardness of mullite at temperatures to 1000 °C, J. Am. Ceram. Soc. 72 (1989) 1739–1740. [4] J.W. MacCauley, Structure and Properties of Aluminum Nitride and AlON Ceramics, US Army Research Lab Aberdeen Proving Ground MD Weapons and Materials Research Directorate, 2002, pp. 20. [5] A. Lejus, Sur la formation à hautes températures de spinelles non stochiométriques et de phases dérivées (On the high temperature formation of non stoechiometric spinels and derived phases), Rev. Int. Hautes Temp. Refract. 1 (1964) 53–95. [6] I. Adams, T.R. Aucoin, G.A. Wolf, Luminescence in the system Al2O3–AlN, J. Electrochem. Soc. 109 (1962) 1050–1054. [7] T. Sakai, Hot-pressed oxynitrides in the system AlN–Al2O3, in: D. Kolar, S. Pejovnik, M.M. Ristic (Eds.), Sintering – Theory and Practice, Materials Science Monographs, Elsevier, Amsterdam, 1982, pp. 591. [8] S.F. Bartman, G.A. Slack, Al10N8O3 and Al9N7O3, two new repeated-layer structures in he Al2O3 system, Acta Crystallogr. B 35 (1979) 2281–2283. [9] D. Michel, Contribution à l’étude de phénomènes d’ordonnancement de défauts dans des monocristaux de matériaux réfractaires à base d’alumine et de zircone (A contribution to the study of defects organisation in monocristals of alumina and zirconia based refractories), Rev. Int. Hautes Temp. Refract. 9 (1972) 225–242. [10] P. Bassoul, A. Lefevre, J.C. Gilles, Phases epsilon obtenues par décomposition des spinelles non stoechiométriques dans les systèmes Ga2O3–MgO, Al2O3– NiO, Al2O3–AlN et Al2O3–NiO (Epsilon phases obtained by decomposition of non stoechiometric spinels in Ga2O3–MgO, Al2O3–NiO, Al2O3–AlN and Al2O3– NiO systems), Mater. Res. Bull. 11 (1976) 11–14. [11] J.W. MacCauley, N.D. Corbin, High temperature reactions and microstructures in the Al2O3–AlN system, in: F.L. Riley (Ed.), Progress in Nitrogen Ceramics, Proc. NATO Adv. Study Institute on Nitrogen Ceramics, held at Univ. of Sussex, 27 July to 7 August, 1981, Martinus Nijhoff, Boston, 1983, pp. 111–118. [12] P. Goursat, P. Goeuriot, M. Billy, Contribution à l’étude du système Al/O/N. IRéactivité de l’oxynitrure d’aluminium’’(A contribution to the study of the Al/ O/N system, I – aluminum oxynitride reactivity), Mater. Chem. 1 (1976) 131– 149. [13] P. Goursat, M. Billy, P. Goeuriot, Contribution à l’étude du système Al/O/N. II: Rétention d’azote dans les produits d’oxydation de l’oxynitrure d’aluminium (A contribution to the study of the Al/O/N system, II – nitrogen retention in the oxydation products of aluminum oxynitride), Mater. Chem. 6 (1981) 81–93. [14] D. Goeuriot, Réactivité, frittage et caractérisation des céramiques dans les systèmes alumine-oxynitrure d’aluminium gamma et nitrure de bore (Reactivity, sintering and characterisation of the gamma aluminum oxynitride–boron nitride ceramic systems), Ph.D. thesis, Lyon, Université Claude Bernard – Lyon I, 1987, pp. 319. [15] P. Lefort, G. Ado, M. Billy, Oxydation de l’oxynitrure d’aluminium pulvérulent dans l’oxygène au-dessus de 1100 °C (Oxydation of aluminum oxynitride powders in oxygen above 1100 °C), J. Phys. 47 (1989) 229–237. [16] J. Zheng, B. Forslund, On the thermal stability and oxidation behaviour of aluminum oxynitride (AlON) powder, Chin. J. Mater. Res. (1998) 497–506. [17] X. Wang, S. Du, L. Wenchao, Kinetic studies of the oxidation of c-aluminum oxynitride, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 33 (2002) 201–207. [18] D. Dubots, P. Faure, Direct nitridation process of metals having a low melting point, European Patent EP 0494 129G, 1992. [19] E. Xolin, Etude de propriétés physiques et mécaniques de matériaux granulaires cohésifs: application aux meules abrasives à liants vitreux, Ph.D. thesis, INSA-Lyon, 2005, pp. 198. [20] W. Jander, Reaktionen im festen Zustande bei hoheren Temperaturen (Reactions in solid state at high temperature), Z. Anorg. Allg. Chem. 163 (1927) 1. [21] A.M. Ginstling, B.I. Brounshtein, Concerning the diffusion kinetics of reactions in spherical particles, Zh. Prikl. Khim. 23 (1950) 1327–1338. [22] R.E. Carter, Kinetic model for solid-state reactions, J. Chem. Phys. 34 (1961) 2010–2015. [23] J.H. Sharp, G.W. Brindley, B.N. Narahari Achar, Numerical data for some commonly used solid state reaction equations, J. Am. Ceram. Soc. 49 (1966) 379–382. [24] J. Jach, The thermal decomposition of NaBrO3. Part I unirradiated material, J. Phys. Chem. Solids 24 (1963) 63–73. [25] B.V. Erofeyev, A generalized equation of chemical kinetics and its application in reactions involving solids, C.R. (Dokl.) Acad. Sci. URSS 52 (1946) 511–514. [26] A. Shimizu, Suitability of the kinetic model for estimation of powder reaction rate, Powder Technol. 100 (1998) 24–31.