Directed melt oxidation and nitridation of aluminium alloys: a comparison B S.S. Daniel and V.s. R Murthy Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur 208016, India Received 4 September 1995; accepted 21 September 1995 Directed melt oxidation and nitridation are recently developed in-situ methods for the formation of metal-ceramic composites. Both processes are based on liquid-gas reaction, wherein a nolten alloy is reacted with a gaseous species, i.e. oxygen or nitrogen, to form A2O,/Al or AIN/A composite microstructures Composite formation is controlled by the alloy composition partial between these two processes, there are differences in the microstructural development. Furthe pressure of gaseous species, processing tomperature and time. Although there are similarities the growth rates of these composites are accelerated using filler materials which provide sites for secondary nucleation and in turn compete with the primary growth of the ceramic phase. the mechanical properties of Al,O Al and AIN/Al are comparable, but AIN/Al exhibits higher thermal conductivity and sensitivity to moisture. Finally, the applications and limitations of these composites are presented Keywords: oxidation; nitridation; in-situ composite Introduction reactions, there are differences between the two. First To meet the demand of recent high-technology applica- in the former method the reacting gas is injected into tions, a large number of metal-matrix composites the molten pool and the reaction time is relatively short (MMCs)were developed adopting suitable processing to control the process effectively. In the latter the reac- methodologies-3. These conventional processing tion is a continuous process over an extended period of methods, however, have some limitations, i.e. residual time. Second, by modifying both alloy and gas compo- microporosity, uneven distribution of reinforcing mate- sition, a large number of reinforcing constituents can be rials(inhomogeneity), non-wetting of reinforcements, formed On the other hand, in direct melt reactions a matrix-rcinforccment interfacc microstructure control singular reinforcing phase is possible. Third, by blowing and its cleanliness. etc. To overcome these barriers. gases, fine(10-20 um) and isolated particles are formed several innovative processing methods have been devel- and, beyond a certain percentage, higher-volume frac oped, wherein the reinforcements are formed by in-sitt tions of dis ssible, whereas, in the reactions. These processes are broadly classified based DIMOX/PRIMEX processes the ceramic phases that on their reactant phases, i.e. liquid-gas, liquid-solid and re formed are large in percentage terms and moreover solid-solid reactions. In liquid-gas reaction processing the metal and ceramic phases are interconnected there one method of forming a reinforcing phase is by injec- is also a greater fiexibility in the volume fraction of rein tion of a reactive gas into a liquid alloy reservoir. forcing phase, i.e. both metal and ceramic matrix Depending on the alloy and gas composition, one or composites are feasible by controlling process variables more phases are formed+6. Another technique which en recently developed is again based on The basic principle liquid-gas reaction, but uses a different approach. In In directed melt composite reactions a molten alloy is this process(developed by Lanxide Corporation, USA) reacted with a gaseous species (sometimes a gaseous the liquid alloy is treated in a static/flowing gas en- mixture)to develop a metal-ceramic composite AIN/Al composites are formed as the reaction product. nitriding, ALO,Al or AIN/Al composites are obtain.or vironment. Depending on the atmosphere, AlOyAl or Depending on the atmosphere, whether oxidizing or hese processes are now commonly referred to as respectively(Figure 1). Under suitable conditions a directed melt oxidation(DIMOX), and direct melt reaction product initially forms on the surface of th Although both processes are based on liquid-gas uous wich y and the product grows outward by contin- nitridation(PRIMEX) respectively molten all kingof liquid alloys through microchannels that are present within the reaction product Thus the Correspondence to v.S.R. Murth final composite contains an interconnected and inter. penetrating network of metal and ceramic. Further 261-306995/03015507 Materials Desian Volume 16 Number 3 1995 155
Directed melt oxidation and nitri~ation of a~uminium alloys: a comparison B. S. S. Daniet and V. S. R. gushy Department of Materials and Metalfurgical Engineering, Indian Institute of Tec~ffology, Ka~~ur ZU8U76, vodka Received 4 September 7995; accepted 21 September 1995 Directed melt oxidation and nitridation are recently developed in-situ methods for the formation of metal-ceramic composites. Both processes are based on liquid-gas reaction, wherein a molten alloy is reacted with a gaseous species, i.e. oxygen or nitrogen, to form AI&/AI or AINfAl composite microstructures. Composite formation is controlled by the alloy composition, partial pressure of gaseous species, processing temperature and time. Although there are similarities between these two processes, there are differences in the microstructural development. Further, the growth rates of these composites are accelerated using filler materials which provide sites for secondary nucleation and in turn compete with the primary growth of the ceramic phase. The mechanical properties of AI,OdAI and AIN/AI are comparable, but AIN/AI exhibits higher thermal conductivity and sensitivity to moisture. Finally, the applications and limitations of these composites are presented. Keywords; oxidation; nitridation; in-situ composite Introduction To meet the demand of recent high~technology applications, a large number of metal-matrix composites (MMCs) were developed adopting suitable processing methodologies’“. These conventional processing methods, however, have some limitations, i.e. residual microporosity, uneven distribution of reinforcing materials (inhomogeneity), non-wetting of reinforcements, matrix-reinfor~ment interface mi~rost~cture control and its cleanliness, etc. To overcome these barriers, several innovative processing methods have been developed, wherein the reinforcements are formed by in-situ reactions. These processes are broadly classified based on their reactant phases, i.e. liquid-gas, liquid-solid and solid-solid reactions4. In liquid-gas reaction processing, one method of forming a reinforcing phase is by injection of a reactive gas into a liquid alloy reservoir. Depending on the alloy and gas composition, one or more phases are formed4”. Another technique which has been recently developed is again based on liquid-gas reaction, but uses a different approach. In this process (developed by Lanxide Corporation, USA) the liquid alloy is treated in a static/flowing gas environment. Depending on the atmosphere, AI,O,/Al or AIN/AI composites are formed as the reaction product. These processes are now commonly referred to as directed melt oxidation (DIMOX)7,S and direct melt nitridation (PRIMEX)‘,“, respectively. Although both processes are based on liquid-gas Correspondence to V. S. R. Murthy 026%3069/95/030155-07 reactions, there are differences between the two. First, in the former method the reacting gas is injected into the molten pool and the reaction time is relatively short to control the process effectively. In the latter the reaction is a continuous process over an extended period of time. Second, by modifying both alloy and gas composition, a large number of reinforcing constituents can be formed. On the other hand, in direct melt reactions a singular reinforming phase is possible. Third, by blowing gases, fine (lo-20 pm) and isolated particles are formed and, beyond a certain percentage, higher-volume fractions of dispersoids are not possible, whereas, in the DIM~~RIMEX processes the ceramic phases that are formed are large in percentage terms and, moreover, the metal and ceramic phases are interconnected. There is also a greater flexibility in the volume fraction of reinforcing phase, i.e. both metal and ceramic matrix composites are feasible by controlling process variables. The basic principle In directed melt composite reactions a molten alloy is reacted with a gaseous species (sometimes a gaseous mixture) to develop a metal-ceramic composite. Depending on the atmosphere, whether oxidizing or nitriding, AI,OJAl or AlN/Al composites are obtained respectively (~~g~~~ I). Under suitable conditions, a reaction product initially forms on the surface of the molten alloy and the product grows outward by continuous ‘wicking’ of liquid alloys through microchannels that are present within the reaction product. Thus the final composite contains an interconnected and interpenetrating network of metal and ceramic. Further, Materials St Desian Volume 76 Number 3 1995 155
Directed melt oxidation and nitridation in aluminium alloys: B S S. Daniel and V.S. R Murthy Vapor phase Oxidising Nitriding Atmosphere Metal-Ceramic Alloy AlO Crucible Vapor phase B Metal-c。rami Composit Molten Alloy A20 Figure 1 Schematic of outward growth of ceramic/ metal reaction product(A)into free space and( B)through filler material. The corresponding microstructures in oxidizing and nitriding atmospheres are show these in-situ reactions can be utilized to obtain 'multi Additionally, these elements also improve the wettability phase composites by placing filler materials in the path of liquid alloys during wicking. In addition to the volatile of the growth direction(Figure 1), lo, The objective of species, elements such as Si, Ge and Sn are also added to using filler materials is mainly to increase growth rates, in control the reaction kinetics". In oxidation experiments addition to modifying the properties. a wide range of there is a choice of adding these elements either in materials are produced using different filler materials. elemental form (in the liquid alloy) or as surface Some examples are SiC/AlO Al, Al,O AlO,Al, dopants, Addition of surface oxides is believed to TiB,AIN/AL, SiNAIN/AL, etc,. All composit reduce the incubation period obtained by gas-metal reactions exhibit not only good In contrast, in nitridation experiments both II and IIB mechanical properties(KIC, stiffness, wear and corrosion elements are added to the alloy. In this process, group II resistance)but also improved electrical and thermal elements not only improve the wettability but also act as (shock resistance) properties compared to their ceramic a gettering agent in maintaining the oxygen partial pres- counterparts. The potential advantages of this method sure below a critical level. In the dimOx process are low processing cost, simplicity, near-net shape and AlOy Al composite growth was reported in binary alloy lexibility in filler material selection like Al-Mg, Al-Zn and Al-Li alloys where there are few limitations in nitridation For instance, AIN/Al composite growth was not seen in Al-Mg or Al-Si Process variables binary alloys up to 1450oC when a 4oC/min heating rate When pure aluminium is oxidized a thin oxide film is Al-Si-Mg exhibited accelerated weight gain beyond formed on the surface preventing further oxidation of the 1200oC19. In contrast, Scholz and Greil2D,21 reported metal. The oxidation behaviour of pure aluminium is conversion of Al to AIN in Al-2.3Li and Al-2.5Mg drastically changed with small amounts of volatile alloys. Additionally, in nitridation, small quantities of elements such as Mg, Zn, Li and Na-3. Initially, these iron were found to aid the infiltration whereas in oxida elements diffuse at a rapid rate to form a porous oxide tion, iron additions give rise to undesirable intermetallic on the surface. However, at a later stage, the surface which in turn affect the mechanical properties". In both oxide dissociates under a concentration gradient and processes, nickel addition was found to be beneficial helps to maintain the Mg level in the liquid reservoir because it refines the composite structure 156 Materials design Volume 16 Number 3 1995
Vapor Phase A Metal-C~amic Composite A1203 Crucible Vapor Phase B Filler Material Metal-Ceramic Composite v ‘AI203 Crucible Directed melt oxidation and nitridation in ~lu~iniu~ alloys: B. S. S. t)anje/ and V. S, R. Mu&y Nitriding Atmosphere Figure 1 Schematic of outward growth of ceramicl metal reaction product (A) into free space and (B) through filler material. The co~esponding microstructures in oxidizing and nitriding atmospheres are shown these in-situ reactions can be utilized to obtain ‘multiphase’ composites by placing filler materials in the path of the growth direction (F&z 1)‘,‘*. The objective of using filler materials is mainly to increase growth rates, in addition to modifying the properties. A wide range of materials are produced using different filler materials. Some examples are SiC/AI,O,/Al, Al,O,/Al,O,/Al, TiB~AlN/~, Si~N~AlN/Al, etc7,‘*. All composites obtained by gas-metal reactions exhibit not only good mechanical properties (&, stiffness, wear and corrosion resistance) but also improved electrical and thermal (shock resistance) properties compared to their ceramic counte~a~, The potential advantages of this method are low processing cost, simplicity, near-net shape and ~exibility in filler material selection. Process variables A flay selection When pure al~inium is oxidized a thin oxide film is formed on the surface preventing further oxidation of the metal. The oxidation behaviour of pure aluminium is drastically changed with small amounts of volatile elements such as Mg, Zn, Li and Na”-‘3. initially, these elements diffuse at a rapid rate to form a porous oxide on the surface. However, at a later stage, the surface oxide dissociates under a concentration gradient and helps to maintain the Mg level in the liquid reservoir’4. Additionally, these elements also improve the wettability of liquid alloys during wicking. In addition to the volatile species, elements such as Si, Ge and Sn are also added to control the reaction kineticP. In oxidation experiments there is a choice of adding these elements either in elemental form (in the liquid alloy) or as surface dopants”,“. Addition of surface oxides is believed to reduce the in~bation period. In contrast, in nitridation experiments both II and IIB elements are added to the alloy. In this process, group II elements not only improve the wettability but also act as a gettering agent in maintaining the oxygen partial pressure below a critical level. In the DIMOX process, Al,03/Al composite growth was reported in binary alloys like Al-M& Al-Zn and Al-Li alloys@ where there are few limitations in nitridation. For instance, AlN/Al composite growth was not seen in Al-Mg or Al-Si binary alloys up to 1450°C when a 4Wmin heating rate was employed. However, ternary alloys such as Al-Si-Mg exhibited acc&erated weight gain beyond 1200°C’9. In contrast, Scholz and Grei120,2’ reported conversion of Al to AlN in Al-2.3Li and Al-2.5Mg alloys. Additionally, in nitridation, small quantities of iron were found to aid the in~ltration” whereas in oxidation, iron additions give rise to undesirable intermetallics, which in turn affect the mechanical properties2’. In both processes, nickel addition was found to be beneficial because it refines the composite structurez3. 156 Materials & Design Volume 16 Number 3 1995
Directed melt oxidation and nitridation in aluminium alloys: B. S.s. Daniel and V. s R Murthy equilibrium thermodynamics, it is established that an oxygen partial pressure of -10 Pa is required for AIN formation. This is never achieved, but fortunately in the presence of local gettering agents such as Mg, Li or Na introduced as alloying additives, kinetics overrules Addition of H, to N,(forming gas) hel oxygen activity low and further improves the kinetics of the process,. Better wetting of the preform surface by the liquid metal is anticipated in a nitrogen atmosphere Composite growth Nucleation and growth mechanisms during high temperature oxidation of liquid aluminium alloys have been studied, 4: 20.27. In the early stages, when the alloy is held at above the liquidous temperature a duplex oxide(MgO MgAl O4) layer forms on the surface These oxide layers are porous and contain intercon nected microdiscontinuities. The liquid metal from the underneath reservoir wicks through these channels and emerges as small nodules on the surface. Many such Aloy Reservo nodules coalesce to form cauliflower, type of colony on the surface. Finally, several such colonies form planar oxidation front. At each stage, nodules are covered with a duplex layer and magnesium (or any Figure 2 Cross-sectional volatile species) in the alloy reservoir depletes with time When the sium content in the alloy drops belot a threshold value (i.e. 0.3 wt% Al2O3)columnar crystals Temperature and time nucleate and grow to several tens of microns uniter In addition to chemistry of the melt, control of the ruptedly (Figure 2). ALO, develops mainly as a with tion kinetics is established by varying temperature growth direction maintained parallel to the c-axis of time to arrive at the desired microstructure. At temper Al,o, " 1. These Al,O, crystals grow in an interconnected atures below 950@C selective conversion of magnesium manner, but appear isolated in two-dimensional to MgO-MgAl2O, takes place. On the other hand, sections. Finally, when there is a bulk growth temperatures above [400c lead to formation of a ALOyAl composite, an oxygen gradient builds porous material. Hence, process temperatures are across the layers and Mgo and MgA1 O in the layers confined between 1000oC and 1300%C .In contrast. underneath dissociate to give Mg", Al*and O-ions conversion of Al to aIN can be achieved at relatively Mg*ions diffuse to the surface and maintain the non lower temperatures(2750C)when compared to oxida- protective Mgo layer, whereas AP*and 02-contribute tion. Further, in nitridation, temperature has a more to growth of Al,03 3, 4.28 pronounced effect on structure; at lower temperatures, The number of steps that are involved in aIN/Al predominantly metal-matrix composites are obtained, composite growth are smaller compared to oxidation whereas at higher temperatures, metal-dispersed Here no complex surface oxides are formed. However, eramic matrix are generated. Such flexibility is not the basic processes such as nodule formation by wicking observed in oxide systems, where the final microstruc- and the conversion of these nodules to composite struc ture is ceramic dominant. To some extent, temperature ture ture are seen(Figure 3). The specific volume ratio of so dependent on the the nitride layer(VAIN/VA=1.26)is greater than unity composition of the alloy and gaseous species and is expected to form a protective nitride surface layer. Magnesium present in the alloy is reported to Atmosp have a catalytic effect on nitride formation such that Both reaction processes are sensitive to partial pressure helps to transfer the surface reaction into a volume of oxygen.4. Nagelbergreported a Po4 dependency reaction". Based on the kinetics of composite forma- for the oxidation of Al-10Si-3Mg alloy at a weight gain tion, Aluminium to aIn conversion is broadly classified of 0. 2 g/cm, whereas no significant influcnce of oxygen into four reaction domains. They are: (1)passivating partial pressure was noticed in a complex alloy". Thus surface nitridation(observed in many binary alloys), oxidation reaction kinetics appears to be largely depen-(2) n-controlled volume nitridation, ( 3)volume dent on alloy composition. On the other hand, for nitric involving outward growth of AIN-AI ture-free nitrogenous atmosphere is essential. From AIN-Al microstructures in domains(3)andya% effective conversion of Al to AIN, an oxygen and mois- and (4) break-away nitridation. The morphologies Materials Design Volume 16 Number 3 1995 157
Figure 2 Cross-sectional microstructure depicting various phases of directed melt oxidation. Inset shows thr~-dimensional view of A&O, crystals Temperature and time In addition to chemistry of the melt, control of the reaction kinetics is established by varying temperature and time to arrive at the desired microst~cture. At temperatures below 950°C selective conversion of magnesium to Mg~MgAl~~~ takes place. On the other hand, temperatures above 1400°C lead to formation of a porous material. Hence, process temperatures are confined between 1000°C and 1300°C. In contrast, conversion of Ai to AlN can be achieved at relatively lower temperatures (2750°C) when compared to oxidation”. Further, in nitridation, temperature has a more pronounced effect on structure; at Iower temperatures, predominantly metal-matrix composites are obtained, whereas at higher temperatures, metal-dispersed ceramic matrix are generated. Such flexibility is not observed in oxide systems, where the final microstructure is ceramic dominant. To some extent, temperature and time combinations are also dependent on the composition of the alloy and gaseous species. Atmosphere Both reaction processes are sensitive to partial pressure of oxygen14~24. Nagelberg” reported a pg; dependency for the oxidation of Al-IOSi-3Mg alloy at a weight gain of 0.2 g/cm*, whereas no significant influence of oxygen partial pressure was noticed in a complex alloy”. Thus oxidation reaction kinetics appears to be largely dependent on alloy composition. On the other hand, for effective conversion of Al to AIN, an oxygen and moisture-free nitrogenous atmosphere is essential. From equilibrium the~odynamics, it is es~blished that an oxygen partial pressure of -lOmBPa is required for AIN fo~ation, This is never achieved, but fortunately in the presence of local gettering agents such as Mg, Li or Na introduced as alloying additives, kinetics overrules. Addition of H2 to N2 (forming gas) helps to keep the oxygen activity low and further improves the kinetics of the process’. Better wetting of the preform surface by the liquid metal is anticipated in a nitrogen atmosphere2? Composite growth Nucleation and growth mechanisms during high temperature oxidation of liquid aluminium alloys have been studied*3,~~26,“. In the early stages, when the alloy is held at above the liquidous temperature, a duplex oxide (MgO + MgA1,03 layer forms on the surface. These oxide layers are porous and contain interconnected microdiscontinuities. The liquid metal from the underneath reservoir wicks through these channels and emerges as small nodules on the surface. Many such nodules coalesce to form a ‘cauliflower’ type of colony on the surface. Finally, several such colonies form a planar oxidation front. At each stage, nodules are covered with a duplex layer and magnesium (or any volatile species) in the alloy reservoir depletes with time. When the magnesium content in the alloy drops below a threshold value (i.e. 0.3 wt% Al,OJ columnar crystals nucleate and grow to several tens of microns uninterruptedly (Figure 2). A&O, develops mainly as OL with growth direction maintained parallel to the c-axis of A1203”“. These A&O, crystals grow in an interconnected manner, but appear isolated in two-dimensional sections. Finally, when there is a bulk growth of AI,O,/Al composite, an oxygen gradient builds up across the layers and MgO and MgA1204 in the layers underneath dissociate to give Mg2’, A13+ and 0” ions. Mg” ions diffuse to the surface and maintain the nonprotective MgO layer, whereas AP’ and 0”” contribute to growth of Al20313,24? The number of steps that are involved in AlN/Al composite growth are smaller compared to oxidation. Here no complex surface oxides are formed. However, the basic processes such as nodule formation by wicking and the conversion of these nodules to composite structure are seen (Figure 3)29. The specific volume ratio of the nitride layer ( VxlN/VA, = 1.26) is greater than unity and is expected to form a protective nitride surface layer. Magnesium present in the alloy is reported to have a catalytic effect on nitride formation such that it helps to transfer the surface reaction into a volume reaction2’. Based on the kinetics of composite formation, Aluminium to AlN conversion is broadly classified into four reaction domains. They are: (1) passivating surface nitridation (observed in many binary alloys), (2) diffusion-controlled volume nitridation, (3) volume nitridation involving outward growth of AlN-Al and (4) break-away nitridation. The morphologies of AIN-Al mi~rostructures in domains (3) and (4) are Materials & Design Volume 16 Number 3 1995 157
Directed melt oxidation and nitridation in aluminium alloys: B S S. Daniel and V.S. R Murthy A B AIN/AI Allo 100 C D 20 um 10m Figure 3 SEM micrographs showing various domains of nitride formation. (A)Surface nitridation and globule formation in the early stages of nitridation.(B)AIN/Al composite microstructure. (C)Deep etched sample showing morphology of interconnected AIN. (D) Ain crystal growth Table 1 Properties of Al2O, and AIN AIN Corundum Wurtzite 4.76;c1299a3.l1;c498 N】 RIDAIJON 140-170 TCE(25500C)(×10◆C Dielectric constant at 1 MHz Flexural strength(MPa) 5 Fracture toughness(MPa vm)2-3 Elastic modulus(GPa) 320-370 300-310 (15 N scale) 95.5 94.5 computed for different alloys during oxidation varies from 89 to 400 kJ/mol 1, 4 263. A typical growth for the Figure 4 Schematic diagram showing weight gain rates during DiMOX process can be up to 5-8 mm/day, but it can be further improved to 30 mm/day when preforms are shown in Figure 3. The AIN formed in the composite used32. The activation energy for AIN formation is 100 has a wurtzite structure with the growth direction kJ/mol and the growth rates are at least three orders oriented along the [0001] direction". The liquid metal higher compared to oxidation(Table 1) exhibits good wettability with aiN, and finds it easy to wick through the AIN crystals during growth. In oxida- Composites using filler materials tion, wicking is due to capillary action aided by magne- In the direct melt infiltration of composites, the growth in an oxide of reaction product(Al, O, /Al or AIN/AD)is restricted or system is expected to be slower channelled within the cavities of a filler material. The Both in nitridation and oxidation, different stages in filler material can be loose powder (or fibres )or owth are distinctly visible in themogravimetric analy- sintered porous ceramic body, which is usually in the sis (TGa)(Figure 4). The activation energy values shape of the final product. During infiltration, limited 158 Materials Design Volume 16 Number 3 1995
Directed melt oxidation and nitfidation in a/u~i~iu~ a/fop: B. S. S. Daniel and V. S. R. hubby Figure 3 SEM micrographs showing various domains of nitride formation. (A) Surface nitridation and globule formation in the early stages nitrida .tion. (B) AlN/Al composite microstructure. (C) Deep etched sample showing morphology of inter~nnected AIN. (D) AIN crystal grov during break-away nitridatiin MgA$O~ Formation Tim4 - Figure 4 Schematic diagram showing weight gain rates during directed melt oxidation and nitr~dation shown in Figure 3. The AlN formed in the composite has a wurtzite structure with the growth direction oriented along the [OOOI] direction”. The liquid metal exhibits good wettability with A1N3’, and finds it easy to wick through the AlN crystals during growth. In oxidation, wicking is due to capillary action aided by magnesium-induced wetting. Hence, liquid rise in an oxide system is expected to be slower. Both in nitridation and oxidation, different stages in growth are distinctly visible in themogravimetric analysis (TGA) (Figure 4). The activation energy values of vth Table 1 Properties of A&O, and AIN Property AR AIN Crystal structure Corundum Wurtzite Lattice parameter (A) a 4.76; c 12.99 a 3.11; c4.98 Theoretical density &m/cm’) 3.98 3.26 Thermal conductivity (Wm.‘K-l) 26 14cL170 TCE (25-SOO’C) (x lO?‘C) 7.1 4.19 Dielectric constant at 1 MHz 9.5 Resistivity at RT (&rn) >10J4 ;Oli Flexural strength (MPa) 400 280-320 Fracture toughness (MPa t’;;) 2-3 Elastic modulus (GPa) 320-370 300-3 10 Rockwell hardness (I.5 N scale) 95.5 94.5 computed for different alloys during oxidation varies from 89 to 400 kJ/mo111~t4~~~31. A typical growth for the DIMOX process can be up to 5-8 mm/day, but it can be further improved to 30 mm/day when preforms are used32. The activation energy for AIN formation is -100 kJ/moI and the growth rates are at least three orders higher compared to oxidation (Table 1)9,20. Composites using filler materials In the direct melt in~ltration of composites, the growth of reaction product (Al,O,/Al or AlN/AI) is restricted or channelled within the cavities of a filler material. The filler material can be loose powder (or fibres) or a sintered porous ceramic body, which is usually in the shape of the final product. During infiltration, limited 158 Materials & Design Volume 16 Number 3 1995
Directed melt oxidation and nitridation in aluminium alloys: B. S.S. Daniel and V.S. R Murthy oxygen-induced reactions. To reduce the degradation and imt bility between the matrix and come essentia Nicalon SiC fibre preforms are coated to prevent surface oxidation and minimize the interfacial reactions33 During infiltration, filler materials act as secon nucleation sites for the reaction product(AlO, or AIN and refine the composite microstructure (Figure 5), while maintaining a growth direction similar to that of primary growth 22. Secondary nucleation seen in oxide systems is attributed to the reduction of oxide coatings Sic that are inherently present on the filler material (e.g. 5 ur Sio2 or SiC) or that which are formed(ZnO, Mgo vapour deposits)during infiltration. For instance, silica B present on SiC is reduced by the advancing aluminium 3SO2+4A→3Si+2Al2O3 and Al,O, tends to nucleate on the surface of Sic par- ticles. The reaction by-product, silicon, occupies portion of channel space and in extreme cases growth elements(e.g. Zn)that are added to the liquid melt form th front. The deposited Zno is 5 um reduced by the advancing aluminium to give rise to AlO, nucleation Figure 5 Secondary nucleation of (A)AL, 0, on SiC particulate and 3Zn0+ 2Al-Al2O3+ 3Zn (B)AIN oI (Nicalon) fibi In AlO, filler materials, newly formed oxide reactions between the preform and liquid alloys are epitaxial growth, whereas on Sic particles,no nd to enhance the lographic matching was evident". In nitridation wettability. In nitridation, adverse reactions are rather secondary nucleation is verge. limited, however, in an oxidative atmosphere the rein- nous, but a clear mechanism relating to their surface forcements can lose their basic properties due to chemistry is not known Table 2 Mechanical properties of direct melt oxidized composites,> Property Al,O Al composite SiCPAlOyA Al,OAl AlO/Al Uniaxial 2D, 12HSW Youngs modulus(GPa) Shear modulus(GPa) Flexural strength 720±150 540± (3-point bend test) (WM 20)(WM 22)(max 880) (max 620) (MPa vm (max 29) (max 16) 414-2100 0.1370.312 Vickers hardness(GPa 141-15.0 at700°C l18-39.6 Coeff. of thermal expansion 9.3-11.0 Dielectric constant o C-axis of Al O, ⊥ to c-axis of Al2O3 4 awith variation in residual metal content Weibull modulu Materials Design Volume 16 Number 3 1995 159
Directed melt oxidation and nitridatjo~ in aiuminium alloys: f3. S. S. Daniel and V. S. R. Murthy Figure 5 Secondary nucleation of (A) A&O, on Sic particulate and (B) AlN on Sic (Nicalon) fibres reactions between the preform and liquid alloys are inevitable and these reactions tend to enhance the wettability. In nitridation, adverse reactions are rather limited, however, in an oxidative atmosphere the reinforcements can lose their basic properties due to oxygen-induced reactions. To reduce the degradation and improve the compatibility between the matrix and the fibre, coatings become essential. For instance, Nicalon Sic fibre preforms are coated to prevent surface oxidation and minimize the interfacial reactions33. During infiltration, filler materials act as secondary nucleation sites for the reaction product (A&O, or AlN) and refine the composite microst~cture (r;igure 5), while maintaining a growth direction similar to that of primary growth iLz2 . Secondary nucleation seen in oxide systems is attributed to the reduction of oxide coatings that are inherently present on the filler material (e.g. Si02 or Sic) or that which are formed (ZnO, MgO vapour deposits) during infiltration. For instance, silica present on Sic is reduced by the advancing aluminium via the displacement reaction 3SQ + 4Al-+ 3Si + 2Al,O, (1) and Al,O, tends to nucleate on the surface of Sic particle.?. The reaction by-product, silicon, occupies a portion of channel space and in extreme cases growth is inhibited due to a ‘choking’ effect3’. Further, the solute elements (e.g. Zn) that are added to the liquid melt form a vapour and are deposited on the filler material surface ahead of the growth front. The deposited ZnO is reduced by the advancing aluminium to give rise to A&O, nucleation23. 3Zn0 f 2Al+ Al& + 3Zn (2) In A&O, filler materials, newly formed oxide exhibits epitaxial growth, whereas on Sic particles, no crystallographic matching was evident34. In nitridation, secondary nucleation is expected to be simply heterogenous, but a clear mechanism relating to their surface chemistry is not known. Table 2 Mechanical properties of direct melt oxidized composites’“*3~‘,4~ Property Al,O,/AI composite without preforma Young’s modulus (GPa) 88-304 Shear modulus @Pa) 41-123 Flexural strength 45-525 (4-point) (MPa) (3-point bend test) Fracture toughness 2.9-9.5 Wa ~‘3 Compressive strength 414-2100 @@a) Poisson’s ratio 0.137-0.312 Vickers hardness (GPa) 1.41-15.0 at 700°C - Thermal conductivity 1 t s-39.6 (Wm-’ K-l) Coeff. of thermal exapansion 9.3-I 1 .o (X 10*/K) Dielectric constant [j to c-axis of A&O, 8.0 I to c-axis of A&O3 6.4 With variation in residual metal content bWeibull modulus 403d Al,OrtAl 301 312 (WMb 20) 5.9 - 8.3 4.9 36 9.0 SiCd SiC~Al~O~lAl Al~O,IAI Uniaxial 2D, I2HSW 324 - 334 720 k 150 540 f 60 (WM 22) (max 880) (max 620) 6.9 27 k3 15 + 1 (max 29) (max 16) - - 6.4 3.7 82 7.5 Materials & Design Volume 16 Number 3 1995 159