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The influence of oxygen impurities on the formation of AlN–Al composites by infiltration of molten Al–Mg
Materials Science and Engineering A337 (2002) 134 /139 www.elsevier.com/locate/msea
The influence of oxygen impurities on the formation of AlN Al composites by infiltration of molten Al Mg /
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S. Swaminathan, B. Srinivasa Rao *,1, V. Jayaram Department of Metallurgy, Indian Institute of Science, Bangalore 560012, India Received 14 August 2001; received in revised form 17 December 2001
Abstract The influence of residual oxygen in nitrogen on the formation of AlN /Al matrix by reactive infiltration has been investigated. Increasing the oxygen content from 10 ppm upwards decreased the nitride content in the matrix from 64 to 6%. Based on the analysis of the availability of oxygen at the Al-melt/gas interface, three distinct scenarios have been proposed: (i) at lowest values, oxygen does not interfere with either infiltration or nitridation reaction; (ii) at intermediate values, nitridation is suppressed, however infiltration continues; and (iii) at a critical upper value, the melt passivates without any infiltration. This phenomenon offers control of the AlN/Al ratio in the matrix and the possibility of creation of microstructural gradients by the appropriate choice of gas mixtures. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Infiltration; Residual oxygen; Nitridation; AlN
1. Introduction The formation of ceramic particulate composites with AlN /Al matrices with interesting mechanical and thermal applications is made possible by the nitridation and simultaneous infiltration of molten Al alloys into a preform. It is known that during the process of nitridation of Al /Mg alloys the degree of conversion of Al to AlN is a strong function of temperature and alloy composition. Two different types of experiments have been conducted. In the first, alloys have been exposed to nitrogen containing atmospheres in the absence of a preform. The increase in AlN content at 1000 8C with Mg in the range 1/10% and with the Mg:Si ratio has been determined [1]. In addition, it was pointed out that the degree of conversion also depended on oxygen impurities in static atmosphere. For Al / 2Mg /0.8Si, the fractional conversion to AlN diminished * Corresponding author. Present address: Technische Universita¨t Darmstadt, FB Materialwissenschaft, FG NichtmetallischAnorganischewerkstoffe, Petersen Straße 23, 64287 Darmstadt, Germany. Tel.: 49-6151-16-6318; fax: 49-6151-16-6314 E-mail address: [email protected] (B. Srinivasa Rao). 1 Presently a Humboldt Fellow.
from 95% in N2 /0.7% H2 (0.1 ppm oxygen), to 16% with pure nitrogen (3 ppm oxygen), while oxygen contents of /1% brought the level of nitridation to :/3% and lower [1]. In contrast, the practically more useful situation is when the alloy is simultaneously infiltrating into a ceramic preform, such as alumina [2 /6]. Here, it has been shown that Mg plays a key role in reducing the oxygen partial pressure to enable the alloy to infiltrate by capillary action. While alloyed Mg content is important, in this regard, it was also shown [5,6] that gettering the incoming nitrogen with a molten Al /Mg alloy was even more effective in ensuring that passivation of the infiltration did not occur. This phenomenon applies in the low temperature regime (850 /950 8C) in which the extent of nitridation is small and equally at higher temperatures at which AlN becomes a significant phase. At the same time, it has been shown that the transition temperature from metalrich to ceramic-rich composites increases slightly when commercial nitrogen is used instead of a nitrogen / hydrogen mixture. The purpose of the present paper is to report a much more substantial variation in nitride content that may be achieved during infiltration by introducing oxygen impurities and to examine the possible reasons for this sensitive behaviour.
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 0 0 2 - 3
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2. Experimental procedure Fused alumina particulate of 53 mm size were used as the dispersion into which infiltration was carried out. Preforms of the dimensions 80 /40 /5 mm were made by cold pressing in a steel die by applying a pressure of :/10 tons. These cold compacts were heated at 3 8C min 1 to 800 8C to avoid cracking of the preform during binder removal. After 800 8C the preforms were heated to 1600 8C and held for 1 h in order to obtain sufficient mechanical strength for handling purposes. Commercial purity aluminium and magnesium were used to prepare Al /2 wt.% Mg and Al /5 wt.% Mg alloys. Aluminium was melted in air in a clay graphite crucible at a temperature of :/750 8C and the melt was degassed by using hexachloroethane before and after magnesium was added to the melt. Magnesium was wrapped in aluminium foil in order to minimise evaporative losses. An excess of 10% (over and above the desired value) of magnesium was added to compensate for any residual evaporation. The alloys were cast into plates and homogenised at a temperature of 400 8C for 24 h and cut to the final size for infiltration. Alloy composition was determined using atomic absorption spectroscopy and inductively coupled plasma spectroscopy (ICP). Mg contents in these alloys were found to be 1.95 and 5.01% for the Al /2Mg and Al /5Mg, respectively. Alumina trays of dimensions 84 /42 /30 mm were used as crucibles. The insides of the crucibles were lined with graphite foil in order to facilitate easy removal of the composite and residual metal after the experiment. Infiltration experiments were carried out in a high temperature tube furnace (Barnstead Thermolyne). The dimensions of the composites fabricated were 80 /40 / 5 mm. Alumina preforms were kept on the top of the Al /2Mg alloy billets in the alumina crucible. The incoming gas was passed over an alloy of Al /5Mg that acted as an oxygen getter before the gas reached the melt-preform assembly. The getter was located in the 900 /950 8C temperature zone of the same tube. The effectiveness of the getter could be gauged by the formation of MgO whiskers on the wall of the furnace tube, downstream of the getter. However, no measurements were made of the oxygen partial pressure. The tube was evacuated to a pressure of 0.05 mbar and then back-filled with N2 /2% H2 (10 ppm oxygen) or commercial nitrogen (5000 ppm oxygen) before the start of experiment. Gas was passed at a rate of 250 ml min 1 throughout the experiment at a pressure slightly above one atmosphere and bubbled out through a column of oil before being vented to the atmosphere. The furnace was heated at a rate of 4 8C min 1 to 975 8C and held there for 3 h to ensure complete infiltration of the preform.
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After infiltration the crucibles were cut longitudinally with a high speed precision saw using a diamondwafering blade for optical metallography. The sectioned samples were polished using standard metallographic procedures to 1 mm diamond finish. The polished samples were examined under an optical microscope to determine the distribution of various phases in the composite. For scanning electron microscopy, the polished samples were etched with concentrated Keller’s reagent in order to remove Al and the samples were coated with gold in order to make the surface electrically conducting. The matrix hardness was measured by a SHIMADZU HMV-2000 tester at test loads of 100 g and a hold period of 10 s. Care was taken to ensure that the indentation fell within the matrix and that there was no interference with adjacent particles. An average of ten to 15 readings was taken as a measure of the representative hardness of the matrix. The volume fraction of AlN present in the composite was determined by decomposing it with sodium hydroxide solution to evolve ammonia which is quantitatively analysed by titrating against a strong acid.
3. Results and discussion The microstructures of the composites fabricated at 975 8C are shown in Figs. 1 and 2. As shown in Table 1, the first three runs, all carried out in N2 /2% H2, indicated wide scatter in the AlN content which was, however, consistent with the microstructures (Fig. 1) and hardnesses. It was also noted that the microstructures were uniform throughout the composite and that the infiltration was therefore reproducible across the section and along the length. It transpired that in the first two experiments in which the AlN content was as low as 6%, that the O-rings that sealed the tube furnace had given way during previous experiments. Degradation of the seals was evident when attempts were made, unsuccessfully, to fully evacuate the system after these two runs. During subsequent runs, the seals were replaced more frequently and, accordingly, the AlN content obtained increased to 64.5% (No. 3 in Table 1). The experiment was repeated for reproducibility at the same temperature with a new seal and a nitride content of 62 vol.% was obtained. Consequently, this value is expected to be representative of the maximum achievable at that temperature with N2 /2% H2. The AlN /Al matrix in the above cases is divided into two parts: a portion adjacent to the particle surfaces that is rich in AlN (Al/AlN in Fig. 1B,C) and the inter-particulate region that contains metal pockets as shown in Fig. 1(B) and Fig. 2(B). The relative microstructural scales may be seen more clearly in the etched scanning electron micrographs in Fig. 2. The hardness of the metal
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Fig. 1. Microstructures of the samples processed at 975 8C depicting the influence of different amounts of oxygen impurities on the amount of AlN formed in the matrix. (A) 6% AlN; AlN is mostly dispersed as isolated particles in Al-alloy matrix as indicated with an arrow. (B) 34% AlN; formation of AlN around the Al2O3 particles and starts extending into the inter-particle pores is the dominant feature at this volume fraction. However, large metal channels can still be seen (marked with M; white in colour). (C) 64.5% AlN; almost all of the inter-particle regions are covered with AlN (dark in colour) interspersed with fine Al-alloy channels (white in colour).
pockets changes little, suggesting that its nitride content is minimal. The mechanisms of alloy infiltration and nitridation are, at this time, not fully understood. However, certain physical features of the process have been established [7,8]. The melt tends to preferentially climb along alumina particle surfaces and commences nitriding outwards from the surface along an irregular front into the inter-particle pores. It is possible for residual porosity in the portions infiltrated earlier (far below the maximum extent of melt penetration), to then become isolated from the nitrogen diffusing in from the outside. These regions then fill up with advancing alloy without reacting significantly with nitrogen and lead to the large metal-rich pockets that are seen in Figs. 1 and 2. Thus, the appearance of the microstructure is completely consistent with experiments [7,8] that are carried out as a function of temperature (after ensuring the absence of air leaks) in the range 950/1100 8C in which the total amount of AlN increases, both by the expansion of the composite Al/AlN zone adjacent to the particulate at the expense of the alloy pockets as well as by an increase in the fraction of AlN within these regions. The amount of oxygen leaked into the system in the present instance is not known. However, to confirm its role, an experiment was conducted at 975 8C with commercial purity nitrogen, which has a residual oxygen content of :/ 5000 ppm and which yielded a composite with a nitride content of 29 vol.% (Fig. 1) that is close to the result of one of the earlier runs in N2 /2% H2, which is known to have been performed with a leaking seal (Table 1, No. 2). Thus, the behaviour with decreasing oxygen is similar to that displayed with increasing temperature, i.e. an increase in the nitride content of the composite. In order to understand the role of oxygen, the suggested mechanisms of nitrogen uptake by the melt are now briefly reviewed. It has been suggested [9] that nitridation proceeds by a two-step reaction, beginning with the nitridation of Mg in the vapour phase to Mg3N2, followed by reduction to AlN by the advancing melt. This sequence is similar to that which is now accepted for directed melt oxidation, in which MgO or MgAl2O4 plays a similar intermediate role. However, the presence of Mg3N2 at the composite surface has not yet been established. It has been shown [10] that if the oxides are suppressed by adequate gettering, that AlN is more stable than Mg3N2 in the alloys of interest. However, this balance could be overturned if the Mg activity was substantially increased in the vapour phase, thereby making it more available for reaction with the incoming gas. At this point in time, this issue is unresolved, i.e. does nitrogen combine directly with aluminium in the melt or via an intermediate compound. What has been believed, however, is that residual oxygen levels need to be sufficiently lowered so that oxides (MgO or MgAl2O4) do not form and terminate
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Fig. 2. SEM photographs after etching of the samples. (A) 6% AlN and (B) 64.5% AlN in the matrix. AlN formed around the Al2O3 particles, fine metal channels and a metal pocket (marked with P) that has been etched away can be seen clearly in the case of 64.5% AlN sample.
the infiltration. (The system may then choose to grow an oxide/metal composite, as in directed melt oxidation, [11] if the temperature and oxygen partial pressures are sufficiently high.) The results presented here show that there is an additional degree of complexity in the infiltration. Despite the reduced extent of nitridation Table 1 Hardness of and AlN content in the matrix S. No. Gas
AlN (vol.%)
Hardness (GPa) Nitride region
a
1. 2.a 3. 4. 5. a b
N2 /2%H2 N2 /2%H2 N2 /2%H2 Repeat of 3 N2b
6 34 64.5 62.0 29.0
/ 3.890.6 6.491.1
In the presence of oxygen impurities. Contains 0.5% oxygen.
Metal pockets 0.6890.06 0.790.07 0.890.2
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in some of the runs, the oxygen levels are sufficiently low that infiltration is still able to occur. Thus, one of the key presumptions in our prior understanding of the pressureless infiltration of Al-alloys, i.e. the need to maintain extremely low oxygen partial pressures, needs to be modified. Now one has to interpret these results in a framework that allows minute amounts of oxygen to permit capillary advance of the metal at rates that are not significantly altered, but effectively poisons the nitridation reaction, even though, for all practical purposes, the partial pressure of nitrogen is one atmosphere and that of oxygen is several orders of magnitude lower. We now examine the issues relating to the thermodynamics and rate of oxygen transport in this system. It has been pointed out [1] that the oxygen partial pressures under which AlN is stabilised in preference to Al2O3 are extremely low and in the vicinity of 10 20 atmosphere. At these pressures, the number of oxygen molecules available within the entire preform would be only / unity. This value becomes even lower if one assumes that Mg vapour from the getter is in equilibrium with MgO and oxygen. Similar or lower numbers are obtained when examining the competition between other equilibria, e.g. those involving MgO and Mg3N2. Thus, it is likely that thermodynamic considerations do not play as important a role as the availability of the appropriate gas molecules to enable a particular equilibrium to be achieved. In particular, it is likely that while experiments carried out in closed systems might conceivably approach such equilibria, the situation in flowing gas, which is the practically relevant one for the fabrication of matrices that contain AlN, is certain to be dominated by kinetic availability. To analyse the problem of oxygen arrival at the melt-atmosphere interface, it is assumed that the flux is determined by the partial pressure difference between the interface and the outside of the preform. In addition, the ratio of the oxygen flux to the rate at which the melt advances through the preform determines the net uptake of oxygen by the alloy. We assume that whether the surface is passivated or not can be defined with respect to two limits of oxygen content. The lower limit is one at which the bulk level of oxygen in the melt, enhanced by surface segregation, can lead to strong binding between Al and O atoms, thereby preventing any nitrogen uptake. Beyond the upper limit, precipitation of oxides takes place. Between the lower and upper limits, infiltration is possible, but with a gradual decrease in the extent of nitridation, i.e. the composite becomes progressively metal rich as the oxygen content rises. Let Jg be the flux of oxygen arriving at the surface of melt, D Dc /Dx , where Dc is the concentration difference between furnace atmosphere and melt surface, D is the gas phase diffusivity and Dx is the preform thickness. Assume that melt surface is at 0, i.e. everything that
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arrives there is instantaneously adsorbed. Then Dc / cfurnace /P /RT, where P is the partial pressure of oxygen in furnace. Let us suppose that poisoning occurs when the average O:Al ratio exceeds f. For this to happen, before the interface moves up by one monolayer, the number of oxygen atoms arriving per unit area must exceed f or fN moles oxygen per area, where N is the number of moles of aluminium per unit area. Thus, the limiting flux is given by: Jt /fN , where t is the time taken for melt to move up a layer /y /vm, where y is the height of one monolayer and vm is the speed of infiltration. Therefore, the limiting value of pressure is given by: DP=(RTDx) (vm fN)=y or P (fNRTvm Dx)=(Dy) At 1248 K, Dx is the preform thickness :/0.01 m and D /10 4 m2 s 1 [12], with oxygen ions of 0.28 nm diameter and an experimentally observed infiltration rate of :/2 mm h 1: y 2:8 10 10 m vm 0:002=3600 5:6 10 7 m s 1 Let oxygen adsorb as a close-packed array with a spacing of 0.28 nm. The number of moles of oxygen atoms per unit area is {(2.8)2(0.866)(10 20)(6) (1023)} 1 /2.45 /10 5 and the number of molecules per unit area is N /10 5 per m2. Since the oxygen ion radius (0.14 nm) and aluminium atom radius are the same, the number of moles of each per unit area is the same for normalisation, assuming that Al is close-packed at 1248 K. P f f(10 5 )(10375)(5:6 10 7 )(0:01)g=f(2:8 10 10 ) (10 4 )g : 2 10 4 f Pascals It seems likely that the critical oxygen concentration that can be tolerated before all infiltration ceases is the solubility limit, beyond which spinel or alumina would be precipitated. Thus, the limit for poisoning of nitridation is expected to be lower than this solubility limit. It will be even further lowered, if surface segregation is present. Some evidence is available in the literature to support the argument that the Al melt is partially covered with oxygen before Al2O3 precipitates out. Goumiri and Joud [13] studied the initial stages of oxidation and its effect on the surface tension of Al by auger electron spectroscopy. It was reported that the surface tension of Al decreases strongly at initial oxidation and stays rather constant for coverage up to one monolayer. The decrease in the surface tension of the melt was attributed to the formation of oxide islands up to approximately saturation coverage. There was no
description of the steps involved before Al2O3 precipitates out. Stucki et al. [14] studied the initial oxidation of Al. When a solid Al is exposed to oxygen, there is chemisorption up to 20 Langmuirs and then the oxide peaks start to appear in the auger signal. In contrast, when the liquid is exposed, there was no chemisorption, but the oxygen appears to go into the solution until 1000 Langmuirs and then the oxide signal appears directly. This tells us that liquid Al can dissolve more oxygen than solid Al, however, there was no indication of what happens at the melt surface after the melt is saturated with oxygen and before Al2O3 starts to precipitate. An estimate of the solubility of O in pure Al is 10 6 [15]. Thus, if we use f /10 6, then we get P /0.02 Pa and assuming a total pressure of 1 atm or 105 Pa, the oxygen level is 2 /10 7 Pa or 0.2 ppm. In the absence of more detailed information on solubilities or the actual partial pressures prevailing, it is difficult to be more quantitative. The above estimate is intended to demonstrate that partial pressures much higher than those predicted from thermodynamic equilibrium can be tolerated, given the kinetic constraints when infiltration is proceeding simultaneously. Evidence supporting the inhibiting effect of O2 on AlN growth comes from the literature on the growth of group-III nitride single crystals (AlN, GaN and InN) by dissolution of nitrogen in the melts (Al, Ga and In) under high nitrogen pressure [16]. Quantum mechanical calculations have provided some insight into the mechanisms of nitrogen dissolution into these metals [17]. It was shown that the adsorption of N2 on the surface of Al melt leads to the dissociation of N2 molecule. The energy barrier for this dissociation is around 3.2 eV, which is less than half of the dissociation energy of N2, which is 9.8 eV. In contrast, O2 dissociates and then adsorbs onto the melt surface without any energy barrier. This points to the fact that the growth of defect free AlN or GaN is difficult since any oxygen in the processing environment dissolves in the melt very easily. In the similar way, it is expected that the presence of oxygen impurities during the infiltration of Al /Mg into Al2O3 preform inhibit the growth of AlN mainly by adsorption onto the Al melt surface. Strictly speaking, one has to consider the effect of Mg on the adsorption of O2 onto the surface of Al melt since Al /2% Mg is used for infiltration. At this moment, it is unclear whether Mg helps or hinders the formation of AlN. However, its role in keeping the melt surface free of any passivating oxides, such as alumina or spinel, is well studied [5,6]. From the above analysis, one can then envisage three distinct scenarios of infiltration that are delineated by the arrival rate of oxygen at the interface. At the lowest values, oxygen dissolves in the melt in small amounts and does not interfere with either infiltration or the reaction between Al and N2. At intermediate values,
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local regions of the surface experience a sufficiently high oxygen content (possibly enhanced by surface segregation), such that Al and O atoms preferentially bind to one another. Thus, nitride formation is suppressed but infiltration is not seriously affected, thus leading to a metal-rich composite. At a critical upper value of oxygen content, the surface passivates through the formation of a monolayer of spinel or MgO. Infiltration of metal now ceases entirely and the reaction front has to advance by mechanisms which are closer to those encountered in directed melt oxidation [11,18/20].
4. Conclusions The residual oxygen in the N2 strongly influences the amount of AlN formed during the infiltration of Al /Mg alloys into Al2O3 preforms in N2. The amount of AlN increases with decreasing residual oxygen content. Oxygen appears to inhibit the reaction between Al and N2, mainly due to the relative ease with which it is adsorbed on to Al. The control of ceramic to metal ratio and the creation of microstructural gradients may now be more readily accomplished through the use of appropriate gas mixtures of N2 and O2 compared to the alternative, slower method of cycling the process temperature.
Acknowledgements Financial support for this work has been provided by the Volkswagen Foundation, Germany.
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References [1] H. Scholz, P. Greil, J. Mater. Sci. 26 (1991) 669. [2] M.K. Aghajanian, M.A. Rocazella, J.T. Burke, S.D. Keck, J. Mater. Sci. 26 (1991) 447. [3] D. Creber, S.D. Poste, M.K. Aghajanian, T.D. Claar, Ceram. Eng. Sci. Proc. 9 (1988) 975. [4] N. Nagendra, B.S. Rao, V. Jayaram, Mater. Sci. Eng. A269 (1999) 26. [5] B. Srinivasa Rao, Ph.D. Thesis, December, 1999, Department of Metallurgy, Indian Institute of Science, Bangalore, India. [6] B. Srinivasa Rao, V. Jayaram, Acta Mater. 49 (13) (2001) 2373. [7] S. Swaminathan, B. Srinivasa Rao, V. Jayaram, Acta Mater, in review. [8] S. Swaminathan, M.Sc. (Eng.) Thesis, November, 2000, Department of Metallurgy, Indian Institute of Science, Bangalore, India. [9] G.H. Schiroky, D.V. Miller, M.K. Aghajanian, A.S. Fareed, Key Eng. Mater. 127 /131 (1997) 141. [10] B. Srinivasa Rao, V. Jayaram, J. Mater. Res. 16 (10) (2001) 2906. [11] M.S. Newkirk, A.W. Urquhart, H.R. Zwicker, E. Breval, J. Mater. Res. 1 (1986) 81. [12] D.R. Lide (Ed.), Handbook of Chemistry and Physics, seventyseventh edition, CRC Press, 1996 /1997, 6 /218. [13] L. Goumiri, J.C. Joud, Acta Metall. 30 (1982) 1397. [14] F. Stucki, M. Erbudak, G. Kostorz, Appl. Surf. Sci. 27 (1987) 393. [15] L.F. Mandolfo, Aluminium Alloys: Structure and Properties, Butterworths, London, UK, 1976. [16] S. Krukowski, Z. Romanowski, I. Grzegory, S. Porowski, J. Cryst. Growth 189/190 (1998) 159. [17] S. Porowski, J. Nitride Semicond. Res. 4 (1999) 4. [18] O. Salas, H. Ni, V. Jayaram, K.C. Vlach, C.G. Levi, R. Mehrabian, J. Mater. Res. 6 (9) (1991) 1964. [19] O. Salas, V. Jayaram, K.C. Vlach, C.G. Levi, R. Mehrabian, J. Am. Ceram. Soc. 78 (1995) 609. [20] A.S. Nagelberg, S. Antolin, A.W. Urquhart, J. Am. Ceram. Soc. 75 (2) (1992) 455.
The influence of oxygen impurities on the formation of AlN Al composites by infiltration of molten Al Mg /
/
S. Swaminathan, B. Srinivasa Rao *,1, V. Jayaram Department of Metallurgy, Indian Institute of Science, Bangalore 560012, India Received 14 August 2001; received in revised form 17 December 2001
Abstract The influence of residual oxygen in nitrogen on the formation of AlN /Al matrix by reactive infiltration has been investigated. Increasing the oxygen content from 10 ppm upwards decreased the nitride content in the matrix from 64 to 6%. Based on the analysis of the availability of oxygen at the Al-melt/gas interface, three distinct scenarios have been proposed: (i) at lowest values, oxygen does not interfere with either infiltration or nitridation reaction; (ii) at intermediate values, nitridation is suppressed, however infiltration continues; and (iii) at a critical upper value, the melt passivates without any infiltration. This phenomenon offers control of the AlN/Al ratio in the matrix and the possibility of creation of microstructural gradients by the appropriate choice of gas mixtures. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Infiltration; Residual oxygen; Nitridation; AlN
1. Introduction The formation of ceramic particulate composites with AlN /Al matrices with interesting mechanical and thermal applications is made possible by the nitridation and simultaneous infiltration of molten Al alloys into a preform. It is known that during the process of nitridation of Al /Mg alloys the degree of conversion of Al to AlN is a strong function of temperature and alloy composition. Two different types of experiments have been conducted. In the first, alloys have been exposed to nitrogen containing atmospheres in the absence of a preform. The increase in AlN content at 1000 8C with Mg in the range 1/10% and with the Mg:Si ratio has been determined [1]. In addition, it was pointed out that the degree of conversion also depended on oxygen impurities in static atmosphere. For Al / 2Mg /0.8Si, the fractional conversion to AlN diminished * Corresponding author. Present address: Technische Universita¨t Darmstadt, FB Materialwissenschaft, FG NichtmetallischAnorganischewerkstoffe, Petersen Straße 23, 64287 Darmstadt, Germany. Tel.: 49-6151-16-6318; fax: 49-6151-16-6314 E-mail address: [email protected] (B. Srinivasa Rao). 1 Presently a Humboldt Fellow.
from 95% in N2 /0.7% H2 (0.1 ppm oxygen), to 16% with pure nitrogen (3 ppm oxygen), while oxygen contents of /1% brought the level of nitridation to :/3% and lower [1]. In contrast, the practically more useful situation is when the alloy is simultaneously infiltrating into a ceramic preform, such as alumina [2 /6]. Here, it has been shown that Mg plays a key role in reducing the oxygen partial pressure to enable the alloy to infiltrate by capillary action. While alloyed Mg content is important, in this regard, it was also shown [5,6] that gettering the incoming nitrogen with a molten Al /Mg alloy was even more effective in ensuring that passivation of the infiltration did not occur. This phenomenon applies in the low temperature regime (850 /950 8C) in which the extent of nitridation is small and equally at higher temperatures at which AlN becomes a significant phase. At the same time, it has been shown that the transition temperature from metalrich to ceramic-rich composites increases slightly when commercial nitrogen is used instead of a nitrogen / hydrogen mixture. The purpose of the present paper is to report a much more substantial variation in nitride content that may be achieved during infiltration by introducing oxygen impurities and to examine the possible reasons for this sensitive behaviour.
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 0 0 2 - 3
S. Swaminathan et al. / Materials Science and Engineering A337 (2002) 134 /139
2. Experimental procedure Fused alumina particulate of 53 mm size were used as the dispersion into which infiltration was carried out. Preforms of the dimensions 80 /40 /5 mm were made by cold pressing in a steel die by applying a pressure of :/10 tons. These cold compacts were heated at 3 8C min 1 to 800 8C to avoid cracking of the preform during binder removal. After 800 8C the preforms were heated to 1600 8C and held for 1 h in order to obtain sufficient mechanical strength for handling purposes. Commercial purity aluminium and magnesium were used to prepare Al /2 wt.% Mg and Al /5 wt.% Mg alloys. Aluminium was melted in air in a clay graphite crucible at a temperature of :/750 8C and the melt was degassed by using hexachloroethane before and after magnesium was added to the melt. Magnesium was wrapped in aluminium foil in order to minimise evaporative losses. An excess of 10% (over and above the desired value) of magnesium was added to compensate for any residual evaporation. The alloys were cast into plates and homogenised at a temperature of 400 8C for 24 h and cut to the final size for infiltration. Alloy composition was determined using atomic absorption spectroscopy and inductively coupled plasma spectroscopy (ICP). Mg contents in these alloys were found to be 1.95 and 5.01% for the Al /2Mg and Al /5Mg, respectively. Alumina trays of dimensions 84 /42 /30 mm were used as crucibles. The insides of the crucibles were lined with graphite foil in order to facilitate easy removal of the composite and residual metal after the experiment. Infiltration experiments were carried out in a high temperature tube furnace (Barnstead Thermolyne). The dimensions of the composites fabricated were 80 /40 / 5 mm. Alumina preforms were kept on the top of the Al /2Mg alloy billets in the alumina crucible. The incoming gas was passed over an alloy of Al /5Mg that acted as an oxygen getter before the gas reached the melt-preform assembly. The getter was located in the 900 /950 8C temperature zone of the same tube. The effectiveness of the getter could be gauged by the formation of MgO whiskers on the wall of the furnace tube, downstream of the getter. However, no measurements were made of the oxygen partial pressure. The tube was evacuated to a pressure of 0.05 mbar and then back-filled with N2 /2% H2 (10 ppm oxygen) or commercial nitrogen (5000 ppm oxygen) before the start of experiment. Gas was passed at a rate of 250 ml min 1 throughout the experiment at a pressure slightly above one atmosphere and bubbled out through a column of oil before being vented to the atmosphere. The furnace was heated at a rate of 4 8C min 1 to 975 8C and held there for 3 h to ensure complete infiltration of the preform.
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After infiltration the crucibles were cut longitudinally with a high speed precision saw using a diamondwafering blade for optical metallography. The sectioned samples were polished using standard metallographic procedures to 1 mm diamond finish. The polished samples were examined under an optical microscope to determine the distribution of various phases in the composite. For scanning electron microscopy, the polished samples were etched with concentrated Keller’s reagent in order to remove Al and the samples were coated with gold in order to make the surface electrically conducting. The matrix hardness was measured by a SHIMADZU HMV-2000 tester at test loads of 100 g and a hold period of 10 s. Care was taken to ensure that the indentation fell within the matrix and that there was no interference with adjacent particles. An average of ten to 15 readings was taken as a measure of the representative hardness of the matrix. The volume fraction of AlN present in the composite was determined by decomposing it with sodium hydroxide solution to evolve ammonia which is quantitatively analysed by titrating against a strong acid.
3. Results and discussion The microstructures of the composites fabricated at 975 8C are shown in Figs. 1 and 2. As shown in Table 1, the first three runs, all carried out in N2 /2% H2, indicated wide scatter in the AlN content which was, however, consistent with the microstructures (Fig. 1) and hardnesses. It was also noted that the microstructures were uniform throughout the composite and that the infiltration was therefore reproducible across the section and along the length. It transpired that in the first two experiments in which the AlN content was as low as 6%, that the O-rings that sealed the tube furnace had given way during previous experiments. Degradation of the seals was evident when attempts were made, unsuccessfully, to fully evacuate the system after these two runs. During subsequent runs, the seals were replaced more frequently and, accordingly, the AlN content obtained increased to 64.5% (No. 3 in Table 1). The experiment was repeated for reproducibility at the same temperature with a new seal and a nitride content of 62 vol.% was obtained. Consequently, this value is expected to be representative of the maximum achievable at that temperature with N2 /2% H2. The AlN /Al matrix in the above cases is divided into two parts: a portion adjacent to the particle surfaces that is rich in AlN (Al/AlN in Fig. 1B,C) and the inter-particulate region that contains metal pockets as shown in Fig. 1(B) and Fig. 2(B). The relative microstructural scales may be seen more clearly in the etched scanning electron micrographs in Fig. 2. The hardness of the metal
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Fig. 1. Microstructures of the samples processed at 975 8C depicting the influence of different amounts of oxygen impurities on the amount of AlN formed in the matrix. (A) 6% AlN; AlN is mostly dispersed as isolated particles in Al-alloy matrix as indicated with an arrow. (B) 34% AlN; formation of AlN around the Al2O3 particles and starts extending into the inter-particle pores is the dominant feature at this volume fraction. However, large metal channels can still be seen (marked with M; white in colour). (C) 64.5% AlN; almost all of the inter-particle regions are covered with AlN (dark in colour) interspersed with fine Al-alloy channels (white in colour).
pockets changes little, suggesting that its nitride content is minimal. The mechanisms of alloy infiltration and nitridation are, at this time, not fully understood. However, certain physical features of the process have been established [7,8]. The melt tends to preferentially climb along alumina particle surfaces and commences nitriding outwards from the surface along an irregular front into the inter-particle pores. It is possible for residual porosity in the portions infiltrated earlier (far below the maximum extent of melt penetration), to then become isolated from the nitrogen diffusing in from the outside. These regions then fill up with advancing alloy without reacting significantly with nitrogen and lead to the large metal-rich pockets that are seen in Figs. 1 and 2. Thus, the appearance of the microstructure is completely consistent with experiments [7,8] that are carried out as a function of temperature (after ensuring the absence of air leaks) in the range 950/1100 8C in which the total amount of AlN increases, both by the expansion of the composite Al/AlN zone adjacent to the particulate at the expense of the alloy pockets as well as by an increase in the fraction of AlN within these regions. The amount of oxygen leaked into the system in the present instance is not known. However, to confirm its role, an experiment was conducted at 975 8C with commercial purity nitrogen, which has a residual oxygen content of :/ 5000 ppm and which yielded a composite with a nitride content of 29 vol.% (Fig. 1) that is close to the result of one of the earlier runs in N2 /2% H2, which is known to have been performed with a leaking seal (Table 1, No. 2). Thus, the behaviour with decreasing oxygen is similar to that displayed with increasing temperature, i.e. an increase in the nitride content of the composite. In order to understand the role of oxygen, the suggested mechanisms of nitrogen uptake by the melt are now briefly reviewed. It has been suggested [9] that nitridation proceeds by a two-step reaction, beginning with the nitridation of Mg in the vapour phase to Mg3N2, followed by reduction to AlN by the advancing melt. This sequence is similar to that which is now accepted for directed melt oxidation, in which MgO or MgAl2O4 plays a similar intermediate role. However, the presence of Mg3N2 at the composite surface has not yet been established. It has been shown [10] that if the oxides are suppressed by adequate gettering, that AlN is more stable than Mg3N2 in the alloys of interest. However, this balance could be overturned if the Mg activity was substantially increased in the vapour phase, thereby making it more available for reaction with the incoming gas. At this point in time, this issue is unresolved, i.e. does nitrogen combine directly with aluminium in the melt or via an intermediate compound. What has been believed, however, is that residual oxygen levels need to be sufficiently lowered so that oxides (MgO or MgAl2O4) do not form and terminate
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Fig. 2. SEM photographs after etching of the samples. (A) 6% AlN and (B) 64.5% AlN in the matrix. AlN formed around the Al2O3 particles, fine metal channels and a metal pocket (marked with P) that has been etched away can be seen clearly in the case of 64.5% AlN sample.
the infiltration. (The system may then choose to grow an oxide/metal composite, as in directed melt oxidation, [11] if the temperature and oxygen partial pressures are sufficiently high.) The results presented here show that there is an additional degree of complexity in the infiltration. Despite the reduced extent of nitridation Table 1 Hardness of and AlN content in the matrix S. No. Gas
AlN (vol.%)
Hardness (GPa) Nitride region
a
1. 2.a 3. 4. 5. a b
N2 /2%H2 N2 /2%H2 N2 /2%H2 Repeat of 3 N2b
6 34 64.5 62.0 29.0
/ 3.890.6 6.491.1
In the presence of oxygen impurities. Contains 0.5% oxygen.
Metal pockets 0.6890.06 0.790.07 0.890.2
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in some of the runs, the oxygen levels are sufficiently low that infiltration is still able to occur. Thus, one of the key presumptions in our prior understanding of the pressureless infiltration of Al-alloys, i.e. the need to maintain extremely low oxygen partial pressures, needs to be modified. Now one has to interpret these results in a framework that allows minute amounts of oxygen to permit capillary advance of the metal at rates that are not significantly altered, but effectively poisons the nitridation reaction, even though, for all practical purposes, the partial pressure of nitrogen is one atmosphere and that of oxygen is several orders of magnitude lower. We now examine the issues relating to the thermodynamics and rate of oxygen transport in this system. It has been pointed out [1] that the oxygen partial pressures under which AlN is stabilised in preference to Al2O3 are extremely low and in the vicinity of 10 20 atmosphere. At these pressures, the number of oxygen molecules available within the entire preform would be only / unity. This value becomes even lower if one assumes that Mg vapour from the getter is in equilibrium with MgO and oxygen. Similar or lower numbers are obtained when examining the competition between other equilibria, e.g. those involving MgO and Mg3N2. Thus, it is likely that thermodynamic considerations do not play as important a role as the availability of the appropriate gas molecules to enable a particular equilibrium to be achieved. In particular, it is likely that while experiments carried out in closed systems might conceivably approach such equilibria, the situation in flowing gas, which is the practically relevant one for the fabrication of matrices that contain AlN, is certain to be dominated by kinetic availability. To analyse the problem of oxygen arrival at the melt-atmosphere interface, it is assumed that the flux is determined by the partial pressure difference between the interface and the outside of the preform. In addition, the ratio of the oxygen flux to the rate at which the melt advances through the preform determines the net uptake of oxygen by the alloy. We assume that whether the surface is passivated or not can be defined with respect to two limits of oxygen content. The lower limit is one at which the bulk level of oxygen in the melt, enhanced by surface segregation, can lead to strong binding between Al and O atoms, thereby preventing any nitrogen uptake. Beyond the upper limit, precipitation of oxides takes place. Between the lower and upper limits, infiltration is possible, but with a gradual decrease in the extent of nitridation, i.e. the composite becomes progressively metal rich as the oxygen content rises. Let Jg be the flux of oxygen arriving at the surface of melt, D Dc /Dx , where Dc is the concentration difference between furnace atmosphere and melt surface, D is the gas phase diffusivity and Dx is the preform thickness. Assume that melt surface is at 0, i.e. everything that
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arrives there is instantaneously adsorbed. Then Dc / cfurnace /P /RT, where P is the partial pressure of oxygen in furnace. Let us suppose that poisoning occurs when the average O:Al ratio exceeds f. For this to happen, before the interface moves up by one monolayer, the number of oxygen atoms arriving per unit area must exceed f or fN moles oxygen per area, where N is the number of moles of aluminium per unit area. Thus, the limiting flux is given by: Jt /fN , where t is the time taken for melt to move up a layer /y /vm, where y is the height of one monolayer and vm is the speed of infiltration. Therefore, the limiting value of pressure is given by: DP=(RTDx) (vm fN)=y or P (fNRTvm Dx)=(Dy) At 1248 K, Dx is the preform thickness :/0.01 m and D /10 4 m2 s 1 [12], with oxygen ions of 0.28 nm diameter and an experimentally observed infiltration rate of :/2 mm h 1: y 2:8 10 10 m vm 0:002=3600 5:6 10 7 m s 1 Let oxygen adsorb as a close-packed array with a spacing of 0.28 nm. The number of moles of oxygen atoms per unit area is {(2.8)2(0.866)(10 20)(6) (1023)} 1 /2.45 /10 5 and the number of molecules per unit area is N /10 5 per m2. Since the oxygen ion radius (0.14 nm) and aluminium atom radius are the same, the number of moles of each per unit area is the same for normalisation, assuming that Al is close-packed at 1248 K. P f f(10 5 )(10375)(5:6 10 7 )(0:01)g=f(2:8 10 10 ) (10 4 )g : 2 10 4 f Pascals It seems likely that the critical oxygen concentration that can be tolerated before all infiltration ceases is the solubility limit, beyond which spinel or alumina would be precipitated. Thus, the limit for poisoning of nitridation is expected to be lower than this solubility limit. It will be even further lowered, if surface segregation is present. Some evidence is available in the literature to support the argument that the Al melt is partially covered with oxygen before Al2O3 precipitates out. Goumiri and Joud [13] studied the initial stages of oxidation and its effect on the surface tension of Al by auger electron spectroscopy. It was reported that the surface tension of Al decreases strongly at initial oxidation and stays rather constant for coverage up to one monolayer. The decrease in the surface tension of the melt was attributed to the formation of oxide islands up to approximately saturation coverage. There was no
description of the steps involved before Al2O3 precipitates out. Stucki et al. [14] studied the initial oxidation of Al. When a solid Al is exposed to oxygen, there is chemisorption up to 20 Langmuirs and then the oxide peaks start to appear in the auger signal. In contrast, when the liquid is exposed, there was no chemisorption, but the oxygen appears to go into the solution until 1000 Langmuirs and then the oxide signal appears directly. This tells us that liquid Al can dissolve more oxygen than solid Al, however, there was no indication of what happens at the melt surface after the melt is saturated with oxygen and before Al2O3 starts to precipitate. An estimate of the solubility of O in pure Al is 10 6 [15]. Thus, if we use f /10 6, then we get P /0.02 Pa and assuming a total pressure of 1 atm or 105 Pa, the oxygen level is 2 /10 7 Pa or 0.2 ppm. In the absence of more detailed information on solubilities or the actual partial pressures prevailing, it is difficult to be more quantitative. The above estimate is intended to demonstrate that partial pressures much higher than those predicted from thermodynamic equilibrium can be tolerated, given the kinetic constraints when infiltration is proceeding simultaneously. Evidence supporting the inhibiting effect of O2 on AlN growth comes from the literature on the growth of group-III nitride single crystals (AlN, GaN and InN) by dissolution of nitrogen in the melts (Al, Ga and In) under high nitrogen pressure [16]. Quantum mechanical calculations have provided some insight into the mechanisms of nitrogen dissolution into these metals [17]. It was shown that the adsorption of N2 on the surface of Al melt leads to the dissociation of N2 molecule. The energy barrier for this dissociation is around 3.2 eV, which is less than half of the dissociation energy of N2, which is 9.8 eV. In contrast, O2 dissociates and then adsorbs onto the melt surface without any energy barrier. This points to the fact that the growth of defect free AlN or GaN is difficult since any oxygen in the processing environment dissolves in the melt very easily. In the similar way, it is expected that the presence of oxygen impurities during the infiltration of Al /Mg into Al2O3 preform inhibit the growth of AlN mainly by adsorption onto the Al melt surface. Strictly speaking, one has to consider the effect of Mg on the adsorption of O2 onto the surface of Al melt since Al /2% Mg is used for infiltration. At this moment, it is unclear whether Mg helps or hinders the formation of AlN. However, its role in keeping the melt surface free of any passivating oxides, such as alumina or spinel, is well studied [5,6]. From the above analysis, one can then envisage three distinct scenarios of infiltration that are delineated by the arrival rate of oxygen at the interface. At the lowest values, oxygen dissolves in the melt in small amounts and does not interfere with either infiltration or the reaction between Al and N2. At intermediate values,
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local regions of the surface experience a sufficiently high oxygen content (possibly enhanced by surface segregation), such that Al and O atoms preferentially bind to one another. Thus, nitride formation is suppressed but infiltration is not seriously affected, thus leading to a metal-rich composite. At a critical upper value of oxygen content, the surface passivates through the formation of a monolayer of spinel or MgO. Infiltration of metal now ceases entirely and the reaction front has to advance by mechanisms which are closer to those encountered in directed melt oxidation [11,18/20].
4. Conclusions The residual oxygen in the N2 strongly influences the amount of AlN formed during the infiltration of Al /Mg alloys into Al2O3 preforms in N2. The amount of AlN increases with decreasing residual oxygen content. Oxygen appears to inhibit the reaction between Al and N2, mainly due to the relative ease with which it is adsorbed on to Al. The control of ceramic to metal ratio and the creation of microstructural gradients may now be more readily accomplished through the use of appropriate gas mixtures of N2 and O2 compared to the alternative, slower method of cycling the process temperature.
Acknowledgements Financial support for this work has been provided by the Volkswagen Foundation, Germany.
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