Composites Science and Technology 49 (1993) 1-12

ON THE CASTABILITY OF A1-Si/SiC PARTICLE- REINFORCED METAL-MATRIX COMPOSITES: FACTORS

AFFECTING FLUIDITY AND SOUNDNESS

A. M. Samuel, H. Liu & F. H. Samuel* Ddpartement des Sciences Appliqudes, Universit~ du Qudbec ?z Chicoutimi, Chicoutimi, Qudbec, Canada, G 7tt 2BI

(Received 24 January 1992; revised version received 24 September 1992; accepted 16 October 1992)

Abstract This paper presents results of studies carried out to determine the effect of melt and mold temperatures on the castability of four AI-Si/SiCo, ~ reinforced metal- matrix composites containing two levels of silicon (7 and lO%w 0 and two levels of SiC (10 and 20%vol.), in terms of the fluidity and soundness. These were assessed by monitoring the Al4C3 formation, SiC distribution, porosity volume fraction and melt cleanli- ness (in terms of the oxide content) in specimens prepared under different melting and casting condi- tions. The results show that a low silicon content coupled with a high SiC level accelerates the formation of AI4C3 which is detrimental to the fluidity and hence castability of the composite alloy. Increasing the silicon level from 7 to 10%wt improves the castability through a significant decrease in AI4C3 content. Increasing the SiC content from 10 to 20%vol. results in a relatively homogeneous distribution of the particles within the matrix, even at low cooling rates of about lOCs-1. The presence of oxides in an otherwise fluid composite melt considerably reduces the castability.

Keywords: metal-matrix composites, particulate com- posites, castability, fluidity, AI4C 3 formation, SiC distribution, porosity distribution, aluminum alloy matrix

INTRODUCTION

The last forty years have witnessed many research efforts undertaken in the development of metal-matrix composites as new engineering materials with better

*To whom correspondence should be addressed.

Composites Science and Technology 0266-3538/93/$06.00 1993 Elsevier Science Publishers Ltd.

properties, lightweight characteristics and competitive costs than those of traditional materials. Some of the most successful have been metal-matrix composites with an aluminum or aluminum alloy base reinforced by SiC powders or short alumina fibres. 1'2 Both solid- and liquid-phase processing methods have been used, the latter having the advantage that the fluidity of the metal allows for the use of a wide range of reinforcements and the capability of producing near-net-shape castings. In the case of particle- reinforced composites, the reinforcement particles influence the solidification process in various ways, through their settling in the melt, through chemical reactions with the matrix, and as a result of the particle pushing that occurs by the impingement of the solidification growth front on these particles. In addition, the particles can also act as nucleation sites under favourable conditions. 3 Basically, the reinforce- ment particles affect the fluidity and hence the castability of the composites in one or more of the above-mentioned ways.

The present work was carried out to investigate the interrelated aspects of castability, fluidity and soundness in AI-Si/SiCtp ) metal-matrix composites obtained from remelted Duralcan F3A (Al-7%wt Si) and F3S (Al-10%wt Si) composite alloys reinforced with 20% vol. SiC. For comparison, castings from F3A- and F3S-10%vol. SiC composite alloys were also studied. The effect of melt and mold temperatures on the castability were investigated by means of fluidity tests. Specimens obtained from these tests and from as-cast test bars were microstructurally examined for porosity content, oxide formation, SiC volume fraction and distribution, and aluminum carbide formation, in order to determine the manner and the extent to which these parameters affect the fluidity. The technical part of this investigation focuses on factors such as mold temperature and stirring mode/speed that can control the soundness of the final composite castings.

2 A .M. Samuel, H. Liu, F. H. Samuel

Table 1. Chemical composition of AI-Si/SiC~p~ metal-matrix composites

Composite Standard designation equivalent

Chemical composition of AI-Si matrix alloy

Si Fe Cu Mn Mg Ti Ni Sr

SiC (volume %)

F3A.20S A1-7Si/20SiC 7.17 0.11 0-005 0.004 0.41 0.10 0-004 F3S.20S AI-10Si/20SiC 9.16 0.15 0-012 0.004 0.62 0.09 0-004 F3A.10S A1-7Si/10SiC 7.45 0.13 0-019 0.003 0-40 0.10 -- F3S.10S AI-10Si/10SiC 9.46 0-14 0-004 0-003 0-58 0.11 --

0.013 18-83 0-013 19-98 0.014 11.30 0-013 11.34

EXPERIMENTAL DETAILS

The chemical compositions of the four composites used in the present work are given in Table 1. In the Duralcan system of designation, an F3A.xxS compos- ite contains 7%wt silicon, whereas an F3S.xxS composite contains 10%wt silicon. The SiC content (in %vol.) is given thereafter by the xxS, with 20S and 10S representing 20 and 10%vol. SiC, respectively. Thus an F3A.20S composite designates an Al- Si/SiC(p) composite that contains 7%wt silicon and 20%vol. SiC. Simplified designations are given in the 'Standard Equivalent' column (column 2) in Table 1 (see Ref. 2 for details). The 7 and 10%wt (for silicon) and 10 and 20%vol. (for SiC) are the standard specifications. However, the actual values can differ from ingot batch to ingot batch, as is observed in Table 1. Nevertheless, the composite is still referred to by its standard specification/designation. This procedure has been followed throughout in the present work. The composites were received in the form of 12 kg ingots. They were cut into small pieces (about 3 in thick) and melted in 7 kg capacity silicon carbide crucibles, using an electrical resistance heating furnace. When the material was melted, continuous mechanical stirring was initiated with a special graphite impeller designed in our laboratory.

Fluidity test measurements were carried out for melts held at various temperatures ranging from 725C to 850C, using a 4210 Ragone fluidity tester (supplied by George Fischer Foundry Systems, Inc., Michigan, USA), where the length of solidified tube indicated the corresponding fluidity. 4 Fluidity is measured in terms of the distance a stream of molten metal will flow in a spiral or channel before it is stopped by solidification. In the Ragone fluidity tester, the spiral channel is replaced by a tube (thin-walled stainless steel) bent at one end for the introduction of metal. The pressure head is provided by an adjustable partial vacuum to one end of the tube. The vacuum measured by a gauge in mm Hg can be converted to the effective metal head (in imperial units Hg) through a given standard formula to compare the mold filling abilities of different metals. This test is sensitive to metal properties and unlike the spiral test, is independent of

pouring speed. 4 In the present work, in one set of fluidity test measurements, the melts were heated to 850C, cooled down to 725C and then tested for their fluidity. In every case, the melt was stirred for about 20-30 min prior to testing.

Specimens for metallographic observations were taken from the solidified tubes obtained from the fluidity tests. In each case, both longitudinal and transverse sections were selected for examination, from the central part of the solidified tube lengths (rods). From the same rods, specimens were also prepared for X-ray analysis, to determine mainly the presence of A14C3 (as well as that of silicon and SiC) and in what amount. The X-ray diffraction analysis was carried out employing standard powder diffractometry procedures, using Cu K~ radiation. 5

The metallographic specimens were polished using a technique specially developed for such composites. The polished specimens were then examined optically for A14C3, porosity, oxide content, and SiC distribu- tion. The SiC and porosity volume fractions were measured by means of a Leco 2001 Image Analyser system attached to the optical microscope (Olympus PMG3). For the latter measurements, the number of fields ('field' representing an area of 1.61242 x 105 sq. microns at 200 magnification in the optical micro- scope) examined were chosen so as to cover the entire area of the specimen. That is to say, the entire surface of the specimen is traversed in a regular, systematic fashion and, depending upon the judgement of the observer, a suitable number of fields are chosen for which the SiC or porosity volume fraction is noted. The image analysis gives the average value of the measured volume fraction obtained over the number of fields taken. Using the point count method, 6 the A14C3 content was measured from optical micrographs (taken at 200x magnification). The ALC3 contents reported represent the average taken over at least four micrographs for each specimen.

For the second part of the investigation, test bars were cast from melts corresponding to the four different composites under different conditions. In each case the melt was heated up to 725C, stirred for about 20-30 min and then cast into a permanent Stahl mold maintained at various temperatures ranging

On the castability of AI-Si/SiC particle-reinforced metal-matrix composites 3

from 275 through to 450C. The stirring speed was maintained at 125 rpm in all cases except one, where the speed was increased to 155rpm for a melt prepared from F3S.20S alloy, to study the effect of the increased amount of oxide films expected to be taken up by the melt at the higher, more vigorous speed. It is to be mentioned here that the stirring time of 20-30min and speed of 125 rpm were arrived at following several casting runs made in our laboratory during the course of studies carried out on these composites. Stirring is recommended upon melting in the case of Duralcan composites to avoid SiC particle sedimentation and to ensure a more uniform distribution of these particles. 7 While the 125 rpm stirring was carefully executed to avoid breaking up of the oxide film on the surface of the melt, the purpose of the more vigorous 155 rpm stirring was deliberately to introduce oxide film segments into the melt. Specimens for metallographic observations were cut from the central portions of the test bars, polished as before, and examined for SiC, porosity and oxide contents and distribution. Dendrite arm spacings for these specimens were also measured using the Leco image analyser.

RESULTS AND DISCUSSION

Melt temperature studies

Fluidity measurements Figure 1 represents the fluidity curves obtained for the four composite types. The temperature of interest to the foundryman is 725C (or 750C), the temperature at which most casting operations are performed. In comparing the four composite types, one must consider both the silicon as well as the SiC contents of the respective composites. It is well established 8 that an increasing silicon content increases the fluidity of an AI-Si alloy; the fluidity attains a maximum value around 18% silicon. Thus, in going from an AI-7Si/10/20 SiC (F3A) to an AI-10Si/10/20 SiC (F3S) composite, it is expected that the F3S composite alloy containing a higher amount of silicon would exhibit greater fluidity. This is found to be true for both F3S.10S and F3S.20S composites as compared to their F3A counterparts.

The second factor affecting the fluidity of these composites is their SiC content. From Fig. 1 it is seen that the 20%vol. SiC-containing composites display lower fluidities than the ones containing 10%vol. SiC. Also, in the case of the F3A.20S composite, when the melt is cooled from 850C to 725C, the fluidity is considerably decreased; the metal flow becomes very viscous, to the extent that the metal is almost solid at about 700C, well above the melting point of about 610C for this composite.

50

ao v ~ . . . . . . . .%/ E]"" ~ I*1 U]'"'"" I"1"~'%

~ 25 ~ ,.."~ "'.\ F ,,/" "\ /. \ . f \, ~ 20 ~.." .'.... .1"'~

15 / t

. / . / "

I " / F3S10S F3A10S

10 i~ . f f ~ 0 [] F3$20S F3A20S

B , , , i . . . . i . . . . i . . . . i . . . . i . . . . i . . . . i . . . . i . . . .

650 675 700 725 750 775 800 825 850 875

Melt temperature (o C)

Fig. 1. Fluidity curves obtained for the four AI-Si/SiC composite types studied in the present investigation. The arrows (shown here in one case only) depict the heating (675 to 850C) through cooling (850 down to 725C) cycle. Points A, B and D, E are referred to in the text with respect to

Figs 2 and 6.

The dramatic decrease in fluidity is attributed to the formation of A14C3 due to the reaction between the SiC reinforcement and the metal matrix. 9'1 SiC reacts with the molten aluminum matrix according to the reaction

4 AI + 3 SiC---~ A14C3 -I- 3 Si (1)

The A14C3 forms a stable compound in the melt and the excess silicon alloys in the matrix and increases the silicon level of the alloy. The products of the reaction bring about large increases in the viscosity of the melt that significantly affect the melt fluidity.

Maximum values of fluidity (as measured by the solidified tube length of a Ragone fluidity tester operating at 200 mm Hg) for the four composites are seen to lie between 735 and 775C for the F3A alloys and between 775 and 800C for the F3S alloys, the actual value depending on the SiC content. Thereaf- ter, the values either stay steady or begin to drop, indicating the start of A14C3 formation.

Lloyd & Dewing u and other workers 12'13 have studied the A14C3 reaction with respect to melt temperature, holding time and silicon level in the matrix alloy. Theoretically, silicon levels above 8%at. are required to prevent A14C3 formation. However, practical observations show that the kinetics of the A14C3 formation must also be taken into considera- tion. From a study of holding times and melt temperatures, it is found that AI4C 3 forms rapidly at temperatures above 790 and 895C in the F3A and F3S composite alloys, respectively. The reaction also occurs below these temperatures but at much lower rates. We shall refer to this point in more detail in the

4 A .M. Samuel, H. Liu, F. H. Samuel

section on the X-ray diffraction analysis. For the holding times involved in the present studies (20-30 min), no significant AI4C 3 should be formed, except at high temperatures and in composites containing higher volume fractions of SiC.14 This steep drop in the fluidity in the case of the 20%vol. SiC-reinforced composites at higher temperatures is a result of an effective amount of AI4C 3 being formed. For the 10%vol. SiC-reinforced composites, however, the fluidity remains relatively constant through the range 750-850C, once maximum fluidity is achieved, indicating that A14C3 formation is taking place, but at a very low rate. It is thus a combined effect of the silicon level and the SiC content that results in maximum fluidity in the F3S.10S alloy and minimum fluidity in the F3A.20S alloy, Fig. 1.

Lloyd ]4 has studied the fluidity of A356 and A356 alloy containing 10 and 20%vol. SiC at different temperatures. Two observations may be made from a comparison of his results and those shown in Fig. 1: (i) the trend of fluidity versus melt temperature is essentially the same, showing thereby that (ii) the method of determining the fluidity (spiral versus Ragone) does not affect the end results in any way. Lloyd carried out measurements up to 750C, which is the practical temperature to which composite melts are normally heated. In the present work, the heating was extended to 850C to observe the effect of superheating on the A14C3 formation.

The SiC-AI reaction is irreversible and once A14C3 formation is initiated at higher temperatures, lowering the melt temperature will only increase the viscosity and decrease the fluidity of the melt due to the bridging effect of the A14C3 and silicon particle formation on the SiC particles. Also, the time interval between heating and cooling down to the lower melt temperature (about 2 h) will permit further reaction to take place. Thus, the fluidity is much lower for such melts (heating ~ cooling). Figure 2 shows two castings obtained from F3A.20S composite melts represented by points A and B in Fig. 1. The contrast in surface topography of the two castings, resulting from the difference in fluidities of their respective melts, is easily observed: cast A (obtained from a much more fluid melt) displays a smooth surface indicative of high fluidity compared to the rather wrinkled surface of cast B obtained from an extremely viscous melt.

Metallography Details regarding the preparation of metallographic specimens used for various purposes have already been given in the Experimental Details section. The specimens prepared from the fluidity test (solidified tube) rods were used to determine quantitatively the amount of AI4C3 formed in the four composites as a function of melt temperature and casting conditions.

Fig. 2. Castings obtained from F3A.20S composite melt corresponding to points A (725C) and B (850-->725C, cooled) in Fig. 1. Note the contrast in surface topography of the two castings, resulting from differences in fluidity of

their respective melts.

The A14C3 volume fractions determined for these specimens are displayed in Fig. 3. It is evident that increasing the silicon content to 10%wt (i.e. F3A versus F3S composites) results in a significant reduction in the amount of AI4C3 formed. Increasing the SiC content in the case of F3A alloy aids in increasing the reaction rate and AI4C3 formation (except for 725C), whereas in the case of the F3S composite, the reaction is not affected as much. This may be understood from considerations of the theoretical silicon levels required to prevent AI4C3 formation: approximately 12%at. silicon is required for melt temperatures of 850C. In the case of the F3A alloy, this level is always much below the required value, even with the additional silicon formed from the SiC dissociation. For the F3S alloy,

4

o ~3 q

2

. . . . . . . . . ,850H

F3A, 725 C

~ F 3 S , 725 C

S 10 15 20 SIC ~t lcu l l~ vo l%

Fig. 3. ALC3 volume fractions obtained from fluidity test specimens for F3A (AI-7Si) and F3S (Al-10Si) composites, as a function of their SiC content (H: heating; C: cooling).

On the castability of AI-Si/SiC particle-reinforced metal-matrix composites 5

platelets, bridging the SiC particles at/between which the reaction is taking place. Figure 4b shows the microstructure obtained from the sample correspond- ing to cast B in Fig. 2. The bridging effect of the A14C3 particles is seen in this case and explains the sluggish characteristics of the cast.

Fig. 4a. Optical micrograph obtained from F3A.20S composite at 725C melt temperature. Arrows indicate the initiation of the ALC3 reaction at the very edges of the SiC

particles (1500x).

however, which already contains about 10% silicon, this level is expected to be reached much sooner. This explains the difference in the slopes observed for the two composite types. Upon cooling from 850 to 750C over 2 h (giving more time for the reaction to take place), the volume fractions are seen to increase slightly, which suggests that the reaction has reached its equilibrium state with respect to time. These results also indicate that the reaction is more temperature-dependent than time-dependent.

Figure 4a is the optical micrograph obtained from F3A.20S alloy at 725C melt temperature. The initiation of the AI4C3 reaction can be observed at the edges of the SiC particles as indicated by the arrows. As the melt temperature is increased, the reaction proceeds faster, the AI4C3 now clearly visible as black

Fig. 4b. Optical micrograph obtained from F3A.20S composite at 850--* 725C cooled melt temperature (cast B, Fig. 2). The AI4C3 is now visible as black platelets bridging

the SiC particles (1500x).

X-ray diffraction analysis X-ray diffraction patterns were recorded for specim- ens prepared from the fluidity test rods obtained from the four composites at 725 and 850C melt temperatures, using Cu K~ radiation.

The X-ray diffraction pattern for F3A.20S compos- ite taken from the 850C---725C (cooling) melt temperature is displayed in Fig. 5, while the patterns obtained for this composite from 725 and 850C as well as that shown in Fig. 5 are analyzed in Table 2. These were chosen as they exhibited the most interesting features. The other patterns followed a similar trend.

The main lines that appear in these patterns are those of aluminum, silicon, SiC and AI4C3, as depicted in Table 2. The relative intensities of the peaks are measured taking the Aim line (at a d value of 2.338 A) as having 100% intensity. Thus, the values recorded in the I/I1 column represent percentage intensities with respect to this line.

The main silicon lines or peaks are (111), (220), (311) and (331) with I/Ii values of 100%, 60%, 35% and 13%, respectively for the standard pattern.

In the case of SiC, ~5 there are two forms that occur: the low temperature cubic form or fl SiC and the high temperature hexagonal (or rhombohedral) form or tr SiC that exhibits several polytypes, chiefly 4H, 6H and 15R, the fl SiC form being the 3C polytype. (Polytypism is a phenomenon that may be defined as a one-dimensional case of polymorphism and polytypes are the various crystallographic structures exhibited by the same material, with two of the three dimensions the same, the variation along the third resulting in the different polytypic structures.) The standard patterns corresponding to these four most commonly observed polytypes have several lines that overlap (i.e. they occur at the same d values), see Table 2. Ruska et al. 16 have shown that it is possible to calculate/determine the polytype fractions directly from peak intensities. No quantitative estimates were attempted in the present case, as even when pure materials are used, experimental determination of these polytypes by X-ray diffraction is quite a problem. 17 However, based on the fact that both a~ and fl SiC phases were observed, all polytypes corresponding to a particular d value are quoted in column 1 of Table 2.

The hexagonal AI4C3 phase exhibits six main lines corresponding to (110), (012), (107), (00-12), (1 1.12) and (1 0-10) planes, with relative intensities of 100%,

6 A .M. Samuel, H. Liu, F. H. Samuel

0 0 v

U

fO t II ~ _ "~. -

~ ~ .-.~ o~ , A II r_ u

"~ , : II I1 , mO~ u ,'7- ,.., f i f ] / ~u h

0 U m ~ . _ tO

1

5

U

5

80 70 60 50 40 30

20

Fig. 5. X-ray diffraction pattern recorded for F3A.20S composite specimen taken from 850--* 725C melt temperature, using Cu K~ radiation.

60%, 60%, 20%, 20% and 10%, respectively for the standard pattern.

Analyzing Table 2, the following observations may be made. At 725C melt temperature, no lines for A14C3 are observed, indicating that at this temperature no reaction is taking place between the SiC and the molten aluminum matrix.

At 850C, the reaction is in progress and all the main mlaC3 reflections are observed albeit with very small intensities. An interesting observation with respect to these reflections is that they are always spread over a range of d values, around the main value at which the reflection is supposed to occur (see the thick black arrows in Fig. 5). This is probably due to the fact that the AI4C3, in the process of formation, occurs in various non-stoichiometric forms before the phase stabilizes to the stoichiometric AI4C3 form. With reference to the A14C3 peak intensity versus holding time curves shown in Fig. 2 of the paper by Lloyd, ~4 the short holding time (30min) used in the present work falls in the region where the reaction is still in progress and is only expected to stabilize much later. This results in the observed spread in d values for A14C3, indicating the various stages of its formation.

When the melt is cooled down to 725C from 850C, essentially the same features are observed. However, a slight increase in the peak intensities is noted, which is the result of the increase in reaction over the extra time of about 2 h which the melt requires to cool down to 725C. That the SiC-AI reaction is taking place is also affirmed by the corresponding increase in the intensities of the silicon lines on progressing from 725C to the last case.

With regard to the SiC phase, as mentioned earlier, both te and fl phases were observed, notably through the 2.628 ~ and 1-67/~ d values for the tr SiC phase for which no corresponding fl SiC line exists. Similarly, the line at 2.17/~ for fl SiC indicates the presence of this phase.

The other d values can correspond to both o: and fl phases. Thus, all possible polytypes occurring at these d values have been mentioned in Table 2, column 1, since no quantitative estimates were attempted. However, it is the authors' opinion that both phases could and apparently do exist, based on the facts concerning SiC polytypes and polytypes in general. It has been reported that both fl---~tr and a:--~fl transformations take place in the presence of impurity phases. TM In particular, the presence of non-soluble phases like iron leads to the c~--~fl transition. Normally, the cubic fl SiC is known to be a low-temperature phase, stable below 2200C, above which only the a: phase is supposed to be stable. But under specific conditions of pressure and temperature, the treatment of powdered SiC in inert atmospheres (e.g. nitrogen or argon) can result in an a:--~fl transformation or vice versa . 19 The polytypes of a substance have very nearly the same internal energies and can thus occur together under identical conditions of growth.

While it is well known that stacking faults play an important role in polytype formation, other factors like impurities, pressure, temperature, dislocation content and supersaturation can also affect the phenomenon. ~5 The SiC-particle-reinforced compos- ites used in the present work were prepared using a proprietary process to incorporate the SiC particles

On the castability of Al-Si/SiC particle-reinforced metal-matrix composites

Table 2. Analysis of X-ray diffraction patterns obtained for F3A.20S composite from fluidity test samples taken at various melt temperatures

Melt temperature (C) Phase and Standard corresponding pattern hkl reflection values 725 850 850--* 725 (cooled)

Phase hkl d(A ) I/It d(A ) I/It b d(A ) I/lt b d(A ) I / I t b

Si 111 3.138 100 3.144 9.4 Al4C3 012 2-80 60

a~ SiC: 6H " 2-61 75 2.628 8.7 ot SiC: 4H 004 2-52 55 2-52 17-9

6H a 2.51 85 15R " 2.51 80

fl SiC: 3C 111 2.51 100 A1 111 2.338 100 2.338 100

AI4C3 107 2.23 60 a~ SiC:6H ~ 2.19 50

15R ~ 2.19 10 fl SiC:3C 200 2.17 20 2.179 3-4

A14C3 00"12 2"07 20 AI 200 2-024 47 2"026 61 "3 Si 220 1"920 60 1"919 5"3

AI4C3 10.10 1.88 10 tr SiC : 6H " 1.67 35

A14C3 110 1"66 100 Si 311 1"638 35 1"637 2"7

a~ SiC: 4H 110 85 1-541 12.1 6H a 100 15R " 1-54 100

fl SiC : 3C 220 63 AI 220 1-431 22 1.433 24.7

a~ SiC:4H 114 100 1.315 6.9 6H a 1.31 100 15R a 90

fl-SiC : 3C 311 50 A14C3 11.12 1.295 20

a~ SiC : 4H 107 1-26 45 1.258 0.6 6H a 1.25 35 15R a 1-26 20

fl SiC: 3C 222 1.26 5 Si 331 1.246 13 1.245 0-8 AI 311 1-221 24 1.221 28.1 AI 222 1-169 7 1.168 4.8

3.138 9.5 3.143 11-7 2.823 1.6 2-817 1-3 2-623 6.0 2.626 8.3 2.524 14.5 2.517 1-6 2.513 14.9

2"338 100 2.249 1.6

2.171 4-8 2.106-2.096 1"6-0-8

2-025 65.32 1.924 8.5 1.919 9-1

1.871-1.857 2.4-2.3 1.679 1.4 1.669 1.8 1.637 6.5 1"542 7.3 1.540 8.9

2.338 100 2.250-2.237 2.1-2.7

2.192 5.0 2" 185 6.0 2.176 7"5

2-089-2.084 1-7-1.4 2-027 67"9 1-919 8-3

1-876-1-849 1"7-2-9&4.2 1.678 1.9

1.667-1.655 1-9-1 "3 1.638 6.7 1.542 1-6

1-431 22-2 1-433 34.6 1-314 8"9 1.315 17-5

1.290-1-288 2-02 1.291-1.259 2.5-2.1 1"259 2" 1

1-246-1.242 2.8-2.4 1.247 2.71 1.220 42-7 1.222 46.3 1-170 7-5 1-170 10.4

For hkl values for 6H and 15R please see Thibault, Am. Min., 29 (1944) hi/It values are given in % taking Alln as 100%

into molten aluminum. In view of this, it is not possible to identify the phases present unless a detailed X-ray analysis is carried out. The authors are inclined to think, therefore, that both phases could be present and that changes could take place from time to time depending on the conditions of the melt and any one of the above mentioned factors coming into play. From the relative decrease in intensity for the oc SiC lines at 2.52 and 1.54 A, however, it is very likely that the SiC-AI reaction is taking place along these planes of the ct SiC phase.

The above observations represent the preliminary results we have obtained in regard to the X-ray diffraction studies being carried out in our laboratory

on these composites. A more detailed study is underway and results will be reported shortly.

Mold temperature studies In the second part of this work, the effect of the mold temperature on the castability of the four composites was investigated. In each case, the composites were melted to 725C (+5C), stirred for 20-30min at 125 rpm and then poured into permanent Stahl molds with mold temperatures kept at 275,300, 350, 400 and 450C, respectively. In one case, for the F3S.20S alloy, to study the effect of increased oxide content in the melt, the melt was stirred more vigorously at

8 A.M. Samuel, H. Liu, F. H. Samuel

Fig. 6a. Casts obtained for F3A.20S composite at different mold temperatures: Cast A: 275C; Cast B: 300C; Cast C: 350C; Cast D: 400C; Cast E: 450C; Cast F: A356 base

alloy, 450C.

155 rpm, to break the oxide film on the surface and cause it to be drawn into the melt.

Casting soundness Figure 6a shows the casts obtained for F3A.20S composite corresponding to different mold tempera- tures. Case F corresponds to A356 base alloy (metal matrix used for the present composites) and a mold temperature of 450C. It can be seen from the figure that a continuous casting below 300C is not obtainable (note cast A at 275C mold temperature). A 300C mold temperature (cast B) indicates insufficient fluidity to produce a solid test bar. The arrow shows the gap in the test bar between the upper and lower sections. The riser is first observed at 400C mold temperature (cast D). Comparison of casts E

and F shows the effect of the presence of 20%vol. SiC on the castability of F3A.20S composite as compared to a cast obtained from the base alloy A356, for the same mold temperature of 450C.

Figure 6b shows the casts obtained under similar conditions for the F3S.20S composite. In this case, the riser is first observed at 300C mold temperature, its size increasing with increasing mold temperature. Cast E represents the cast obtained at 450C after the melt was stirred vigorously, and shows a reduction in the riser size compared to cast D obtained at the same mold temperature. The reduction results from the oxide films introduced into the melt due to the vigorous stirring. The main observations made from Figs 6a and 6b are:

--For the F3A composite, optimum casting is obtained when the mold temperature lies between 400 and 450C, whereas for the F3S composite, casting can be done below 400C.

--Introduction of oxides (through vigorous stirring) considerably reduces the fluidity of the melt (see points D and E in Fig. 1 and corresponding casts D and E in Fig. 6b).

Void formation The void volume fractions (representing mainly gas and shrinkage cavities in the present case--for details on voids and porosity formation see Ref. 20) for F3A.20S and F3S.20S composites measured for metallographic specimens obtained from the central portions of the test bars shown in the previous figure are displayed in Fig. 7. Open symbols represent measurements taken over 75 fields (see Experimental Details section for description of 'field'), whereas closed symbols are measurements taken over 25 fields.

Fig. 6h. Casts obtained for F3S.20S composite at different mold temperatures: Cast A: 300C; Cast B: 350C; Cast C: 400C; Cast D: 450C; Cast E: 450C, containing oxide

contamination from vigorous stirring of the melt.

1.5

P

05

%0 ~oo

I F~s ~0~ ~0s F~ 1

30O 350 400 450

Mold temperature (* C)

Fig. 7. Void volume fractions for F3A.20S and F3S.20S composites obtained from the test bars shown in Fig. 6 at different mold temperatures. Open and closed symbols represent measurements taken for the same specimen over 75 and 25 fields, respectively. See text for details regarding

fields.

On the castability of AI-Si / SiC particle-reinforced metal-matrix composites 9

(3

0

m o

o

O 0 O0 O DO O O

O D 0 o O00 io [] 0

~mm_m~i~i . . .m m ~INNIg IQ I . I .

~d num~r

Fig. 8. Porosity distribution for F3S.20S composite pre- pared from casts D and E of Fig. 6b. See text for details.

Each set of observations is taken at 200x magnifica- tion. The purpose in taking the two sets of observations was to determine the nature of the overall porosity distribution across the specimen surface from a comparison between fewer (25) and more (75) readings taken in each case. It is seen that the F3S composites have a much lower void volume fraction as compared to the F3A composites. Also, increasing the mold temperature improves the soundness of the casting. Typical void patterns obtained for these composites are shown in Fig. 9 (F3A.20S) and Fig. 10 (F3S.20S). It is interesting to note that increasing the number of fields or the area observed for the measurement of the void volume fraction (from 25 to 75) results in a much steeper slope for the A composite on proceeding from 350 to 450C, whereas for the S composite, the slope is almost flat, indicating that the S alloy gives a much better soundness.

Figure 8 shows the porosity distribution in the samples prepared from casts D and E of Fig. 6b, for the F3S.20S composite. Open square symbols represent readings taken from sample E, while closed symbols represent sample D. The field numbers represent fields taken consecutively, traversing from one edge to the centre to the opposite edge of the spherical cross-section of each metallographic speci- men. Localization of oxides in the case of sample E results in the formation of large voids (represented by the random high porosity readings), of the order of a volume fraction of 0.043. The D sample, on the other hand, representing relatively clean material, displays a maximum localized void volume fraction of 0-003. Also, the distribution of voids is much more uniform in this case. Figures 9(a) and 9(b) are optical micrographs taken from F3A.20S composite cast at 300 and 450C mold temperatures, respectively, comparing the void sizes obtained in the two cases. The increase in mold temperature results in a decrease in the void size by about 7-10 times. Figures 10(a) and 10(b) show the difference in void distribution

(a)

(b)

Fig. 9. Optical micrographs taken from F3A.20S composite cast at (a) 300 and (b) 4500C mold temperatures (50), showing the voids (large black areas) obtained in each case.

obtained for F3S.20S for clean (cast D, Fig. 6b) versus oxide-contaminated (cast E, Fig. 6b) material. In Fig. 10(b), the oxide films are always seen to be associated with voids (see inset, taken at a higher magnification). The void size of F3A versus F3S may also be compared from Figs 9(b) and 10(b).

SiC distribution It has been established 2 that the microstructure of an alloy is dependent on the solidification rate or the cooling rate of the casting. This effect becomes important in the case of composites because the reinforcement particle distribution is affected by the dendrite arm spacing (DAS) of the casting. During solidification the SiC particles are rejected at, and pushed ahead of, the solidification front and become trapped by converging dendrite arms in the intercellu- lar regions. Thus the cooling rate during solidification and therefore the DAS influences the distribution of the SiC in the casting obtained. The DAS-cooling

lO A. M. Samuel, H. Liu, F. H. Samuel

(a)

ql

<

I

0 A

Mold t~npor~u~ (e)

Fig. 11. Dendrite arm spacings obtained for the four composite alloys taken at 300 and 450*(:7 mold temperatures. Open and closed symbols represent average readings taken from the centre and edge, respectively, from each

metallographic specimen. See text for details.

(b)

Fig. 10. Optical micrographs comparing the void distribu- tion obtained in F3S.20S specimens prepared from (a) clean (cast D, Fig. 6b) and (b) oxide-contaminated (cast E, Fig. 6b) material (50x). The inset in (a) shows how the oxide

films always appear associated with voids (200).

rate relation is given by the equation

d = Ar a (2)

where d is the DAS, r is the cooling rate, and A and a are constants. 11 In the present work, a was determined to be -0.33 from thermal analysis. At cooling rates typically around 20Cs -], the DAS approximately equals the SiC particle size and hence there is very little particle pushing during solidification. As the cooling rate decreases, the DAS becomes much larger, resulting in a more clustered distribution of the SiC particles. 14

Figure 11 shows the dendrite arm spacing for the four composite alloys taken from specimens obtained from 300 and 450C mold temperatures. Open and closed symbols represent average readings taken from the centre and the edge, respectively, for each specimen. Two conclusions can be drawn from the figure: (i) for a given mold temperature, the A alloys

(AI-7%wt Si), due to their low silicon content, display a coarser DAS as compared to the S alloys (Al-10%wtSi); (ii) the increase in SiC content (20%vol. SiC) in the 20S alloys (F3A.20S and F3S.20S) results in a finer dendrite arm spacing compared to that displayed by the 10S alloys. Several models have been proposed to explain particle entrapment and rejection by a solidifying interface, 2 based on differences in thermal conductivity 22 or thermal diffusivity 23 and energy considerations. 24-26.

Figures 12(a) and (b) show the respective microstructures obtained at the edges of F3S.10S and F3S.20S specimens solidified at a cooling rate of 10-15C s -~ and mold temperature 450C. As can be seen, a wide SiC denuded zone (about 250-300/zm in width) is present in the 10S composite, whereas none exists in the case of the 20S specimen. These effects may be related to the difference in DAS obtained for the two composites, resulting from the difference in their SiC contents. The distribution of the SiC volume fraction measured from the edge to the centre to the other edge for the above specimens is displayed in Fig. 13. It can be seen that the 20S alloy exhibits a more even distribution of the SiC over the specimen cross-section compared to the 10S alloy, due to the absence of the denuded zone. The tendency for SiC segregation towards the centre of the specimen for the 10S alloy is also observed.

SUMMARY

The present work was performed to determine the effect of melt and mold temperatures on the castability of four aluminum metal-matrix composites containing two levels of silicon and two levels of SiC in terms of the fluidity and soundness. These were assessed through a monitoring of the AI4C 3 formation, SiC distribution, porosity volume fraction and oxides

On the castability of Al-Si/SiC particle-reinforced metal-matrix composites 11

(a)

(b)

Fig. 12. Optical micrographs showing the edges of (a) F3S.10S and (b) F3S.20S metallographic specimens, obtained at 450C mold temperature (100x). Compare the wide SiC-free zone in (a) with the relatively uniform SiC

distribution in (b).

~ 15

O ~o

[3 D [] O D (3 rl R ~ D [3 D D O

[]

.

I I I I I 5 10 15 20 25

Field number

Fig. 13. SiC distribution across the specimen cross-section for the F3S.10S and F3S.20S metallographic specimens

shown in Fig. 12. See text for details.

present for specimens prepared under different melting and casting conditions. The results may be summarized as follows:

1. It has been found that a low silicon content (7%wt) coupled with a high SiC level (20%vol.) accelerates the formation of AI4C3 to about 6%vol. at 850C, which is detrimental to the fluidity, and hence, castability of the composite alloy.

2. Increasing the silicon content to 10%wt leads to a significant reduction in the amount of ALC3 (about 2%vol.), even with a 20%vol. SiC content.

3. The F3A.20S composite could not be cast at temperatures below 400C because of the low fluidity, which was not found to be the case for the F3S.20S alloy.

4. For the F3S (10%wt Si) composites, increasing the SiC content to 20%vol. eliminates the presence of the denuded (SiC-free) zone at the specimen periphery. This leads to an homoge- neous distribution of the SiC particles across the casting cross-section even at low cooling rates of 10oc s -1.

5. The improved fluidity of the F3S.20S composite caused by its high silicon content is also reflected by the low porosity content exhibited by this alloy as compared to the higher level observed in the F3A.20S composite.

6. The presence of oxides to a large extent in an otherwise fluid composite melt considerably reduces the castability as was observed for the F3S.20S alloy in the present work.

ACKNOWLEDGMENTS

The authors would like to thank Dr R. Provencher of Arvida R & D Centre, Alcan International Limited for his suggestions concerning this work and the Dural Aluminum Composite Corporation for supplying the SiC-reinforced composites. They would also like to thank Dr D. J. Lloyd of Kingston R & D Centre, Alcan International Limited for helpful discussions during the course of this work. The financial support received from the Natural Sciences and Engineering Research Council of Canada, the Fondation de l'Universit6 du Qu6bec ~ Chicoutimi and the Soci6t6 d'61ectrolyse et de chimie Alcan (SECAL) is gratefully acknowledged.

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12 A. M. Samuel, H. Liu, F. H. Samuel

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