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Reactive synthesis of alumina-boron nitride composites
Acta Materialia 52 (2004) 1823–1835 www.actamat-journals.com
Reactive synthesis of alumina-boron nitride composites G.J. Zhang a
a,*
, J.F. Yang b, M. Ando a, T. Ohji b, S. Kanzaki
b
Synergy Ceramics Laboratory, Fine Ceramics Research Association (FCRA), 2268-1, Shimo-Shidami, Moriyama-ku, Nagoya, Aichi 463-8687, Japan Synergy Materials Research Center, National Institute of Advanced Industrial Science and Technology (AIST), Nagoya, Aichi 463-8687, Japan
b
Received 2 May 2003; received in revised form 11 December 2003; accepted 15 December 2003
Abstract The reactions of aluminum borates (9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 ) with aluminum nitride (AlN) have been used as a new chemical route to synthesize alumina-boron nitride (Al2 O3 –BN) composites. Reaction mechanisms were investigated by TG-DTA and static reaction process. The reactions started at around 1200 °C and completed at around 1500 °C. Soaking at temperatures higher than 1800 °C resulted in the reverse reaction that caused great weight loss. Hot pressing promoted the reactions due to the improved diffusion process. The in situ formed BN phase was in agglomerate shape located at the pockets of Al2 O3 matrix particles and this distribution was suggested to be beneficial to the strength of materials with weak phase dispersoids. The fracture surface analysis demonstrated that the main fracture mode was transgranular, indicating the existence of a strong Al2 O3 network in the in situ synthesized composites. The prepared composites exhibited high strength, low YoungÕs modulus and high strain tolerance. Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Aluminum borate; Aluminum nitride; Alumina; Boron nitride; Reactive hot pressing
1. Introduction Strain tolerance improvement of ceramic materials is important for their applications under thermal and/or mechanical shock conditions. Advanced porous ceramics with controlled pore parameters, such as size, porosity, shape and distribution, demonstrate low elasticity and relatively high strength compared to that of the traditional porous ceramics [1–3]. However, the material strength decreases remarkably at the same time. Our previous results showed that high strength and low elasticity materials can be produced by dispersing a second phase of low YoungÕs modulus such as hexagonal boron nitride (h-BN) with fine particle size and homogeneous distribution in a ceramic matrix. h-BN-containing composites with low elastic modulus, high thermal and mechanical shock resistance and excellent corrosion resistance to molten metals or glass have been or are expected to be used as refractory structural parts in iron and steel or glass industries [4]. Alumina ceramics is one of the most common ceramic *
Corresponding author. Tel.: +81-527390156; fax: +81-527390051. E-mail address: [email protected] (G.J. Zhang).
materials showing excellent comprehensive properties such as high melting point, hardness, strength, wear resistance, and corrosion and oxidation resistance. By the addition of h-BN to alumina, it is possible to get Al2 O3 –BN composites with low elastic modulus, high thermal and mechanical shock resistance and excellent corrosion resistance. In BN composites, synergetic improvement effect of high strength and low YoungÕs modulus is difficult to realize through the conventional process of directly adding BN particles due to the large BN agglomerates or BN platelets possibly preexisting in raw powders [5,6]. These agglomerates or large platelets of BN may act as crack initiation sites during material failure [5] and thus result in low material strength. In addition, hot pressing is usually used to fabricate BN composites due to the poor sinterability. In this case, the obtained composites generally have obvious anisotropic microstructure and properties because of the flake-shape of BN particles. The above-mentioned problems can be solved by in situ synthesis of BN in ceramic matrixes [7,8]. Previously, we prepared various nonoxide-boron nitride composites (NOBNs) such as SiC-BN [5,6,9], Si3 N4 -SiC-BN [10],
1359-6454/$30.00 Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2003.12.021
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AlN-BN, SiAlON-BN and AlON-BN [11,12] composites based on the proposed in situ reactions. The obtained NOBNs demonstrated very fine, homogeneous and isotropic microstructures and as a result high strengths. Coblenz and Lewis [7] obtained Al2 O3 –BN composites with homogeneous and isotropic microstructures from the reaction of B2 O3 and AlN. Li et al. [13] prepared this composite through synthesizing BN in alumina using boric acid and urea as reactants. On the other hand, Kurita et al. [14] investigated the reaction sequence in the AlN–B2 O3 system. The present authors also utilized B2 O3 as the boron source to synthesize SiCBN composite through the reaction system of Si3 N4 – B2 O3 -C [15]. However, there are three main problems during the processing when B2 O3 is used as the boron source. (1) B2 O3 is moisture-sensitive; it absorbs moisture or water at room temperature and gives out H2 O vapor at elevated temperatures. Boric acid (H3 BO3 ) is in general used to replace B2 O3 as the starting material. (2) During the drying process of the mixed powder slurry, in the case of using wet-mixture method, due to the boric acid (water used) segregation occurs easily and hard agglomerates are formed or boric ester (with ethanol used) volatilizes; there is thus difficulty in obtaining mixed powders with homogeneous distribution of reactants or exact chemical compositions. (3) B2 O3 has a low melting point of 450 °C. Although this low melting point brings advantage of a transient liquid phase densification before the chemical reaction [7], the melted B2 O3 will vaporize during the sintering process under a normal pressure condition or even flow out of the dies during hot pressing. For avoiding the above problems of using B2 O3 as boron source, in this work we suggest using aluminum borates as the boron source to synthesize Al2 O3 –BN composites. According to the phase diagram of the Al2 O3 –B2 O3 system as shown in Fig. 1 [16,17], there are two stable aluminum borate compounds, namely 2Al2 O3 B2 O 3 (2AB) and 9Al2 O3 2B2 O3 (9A2B). Their decomposition temperatures are 1035 °C and higher than 1500 °C, respectively. For 2AB, it decompose at 1035 °C to 9A2B and B2 O3 liquid. These two components are water-resistant. Accordingly, synthesis of Al2 O3 –BN composites by using 2AB and 9A2B as reactants is advantageous from the processing viewpoint. The reactions are as follows: 9Al2 O3 2B2 O3 þ 4AlN ¼ 11Al2 O3 þ 4BN
ð1Þ
2Al2 O3 B2 O3 þ 2AlN ¼ 3Al2 O3 þ 2BN
ð2Þ
The calculated BN contents from the two reactions are 13.38 and 22.06 vol%, respectively. Theoretical densities of the composites are 3.734 and 3.587 g/cm3 , respectively, calculated from the density values of 3.96 g/cm3 for Al2 O3 and 2.27 g/cm3 for BN [5]. The two reaction systems are briefly denoted as 9A2B–AlN and 2AB–AlN, and the
Fig. 1. Al2 O3 –B2 O3 phase diagram.
related composites as Al2 O3 –13BN and Al2 O3 –22BN, respectively. In addition, by the combination of the two reactions, synthesis of Al2 O3 –BN composites with BN content among 13.38–22.06 vol% is possible. In this article, reaction thermodynamics, reaction and phase formation mechanisms, reactive hot pressing (RHP), microstructures and mechanical properties of the obtained materials will be discussed.
2. Thermodynamic considerations Because we could not find available thermodynamic data for 9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 , and, in addition these two compounds decompose at different temperatures according to the phase diagram shown in Fig. 1, here we will give a thermodynamic analysis based on B2 O3 . That is, 9Al2 O3 2B2 O3 is treated as 9Al2 O3 + 2B2 O3 , and 2Al2 O3 B2 O3 is treated as 2Al2 O3 + B2 O3 in the following calculation. We think that this analysis is valuable to investigate the phase formation mechanism and phase stability at high temperature. Due to the B2 O3 melting at temperature of 723 K and beginning to evaporate, the reaction equations can be written as B2 O3 ðlÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð3Þ
B2 O3 ðgÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð4Þ
The values of Gibbs free energy of the above reactions at standard condition in the temperature range 800–2200 K are (thermodynamic data are from [18])
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ðDGT °ð3Þ ¼ DH298 ° T DS° ¼ 288193 þ 54:91T ðJÞ DGT °ð4Þ ¼ DH298 ° T DS° ¼ 705589 þ 226:07T ðJÞ These reactions are exothermic. The large negative values of DGT ° (3) and DGT ° (4) in the temperature range of present study indicate that reactions (3) and (4) will occur thermodynamically under standard state. On the other hand, the adiabatic temperatures of the two exothermic reactions are 746 and 1285 K for 9A2B, and 1008 and 1842 K for 2AB, respectively, if these reactions are carried out at 298 K according to the calculation method used in the theoretical analysis of self-propagating high-temperature synthesis (SHS) process [19]. Although there is a possibility of SHS ignition with the thermodynamics (Tad ¼ 1842 K) for reaction (4) in the case of using 2AB as reactant, actually it is impossible because we have not included the enthalpy of vaporization for B2 O3 in the above calculation, namely, the real Tad becomes 904 K. Accordingly, we can conclude that there is no SHS process occurring for the above reactions. On the other hand, under low partial pressure of B2 O3 (g) at high temperature, the above reaction (4) will inversely occur as follows: Al2 O3 þ 2BN ¼ 2AlN þ B2 O3 ðgÞ "
ð5Þ
The value of standard Gibbs free energy of this reaction is the minus of reaction (4), namely DGT °ð5Þ ¼ DGT ° (4). Although the equilibrium temperature at DGT °ð5Þ ¼ 0 is very high (3121 K or 2848 °C), the reaction occurs remarkably above 1700 °C due to the quick evaporation and escape of B2 O3 (g) from the specimen. This reaction was used to synthesize AlON at 1850 °C in argon atmosphere [20]. Accordingly, for inhibiting the occurrence of reaction (5), during the synthesis of Al2 O3 –BN composites based on the reactions (1) and (2) the temperature should be lower than 1700 °C, or some means should be adopted, for example hot pressing, to prevent the escape of B2 O3 (g). Even then, the results in this work showed that the weight loss of the stoichiometric mixed powder compacts according to reactions (1) and (2) was very serious during hot pressing at 1900 °C under 30 MPa in 1.3 atm nitrogen atmosphere.
3. Experimental procedures 3.1. Starting materials 9Al2 O3 2B2 O3 (9A2B, mean particle size D50 ¼ 2.36 lm, BET specific surface 9.89 m2 /g, Shikoku Chemical Co., Marugame-shi, Japan), 2Al2 O3 B2 O3 (2AB, mean particle size D50 ¼ 2.85 lm, BET specific surface 27.91 m2 /g, Shikoku Chemical Co., Marugame-shi, Japan), AlN (F type, mean particle size D50 ¼ 1.23 lm,
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Tokuyama Chemical, Tokyo, Japan) and Al2 O3 (particle size 0.2 lm, Daimei Chemical, Nagano, Japan) were used in this study. The stoichiometric powders, according to reactions (1) and (2), were mixed in ethanol with Al2 O3 balls for 72 h in a plastic bottle and then dried. 3.2. TG-DTA study Thermogravimetry-differential thermal analysis (TGDTA2020 type, Mac Science Co., Tokyo, Japan) was conducted from room temperature (RT) to 1700 °C (limited by apparatus) at a heating rate of 10 °C/min in flowing high purity nitrogen and argon gases of 150 ml/ min. Measurement was performed twice for each condition and excellent repeatability was obtained. The results shown in this report are the average of the two measurements. 3.3. Static reaction process Powder compacts with a dimensions of about 6 mm 12 mm 18 mm (weight about 2 g, relative density about 47%) formed by uniaxially pressing under 60 MPa were soaked on a graphite plate at temperatures from 900 to 1800 °C with a temperature increase rate of 10 °C /min under 1.3 atm pressure of nitrogen and argon atmosphere for 1 h. Then the weight loss and linear shrinkage of the compacts after cooling down to room temperature were measured. X-ray diffraction analysis (XRD, Rigaku 2500V, Rigaku Denki Co., Tokyo, Japan) using Cu Ka radiation was conducted to determine the phase compositions of the obtained reacted compacts after crashed to powders. Microstructures of the fracture surfaces of the reacted compacts were observed by field emission scanning electron microscopy (FESEM, JSM-6330F, JEOL) at 15 kV. The elemental compositions of the particles were qualitatively analyzed by energy-dispersive X-ray (EDX) equipped to the FE-SEM. 3.4. Reactive hot pressing The mixed powders were hot pressed in BN-coated graphite die under 30 MPa at different temperatures for 60 min in 1.3 atm N2 atmosphere. At room temperature the powder compact was pre-pressed under 15 MPa and then the pressure was released. At temperatures of less than 600 °C, a vacuum of about 104 Torr was maintained in the furnace. Then, N2 gas was added into the chamber up to a pressure of 1.3 atm and the pressure used for hot pressing was gradually increased to 30 MPa. The temperature increase rate was 10 °C/min. For comparison an Al2 O3 monolithic ceramic was also prepared by hot pressing at 1350 °C under 30 MPa for 60 min in argon atmosphere.
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3.5. Property evaluations of the RHPed composites After removing the surface layer of the hot pressed blocks with a dimensions of 42 45 6 mm by grinding, the densities of the obtained composites were measured by the water displacement method. Then specimens were cut from the obtained blocks. The specimen surfaces were ground with 600-grade diamond wheel and the four edges were beveled. The three-point bending strength was measured on bars with dimensions of 3 4 42 mm; the span was 30 mm and the crosshead speed was 0.5 mm/min. The strength data were averages of five measurements. The pulse-echo method was used to measure YoungÕs modulus and PoissonÕs ratio [9] in the RHP direction. XRD was conducted to determine the phase compositions of the obtained composites. The fracture surfaces of the composites were observed by SEM (JSM-5600, JEOL, Tokyo, Japan). High resolution transmission electron microscopy (HRTEM) analysis of the microstructures of specimens RHPed at 1600 and 1700 °C for Al2 O3 –13BN composite and 1700 °C for Al2 O3 –22BN composite was conducted at 300 kV in a Hitachi H-9000UHR III. The specimens for HRTEM study were prepared by the usual slicing, polishing and ion etching (with argon ion for 10–15 h).
4. Results and discussion 4.1. TG-DTA investigation The TG-DTA results are shown in Fig. 2. From the DTA curves we can see that there are no apparent exothermic peaks for both the 9A2B–AlN and 2AB– AlN systems in either N2 or Ar atmosphere. This result may be related to the loose compaction of the mixed powders in the TG-DTA study. For any chemical reaction, especially for a solid state one, adequate contact of reactant particles is essential to sustain a fast reaction process. Slow exothermic phenomena are difficult to distinguish from the baseline of DTA curves. On the other hand, however, the curve for 2AB–AlN in Ar atmosphere may indicate an exothermic peak appearance at higher temperature. Furthermore, the difference between the DTA curves in N2 and Ar atmospheres indicates a faster reaction process in Ar than in N2 . Although this difference is not large and the definite reason is not clear, it may relate to the different reactivity of AlN in N2 and Ar atmospheres. This phenomenon should be similar to that in the Si3 N4 –B4 C–C reaction system, where the reaction rate is closely related to the decomposition of the reactant Si3 N4 , and Ar atmosphere is beneficial to this process [21]. The TG curves measured in N2 and Ar atmospheres are a little different especially for the 9A2B–AlN system. In the temperature range from RT to 1700 °C, the total
Fig. 2. TG-DTA curves in N2 or Ar atmosphere: (a) 9A2B–AlN system; (b) 2AB–AlN system.
weight loss are (wt%): 6.4 and 3.4 for 9A2B–AlN, and 1.8 and 0.6 for 2AB–AlN, in N2 and Ar atmospheres, respectively. The weight loss of about 2 wt% for both systems occurred below 1000 °C are considered to be mainly from the dewatering process or the decomposition of the small amount of other absorbed impurities such as organic substances. On the other hand, the weight loss at higher temperature are suggested to be partly related with the vaporization of B2 O3 . For 9A2B– AlN, however, relatively large weight loss appeared over about 1450 °C especially in N2 atmosphere. This result indicates that there are some reactions between the compacts and the N2 atmosphere to cause weight loss. A possible reaction is B2 O3 þ 2:5N2 ¼ 2BN þ 3NO "
ð6Þ
Although this reaction is thermodynamically impossible at standard condition, the removal of the by-product NO gas by the flowing nitrogen atmosphere will promote the reaction. Nevertheless, the above weight loss reaction in nitrogen atmosphere will not have large detrimental effect on the composition and properties of the synthesized composites because of: (1) the low value of weight loss
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by the reaction; (2) the probably non-detrimental effect of the remained reactant AlN, and (3) when the composites are prepared by hot pressing, the reaction (6) with gas phase by-product will be inhibited. Furthermore, if it is necessary the added amount of reactant AlN can be decreased by a small amount that can be calculated out from the B2 O3 amount lost in the weight loss reaction. The reason for the difference in the weight loss behavior of 9A2B and 2AB is not clear, but may be related to the different reaction activity and kinetics. 4.2. Static reaction analysis Fig. 3 shows the weight loss and linear shrinkage of the compacts after soaking at different temperatures in N2 or Ar atmosphere. From Fig. 3(a) we can find that these results are not in accordance with Fig. 2. Namely,
Temperature(˚C ) 0
200
400
600
800 1000 1200 1400 1600 1800 2000
0 -0.5 -1
Weight loss (%)
-1.5 -2 -2.5 (1) (2) (3) (4)
(1) 9A2B-AlN in N2 (2) 9A2B-AlN in Ar (3) 2AB-AlN in N2 (4) 2AB-AlN in Ar
-3 -3.5 -4
(a) -4.5 Temperature(˚C ) 0
200
400
600
800 1000 1200 1400 1600 1800 2000
4 (3) 2 (4) Linear shrinkage (%)
0 -2 -4 -6 -8
(1) 9A2B-AlN in N2 (2) 9A2B-AlN in Ar (3) 2AB-AlN in N2 (4) 2AB-AlN in Ar
(2) (1)
-10
(b) -12
Fig. 3. Weight loss and linear shrinkage of the powder compacts after soaking at different temperatures in N2 and Ar atmosphere: (a) weight loss; (b) linear shrinkage.
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heat treatment for 1 h in Ar atmosphere caused a little larger weight loss than that in N2 atmosphere. There should be different mechanisms for this difference in weight loss behavior for short (in the case of TG-DTA) and long (static reaction process) time period heat treatment. In the static reaction process, the denitrification of the reactant AlN or the newly formed product BN might be the controlling factor of the weight loss behavior, and this process could have been inhibited by the N2 atmosphere. It is reasonable to consider that the denitrification process is time-dependent. On the other hand, reaction (6) might be controlled by surface reaction or diffusion process of gaseous reactants or products, and the weight loss caused by this process should be limited because of the thermodynamic barrier. However, from a whole view of the weight loss behavior, if we treat the weight loss (2%) below 1000 °C to be caused by the dewatering process, the weight loss up to 1800 °C through other above-mentioned mechanisms is only 1.5%. Such a value of weight loss should not be a serious problem during sintering. The linear shrinkage of the 9A2B–AlN and 2AB–AlN is much different and the effect of atmosphere on it is negligible as shown in Fig. 3(b). For 9A2B–AlN system, the compact shrunk from about 1100 °C and a 9% of linear shrinkage demonstrated at 1800 °C. On the other hand, however, for 2AB–AlN system, the compact expanded from about 1000 °C and reached the largest value of expansion (2%) at 1400 °C, then shrank slowly. However, the compact dimension after soaking at 1800 °C still remained about 1% expansion. The shrinkage results should be related with the following three processes, i.e. (1) decomposition of 2AB at 1035 °C for specimen 2AB–AlN. This process may have resulted in the expansion around 1000–1200 °C, (2) phase formation process by the in situ reactions. This process occurs around 1300–1500 °C according to the XRD results (discussed later), (3) densification (sintering) process. Fig. 4 shows the XRD results of the powder compacts soaked at different temperatures for 60 min. For 9A2B–AlN system (Fig. 4(a)), the peaks for Al2 O3 appeared from 1200 °C, indicating the reaction started at this temperature. However, due to the low BN (the calculated content was 13.38 vol%), and the nature of any BN formed at this low temperature was in turbostratic type with low crystallinity, the diffraction peaks for t-BN could not be identified. From about 1500 °C up the peak at 2h ¼ 26:7° for t-BN became identified. The reaction completed at around 1500 °C. However, there was a small amount of AlN remained in the specimen. The reason may be related to the above discussed weight loss mechanisms from that a small part of reactant B2 O3 consumed. When the sample was soaked at 1800 °C, a new phase of aluminum oxynitride Al5 O6 N appeared. Part of this Al5 O6 N
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In fact, reaction (7) is the same with reaction (5). This reaction should be intense at temperatures higher than 1800 °C. We even found that when the specimen was soaked at 1900 °C, the specimen was almost decomposed and there was a great weight loss. This phenomenon indicated that except for reaction (7) other reactions with large weight loss should have occurred. A possible reaction is 3Al2 O3 þ 4BN ¼ 2B2 O3 ðgÞ þ 3Al2 OðgÞ þ 2N2 ðgÞ
Fig. 4. XRD patterns of the powder compacts soaked at different temperatures for 60 min in N2 atmosphere; (a) 9A2B–AlN system; (b) 2AB–AlN system.
was from the reaction of Al2 O3 and the residual AlN. Another possible way is 5Al2 O3 þ 2BN ¼ 2Al5 O6 N þ B2 O3 ðgÞ
ð7Þ
ð8Þ
The above discussion implies that the synthesis temperature of Al2 O3 –BN composites should be around 1700 °C (higher than 1500 °C but lower than 1800 °C). Also, we can conclude that the application temperature of the composites should be lower than 1800 °C even for short period of time. For 2AB–AlN system (Fig. 4(b)), the XRD results demonstrated three differences compared to those for 9A2B–AlN system. (1) 2AB decomposed to 9A2B + B2 O3 at near 1000 °C. (2) Peaks for Al2 O3 phase appeared from 1400 °C that was 200 °C higher than in the 9A2B–AlN system. (3) Peaks for 9A2B phase disappeared at about 1400 °C that was about 100 °C higher than in the 9A2B–AlN system. From Figs. 2 and 3, we can find that there was no exceptional weight loss in the decomposition process of 2AB. It may means that the formed B2 O3 was mainly in liquid state or even if it was in gaseous state at this temperature it was absorbed (‘‘caught’’) tightly by AlN or 9A2B (on the surface of particles). This phenomenon gives us chance to synthesize Al2 O3 –BN composites based on the proposed reaction (2). The existence of B2 O3 coated on the particle surface of 9A2B delivered a boron oxide environment to prevent 9A2B from reacting with AlN (see reactions (1), (9) and (10)) at lower temperature, which moved the appearance of Al2 O3 phase to a higher temperature. The microstructures of the specimens soaked at 1700 °C are shown in Fig. 5. From the particle shapes we notice that there are two types of particles, one is polyhedral with smooth surface and the other is quasispherical with granular rough surface. EDX results pointed out that the polyhedral particles were composed of Al and O (indicating the Al2 O3 phase) and the quasispherical particles were composed of B, N, Al and O (indicating the existence of BN and Al2 O3 ). It was usually easy to find fractured quasi-spherical particles on the fracture surface of the specimens, as shown in Fig. 5(c). From this picture we can realize that the quasispherical particles are actually agglomerates made of many fine (nanosized) particles. Based on the XRD and EDX results, these agglomerates were composed of t-BN and Al2 O3 . By the way, according to the discussion in Section 4.3.3, the structure of these agglomerates might be composed of an Al2 O3 core surrounded by BN.
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4:5ð2Al2 O3 B2 O3 Þ ¼ 9Al2 O3 2B2 O3 þ 2:5B2 O3 ðl; gÞ ð9Þ and the other is in the newly formed 9A2B. The former should mainly be in liquid state at the decomposition temperature of 2AB and would preferentially react with AlN at high temperature B2 O3 ðl; gÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð10Þ
Subsequently the reaction between the newly formed 9Al2 O3 2B2 O3 and AlN occurred at higher temperature. According to the SEM observation of the obtained specimens, we could not find obvious difference in microstructures except for the content of the formed t-BNcontaining agglomerates. This finding may indicate that the synthesis of t-BN is from the same mechanism, possibly from gaseous route. Furthermore, it was suggested that in both reaction systems the polyhedral Al2 O3 particles were the remnant of the 9A2B particles after the removal of B2 O3 . The shapes of these Al2 O3 particles partly maintained the shape features of the starting powders of 9A2B and 2AB. On the other hand, the agglomerates of nanosized t-BN and Al2 O3 were from the products of the reaction between AlN and B2 O3 , as shown in reaction (10). In addition, from Fig. 3(b) and the SEM observation, the pressureless synthesized Al2 O3 –BN composites are porous materials especially for 2AB–AlN system. These materials might be directly used as catalyst supports in their porous state [22,23] or be used as thermal shock resistant parts in their densified state after hot pressing [24]. On the other hand, from the shrinkage results in the static reaction process we can find that dense composites can not be manufactured by simply normal sintering of the mixed powders. 4.3. Reactive hot pressing
Fig. 5. FE-SEM micrographs of the Al2 O3 –BN composites synthesized at 1700 °C. The arrows indicate the agglomerates composed of nanosized t-BN and Al2 O3 : (a) Al2 O3 –BN composite from 9A2B–AlN; (b) Al2 O3 –BN composite from 2AB–AlN; (c) higher magnification of (a), showing the structure of agglomerate composed of nanosized t-BN and Al2 O3 .
For 9A2B–AlN system, the decomposition temperature of 9A2B is higher than the reaction starting temperature of reaction (1) (see Fig. 4(a)), so the reaction sequence should be simply the original reaction (1). On the other hand, for 2AB–AlN system, however, there are two parts of B2 O3 source, one is from the decomposition of 2AB at 1035 °C
4.3.1. X-ray diffraction The XRD patterns of the specimens hot pressed at different temperatures are shown in Fig. 6. For 9A2B– AlN system (Fig. 6(a)), peaks for Al2 O3 phase appeared from 1200 °C, which was the same with that in the pressureless sintered specimen. On the other hand, however, peaks for 9A2B disappeared at 1400 °C, 100 °C lower than that in the pressureless sintering specimen. This indicates that the applied pressure possibly promoted the reaction process by improving the diffusion conditions, for example, by decreasing diffusion distance and providing driving force for material transport. Moreover, from the heightened peak intensity for turbostratic BN (t-BN) phase we can see that the applied pressure also accelerated the formation and crystallization process of BN. If the hot pressing temperature is increased to 1800 °C, similar to the case of pressureless sintering, Al5 O6 N phase appeared from the
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seems to be relative to the percolation threshold of BN in a composite. When the BN flakes form a percolation network, the densification behavior will be obviously worsened due to the impingement occurrence (formation of card-house structure) of grain-growing BN flakes with different basal plane orientation and thus production of pores [9,11]. Fig. 7 shows the densities of the RHPed specimens as a function of the hot pressing temperature. For Al2 O3 – 13BN composite, very high relative density was obtained when the RHP temperature was higher than 1600 °C. On the other hand, however, for Al2 O3 –22BN composite only low-density materials were obtained. This result should implicate that the BN content in the Al2 O3 –13BN composite is under, but that in the Al2 O3 – 22BN composite is above the percolation limit of BN phase in the Al2 O3 matrix. In other words, we can know by inference that continuous network of BN phase has not yet formed in the Al2 O3 –13BN composite, but has formed in the Al2 O3 –22BN composite. This inference will be verified by the microstructure observation in the next section of this article. Here although we will not discuss how to get dense composite with high BN content, to control the grain growth of BN flakes, to destroy the already formed continuous BN network and to increase the slip ability of BN particles in microstructurial level should be the critical points. Two methods might be effective, one is to control the grain growth of h-BN flakes by synthesis conditions including temperature and pressure programs, and the other is to add liquid-phase-forming additives to help the flow and rotation of BN particles at hot pressing temperature like as shown in our previous study on SiC-BN in situ composite [25]. Abreal et al. [26] reported that Y2 O3 addition improved the densification Fig. 6. XRD patterns of the specimens hot pressed at different temperatures: (a) 9A2B–AlN system; (b) 2AB–AlN system.
110 Al2O3-13BN
100 Relative density (TD%)
reaction (7). The formation of gasous B2 O3 by this reaction would be detrimental to the material properties. For 2AB–AlN system, similar to the 9A2B–AlN system, a tendency of accelerated reaction process by hot press can be seen from Fig. 6(b). As discussed above, the coexistence of B2 O3 from the decomposition of 2AB inhibited the reaction between 9A2B and AlN. Namely, in phenomenal the B2 O3 phase stabilized the 9A2B phase to a higher temperature. As will be discussed later, this reaction mechanism will produce great influence on the microstructure development process.
97.29
98.44 99.11
94.38 89.25
90
82.78 78.82
80
75.72
70
77.65
72.31
Al2O3-22BN
59.02
60
1900
1800
1700
1600
1500
1400
1300
1200
55.18
50
1100
4.3.2. Densification h-BN-containing composites usually show poor densification behavior even when being manufactured by hot pressing when the BN content exceeds a limit of the range of 10–20 vol% that is determined by the particle size, shape and distribution of h-BN particles. This limit
Hot press temperature (˚C) Fig. 7. Relative density of the obtained in situ composites.
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of Al2 O3 –BN composites. In addition, it should also be effective by simply increasing the applied pressure of hot pressing. 4.3.3. TEM observation In general the microstructures of in situ synthesized ceramic composites demonstrate some features that are very different from those of conventionally processed ones. These features include special phase distribution, particle size and shape, and unique grain boundary. People can directly produce nanocomposites, whisker/ platelet-reinforced ceramic composites and interpenetrating phase composites through in situ synthesizing processes. In addition, pore structure in the in situ formed materials may also be different from that of the conventionally processed ones. For instance, closed pores may dominate the pore type in in situ composite [12]. The above-mentioned features in microstructural level endow the corresponding materials with some novel properties, such as high strength, toughness, creep resistance, wear resistance or corrosion resistance both at room and high temperatures, and in addition, good thermal shock resistance. The TEM micrographs of the obtained composites are shown in Fig. 8. The hot press direction is vertical in these pictures. In the Al2 O3 –13BN composites hot pressed at both 1600 and 1700 °C, there was a little orientation preference for Al2 O3 phase, which showed a rectangular particle shape with the longer side vertical to the hot press direction. This preferential growth of Al2 O3 particles might resulted from the considerably high temperature hot pressing for Al2 O3 phase, and the BN phase with a low fraction in this composite has not inhibited this growth process. The reaction-formed BN phase was in round shape of agglomerates located at the ‘‘pockets’’ of Al2 O3 particles. For Al2 O3 –13BN composite RHPed at 1700 °C, the reaction, densification and the grain growth of Al2 O3 were possibly too fast at the hot press temperature, some particles of BN phase were distributed at the grain boundaries or inside the grains of Al2 O3 particles. Those BN particles located at the grain boundaries might degrade the material strength. For the Al2 O3 –13BN composite hot pressed at 1600 °C, however, all the BN round agglomerates were perfectly located at those pockets surrounded by Al2 O3 particles. It means that the distribution of the weak phase BN is not continuous in the microstructure. Such a microstructure should have less detrimental effect on the macro-properties (for example, strength) of the material. A good performance in wear resistance may also be expected with this material. For a composite with a weak phase like BN as the dispersoid at a given volume fraction, it is suggested that the round agglomerates of the weak phase with a suitable size might be beneficial to result in the smallest possible effect on degrading the material strength. In
Fig. 8. TEM micrographs showing the agglomerate shape and distribution of the in situ formed BN phase in Al2 O3 matrix: (a) Al2 O3 – 13BN composite RHPed at 1600 °C; (b) Al2 O3 –13BN composite RHPed at 1700 °C; (c) Al2 O3 –22BN composite RHPed at 1700 °C.
other words, a suitable agglomeration and formation of round agglomerates will increase the percolation threshold of the weak phase, and result in an isolated distribution of and an increased free distance between the agglomerates. However, to realize this structure in a large microstructural scale there needs to be careful control of the size, distribution and volume fraction of the weak phase (or pore) in matrix material. The present process of in situ synthesis successfully provides
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a way to manufacture material with such kind of microstructure. Fig. 9 demonstrates a typical agglomerate of BN phase located at an alumina pocket in Al2 O3 –13BN composite RHPed at 1600 °C. The shape feature of this BN agglomerate was very similar to that observed in the pressureless sintered specimen (see Fig. 5(c)). This micrograph seems to reveal the picture of formation process of BN phase from the reactions. The diffusion and BN growth at the radial direction dominated the reaction process. However, there was no predominant orientation relationship between Al2 O3 and BN phases as it can be seen from Fig. 10, which shows an HRTEM micrograph of a typical grain boundary between Al2 O3 and BN phases. Similar microstructural feature was observed in Al2 O3 –22BN composite. We can also find that the interface is very clean with no amorphous phase, indicating a strong interface bonding in this in situ composite. It should be noticed that the round BN agglomerates themselves were fully dense, which was much different with the pore-containing card-house structures in conventionally processed composite. For Al2 O3 –22BN composite (Fig. 8(c)), the reaction formed BN phase showed a very similar round ag-
Fig. 9. A typical agglomerate of BN phase located at the alumina pocket in Al2 O3 –13BN composite RHPed at 1600 °C.
Fig. 10. HRTEM micrograph of an Al2 O3 /BN interface observed in Al2 O3 –13BN composite RHPed at 1600 °C, showing no predominant orientation relationship between Al2 O3 and BN phase.
glomerate shape and distribution to that in the Al2 O3 – 13BN composite except for the different BN fraction. However, this composite remained a high level of porosity and the particle size for Al2 O3 phase was small compared to that in Al2 O3 –13BN composite. In addition, at the microstructure level, although in some areas the BN phase, in round agglomerate shape, was still isolated, a BN phase network has already formed in most area. The small particle size of Al2 O3 might be resulted from the grain growth constraint by this network. On the other hand, however, another strong network formed by Al2 O3 matrix grains also existed in this composite. According to the microstructure features observed by TEM and the XRD results, a microstructure development mechanism was proposed, as shown in Fig. 11. In 9A2B–AlN system, boron oxide in 9A2B moves through solid diffusion or vaporization–solidification process to the surface of AlN particle and then reacts with AlN by mutual diffusions of nitrogen from the AlN core to the outside and oxygen from outside to the core. Consequently, we usually find such core–shell microstructure as shown in Fig. 9. In 2AB–AlN system, the B2 O3 formed by the low temperature decomposition of 2AB coated on all the surfaces of particles (Fig. 11(e)). This B2 O3 coated on the simultaneously formed 9A2B particles prevents the diffusion of boron oxide in 9A2B towards outside. Thus the peaks for Al2 O3 phase appeared at temperature higher than that in 9A2B–AlN system. The densification behavior was poor in this system due to the high BN content. However, the next section will show that even for this composite high strength can be obtained. 4.3.4. Mechanical properties The mechanical properties are listed in Table 1. For Al2 O3 –13BN composite, the specimen RHPed at 1600 °C showed the highest strength. The low strength for composite RHPed at 1500 °C mainly resulted from the low relative density (see Fig. 7). However, the relative low strength for composites RHPed at 1700 and 1800 °C should be related to the microstructure feature (Fig. 8(b)) and the reverse reaction (5). As for the small difference in strength for composites RHPed at 1700 and 1800 °C, it might point out that the reverse reaction (5) was not so serious at the present experimental condition for such a composite with low BN fraction, and on the other hand the increased density (although only a little) gave an improvement in strength for composite RHPed at 1800 °C. For Al2 O3 –22BN composite, the highest strength was obtained when it was RHPed at 1700 °C. From this result we can find that although the density for composite RHPed at 1800 °C was higher, the reverse reaction (5) should have taken a serious role in degrading material strength. The detrimental effect of the reverse
G.J. Zhang et al. / Acta Materialia 52 (2004) 1823–1835
9A2B 9A2B 9A2B
AlN
Al2O3
B 2O 3
Al2O3
Al2O3
Al2O3
9A2B Al2O3
9A2B 9A2B
Al2O3
Al2O3
Al2O3
Al2O3 and BN
Al2O3
Al2O3
Al2O3
(b)
(a)
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BN
Al2O3
Al2O3
(c) B2O3
2AB 2AB
9A2B
AlN 2AB AlN 2AB
9A2B Al2O3
AlN 9A2B AlN
2AB
9A2B
(d)
Al2O3
9A2B
BN
Al2O3
Al2O3 Al2O3
Al2O3Al O Al O 2 3 2 3
(e)
(f)
Fig. 11. Schematic mechanisms for the microstructure development of the in situ Al2 O3 –BN composites: (a)–(c) is for the 9A2B–AlN system; (d)–(f) for the 2AB–AlN system.
Table 1 Mechanical properties of the RHPed Al2 O3 –BN composites Property °Ca °C °C °C
Bending strength (MPa)
YoungÕs modulus (GPa)
Strain to failure (103 )
PoissonÕs ratio
604 74 806 66 694 75 721 61
246 266 260 257
2.46 3.03 2.67 2.81
0.206 0.211 0.206 0.198
Al2 O3 –13BN composite
1500 1600 1700 1800
Al2 O3 –22BN composite
1600 °C 1700 °C 1800 °C
171 6 275 9 254 12
80 109 111
2.14 2.52 2.29
0.219 0.196 0.175
Al2 O3 monolithic
1350 °C
757 45
411
1.84
0.237
a
Hot press temperature.
reaction (5) on strength should become large with the increase of BN content in the composite. Compared with the monolithic Al2 O3 ceramics, the in situ synthesized composites exhibited high strength, low YoungÕs modulus and high strain to failure. In addition, PoissonÕs ratios of the composites were always lower than that of the Al2 O3 monolithic ceramics. As already mentioned above, this high level strength may come from the remained strong network of the Al2 O3 matrix and the distribution feature of the in situ formed BN phase. The weak BN phase in round agglomerate shape and located at the pockets of Al2 O3 particles should have resulted in a small effect on degrading the material strength compared to a BN distribution at the matrix grain boundaries. In the material with this microstructure, the BN agglomerates and closed pores may take a similar effect on the material properties. However, compared with the pores, the BN phase would produce a composite with low friction coefficient, high wear resistance and corrosion resistance. On the other hand, the in situ process gives an opportunity to obtain highly
strong interfacial bonding between both grain boundaries of Al2 O3 /Al2 O3 and Al2 O3 /BN. Fig. 12 reveals the fracture surfaces of the Al2 O3 – 13BN composite RHPed at 1600 °C and the Al2 O3 – 22BN composite RHPed at 1700 °C observed by SEM. The hot press direction is vertical in these pictures. We can find that the dominant fracture mode for both composites was transgranular, indicating a strong bonding at the grain boundaries. Fig. 13 demonstrates an HRTEM picture of a typical grain boundary of Al2 O3 /Al2 O3 grains in Al2 O3 –22BN composite RHPed at 1700 °C. It points out the clean and strongly bonded Al2 O3 grains. Especially, for Al2 O3 –22BN composite, although the remnant porosity was 22.35%, that is, the total volume fraction of the weak BN phase and the porosity combined reached about 45%, a strong Al2 O3 network still existed. A strength of 275 MPa is also very high for Al2 O3 ceramics with a porosity of 45% [27]. Of course, the effect of BN phase and pore on material properties should be different. On the other hand, however, it may imply that by in situ process of
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2.
3.
4.
5.
Fig. 12. Fracture surfaces of the composites observed by SEM indicating the transgranular fracture mode: (a) Al2 O3 –13BN composite RHPed at 1600 °C; (b) Al2 O3 –22BN composite RHPed at 1700 °C.
6.
7.
8.
Fig. 13. HRTEM picture of a typical Al2 O3 /Al2 O3 grain boundary in Al2 O3 –22BN composite RHPed at 1700 °C, pointing out the clean and strongly bonded Al2 O3 grains.
synthesizing BN phase in ceramic matrix, high performance porous ceramic composites can be prepared. 9. 5. Conclusions A new chemical approach using aluminum borates 9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 as boron source was proposed to synthesize Al2 O3 –BN composites in the present study. The following conclusions were obtained. 1. Thermodynamic calculations revealed that the reactions in the 9A2B–AlN and 2AB–AlN systems were
exothermic, but there were no SHS processes occurred at standard conditions. TG results showed that the weight loss in nitrogen atmosphere was a little larger than that in argon atmosphere, possibly resulting from the reaction between B2 O3 and N2 . However, this reaction would not lead to obvious detrimental effect on the composition and properties of the synthesized composites. Due to the reverse reaction between the already formed Al2 O3 and BN, synthesis and subsequent applications of the composites at temperatures higher than 1800 °C should be avoided. The B2 O3 from the decomposition of 2AB at about 1000 °C coated on the surfaces of the AlN particles and the simultaneously formed 9A2B particles, resulting in the appearance of Al2 O3 phase at higher temperature through inhibiting the direct reaction between 9A2B and AlN. SEM observations of the statically reacted specimens demonstrated that the syntheiszed BN phase formed into quasi-spherical agglomerates composing of very fine (nanosized) t-BN and Al2 O3 particles. The polyhedral Al2 O3 particles kept the shape feature of the starting powders of 9A2B and 2AB, but the quasispherical agglomerates were formed by the products of the reaction between AlN and B2 O3 . The applied pressure of hot pressing promoted the reactions between aluminum borates and aluminum nitride, possibly due to improved diffusion processes. Al2 O3 –13BN composite from 9A2B–AlN system with high density of 97–98 TD% was obtained by RHP. On the other hand, however, for Al2 O3 –22BN composite from 2AB–AlN system, density higher than 80 TD% could not be reached under the present RHP conditions owing to the formation of BN networks. TEM observations of the RHPed specimens showed that the in situ formed BN phase was in round agglomerate shape located at the pockets of Al2 O3 matrix particles. It was suggested that such distribution of BN phase was beneficial to the strength of such materials with weak phase dispersoids. Fracture surface observations by SEM revealed that the main fracture mode was transgranular for both Al2 O3 –13BN and Al2 O3 –22BN composites, pointing out the existence of a strong network of the Al2 O3 matrix. The obtained composites showed high strength, low YoungÕs modulus and improved strain tolerance.
Acknowledgements This work has been supported by AIST, METI, Japan, as part of the Synergy Ceramics Project. Part of the work has been supported by NEDO. The authors
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are members of the Joint Research Consortium of Synergy Ceramics.
References [1] Inagaki Y, Ohji T, Kanzaki S, Shigegaki Y. J Am Ceram Soc 2000;83:1807. [2] Yang JF, Zhang GJ, Kondo N, Ohji T. Acta Mater 2002;50:4831. [3] Deng ZY, Fukasawa T, Ando M, Zhang GJ, Ohji T. Acta Mater 2001;49:1939. [4] Paine RT, Narula CK. Chem Rev 1990;90:73. [5] Zhang GJ, Beppu Y, Ohji T, Kanzaki S. Acta Mater 2001;49:77. [6] Zhang GJ, Ohji T. J Am Ceram Soc 2001;84:1475. [7] Coblenz WS, Lewis III D. J Am Ceram Soc 1988;71:1080. [8] Thevenot F, Doche C, Mongeot H, Guilhon F, Miele P, Bonnetot B. J Eur Ceram Soc 1997;17:1911. [9] Zhang GJ, Ohji T. J Mater Res 2000;15:1876. [10] Zhang GJ, Yang JF, Ohji T. Mater Sci Eng A 2002;328:201. [11] Zhang GJ, Yang JF, Ando M, Ohji T. J Am Ceram Soc 2002;85:2938. [12] Zhang GJ, Yang JF, Ohji T, Kanzaki S. Adv Eng Mater 2002;4:15.
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[13] Li YL, Qiao GJ, Jin ZH. Mater Res Bull 2002;37:1401. [14] Kurita K, Zeng ZQ, Takebe H, Morinaga H. Mater Trans JIM 1994;35:258. [15] Zhang GJ, Beppu Y, Ando M, Yang JF, Ohji T. J Am Ceram Soc 2002;85:2858. [16] Gielisse PJM, Foster WR. Nature 1962;195:69. [17] Ray SP. J Am Ceram Soc 1992;75:2605. [18] Chase Jr MW. NIST-JANAF thermochemical tables. 4th ed. New York, USA: American Chemical Society and American Institute of Physics; 1998. [19] Munir ZA. Am Ceram Soc Bull 1988;67:342. [20] Labbe JC, Jeanne A, Roult G. Ceram Int 1992;18:81. [21] Zhang GJ, Yang JF, Ohji T. J Am Ceram Soc 2002;85:2256. [22] Jacobsen CJH. J Catal 2001;200:1. [23] Perdigon-Melon JA, Auroux A, Cornu D, Miele P, Toury B, Bonnetot B. J Organomet Chem 2002;657:98. [24] Goeuriot-Launay D, Brayet G, Thevenot F. J Mater Sci Lett 1986;5:940. [25] Zhang GJ, Yang JF, Deng ZY, Ohji T. J Ceram Soc Jpn 2001;109:45. [26] Abreal A, Goeuriot D, Thevenot F, Lagace M, Gueroult B, Rigard M. J Eur Ceram Soc 1995;15:841. [27] Jayaseelan DD, Kondo N, Brito ME, Ohji T. J Am Ceram Soc 2002;85:267.
Reactive synthesis of alumina-boron nitride composites G.J. Zhang a
a,*
, J.F. Yang b, M. Ando a, T. Ohji b, S. Kanzaki
b
Synergy Ceramics Laboratory, Fine Ceramics Research Association (FCRA), 2268-1, Shimo-Shidami, Moriyama-ku, Nagoya, Aichi 463-8687, Japan Synergy Materials Research Center, National Institute of Advanced Industrial Science and Technology (AIST), Nagoya, Aichi 463-8687, Japan
b
Received 2 May 2003; received in revised form 11 December 2003; accepted 15 December 2003
Abstract The reactions of aluminum borates (9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 ) with aluminum nitride (AlN) have been used as a new chemical route to synthesize alumina-boron nitride (Al2 O3 –BN) composites. Reaction mechanisms were investigated by TG-DTA and static reaction process. The reactions started at around 1200 °C and completed at around 1500 °C. Soaking at temperatures higher than 1800 °C resulted in the reverse reaction that caused great weight loss. Hot pressing promoted the reactions due to the improved diffusion process. The in situ formed BN phase was in agglomerate shape located at the pockets of Al2 O3 matrix particles and this distribution was suggested to be beneficial to the strength of materials with weak phase dispersoids. The fracture surface analysis demonstrated that the main fracture mode was transgranular, indicating the existence of a strong Al2 O3 network in the in situ synthesized composites. The prepared composites exhibited high strength, low YoungÕs modulus and high strain tolerance. Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Aluminum borate; Aluminum nitride; Alumina; Boron nitride; Reactive hot pressing
1. Introduction Strain tolerance improvement of ceramic materials is important for their applications under thermal and/or mechanical shock conditions. Advanced porous ceramics with controlled pore parameters, such as size, porosity, shape and distribution, demonstrate low elasticity and relatively high strength compared to that of the traditional porous ceramics [1–3]. However, the material strength decreases remarkably at the same time. Our previous results showed that high strength and low elasticity materials can be produced by dispersing a second phase of low YoungÕs modulus such as hexagonal boron nitride (h-BN) with fine particle size and homogeneous distribution in a ceramic matrix. h-BN-containing composites with low elastic modulus, high thermal and mechanical shock resistance and excellent corrosion resistance to molten metals or glass have been or are expected to be used as refractory structural parts in iron and steel or glass industries [4]. Alumina ceramics is one of the most common ceramic *
Corresponding author. Tel.: +81-527390156; fax: +81-527390051. E-mail address: [email protected] (G.J. Zhang).
materials showing excellent comprehensive properties such as high melting point, hardness, strength, wear resistance, and corrosion and oxidation resistance. By the addition of h-BN to alumina, it is possible to get Al2 O3 –BN composites with low elastic modulus, high thermal and mechanical shock resistance and excellent corrosion resistance. In BN composites, synergetic improvement effect of high strength and low YoungÕs modulus is difficult to realize through the conventional process of directly adding BN particles due to the large BN agglomerates or BN platelets possibly preexisting in raw powders [5,6]. These agglomerates or large platelets of BN may act as crack initiation sites during material failure [5] and thus result in low material strength. In addition, hot pressing is usually used to fabricate BN composites due to the poor sinterability. In this case, the obtained composites generally have obvious anisotropic microstructure and properties because of the flake-shape of BN particles. The above-mentioned problems can be solved by in situ synthesis of BN in ceramic matrixes [7,8]. Previously, we prepared various nonoxide-boron nitride composites (NOBNs) such as SiC-BN [5,6,9], Si3 N4 -SiC-BN [10],
1359-6454/$30.00 Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2003.12.021
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AlN-BN, SiAlON-BN and AlON-BN [11,12] composites based on the proposed in situ reactions. The obtained NOBNs demonstrated very fine, homogeneous and isotropic microstructures and as a result high strengths. Coblenz and Lewis [7] obtained Al2 O3 –BN composites with homogeneous and isotropic microstructures from the reaction of B2 O3 and AlN. Li et al. [13] prepared this composite through synthesizing BN in alumina using boric acid and urea as reactants. On the other hand, Kurita et al. [14] investigated the reaction sequence in the AlN–B2 O3 system. The present authors also utilized B2 O3 as the boron source to synthesize SiCBN composite through the reaction system of Si3 N4 – B2 O3 -C [15]. However, there are three main problems during the processing when B2 O3 is used as the boron source. (1) B2 O3 is moisture-sensitive; it absorbs moisture or water at room temperature and gives out H2 O vapor at elevated temperatures. Boric acid (H3 BO3 ) is in general used to replace B2 O3 as the starting material. (2) During the drying process of the mixed powder slurry, in the case of using wet-mixture method, due to the boric acid (water used) segregation occurs easily and hard agglomerates are formed or boric ester (with ethanol used) volatilizes; there is thus difficulty in obtaining mixed powders with homogeneous distribution of reactants or exact chemical compositions. (3) B2 O3 has a low melting point of 450 °C. Although this low melting point brings advantage of a transient liquid phase densification before the chemical reaction [7], the melted B2 O3 will vaporize during the sintering process under a normal pressure condition or even flow out of the dies during hot pressing. For avoiding the above problems of using B2 O3 as boron source, in this work we suggest using aluminum borates as the boron source to synthesize Al2 O3 –BN composites. According to the phase diagram of the Al2 O3 –B2 O3 system as shown in Fig. 1 [16,17], there are two stable aluminum borate compounds, namely 2Al2 O3 B2 O 3 (2AB) and 9Al2 O3 2B2 O3 (9A2B). Their decomposition temperatures are 1035 °C and higher than 1500 °C, respectively. For 2AB, it decompose at 1035 °C to 9A2B and B2 O3 liquid. These two components are water-resistant. Accordingly, synthesis of Al2 O3 –BN composites by using 2AB and 9A2B as reactants is advantageous from the processing viewpoint. The reactions are as follows: 9Al2 O3 2B2 O3 þ 4AlN ¼ 11Al2 O3 þ 4BN
ð1Þ
2Al2 O3 B2 O3 þ 2AlN ¼ 3Al2 O3 þ 2BN
ð2Þ
The calculated BN contents from the two reactions are 13.38 and 22.06 vol%, respectively. Theoretical densities of the composites are 3.734 and 3.587 g/cm3 , respectively, calculated from the density values of 3.96 g/cm3 for Al2 O3 and 2.27 g/cm3 for BN [5]. The two reaction systems are briefly denoted as 9A2B–AlN and 2AB–AlN, and the
Fig. 1. Al2 O3 –B2 O3 phase diagram.
related composites as Al2 O3 –13BN and Al2 O3 –22BN, respectively. In addition, by the combination of the two reactions, synthesis of Al2 O3 –BN composites with BN content among 13.38–22.06 vol% is possible. In this article, reaction thermodynamics, reaction and phase formation mechanisms, reactive hot pressing (RHP), microstructures and mechanical properties of the obtained materials will be discussed.
2. Thermodynamic considerations Because we could not find available thermodynamic data for 9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 , and, in addition these two compounds decompose at different temperatures according to the phase diagram shown in Fig. 1, here we will give a thermodynamic analysis based on B2 O3 . That is, 9Al2 O3 2B2 O3 is treated as 9Al2 O3 + 2B2 O3 , and 2Al2 O3 B2 O3 is treated as 2Al2 O3 + B2 O3 in the following calculation. We think that this analysis is valuable to investigate the phase formation mechanism and phase stability at high temperature. Due to the B2 O3 melting at temperature of 723 K and beginning to evaporate, the reaction equations can be written as B2 O3 ðlÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð3Þ
B2 O3 ðgÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð4Þ
The values of Gibbs free energy of the above reactions at standard condition in the temperature range 800–2200 K are (thermodynamic data are from [18])
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ðDGT °ð3Þ ¼ DH298 ° T DS° ¼ 288193 þ 54:91T ðJÞ DGT °ð4Þ ¼ DH298 ° T DS° ¼ 705589 þ 226:07T ðJÞ These reactions are exothermic. The large negative values of DGT ° (3) and DGT ° (4) in the temperature range of present study indicate that reactions (3) and (4) will occur thermodynamically under standard state. On the other hand, the adiabatic temperatures of the two exothermic reactions are 746 and 1285 K for 9A2B, and 1008 and 1842 K for 2AB, respectively, if these reactions are carried out at 298 K according to the calculation method used in the theoretical analysis of self-propagating high-temperature synthesis (SHS) process [19]. Although there is a possibility of SHS ignition with the thermodynamics (Tad ¼ 1842 K) for reaction (4) in the case of using 2AB as reactant, actually it is impossible because we have not included the enthalpy of vaporization for B2 O3 in the above calculation, namely, the real Tad becomes 904 K. Accordingly, we can conclude that there is no SHS process occurring for the above reactions. On the other hand, under low partial pressure of B2 O3 (g) at high temperature, the above reaction (4) will inversely occur as follows: Al2 O3 þ 2BN ¼ 2AlN þ B2 O3 ðgÞ "
ð5Þ
The value of standard Gibbs free energy of this reaction is the minus of reaction (4), namely DGT °ð5Þ ¼ DGT ° (4). Although the equilibrium temperature at DGT °ð5Þ ¼ 0 is very high (3121 K or 2848 °C), the reaction occurs remarkably above 1700 °C due to the quick evaporation and escape of B2 O3 (g) from the specimen. This reaction was used to synthesize AlON at 1850 °C in argon atmosphere [20]. Accordingly, for inhibiting the occurrence of reaction (5), during the synthesis of Al2 O3 –BN composites based on the reactions (1) and (2) the temperature should be lower than 1700 °C, or some means should be adopted, for example hot pressing, to prevent the escape of B2 O3 (g). Even then, the results in this work showed that the weight loss of the stoichiometric mixed powder compacts according to reactions (1) and (2) was very serious during hot pressing at 1900 °C under 30 MPa in 1.3 atm nitrogen atmosphere.
3. Experimental procedures 3.1. Starting materials 9Al2 O3 2B2 O3 (9A2B, mean particle size D50 ¼ 2.36 lm, BET specific surface 9.89 m2 /g, Shikoku Chemical Co., Marugame-shi, Japan), 2Al2 O3 B2 O3 (2AB, mean particle size D50 ¼ 2.85 lm, BET specific surface 27.91 m2 /g, Shikoku Chemical Co., Marugame-shi, Japan), AlN (F type, mean particle size D50 ¼ 1.23 lm,
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Tokuyama Chemical, Tokyo, Japan) and Al2 O3 (particle size 0.2 lm, Daimei Chemical, Nagano, Japan) were used in this study. The stoichiometric powders, according to reactions (1) and (2), were mixed in ethanol with Al2 O3 balls for 72 h in a plastic bottle and then dried. 3.2. TG-DTA study Thermogravimetry-differential thermal analysis (TGDTA2020 type, Mac Science Co., Tokyo, Japan) was conducted from room temperature (RT) to 1700 °C (limited by apparatus) at a heating rate of 10 °C/min in flowing high purity nitrogen and argon gases of 150 ml/ min. Measurement was performed twice for each condition and excellent repeatability was obtained. The results shown in this report are the average of the two measurements. 3.3. Static reaction process Powder compacts with a dimensions of about 6 mm 12 mm 18 mm (weight about 2 g, relative density about 47%) formed by uniaxially pressing under 60 MPa were soaked on a graphite plate at temperatures from 900 to 1800 °C with a temperature increase rate of 10 °C /min under 1.3 atm pressure of nitrogen and argon atmosphere for 1 h. Then the weight loss and linear shrinkage of the compacts after cooling down to room temperature were measured. X-ray diffraction analysis (XRD, Rigaku 2500V, Rigaku Denki Co., Tokyo, Japan) using Cu Ka radiation was conducted to determine the phase compositions of the obtained reacted compacts after crashed to powders. Microstructures of the fracture surfaces of the reacted compacts were observed by field emission scanning electron microscopy (FESEM, JSM-6330F, JEOL) at 15 kV. The elemental compositions of the particles were qualitatively analyzed by energy-dispersive X-ray (EDX) equipped to the FE-SEM. 3.4. Reactive hot pressing The mixed powders were hot pressed in BN-coated graphite die under 30 MPa at different temperatures for 60 min in 1.3 atm N2 atmosphere. At room temperature the powder compact was pre-pressed under 15 MPa and then the pressure was released. At temperatures of less than 600 °C, a vacuum of about 104 Torr was maintained in the furnace. Then, N2 gas was added into the chamber up to a pressure of 1.3 atm and the pressure used for hot pressing was gradually increased to 30 MPa. The temperature increase rate was 10 °C/min. For comparison an Al2 O3 monolithic ceramic was also prepared by hot pressing at 1350 °C under 30 MPa for 60 min in argon atmosphere.
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3.5. Property evaluations of the RHPed composites After removing the surface layer of the hot pressed blocks with a dimensions of 42 45 6 mm by grinding, the densities of the obtained composites were measured by the water displacement method. Then specimens were cut from the obtained blocks. The specimen surfaces were ground with 600-grade diamond wheel and the four edges were beveled. The three-point bending strength was measured on bars with dimensions of 3 4 42 mm; the span was 30 mm and the crosshead speed was 0.5 mm/min. The strength data were averages of five measurements. The pulse-echo method was used to measure YoungÕs modulus and PoissonÕs ratio [9] in the RHP direction. XRD was conducted to determine the phase compositions of the obtained composites. The fracture surfaces of the composites were observed by SEM (JSM-5600, JEOL, Tokyo, Japan). High resolution transmission electron microscopy (HRTEM) analysis of the microstructures of specimens RHPed at 1600 and 1700 °C for Al2 O3 –13BN composite and 1700 °C for Al2 O3 –22BN composite was conducted at 300 kV in a Hitachi H-9000UHR III. The specimens for HRTEM study were prepared by the usual slicing, polishing and ion etching (with argon ion for 10–15 h).
4. Results and discussion 4.1. TG-DTA investigation The TG-DTA results are shown in Fig. 2. From the DTA curves we can see that there are no apparent exothermic peaks for both the 9A2B–AlN and 2AB– AlN systems in either N2 or Ar atmosphere. This result may be related to the loose compaction of the mixed powders in the TG-DTA study. For any chemical reaction, especially for a solid state one, adequate contact of reactant particles is essential to sustain a fast reaction process. Slow exothermic phenomena are difficult to distinguish from the baseline of DTA curves. On the other hand, however, the curve for 2AB–AlN in Ar atmosphere may indicate an exothermic peak appearance at higher temperature. Furthermore, the difference between the DTA curves in N2 and Ar atmospheres indicates a faster reaction process in Ar than in N2 . Although this difference is not large and the definite reason is not clear, it may relate to the different reactivity of AlN in N2 and Ar atmospheres. This phenomenon should be similar to that in the Si3 N4 –B4 C–C reaction system, where the reaction rate is closely related to the decomposition of the reactant Si3 N4 , and Ar atmosphere is beneficial to this process [21]. The TG curves measured in N2 and Ar atmospheres are a little different especially for the 9A2B–AlN system. In the temperature range from RT to 1700 °C, the total
Fig. 2. TG-DTA curves in N2 or Ar atmosphere: (a) 9A2B–AlN system; (b) 2AB–AlN system.
weight loss are (wt%): 6.4 and 3.4 for 9A2B–AlN, and 1.8 and 0.6 for 2AB–AlN, in N2 and Ar atmospheres, respectively. The weight loss of about 2 wt% for both systems occurred below 1000 °C are considered to be mainly from the dewatering process or the decomposition of the small amount of other absorbed impurities such as organic substances. On the other hand, the weight loss at higher temperature are suggested to be partly related with the vaporization of B2 O3 . For 9A2B– AlN, however, relatively large weight loss appeared over about 1450 °C especially in N2 atmosphere. This result indicates that there are some reactions between the compacts and the N2 atmosphere to cause weight loss. A possible reaction is B2 O3 þ 2:5N2 ¼ 2BN þ 3NO "
ð6Þ
Although this reaction is thermodynamically impossible at standard condition, the removal of the by-product NO gas by the flowing nitrogen atmosphere will promote the reaction. Nevertheless, the above weight loss reaction in nitrogen atmosphere will not have large detrimental effect on the composition and properties of the synthesized composites because of: (1) the low value of weight loss
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by the reaction; (2) the probably non-detrimental effect of the remained reactant AlN, and (3) when the composites are prepared by hot pressing, the reaction (6) with gas phase by-product will be inhibited. Furthermore, if it is necessary the added amount of reactant AlN can be decreased by a small amount that can be calculated out from the B2 O3 amount lost in the weight loss reaction. The reason for the difference in the weight loss behavior of 9A2B and 2AB is not clear, but may be related to the different reaction activity and kinetics. 4.2. Static reaction analysis Fig. 3 shows the weight loss and linear shrinkage of the compacts after soaking at different temperatures in N2 or Ar atmosphere. From Fig. 3(a) we can find that these results are not in accordance with Fig. 2. Namely,
Temperature(˚C ) 0
200
400
600
800 1000 1200 1400 1600 1800 2000
0 -0.5 -1
Weight loss (%)
-1.5 -2 -2.5 (1) (2) (3) (4)
(1) 9A2B-AlN in N2 (2) 9A2B-AlN in Ar (3) 2AB-AlN in N2 (4) 2AB-AlN in Ar
-3 -3.5 -4
(a) -4.5 Temperature(˚C ) 0
200
400
600
800 1000 1200 1400 1600 1800 2000
4 (3) 2 (4) Linear shrinkage (%)
0 -2 -4 -6 -8
(1) 9A2B-AlN in N2 (2) 9A2B-AlN in Ar (3) 2AB-AlN in N2 (4) 2AB-AlN in Ar
(2) (1)
-10
(b) -12
Fig. 3. Weight loss and linear shrinkage of the powder compacts after soaking at different temperatures in N2 and Ar atmosphere: (a) weight loss; (b) linear shrinkage.
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heat treatment for 1 h in Ar atmosphere caused a little larger weight loss than that in N2 atmosphere. There should be different mechanisms for this difference in weight loss behavior for short (in the case of TG-DTA) and long (static reaction process) time period heat treatment. In the static reaction process, the denitrification of the reactant AlN or the newly formed product BN might be the controlling factor of the weight loss behavior, and this process could have been inhibited by the N2 atmosphere. It is reasonable to consider that the denitrification process is time-dependent. On the other hand, reaction (6) might be controlled by surface reaction or diffusion process of gaseous reactants or products, and the weight loss caused by this process should be limited because of the thermodynamic barrier. However, from a whole view of the weight loss behavior, if we treat the weight loss (2%) below 1000 °C to be caused by the dewatering process, the weight loss up to 1800 °C through other above-mentioned mechanisms is only 1.5%. Such a value of weight loss should not be a serious problem during sintering. The linear shrinkage of the 9A2B–AlN and 2AB–AlN is much different and the effect of atmosphere on it is negligible as shown in Fig. 3(b). For 9A2B–AlN system, the compact shrunk from about 1100 °C and a 9% of linear shrinkage demonstrated at 1800 °C. On the other hand, however, for 2AB–AlN system, the compact expanded from about 1000 °C and reached the largest value of expansion (2%) at 1400 °C, then shrank slowly. However, the compact dimension after soaking at 1800 °C still remained about 1% expansion. The shrinkage results should be related with the following three processes, i.e. (1) decomposition of 2AB at 1035 °C for specimen 2AB–AlN. This process may have resulted in the expansion around 1000–1200 °C, (2) phase formation process by the in situ reactions. This process occurs around 1300–1500 °C according to the XRD results (discussed later), (3) densification (sintering) process. Fig. 4 shows the XRD results of the powder compacts soaked at different temperatures for 60 min. For 9A2B–AlN system (Fig. 4(a)), the peaks for Al2 O3 appeared from 1200 °C, indicating the reaction started at this temperature. However, due to the low BN (the calculated content was 13.38 vol%), and the nature of any BN formed at this low temperature was in turbostratic type with low crystallinity, the diffraction peaks for t-BN could not be identified. From about 1500 °C up the peak at 2h ¼ 26:7° for t-BN became identified. The reaction completed at around 1500 °C. However, there was a small amount of AlN remained in the specimen. The reason may be related to the above discussed weight loss mechanisms from that a small part of reactant B2 O3 consumed. When the sample was soaked at 1800 °C, a new phase of aluminum oxynitride Al5 O6 N appeared. Part of this Al5 O6 N
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In fact, reaction (7) is the same with reaction (5). This reaction should be intense at temperatures higher than 1800 °C. We even found that when the specimen was soaked at 1900 °C, the specimen was almost decomposed and there was a great weight loss. This phenomenon indicated that except for reaction (7) other reactions with large weight loss should have occurred. A possible reaction is 3Al2 O3 þ 4BN ¼ 2B2 O3 ðgÞ þ 3Al2 OðgÞ þ 2N2 ðgÞ
Fig. 4. XRD patterns of the powder compacts soaked at different temperatures for 60 min in N2 atmosphere; (a) 9A2B–AlN system; (b) 2AB–AlN system.
was from the reaction of Al2 O3 and the residual AlN. Another possible way is 5Al2 O3 þ 2BN ¼ 2Al5 O6 N þ B2 O3 ðgÞ
ð7Þ
ð8Þ
The above discussion implies that the synthesis temperature of Al2 O3 –BN composites should be around 1700 °C (higher than 1500 °C but lower than 1800 °C). Also, we can conclude that the application temperature of the composites should be lower than 1800 °C even for short period of time. For 2AB–AlN system (Fig. 4(b)), the XRD results demonstrated three differences compared to those for 9A2B–AlN system. (1) 2AB decomposed to 9A2B + B2 O3 at near 1000 °C. (2) Peaks for Al2 O3 phase appeared from 1400 °C that was 200 °C higher than in the 9A2B–AlN system. (3) Peaks for 9A2B phase disappeared at about 1400 °C that was about 100 °C higher than in the 9A2B–AlN system. From Figs. 2 and 3, we can find that there was no exceptional weight loss in the decomposition process of 2AB. It may means that the formed B2 O3 was mainly in liquid state or even if it was in gaseous state at this temperature it was absorbed (‘‘caught’’) tightly by AlN or 9A2B (on the surface of particles). This phenomenon gives us chance to synthesize Al2 O3 –BN composites based on the proposed reaction (2). The existence of B2 O3 coated on the particle surface of 9A2B delivered a boron oxide environment to prevent 9A2B from reacting with AlN (see reactions (1), (9) and (10)) at lower temperature, which moved the appearance of Al2 O3 phase to a higher temperature. The microstructures of the specimens soaked at 1700 °C are shown in Fig. 5. From the particle shapes we notice that there are two types of particles, one is polyhedral with smooth surface and the other is quasispherical with granular rough surface. EDX results pointed out that the polyhedral particles were composed of Al and O (indicating the Al2 O3 phase) and the quasispherical particles were composed of B, N, Al and O (indicating the existence of BN and Al2 O3 ). It was usually easy to find fractured quasi-spherical particles on the fracture surface of the specimens, as shown in Fig. 5(c). From this picture we can realize that the quasispherical particles are actually agglomerates made of many fine (nanosized) particles. Based on the XRD and EDX results, these agglomerates were composed of t-BN and Al2 O3 . By the way, according to the discussion in Section 4.3.3, the structure of these agglomerates might be composed of an Al2 O3 core surrounded by BN.
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4:5ð2Al2 O3 B2 O3 Þ ¼ 9Al2 O3 2B2 O3 þ 2:5B2 O3 ðl; gÞ ð9Þ and the other is in the newly formed 9A2B. The former should mainly be in liquid state at the decomposition temperature of 2AB and would preferentially react with AlN at high temperature B2 O3 ðl; gÞ þ 2AlN ¼ Al2 O3 þ 2BN
ð10Þ
Subsequently the reaction between the newly formed 9Al2 O3 2B2 O3 and AlN occurred at higher temperature. According to the SEM observation of the obtained specimens, we could not find obvious difference in microstructures except for the content of the formed t-BNcontaining agglomerates. This finding may indicate that the synthesis of t-BN is from the same mechanism, possibly from gaseous route. Furthermore, it was suggested that in both reaction systems the polyhedral Al2 O3 particles were the remnant of the 9A2B particles after the removal of B2 O3 . The shapes of these Al2 O3 particles partly maintained the shape features of the starting powders of 9A2B and 2AB. On the other hand, the agglomerates of nanosized t-BN and Al2 O3 were from the products of the reaction between AlN and B2 O3 , as shown in reaction (10). In addition, from Fig. 3(b) and the SEM observation, the pressureless synthesized Al2 O3 –BN composites are porous materials especially for 2AB–AlN system. These materials might be directly used as catalyst supports in their porous state [22,23] or be used as thermal shock resistant parts in their densified state after hot pressing [24]. On the other hand, from the shrinkage results in the static reaction process we can find that dense composites can not be manufactured by simply normal sintering of the mixed powders. 4.3. Reactive hot pressing
Fig. 5. FE-SEM micrographs of the Al2 O3 –BN composites synthesized at 1700 °C. The arrows indicate the agglomerates composed of nanosized t-BN and Al2 O3 : (a) Al2 O3 –BN composite from 9A2B–AlN; (b) Al2 O3 –BN composite from 2AB–AlN; (c) higher magnification of (a), showing the structure of agglomerate composed of nanosized t-BN and Al2 O3 .
For 9A2B–AlN system, the decomposition temperature of 9A2B is higher than the reaction starting temperature of reaction (1) (see Fig. 4(a)), so the reaction sequence should be simply the original reaction (1). On the other hand, for 2AB–AlN system, however, there are two parts of B2 O3 source, one is from the decomposition of 2AB at 1035 °C
4.3.1. X-ray diffraction The XRD patterns of the specimens hot pressed at different temperatures are shown in Fig. 6. For 9A2B– AlN system (Fig. 6(a)), peaks for Al2 O3 phase appeared from 1200 °C, which was the same with that in the pressureless sintered specimen. On the other hand, however, peaks for 9A2B disappeared at 1400 °C, 100 °C lower than that in the pressureless sintering specimen. This indicates that the applied pressure possibly promoted the reaction process by improving the diffusion conditions, for example, by decreasing diffusion distance and providing driving force for material transport. Moreover, from the heightened peak intensity for turbostratic BN (t-BN) phase we can see that the applied pressure also accelerated the formation and crystallization process of BN. If the hot pressing temperature is increased to 1800 °C, similar to the case of pressureless sintering, Al5 O6 N phase appeared from the
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seems to be relative to the percolation threshold of BN in a composite. When the BN flakes form a percolation network, the densification behavior will be obviously worsened due to the impingement occurrence (formation of card-house structure) of grain-growing BN flakes with different basal plane orientation and thus production of pores [9,11]. Fig. 7 shows the densities of the RHPed specimens as a function of the hot pressing temperature. For Al2 O3 – 13BN composite, very high relative density was obtained when the RHP temperature was higher than 1600 °C. On the other hand, however, for Al2 O3 –22BN composite only low-density materials were obtained. This result should implicate that the BN content in the Al2 O3 –13BN composite is under, but that in the Al2 O3 – 22BN composite is above the percolation limit of BN phase in the Al2 O3 matrix. In other words, we can know by inference that continuous network of BN phase has not yet formed in the Al2 O3 –13BN composite, but has formed in the Al2 O3 –22BN composite. This inference will be verified by the microstructure observation in the next section of this article. Here although we will not discuss how to get dense composite with high BN content, to control the grain growth of BN flakes, to destroy the already formed continuous BN network and to increase the slip ability of BN particles in microstructurial level should be the critical points. Two methods might be effective, one is to control the grain growth of h-BN flakes by synthesis conditions including temperature and pressure programs, and the other is to add liquid-phase-forming additives to help the flow and rotation of BN particles at hot pressing temperature like as shown in our previous study on SiC-BN in situ composite [25]. Abreal et al. [26] reported that Y2 O3 addition improved the densification Fig. 6. XRD patterns of the specimens hot pressed at different temperatures: (a) 9A2B–AlN system; (b) 2AB–AlN system.
110 Al2O3-13BN
100 Relative density (TD%)
reaction (7). The formation of gasous B2 O3 by this reaction would be detrimental to the material properties. For 2AB–AlN system, similar to the 9A2B–AlN system, a tendency of accelerated reaction process by hot press can be seen from Fig. 6(b). As discussed above, the coexistence of B2 O3 from the decomposition of 2AB inhibited the reaction between 9A2B and AlN. Namely, in phenomenal the B2 O3 phase stabilized the 9A2B phase to a higher temperature. As will be discussed later, this reaction mechanism will produce great influence on the microstructure development process.
97.29
98.44 99.11
94.38 89.25
90
82.78 78.82
80
75.72
70
77.65
72.31
Al2O3-22BN
59.02
60
1900
1800
1700
1600
1500
1400
1300
1200
55.18
50
1100
4.3.2. Densification h-BN-containing composites usually show poor densification behavior even when being manufactured by hot pressing when the BN content exceeds a limit of the range of 10–20 vol% that is determined by the particle size, shape and distribution of h-BN particles. This limit
Hot press temperature (˚C) Fig. 7. Relative density of the obtained in situ composites.
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of Al2 O3 –BN composites. In addition, it should also be effective by simply increasing the applied pressure of hot pressing. 4.3.3. TEM observation In general the microstructures of in situ synthesized ceramic composites demonstrate some features that are very different from those of conventionally processed ones. These features include special phase distribution, particle size and shape, and unique grain boundary. People can directly produce nanocomposites, whisker/ platelet-reinforced ceramic composites and interpenetrating phase composites through in situ synthesizing processes. In addition, pore structure in the in situ formed materials may also be different from that of the conventionally processed ones. For instance, closed pores may dominate the pore type in in situ composite [12]. The above-mentioned features in microstructural level endow the corresponding materials with some novel properties, such as high strength, toughness, creep resistance, wear resistance or corrosion resistance both at room and high temperatures, and in addition, good thermal shock resistance. The TEM micrographs of the obtained composites are shown in Fig. 8. The hot press direction is vertical in these pictures. In the Al2 O3 –13BN composites hot pressed at both 1600 and 1700 °C, there was a little orientation preference for Al2 O3 phase, which showed a rectangular particle shape with the longer side vertical to the hot press direction. This preferential growth of Al2 O3 particles might resulted from the considerably high temperature hot pressing for Al2 O3 phase, and the BN phase with a low fraction in this composite has not inhibited this growth process. The reaction-formed BN phase was in round shape of agglomerates located at the ‘‘pockets’’ of Al2 O3 particles. For Al2 O3 –13BN composite RHPed at 1700 °C, the reaction, densification and the grain growth of Al2 O3 were possibly too fast at the hot press temperature, some particles of BN phase were distributed at the grain boundaries or inside the grains of Al2 O3 particles. Those BN particles located at the grain boundaries might degrade the material strength. For the Al2 O3 –13BN composite hot pressed at 1600 °C, however, all the BN round agglomerates were perfectly located at those pockets surrounded by Al2 O3 particles. It means that the distribution of the weak phase BN is not continuous in the microstructure. Such a microstructure should have less detrimental effect on the macro-properties (for example, strength) of the material. A good performance in wear resistance may also be expected with this material. For a composite with a weak phase like BN as the dispersoid at a given volume fraction, it is suggested that the round agglomerates of the weak phase with a suitable size might be beneficial to result in the smallest possible effect on degrading the material strength. In
Fig. 8. TEM micrographs showing the agglomerate shape and distribution of the in situ formed BN phase in Al2 O3 matrix: (a) Al2 O3 – 13BN composite RHPed at 1600 °C; (b) Al2 O3 –13BN composite RHPed at 1700 °C; (c) Al2 O3 –22BN composite RHPed at 1700 °C.
other words, a suitable agglomeration and formation of round agglomerates will increase the percolation threshold of the weak phase, and result in an isolated distribution of and an increased free distance between the agglomerates. However, to realize this structure in a large microstructural scale there needs to be careful control of the size, distribution and volume fraction of the weak phase (or pore) in matrix material. The present process of in situ synthesis successfully provides
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a way to manufacture material with such kind of microstructure. Fig. 9 demonstrates a typical agglomerate of BN phase located at an alumina pocket in Al2 O3 –13BN composite RHPed at 1600 °C. The shape feature of this BN agglomerate was very similar to that observed in the pressureless sintered specimen (see Fig. 5(c)). This micrograph seems to reveal the picture of formation process of BN phase from the reactions. The diffusion and BN growth at the radial direction dominated the reaction process. However, there was no predominant orientation relationship between Al2 O3 and BN phases as it can be seen from Fig. 10, which shows an HRTEM micrograph of a typical grain boundary between Al2 O3 and BN phases. Similar microstructural feature was observed in Al2 O3 –22BN composite. We can also find that the interface is very clean with no amorphous phase, indicating a strong interface bonding in this in situ composite. It should be noticed that the round BN agglomerates themselves were fully dense, which was much different with the pore-containing card-house structures in conventionally processed composite. For Al2 O3 –22BN composite (Fig. 8(c)), the reaction formed BN phase showed a very similar round ag-
Fig. 9. A typical agglomerate of BN phase located at the alumina pocket in Al2 O3 –13BN composite RHPed at 1600 °C.
Fig. 10. HRTEM micrograph of an Al2 O3 /BN interface observed in Al2 O3 –13BN composite RHPed at 1600 °C, showing no predominant orientation relationship between Al2 O3 and BN phase.
glomerate shape and distribution to that in the Al2 O3 – 13BN composite except for the different BN fraction. However, this composite remained a high level of porosity and the particle size for Al2 O3 phase was small compared to that in Al2 O3 –13BN composite. In addition, at the microstructure level, although in some areas the BN phase, in round agglomerate shape, was still isolated, a BN phase network has already formed in most area. The small particle size of Al2 O3 might be resulted from the grain growth constraint by this network. On the other hand, however, another strong network formed by Al2 O3 matrix grains also existed in this composite. According to the microstructure features observed by TEM and the XRD results, a microstructure development mechanism was proposed, as shown in Fig. 11. In 9A2B–AlN system, boron oxide in 9A2B moves through solid diffusion or vaporization–solidification process to the surface of AlN particle and then reacts with AlN by mutual diffusions of nitrogen from the AlN core to the outside and oxygen from outside to the core. Consequently, we usually find such core–shell microstructure as shown in Fig. 9. In 2AB–AlN system, the B2 O3 formed by the low temperature decomposition of 2AB coated on all the surfaces of particles (Fig. 11(e)). This B2 O3 coated on the simultaneously formed 9A2B particles prevents the diffusion of boron oxide in 9A2B towards outside. Thus the peaks for Al2 O3 phase appeared at temperature higher than that in 9A2B–AlN system. The densification behavior was poor in this system due to the high BN content. However, the next section will show that even for this composite high strength can be obtained. 4.3.4. Mechanical properties The mechanical properties are listed in Table 1. For Al2 O3 –13BN composite, the specimen RHPed at 1600 °C showed the highest strength. The low strength for composite RHPed at 1500 °C mainly resulted from the low relative density (see Fig. 7). However, the relative low strength for composites RHPed at 1700 and 1800 °C should be related to the microstructure feature (Fig. 8(b)) and the reverse reaction (5). As for the small difference in strength for composites RHPed at 1700 and 1800 °C, it might point out that the reverse reaction (5) was not so serious at the present experimental condition for such a composite with low BN fraction, and on the other hand the increased density (although only a little) gave an improvement in strength for composite RHPed at 1800 °C. For Al2 O3 –22BN composite, the highest strength was obtained when it was RHPed at 1700 °C. From this result we can find that although the density for composite RHPed at 1800 °C was higher, the reverse reaction (5) should have taken a serious role in degrading material strength. The detrimental effect of the reverse
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9A2B 9A2B 9A2B
AlN
Al2O3
B 2O 3
Al2O3
Al2O3
Al2O3
9A2B Al2O3
9A2B 9A2B
Al2O3
Al2O3
Al2O3
Al2O3 and BN
Al2O3
Al2O3
Al2O3
(b)
(a)
1833
BN
Al2O3
Al2O3
(c) B2O3
2AB 2AB
9A2B
AlN 2AB AlN 2AB
9A2B Al2O3
AlN 9A2B AlN
2AB
9A2B
(d)
Al2O3
9A2B
BN
Al2O3
Al2O3 Al2O3
Al2O3Al O Al O 2 3 2 3
(e)
(f)
Fig. 11. Schematic mechanisms for the microstructure development of the in situ Al2 O3 –BN composites: (a)–(c) is for the 9A2B–AlN system; (d)–(f) for the 2AB–AlN system.
Table 1 Mechanical properties of the RHPed Al2 O3 –BN composites Property °Ca °C °C °C
Bending strength (MPa)
YoungÕs modulus (GPa)
Strain to failure (103 )
PoissonÕs ratio
604 74 806 66 694 75 721 61
246 266 260 257
2.46 3.03 2.67 2.81
0.206 0.211 0.206 0.198
Al2 O3 –13BN composite
1500 1600 1700 1800
Al2 O3 –22BN composite
1600 °C 1700 °C 1800 °C
171 6 275 9 254 12
80 109 111
2.14 2.52 2.29
0.219 0.196 0.175
Al2 O3 monolithic
1350 °C
757 45
411
1.84
0.237
a
Hot press temperature.
reaction (5) on strength should become large with the increase of BN content in the composite. Compared with the monolithic Al2 O3 ceramics, the in situ synthesized composites exhibited high strength, low YoungÕs modulus and high strain to failure. In addition, PoissonÕs ratios of the composites were always lower than that of the Al2 O3 monolithic ceramics. As already mentioned above, this high level strength may come from the remained strong network of the Al2 O3 matrix and the distribution feature of the in situ formed BN phase. The weak BN phase in round agglomerate shape and located at the pockets of Al2 O3 particles should have resulted in a small effect on degrading the material strength compared to a BN distribution at the matrix grain boundaries. In the material with this microstructure, the BN agglomerates and closed pores may take a similar effect on the material properties. However, compared with the pores, the BN phase would produce a composite with low friction coefficient, high wear resistance and corrosion resistance. On the other hand, the in situ process gives an opportunity to obtain highly
strong interfacial bonding between both grain boundaries of Al2 O3 /Al2 O3 and Al2 O3 /BN. Fig. 12 reveals the fracture surfaces of the Al2 O3 – 13BN composite RHPed at 1600 °C and the Al2 O3 – 22BN composite RHPed at 1700 °C observed by SEM. The hot press direction is vertical in these pictures. We can find that the dominant fracture mode for both composites was transgranular, indicating a strong bonding at the grain boundaries. Fig. 13 demonstrates an HRTEM picture of a typical grain boundary of Al2 O3 /Al2 O3 grains in Al2 O3 –22BN composite RHPed at 1700 °C. It points out the clean and strongly bonded Al2 O3 grains. Especially, for Al2 O3 –22BN composite, although the remnant porosity was 22.35%, that is, the total volume fraction of the weak BN phase and the porosity combined reached about 45%, a strong Al2 O3 network still existed. A strength of 275 MPa is also very high for Al2 O3 ceramics with a porosity of 45% [27]. Of course, the effect of BN phase and pore on material properties should be different. On the other hand, however, it may imply that by in situ process of
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2.
3.
4.
5.
Fig. 12. Fracture surfaces of the composites observed by SEM indicating the transgranular fracture mode: (a) Al2 O3 –13BN composite RHPed at 1600 °C; (b) Al2 O3 –22BN composite RHPed at 1700 °C.
6.
7.
8.
Fig. 13. HRTEM picture of a typical Al2 O3 /Al2 O3 grain boundary in Al2 O3 –22BN composite RHPed at 1700 °C, pointing out the clean and strongly bonded Al2 O3 grains.
synthesizing BN phase in ceramic matrix, high performance porous ceramic composites can be prepared. 9. 5. Conclusions A new chemical approach using aluminum borates 9Al2 O3 2B2 O3 and 2Al2 O3 B2 O3 as boron source was proposed to synthesize Al2 O3 –BN composites in the present study. The following conclusions were obtained. 1. Thermodynamic calculations revealed that the reactions in the 9A2B–AlN and 2AB–AlN systems were
exothermic, but there were no SHS processes occurred at standard conditions. TG results showed that the weight loss in nitrogen atmosphere was a little larger than that in argon atmosphere, possibly resulting from the reaction between B2 O3 and N2 . However, this reaction would not lead to obvious detrimental effect on the composition and properties of the synthesized composites. Due to the reverse reaction between the already formed Al2 O3 and BN, synthesis and subsequent applications of the composites at temperatures higher than 1800 °C should be avoided. The B2 O3 from the decomposition of 2AB at about 1000 °C coated on the surfaces of the AlN particles and the simultaneously formed 9A2B particles, resulting in the appearance of Al2 O3 phase at higher temperature through inhibiting the direct reaction between 9A2B and AlN. SEM observations of the statically reacted specimens demonstrated that the syntheiszed BN phase formed into quasi-spherical agglomerates composing of very fine (nanosized) t-BN and Al2 O3 particles. The polyhedral Al2 O3 particles kept the shape feature of the starting powders of 9A2B and 2AB, but the quasispherical agglomerates were formed by the products of the reaction between AlN and B2 O3 . The applied pressure of hot pressing promoted the reactions between aluminum borates and aluminum nitride, possibly due to improved diffusion processes. Al2 O3 –13BN composite from 9A2B–AlN system with high density of 97–98 TD% was obtained by RHP. On the other hand, however, for Al2 O3 –22BN composite from 2AB–AlN system, density higher than 80 TD% could not be reached under the present RHP conditions owing to the formation of BN networks. TEM observations of the RHPed specimens showed that the in situ formed BN phase was in round agglomerate shape located at the pockets of Al2 O3 matrix particles. It was suggested that such distribution of BN phase was beneficial to the strength of such materials with weak phase dispersoids. Fracture surface observations by SEM revealed that the main fracture mode was transgranular for both Al2 O3 –13BN and Al2 O3 –22BN composites, pointing out the existence of a strong network of the Al2 O3 matrix. The obtained composites showed high strength, low YoungÕs modulus and improved strain tolerance.
Acknowledgements This work has been supported by AIST, METI, Japan, as part of the Synergy Ceramics Project. Part of the work has been supported by NEDO. The authors
G.J. Zhang et al. / Acta Materialia 52 (2004) 1823–1835
are members of the Joint Research Consortium of Synergy Ceramics.
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