Synthesis of in-situ reinforced Al composites from AlSiMgO precursors

Acta mater. Vol. 45, No. 10, pp. 40674076, 1997 6 1997Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved

Pergamon

PII: S1359-6454(97)00085-Z

SYNTHESIS

Printed in Great Britain 1359-6454/97 $17.00 + 0.00

OF IN-SITU REINFORCED Al COMPOSITES FROM Al-Si-Mg-0 PRECURSORS M. HANABE and P. B. ASWATH

Materials Science and Engineering Program, University of Texas at Arlington, Arlington, TX 76019. U.S.A. (Received 17 September

1996; accepted 20 February 1997)

Abstract-Al matrix composites reinforced with micro-composite Al203/Al particles were synthesized by reactive infiltration of molten Al into preforms of particulate Si02 or Mg + SiO2 mixtures at 1075°C. Displacement reactions between silica and magnesium containing oxides lead to in-situ formation of the reinforcements and also aid the infiltration of the melt. In the presence of Mg, it was found that the consistency of infiltration was better and that the transformation of silica to alumina involved intermediate displacement reactions unlike the single step reaction without Mg in the preform. It was observed that the morphology and size scale of the micro-composite A120,/Al particles were affected by the presence of Mg. Without Mg a finer scale AlzO,/Al microstructure with a tendency to be elongated in the growth direction formed, while a coarser morphology with interconnectivity in both phases developed from Mg + Si02 preforms. Potential for such microstructural variations of the reinforcements permits control of the overall mechanical properties of the composite. % 1997 Acta Mera/lurgica Inc.

INTRODUCTION In recent years displacement

reactions between liquid metal and ceramic oxides have been used to fabricate ceramic and metal matrix composites. Such reactions may be of the type 4M + 3Si02 = 2M20,+ 3% where M is a trivalent metal, the thermodynamic criterion being that at the processing conditions the Gibbs free energy of the reaction is negative. While it is an irony that such reactions are a source of reinforcement/matrix degradation in composites, a judicious choice of reactants will result in the in-situ processing of thermodynamically compatible and technically important matrices and reinforcements. More specifically, the application of this concept to make A1203 and Al composites has been extensive. In the DIMOX process, where an Al,O,/Al composite is grown from an Al alloy surface at temperatures well above its liquidus, reactions between oxides like MgO, MgA1204 and ZnO and liquid Al are the source of the composite [l-4]. Alloying elements like Mg and Zn help in the formation of these oxides on the surface of the melt. While these oxides are cyclically reduced and generated by the Al alloy during the process, solid oxide precursors, primarily silica and mullite, in the shape of the component required, have also been used to make alumina matrix composites. Reactions between silica and/or mullite and molten Al have been an issue in the foundry industry for a long time [5-71. Brondyke [S] showed that exposure of commercial alumina-silica refractories with even up to 99% Al203 content to molten Al resulted in the formation of Al,Oz and metallic Si. In a related work 4067

on the reaction between vitreous silica and molten Al, Standage and Gani [6] found three phases of alumina with q-A1203 and 8-A1203 as primary products and c(-AI~O~ as a secondary product formed by phase change from 8-AlzO,. The results of Prabriputaloong and Piggott [7] indicated that initiation of the reaction between silica and Al was delayed by the native oxide skin on Al, which prevented direct Al/SiOI contact. Devereux demonstrated that these displacement reactions can be beneficially applied to form A120,/Al-Si composites [8]. In a carefully designed set-up, SiO2 (glass) specimens were immersed in an Al bath for sufficient times. Three dimensionally interconnected Al203/AI-Si composites were formed and it was found that the Si content in Al affected the morphology of the alumina. More recently Breslin et al. [9] and Loehman et al. [lo] converted amorphous SiOt and mullite preforms, respectively, into A&Ox/Al composites using the immersion technique. A remarkable advantage of these processes is that by carefully controlling the density of the precursors and amount of Al, the shape of the precursor can be faithfully reproduced in the final composite; thus net shape fabrication is feasible. Dhandapani et al. [1 1] while infiltrating Sic particulate preforms with A1203/Al also observed that interaction between the surface SiOz on the Sic and liquid Al led to higher metal contents in the composite. Very recently, the present authors demonstrated that such displacement reactions can be applied to the processing of in-situ reinforced Al matrix composites [ 121. In their process a porous amorphous silica

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particulate preform was infiltrated by molten Al. The ensuing displacement reaction between the silica particles and Al leads to the formation of the reinforcements (micro-composite AlzOiiAl particulates) in-situ and also aids in the infiltration of the metal into the preform. In addition to the inherent advantages of low cost and simplicity of the process, it is also believed that by controlling the distribution and morphology of the reinforcements (being micro-composites themselves), the overall mechanical properties of the composite can be varied. While still saving the attractiveness of single step processing. this article will show that by controlling the composition of the preform, multiple displacement reactions between different oxides and Al can be promoted during processing which affect the infiltration and morphology of the A120j/Al particles. Composite formation will be discussed mostly qualitatively based on the information available on the thermodynamics and wetting of Al/solid oxide systems and experimental techniques of optical and electron microscopy, energy dispersive and powder diffraction X-ray analysis.

EXPERIMENTAL

PROCEDURE

The process of making the in-situ reinforced Al matrix composite involves reacting a SiOz particulate preform contained in a quartz tube with molten Al at temperatures between 1000 and 1lOO’C by immersing it in the Al melt as shown in Fig. 1 [12]. The melt infiltrates the preform by reacting with the SiOz and the Si released from the reaction diffuses out of the preform and into the bulk melt. The immersion

Si wafer

/

II

Mg powder

\

SiOZ

Fig. 1. Schematic of the experimental set-up to synthesize in-situ reinforced Al matrix composite. (i) SiO2 particles packed into a quartz tube with Mg powder at the ends for oxygen gettering and (ii) 3 wt% Mg powder is mixed in with 502

particles.

Al COMPOSITES

technique, however, affects the consistency of infiltration. Upon immersion the entrapped oxygen from the preform reacts with liquid Al at the open ends of the quartz tube and forms a passivating native oxide which prevents direct Al/SiO, contact. Infiltration occurs only when Al permeates through the cracks developed in the aluminum oxide layer owing to thermal/mechanical stresses and comes in contact with the SiOz particles. During the course of the experiments, it was determined that the consistency of the infiltration was also better if the molten Al came into contact initially with a material that would subsequently dissolve in Al as opposed to a porous Si02 preform, and also if an inert atmosphere was maintained within the preform. In order to maintain the simplicity of the process and retain the high rates of composite formation offered by the immersion technique, elaborate set-ups were eliminated by developing a self generating vacuum system within the quartz tube itself. This system had Mg powder (used in two different ways as detailed later) along with crystalline SiOz (quartz) particles inside a tube with open ends. The open ends of the tube were sealed with Si wafers (0.2 mm thick) using very small amounts of refractory cement. Upon immersion of this preform into molten Al at 750°C Al comes into contact with Si and dissolves it. While this is in progress, Mg by virtue of its excellent oxygen gettering ability reduces the oxygen partial pressure inside the tube. Hence, molten Al comes into contact with an oxygen free preform after most of the Si from the wafer has diffused away from the end of the tubes. In this process it is assumed that after the above-mentioned gettering process and for the remainder of the processing time, the only source of oxygen for the reaction is from the preform. The solubility of oxygen in molten Al is known to be negligible. Magnesium. as mentioned above, was packed inside the tube in two ways. In the first method, it was used only as a gettering agent by packing it at the two ends of the quartz tube to about l-2 mm thickness (see Fig. I ). In the second method, Mg powder (~44 pm), 3% by weight of SiOz, was mixed thoroughly with the SiO, particles. This amount was greater than that required for gettering and was added to investigate its effect on the mechanism of composite formation. The experiment involved immersing the preforms in a bath of molten Al held in a clay graphite crucible at 750°C by placing preheated alumina rods on them. Thereafter the furnace temperature was raised to 1075’C. Samples were held for different times at this temperature to study the nucleation of A1203 and its morphology. After the required holding time, the temperature was lowered to 750°C and the samples removed. The raw materials used in this study together with their characteristics are presented in Table 1. It is to be noted here that the size of all the preforms (9 mm dia. x 12 mm length), the amount of

HANABE

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Table 1. Raw material propertles Material

Particle size

Assay

Al Mg SiO* (low quartz)

lumps <44pm IO&l25 pm

99.1% 99.8%

REINFORCED

4069

porosity is eliminated only after most of the Si has diffused out. The silicon concentration in the finished composite varied typically between 2 and 5 wt% and this could be varied by changing the ratio of the volume of the preform to the volume of bulk Al. AI-Si02-Mg

Al and Mg and the ratio of the volume of the preform to the volume of bulk Al were kept constant. The composites were subsequently cut on a low speed diamond saw and the cross-sections prepared for optical and scanning electron microscopy. Energy dispersive analysis using X-rays was conducted on a Cambridge 120 Stereoscan Scanning Electron Microscope using a Kevex Super Dry Detector. Elemental X-ray dot mapping was performed to understand the distribution of the various elements involved in the displacement reactions. Powder X-ray diffraction was carried out on selected samples using a Phillips Powder X-ray Diffractometer with CuKoc radiation in the 20 range of 10-80” to study the evolution of different phases as a function of reaction time. In order to accentuate the X-ray intensities from other low volume fraction phases, the Al signal was eliminated by dissolving equal volumes of reacted samples in dilute HCl to leach out the Al. Residue from the acid leaching was washed thoroughly with water and acetone and then powdered.

Al COMPOSITES

g’stem

In the second set of experiments, Mg (3 wt% of SiOZ) was incorporated into the preform to study its effect on the composite forming mechanism in addition to its role as an infiltration initiator. The first step in this study was to determine if there was any interaction between Mg and SiOZ even before the infiltration of Al. For this, Mg + Si02 mixtures in the same proportion as those employed for composite formation were encapsulated in a quartz tube with one end open. The other end of the tube was sealed off with a 5 mm thick quartz disc which eliminated the Al/preform interaction for at least 2 h. This tube

RESULTS Al-Si02

system

This section describes the results of the infiltration and displacement reaction in the absence of Mg. Figure 2(a) is a back scattered electron (BSE) image, Si and Al elemental dot maps of a representative cross-section of the preform which was immersed in molten Al for 15 min at 1075°C. These figures show partially reacted SiO2 particles which have an outer shell of A1203/Al composite and an inner core of unreacted SiO, and some of the smaller SiOZ particles which have been completely converted to A120,/A1 particles. Silicon displaced by the reaction surrounds the particles. Also evident in this figure and in other samples examined after such short holding times were the large porosities even though the Al had infiltrated the complete length of the preform. Figure 2(b) shows the cross-section of a preform after 1 h at 1075°C. From the elemental X-ray images of Si and Al it is clear that the original Si02 particles have completely transformed to A120j/Al and most of the elemental Si has diffused out of the composite into the bulk Al melt. As the hold time at 1075°C increases, the Si content in the matrix drops with a concurrent decrease in the porosity of the composite. Figure 3 developed as a result of X-ray diffraction of powders shows the evolution of phases as a function of infiltration/reaction time at 1075°C. Thus, although the infiltration and complete transformation of SiO: to A1203 is quite rapid, a substantial amount of

@I Fig. 2. Back scattered electron image, Si and Al elemental X-ray dot maps of the composite synthesized from SiO2 particulate preform without Mg powder at 1075°C: (a) after 15 min of processing, shows partially and fully reacted SiO2 particles in a Si enriched Al matrix and (b) after 60 min of processing, shows fully reacted SiOz particles yielding AhO,/AI microcomposite particles in an Al matrix. A: transformed region made up of Al203 and Al channels; S: Si displaced by the reaction which has not yet diffused away; Q: unreacted SiO2; M: Al matrix.

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Al COMPOSITES

100

Ii& 0

3 b .g tij E .E E .F 0

20

40

60

80

loo

120

3

10

!i 0 0

1

A

0.1

A

Reaction Time (minutes)

Fig. 3. Evolution of phases in the composite as a function of time in the AI/SiOz system. X-ray diffraction intensities from the strongest line of each phase were used to calculate the percentage relative intensities. The signal from Al was eliminated by leaching out the phase.

was then immersed in molten Al and held at 1075°C for two different times, viz. 15 and 30 min. Heat treatments beyond these durations were considered unnecessary since the results of infiltration indicated that it took only 1.5min for Al to completely infiltrate

the preform, thus rendering Mg + Si02 interactions beyond 30 min irrelevant. X-ray diffraction patterns of powdered samples after this heat treatment revealed small amounts of Mg,SiO., and Si. Figure 4 is a BSE image of the cross-section of a preform which was immersed in Al for 15 min at 1075°C along with elemental X-ray images for Si, Mg and Al. Similar to the system described in the previous section, localized porosity was evident, while the Al had infiltrated the entire length of the preform. Comparing both the Al and Mg elemental X-ray maps, it is evident that an interfacial layer is seen to develop on the SiOZ particle in contact with molten Al. Energy dispersive spectroscopy (EDS) of this layer indicated an Mg:Al ratio close to 1:2, which

Fig. 4. Back scattered electron (BSE) image Al, Si and Mg elemental X-ray dot maps of the composite synthesized from SiOl particles mixed with 3 wt% Mg powder after 15 min of processing time at 1075”C, showing Mg spine1 at the boundary of the SiOl particles. M: Al rich matrix; Q: unreacted SiOl; S,: MgAb04.

0.01 1

10 100 Reaction Time(minutes)

Fig. 5. Average Si and synthesized by reacting preform at 1075°C. Each least 10 readings taken

1000

Mg content in the Al matrix a SiOl + 3 wt% Mg particulate point represents an average of at at different locations by EDS.

suggests that this reaction product is MgA1204. X-ray diffraction of powdered samples also confirmed the presence of spine]. The average composition of the alloy (measured by a semi-quantitative EDS routine) as a function of reaction time at 1075°C is shown in Fig. 5. The concentration of Si steadily declines with time. As the displacement reaction proceeds the Mg and Al displace Si from the SiOZ and eventually the Si diffuses away into the molten Al melt. The concentration of Mg also shows a steady decline with reaction time. As the reaction front moves into the Si02 particle for prolonged holding, two scenarios develop. In the first case, some of the smaller particles in the preform transform completely into spinel, while those that are close to the average particle size (100 pm) react with a lean Al-Mg alloy and form an Al rich oxide (henceforth referred to as AO) with some Mg and Si. Figure 6(a) depicts the latter case, where the set of elemental X-ray maps confirm the presence of Si and Mg within a predominantly Al oxide. Also seen in this figure is the outer spine1 layer. The next step in the transformation is shown in Fig. 6(b) where coarse AlgO channels are seen nucleating at the AO/spinel interface. Alumina grows into A0 while the spine1 layer remains more or less intact. Two features are evident from this set of micrographs. While the A0 adheres fairly well to the outer spine1 layer, alumina shrinks away from the spine1 leaving behind a gap which is subsequently f&d up by Al. Secondly, there is a build-up of Si at the AO/alumina, alumina/spine1 and spinel/matrix interfaces. Transformation to alumina is rather rapid and occurs at the expense of the AO. In fact, within the first 30 min almost all the particles transform completely into alumina channels (Fig. 7). The spine1 layer, however, persists for much longer times. Figure 8 which shows the percentage relative intensity of the

HANABE and ASWATH:

SYNTHESIS OF IN-SITU REINFORCED

407 1

Al COMPOSITES

phases (calculated from X-ray diffraction of powdered samples) as a function of the reaction time lends support to this fact. The MgA1204 itself has fine pockets of metal interspersed in it as seen in Fig. 9. It should be noted here that once the Al infiltrated the preform there was no evidence of magnesium silicates or aluminosilicates. Also, no other phases of alumina, except cr-alumina, were detected. Morphology

of A1203/Al; effect of Mg

An attractive outcome of this study was the possibility of changing the morphology of the reinforcement by varying the composition of the preform. The concept of incorporating other phases like Sic along with SiOt in the preform was introduced in the earlier work [12], but the SIC did not participate in either the infiltration or reinforcement forming mechanism. By including an active

(4

(W Fig. 6. Cross-section of a SiOz particle from a Mg + SiO? preform which was reacted with Al at 1075°C along with the corresponding Al, Si and Mg X-ray maps: (a) shows the transformation to an Al rich oxide(A0) which occurs during the first 30 min of the reaction time as the metal front advances beyond the outer spine1 layer. Note the presence of Si and Mg within AO. (b) Nucleation of a-ALO, channels at the spinehA interface after 30min of reaction time. Notice the enrichment of Si at AO/spinel, spinel/alumina and spinel/matrix interfaces. M: Al rich matrix; S,: MgALOr; AO: Al rich oxide; A: Al,O,/Al channels.

Fig. 7. Cross-section of a particle showing the almost complete transformation of A0 into a-Al203 and Al channels after 1h of reacting SiOs + 3 wt% Mg with Al at 1075’C.

ingredient like Mg in the preform a two step displacement reaction involving Si02 first, and then AO/spinel to yield AlI03/A1 can be forced, as discussed in the previous section. The influence of such a route on the morphology of A1,03/Al is clearly seen in Fig. 10. Figure 10(a) is a representative cross-section towards the centre of a single Si02 particle transformed into alumina without Mg in the preform. Microscopically distinct regions of A120,/Al channels separated by a metal layer are seen here. These represent growth fronts which originated from different sides of the silica particle and have converged at the center. Figure 10(b) and (c) shows higher magnification images of the representing growth directions parallel and normal to the plane of the paper. Comparison of these shows that the AhOJAl are elongated in the growth direction. These figures also indicate a higher degree of interconnectivity in the alumina. A1203/Al formed from Mg treated preform [Fig. 10(d)], however, did not show such preferential alignments. Instead, a coarser morphology with interconnectivity in both phases

L 50

100

150

Reacth

Time(minute8)

200

250

D

Fig. 8. Relative intensity of different phases in the composite as a function of processing time at 1075°C in the Al-SiOrMg system when the SiO2 particles were mixed with 3 wt% Mg before processing. Aluminum was leached out to accentuate the intensities of other phases.

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Al COMPOSITES

Fig. 9. High magnification micrograph of a particle showing A1201 channels and spine1 which has pockets of metal interspersed in it.

developed in the presence of Mg. In addition, the alumina which formed from a Mg treated preform coarsened as a function of time. Figure 1l(a) and (b) shows two Alz03 particles which were originally SiOz. The effect of processing time is clearly evident from this figure, where after 14 h of processing the alumina has coarsened. The presence of Mg and the prolonged processing time also leads to the penetration of individual alumina grains by Al. This is clearly seen both in Fig. 11(b) and bright field TEM (Fig. 12), where the metal has penetrated the grain boundaries of some finer alumina grains. DISCUSSION Development of various intermediate compounds observed during this study and their eventual

(b) Fig. 11. Morphology of AIzO,/AI particles in composites made from preforms of Mg + SiOz; (a) 30min of processing: note the interconnectivity and channel size of alumina; (b) 14 h of processing: alumina has coarsened and some of it has also been dispersed in 41.

conversion to Al*Oj/Al indicate their effect on both the infiltration process and also on the morphology of the final A1203/Al particles. Thus any analysis of the results will have to include first the examination

Fig. 10 Electron micrographs of an AL03/Al particle illustrating the effect of Mg on the morphology of the particle. (a) Growth fronts of AIZOJ/A1 from different sides of SiOl formed in the absence of Mg converging at the centre of the original particle. Regions p and n are magnified in (b) and (c) which show growth directions parallel and normal to the plane of the micrograph, respectively. (d) Coarse channels ofAhOs/Al formed by the incorporation of Mg in the preform. Note that magnification in (b), (c) and (d) is the same and the metal was etched out for better contrast.

Fig. 12. Bright field TEM of an ALOX/AI from a Mg treated preform showing penetration of ALO? grains by Al.

HANABE and ASWATH:

SYNTHESIS OF IN-SITU REINFORCED

Al COMPOSITES

4073

Mullite + Si + Al203

Si

Al

Wt.% Si

(log Po2 = -32.2 at.m/ (log Paz = -34.96 atm) (log Po2 = -33.5

A1,Mg + Al,

Al 0.11

3

lo

Mg Wt.% Mg

(b) Fig. 13 Calculated isothermal section of (a) AI-%-O and (b) AI-Mg-0 phase diagrams at 1075°C. Phase fields are enlarged for clarity; * and ??indicate the nominal initial and final Mg content in the matrix. respectively.

of the AI-Sl-0 and Al-Mg-0 phase equilibria and then the issues pertaining to the wetting of ceramic oxides by liquid Al.

Phase relations in Al-Si-0

and AI-Mg-0

systems

Both these ternary systems have been widely studied to aid the foundry and the ceramic/metal composite processing industries and will be reviewed again here.

In the Al-Sl-0 system reactions between SiOz and mullite with liquid Al above 800°C are thermodynamically favored with the reaction product being alumina. Although cc-alumina is the most stable phase, there has been evidence of q- and o-phases of alumina [6]. Figure 13(a) is a calculated Al-Si-0 phase diagram at 1075”C, which is the reaction temperature for Si02 and Al in the present case. It is seen here that both Si02 and mullite are stable with

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HANABE and ASWATH: SYNTHESIS OF IN-SITU REINFORCED Al COMPOSITES

only pure Si and that trace amounts of Al are enough to stabilize A&O? according to the following reaction: 4Al+

3Si02 = 2A120, + 3Si.

(1)

Even if mullite or any alumina-silica solutions form as intermediate products, they would eventually transform to alumina. Diffusion of Si out of the composite and into the bulk melt owing to a concentration gradient helps in maintaining the matrix composition within the Al-Si + A&O, phase field of the diagram. This aspect provides the potential of varying the Si content of the matrix by adjusting the proportions of SiOZ and Al to achieve the desired wear resistance properties of the composite. The above-described interaction represents the scenario in the preform away from the open ends of the quartz tube where Mg was packed to initiate melt infiltration. Upon dissolution of this Mg layer into the advancing Al melt its concentration continues to drop owing to its diffusion away from the front and into the bulk melt, thus negating its effect on the Al-Si02 interaction in the interior of the preform. However, in the regions near the Mg layer and also in the preforms containing Mg and Si02, Mg alters by introducing the binary AI-Si02 interaction additional intermediate displacement reactions. Evidence of a Mg,SiO,, in an uninfiltrated preform points to the fact that Mg reacts with the SiOl particle at its surface according to the following reaction: 2Mg + 2Si02 = Mg2SiOl + Si.

(2)

However, on coming into contact with Al when the melt infiltrates the preform, the silicate is highly unstable and immediately converts to MgA1204 as follows: 2Al+ MgzSiOl = MgA1204 + Si + Mg.

(3)

Spine1 can also form when the Al-Mg advances into the SiOZ particle in the following manner: 2Al + MgtAll+ 2SiO2 = MgAhO., + 2Si.

(4)

Again, thermodynamic calculations show that the equilibrium conditions for both reactions (2) and (3) define a 100% Si; thus both reactions can proceed in the forward direction whenever Si diffuses away from the reaction front. The reason for Mg,SiOl not being a direct source of Al2O3 through the reaction 8Al + 3Mg2Si0, = 4A1203 + 6Mg + 3Si

(5)

can be seen by considering the calculated Al-Mg-0 phase diagram at 1075°C [Fig. 13(b)], the magnitude of the free energy changes for reactions (4) and (5) and also the matrix Mg content after 15 min reaction time, which varied between 0.5 and 1% (Fig. 5). The phase diagram indicates that A1203 is stable only when the Mg content in the alloy is below 0.11%. Also the magnitudes of the free energy changes for reactions (4) and (5) being - 1352.18 and

- 63.10 kcals, respectively, indicate the higher stability of MgA1204 over A1201. Calculation of the Al-Si-Mg-0 diagram was not attempted because the interaction parameters of Al, Si and Mg in their solution is not very well known. It has, however, been reported that Si in Al-Mg tends to reduce the activity of Mg [2]. The development of AO, depicted in Fig. 6(a), is the result of the interaction between the SiO, and an Al-Si-Mg alloy which does not have enough Mg to sustain reaction (4) beyond the 10 pm of spine1 already formed on the surface of the SiO*. The presence of Mg and Si within A0 along with the fact that the nucleation of Al,O, [Fig. 6(b)] is accompanied by an enrichment of Si in the metal channels leads to the belief that A0 is initially stabilized by Si and Mg. Eventual transformation of this oxide to alumina is controlled by the diffusion of Si and Mg owing to a compositional gradient set-up in the Al matrix across the spine1 layer. The spine1 remains stable as long as the Mg wt% in the matrix is above 0.11% as predicted by the calculated phase diagram in Fig. 13(b). However, depletion of Mg from the matrix (Fig. 5) shifts the equilibrium to the Al, Mg + AIZO1phase field and the conversion of spine1 occurs in the following manner: 2Al+ 3 MgA1204 = 4A1201+ 3Mg.

(6)

Free energies of formation of all the oxides described in the different reactions were obtained from JANAF tables [ 131.The Al-Si-0 phase diagram was calculated on the basis of a description of this system in Ref. [14] and the Al-Si activities were assessed from solution data published in Ref. [15]. The Al-Mg-0 system was calculated using the Al-Mg solution data from Ref. [ 161and also from the description of the system in Refs [2] and [17]. Reactive

wetting in Al-Si-Mg-0

system

The development of the Al matrix in the present study occurs without the aid of external pressure. Such infiltration requires spontaneous wetting of the preform particles by liquid Al. Numerous studies of wetting in liquid metal/ceramic systems have indicated that very low contact angles can be achieved if an interfacial reaction occurs between the ceramic and the liquid metal [18]. Under such conditions the contact angles dramatically decrease until they assume an equilibrium value in contact with the interfacial product. Aksay et al. [19] examined the role of the interfacial free energy in determining the wettability, and proposed that the solid-liquid interfacial tension ysmidecreases by an amount equal to the magnitude of the change in the free energy owing to a chemical reaction (A@‘) at the interface according to

where, ysl and yi are final and initial interfacial tensions, respectively, of the solid-liquid interface.

HANABE and ASWATH:

SYNTHESIS OF IN-SITU REINFORCED Al COMPOSITES

The concept of reactive wetting has been demonstrated in both the Al-Si-0 and Al-Mg-0 systems. Marumo and Pask [20] in their study of the wetting of fused SiOl by liquid Al attributed the initial decrease in the contact angle to the free energy of formation of an Al rich interfacial layer which subsequently decomposed to AllO, upon cooling. The presence of such an interfacial layer between amorphous SiOl and liquid Al was also demonstrated by the present authors in a previous study [12]. Similar trends were observed in the Al-Mg-0 system [ 171. For example, wetting of MgO by Al is accompanied by the formation of an intermediate spine1 layer, while the wetting of the spine1 substrate itself results in an interfacial Al2O3 layer. The relative stabilities of these oxides in contact with Al are determined by the activity of Mg in Al. The aforementioned discussion can be applied very well to the present situation where the negative free energy changes accompanying reaction (1) and reactions (3 and 4) aid in the pressureless infiltration of Al into the preform for both Al-Si-0 and Al-Mg-Si-0 systems. While the theories of reactive wetting as controlled by the magnitude of the interfacial reactions are generally accepted, more recent studies attribute wetting to other factors. Zhou and De Hosson [21] in their study propose that if the change in volume accompanying a displacement reaction is negative then the reaction does not improve wetting. Espie et al. [22] indicate that the final contact angle is dependent on the interfacial product irrespective of the magnitude of the free energy change. However, these two studies fail to account for the initial decrease in contact angle which is controlled by the chemical reaction. When one considers the infiltration of Al into a porous SiO, preform, it is this initial decrease in the contact angle which plays a dominant role over the final contact angle. The initial drop in the contact angle helps in advancing the Al melt from one SiOz particulate layer to the other. This argument is clearly seen by considering Figs 2(a) and (4) where the melt has infiltrated the entire length of the preform forming an interfacial layer of alumina and spinel, respectively. while the core of the particles is still Si02. While reaction induced wetting seems a very plausible mechanism for the infiltration in the present study, the effect of the physical distribution of the SiO, particles must also be considered. In this respect, the work by Yang and Xi [23] may be particularly applicable. Through empirical calculations, they have shown that by considering the particles in a preform as an orderly array of monosized spheres, spontaneous infiltration was highly dependent on the infiltration direction and thus on the nature of packing. While the anisotropic effect of infiltration was hard to notice in our case, the effect of packing density on wetting was, however, clearly observed. In some of the preforms which had a non-uniform

4075

distribution of SiOl, the melt had not completely infiltrated in regions where the volume fractions exceeded - 0.7. Silicon which is released from the reactions and Mg which is incorporated in the preform affect the surface tension of the liquid Al. While the reduction of y,”owing to Si is only slight [17], Mg causes a large decrease in the surface tension of Al [24]. Thus the wetting of the silica particles ahead of the first layer may be even more spontaneous in the presence of Mg owing to this change in the surface tension of Al. Morphological variations in Al~O,/Ai Figure 10 indicated that the morphology and size of Al,OJAI were affected by Mg. At the outset it would seem logical to explain these variations on the basis of the volume changes associated with the different displacement reactions leading up to the final A1203. Other workers [9, 211 have postulated that the negative change in the volume accompanying the conversion of Si02 to AlzOx leads to cracks in the alumina which are subsequently filled up by the metal. Thus the origin of interconnectivity of both the ceramic and metal phases was linked to this volumetric contraction. Volume changes based on molar volumes of compounds for reactions (1), (4) and (6) are _- 25%, - 12.5% and - 14%, respectively. If one assumes that greater volume changes result in a higher density of cracks and thus a finer channel distribution, comparison of Figs 9 and lO(b or c) speaks of the contrary. The formation of spinehmetal accompanied by a volumetric contraction of 14% appears to have a finer structure than the A1203/Al formed from direct interaction of SiOt and Al with a -25% volume change. Furthermore, although the volume changes for reaction (4) and (6) are comparable, the microstructures in Figs 9 and 10(b) show considerable differences in the channel size and interconnectivity between the two phases. In addition, it should also be remembered that inclusion of Mg within the preform resulted in coarsening of the alumina. Thus the available evidence from the present study suggests the fact that volumetric changes by themselves cannot explain the observed morphological variations in channel size and interconnectivity between Al2O3 and Al and that additional mechanisms may be operative. An important aspect to be considered in the reactive infiltration of solid oxide and the development of interconnected phases is the dihedral angle cos @ = l/2(y,,/ysJ [25]. This angle determines the ability of the liquid Al alloy to reactively infiltrate the silica particle itself and also permit wicking of liquid Al through the AlzOj channels to the reaction front. Any liquid alloy can completely penetrate oxide grain boundaries and form a continuous phase if @ = 0” or when the grain boundary energy (ysq) is twice the solid-liquid interfacial energy, ~~,. However, if CDis greater than zero the extent of penetration decreases,

4076

HANABE

and ASWATH:

SYNTHESIS

250

OF IN-SITU -I

200 a E f 150 ol 5 R Tii 100 ?I 3 IL 50

0 IOOpm 2OOpm 100pm (big treated) Characteristics of micro-composite reinforcements

Fig. 14. Preliminary flexural strength data for various Al composites in the as-processed condition. Note that the particle sizes are those of the original SiOl.

and when Q is larger than 90” then the alloy will no longer be continuous but remain as isolated pockets. Formation of alumina in the absence of Mg would probably lead to a situation where the orientation of alumina grains may be such that penetration of alumina grain boundaries becomes less feasible owing to grain boundary energies being lower than the Al-AlI03 interfacial energies. This could result in the microstructure observed in Fig. 10(b) and (c) where the alumina appears to be more continuous. Coarsening of alumina in the Mg treated preforms (Fig. 11) indicates the effect of Mg in controlling the ionic diffusion, although the exact mechanism of this phenomenon is not known at this time. Grain boundary penetration, however, may be due to the fact that Mg, being more electropositive, migrates to the AI-Alz03 interfaces and lowers the ysl. While the implications of such morphological variations in the reinforcement is under investigation, preliminary results have indicated that a Mg treated preform results in a composite with higher flexural strengths (Fig. 14). These results were obtained from as-processed composites without any heat treatment. CONCLUSIONS (1) Al matrix composites reinforced with Alz03/Al particles were prepared by reacting either Si02 or Mg + SiOz particulate preforms with liquid Al at 1075°C. The consistency of infiltration was better when Mg was used as an infiltration initiator. (2) It was found that the in-situ nucleation of the Al,O,/Al particles was a result of a single displacement reaction between Al and Si02 in the absence of Mg. In Mg containing preforms, more than one displacement reaction was observed. It is believed here that the negative free energy changes associated with the various displacement reactions aid in the infiltration of the Al melt into the preform.

REINFORCED

Al COMPOSlTES

(3) The A1203/Al particles that form as a result of a direct Al/SiO, interaction have a finer microstructure, with the alumina showing a tendency to be elongated in the growth direction, while those that form when a Mg treated preform is reacted with Al have a coarser morphology with both phases being interconnected. Grain coarsening and penetration by Al was also observed with the Mg containing preforms. Acknowledgements-Support provided by the Mechanical and Aerospace Engineering Department and the Materials Science and Engineering Program at the University of Texas at Arlington is gratefully acknowledged. Helpful discussion with Dr Sriram Rangarajan is greatly appreciated.

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