10 Kent Ridge Crescent, Singapore 0511. (Received September 18, 1995; Accepted February 5, 1996;). (Refereed). ABSTRACT. Sintering temperature is one of ...
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Materials Research Bulletin, Vol. 31, No. 5, pp. 453-464, 1996 Copyright O 1996 Elsevier Science Ltd Printed in the USA. All fights reserved 0025-5408/96 $15.00 + .00
PII S0025-5408(96)00024-4
INFLUENCE OF SINTERING PROCESS ON THE MECHANICAL PROPERTY AND MICROSTRUCTURE OF BALL MILLED COMPOSITE COMPACTS
L. Lti, M.O. Lai and G. Li Department of Mechanical & Production Engineering National University of Singapore 10 Kent Ridge Crescent, Singapore 0511
(Received September 18, 1995; Accepted February 5, 1996;) (Refereed)
ABSTRACT Sintering temperature is one of the critical parameters in the sintering of composites. It was found that the density of a sintered metal matrix composite compact at relatively lower sintering temperature could only be slightly changed. Although compaction of the matrix metal (A1-4.5wt.%Cu) could be densified from an original 70% to 96% of its theoretical density by liquid phase sintering (LPS), density of a composite compact could not be increased at the same sintering temperature. The composite compacts could be densified up to only about 91% of their theoretical density if 21.6% of liquid phase sintering was employed. The temperature employed was much higher than that used in the liquid phase sintering of its base metal. Ktc measurement of the compacts revealed an increase of about 50 % if high temperature sintering was used. Fracture surface analysis showed an incompletely sintered structure due mainly to high viscosity of the sintered particles. It is, therefore, suggested that the temperature for sintering metal matrix composites should be increased in order to lower the viscosity. A simple model is here proposed to predict the amount of the liquid phase needed whereby the sintering temperature can be chosen accordingly. KEYWORDS: A. composites, D. mechanical properties, D. diffusion. INTRODUCTION The incorporation of SiC particulates in AI or its alloy matrices results in increase in the values of the elastic modulus and yield strength by an Orowan strengthening effect and an 453
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increase in dislocation density of the A1 matrix due to the difference in the coefficient of thermal expansion between the SiC and the matrix (1). The physical and mechanical properties of MMCs can therefore be tailored to specific needs in a range of structural and nonstructural applications. The major difficulty in producing A1 matrix composites using foundry route, however, is to disperse the fine reinforcement such as SiC particulates. To overcome the inhomogeneous distribution of reinforcement, powder metallurgy has been widely used (2, 3). In recent years, novel technique such as mechanical alloying (MA) has been investigated in the fabrication of MMCs (4, 5). However, it has been found that the sintering process used for ball milled (BM) or mechanical alloyed compacts is one of the problems to be overcome. In general, hardness of the MA or BM powder particles is increased due to work hardening making cold compact more difficult. Sintering a compacted sample can generally be carried out when it is heated to a temperature of approximately over one-half of its absolute melting temperature. However, the sintering process can be improved if the compact is sintered in a partial liquid phase. Densification occurs by the combination of wetting, liquid flow, particle rearrangement and solution reprecipitation in the partial liquid particles (6). This sintering is generally very rapid in the presence of a liquid phase. In liquid phase sintering (LPS), rapid densification takes place followed by a continual decrease in densification rate (7). Above 96% of theoretical density can be attained within 15 rain of sintering if a suitable sintering temperature is selected (8). There is an optimal sintering temperature range above the solidus temperature where a high sintered density, low degree of microstructural coarsening and minimal compact slumping can be obtained. Hence, sintering temperature is of fundamental importance (9, 10). Besides the influence of sintering temperature, green density and initial particle size can also have profound effects on densification. Coarse particle size and high green density offset the favorable effects of the liquid. Enhancement of LPS is generally the result of increased driving force through physical or chemical treatments. It is attributed to one or more changes in the fundamental material properties resulting from a special treatment (11). The strongest effects are those associated with changes in the interfacial properties (higher surface energy to lower grainboundary energy). Sintering time has a nominal effect on densification since LPS kinetics is more sensitive to temperature than to sintering time. For solution and reprecipitation shrinkage, the transport path is through the liquid phase. There exists three stages in LPS: the first stage is liquid flow, the second, solution and reprecipitation, and the final, solidphase sintering. At the final stage where solid-state sintering is dominant, the rate of densification is quite low. Therefore, densification is mainly contributed by the first two stages. The maximum liquid phase in LPS is approximately 30 vol.% depending upon the materials used. Beyond this value, little effect can be observed. Too high a volume fraction of liquid phase may lead to slump of the sintered parts. Generally, prolonged sintering is not advantageous because of grain growth at high temperature (12). The objective of the present study is to investigate the influence of BM time and sintering temperature on the fracture toughness of ball milled AI metal matrix composite compacts. THEORETICAL CONSIDERATION Shrinkage and densification of a compact are driven by reduction in surface area during sintering. Two important factors in solid-state sintering have to be considered (13):
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(a) Fine particles sinter much faster than coarse ones because the surface area of the former is higher and the diffusion distances is much smaller. (b) The rate of sintering varies with temperature in exactly the same way as the diffusion coefficient. The rate of densification can therefore be represented by: dp
C
dt
sn
exp(-Q/RT)
(I)
where p is the density, s, the particle size, C and n, constants, Q, the activation energy for sintering, R, gas constant and T, absolute temperature, n is typically about 3, while Q is usually equal to the activation energy for grain boundary diffusion. From Eq.[1], it can be seen that the density of a compact can be increased if sintering temperature is raised. For sintering MMCs, because the composition of metal matrix material is not strongly influenced by the addition of SiC particulates, only the matrix metal is considered in LPS. FIG. 1 schematically shows a phase diagram. If a compact with composition Xa is heated up to temperature Ts, partial liquid phase starts to form. The amount of liquid phase is increased if the temperature of the compact is above T,. The volume fraction of solid phase at any temperature in the two phase region can be calculated according to the following equation (6):
x,-x x,-x
(2)
Heterogeneous sites are preferred for liquid nucleation. Therefore, grain boundaries and interdendritic regions will be the first to exhibit liquid formation since the liquid must form a film surrounding the solid phase, wetting is the initial requirement. Secondly, the liquid must have a solubility for the solid (14). Diffusion for the solid atoms dissolved into the liquid should be high enough to ensure a rapid sintering. During heating, the liquid phase will initially be insufficient to coat all of the grain boundaries. Hence, to maintain the minimum boundary thickness, a fractional coverage Fc of the grain boundary is required.
Tm
L
Ti ¢D
S
X, X,
X~ Composition
FIG. 1 A schematic binary phase diagram.
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Accordingly, some of the liquid is located at the interior of the grain and may not play a useful role in particle decomposition. According to German (6), the fractional solid content can be expressed by: = 1 -
[(g, f F ) / 2 ( 1
-
F,)g,~]
(3)
where gs is 26.78, gv, 11.31, Fc, the fractional coverage of liquid, Fj, the ratio of the interior liquid to the total liquid, varying from 0 to 1, 6, grain boundary film thickness, and G, the tetrakaidecahedron edge length. There is a critical volume fraction of solid above which the mixture has infinite viscosity. When the solid fraction fs greater than this maximum value, sufficient solid fraction of solid particles gives a rigid structure. If the solid fraction decreases, viscosity will decrease. From a particle point of view, a high sintered density is favoured by increase in the liquid-vapor surface energy, sintering time and temperature, alloying level and liquidus-solidus separation (15). To get high density, 89% of the grain boundaries o f a particle must be coated with the liquid (6). The sintered density ps can be evaluated from: p, = pg/(1 -
AL/Lo) s
(4)
where pg is the green density, AL/Lo, shrinkage during isothermal sintering which can be written as: A L / L o = 0 . 7 5 ? L v t / (Drl)
(5)
where VLv is liquid-vapour surface energy, D, particle size, rl, viscosity of solid-liquid particles and t, sintering time. EXPERIMENTAL PROCEDURE
Materials and Bali Milling Process. Elemental powders of AI, Cu and reinforcement of SiC particulate with nominal composition of AI-4.5wt.%Cu/15wt.%SiC were blended or ball milled under ambient conditions in an ordinary horizontal ball mill for durations varying from 3.6 ks to 360 ks. The mean particle sizes of A1, Cu and SiC particulate used in this study were 120, 50 and 5 ~tm respectively. 400 g of powders were milled at a rotational speed of 113 rpm to form composite powders at a ball to powder weight ratio of 6.7:1. To minimize contamination such as Fe, ceramic vial with inner diameter o f 200 mm and cylindrical ceramic balls of 21 mm diameter and 21 mm length were used. Preparation of the Compacts. The ball milled powder particles were cold-compacted using a hydraulic press to a pressure of about 408 MPa. After this, the compacts were sintered at 868 K for 7.2 ks or at 898 K for 3.6 ks in an environment of forming gas containing 95%N2 and 5%H2 followed by furnace cooling. From the AI-Cu phase diagram in FIG. 2, the volume fraction of solid phase at two different temperatures can be calculated from Eq.1. The liquid phase is about 5.4% at 868 K and 21.6% at 898 K. A maximum of 26.45 % of liquid phase can be obtained at the temperature of 908 K. Above this temperature, slump occurred. After sintering, all compacts were subjected to a solution treatment at 790 K for 14.4 ks and aged at 443 K for 57.6 ks.
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700 ot+L
600 500
300 200
¢X
/
/ A1
) 1
2
3
5
6
7 18.5
Cu (wL%)
FIG. 2 A1-Cu phase diagram.
FIG. 3 Geometry of the compact tensile specimen.
Fracture Toughness Test and Microstructure Analysis. The sintered and heat-treated compacts were prepared for fracture toughness testing. The dimensions of the compacttension specimen used are given in Fig. 3. Fracture toughness of the peak-aged composites was determined using test procedures in accordance with ASTM E399-88 standard (16) using an Instron machine (model 8501). The crosshead speed was controlled at 0.1 mm/min. Morphologies of the powder particles were analyzed using a JEOL JSM-T30OA scanning electron microscope (SEM) operated at 15 kV. To analyze the cross sectional microstructure of the powder particles, the ball milled powder particles were cold-mounted and then carefully polished. The changes in microstructure of the particles after each duration of BM were examined using the optical microscope and SEM. Fracture surfaces after fracture toughness testing were also examined on the SEM. RESULTS AND DISCUSSION
Mechanical Properties. Fracture toughness of the composites in terms of KQ[MPalTI-1/2], the tentative value of Kzc,was determined in accordance with ASTM E399-88 standard from the following equation (16):
KQ= ~
f
(6)
where PQ is the 5% secant load [kN], B, the specimen thickness [cm], W, the specimen width [cm], a, the crack length [cm] andf(a/w), a geometry dependent constant. The fracture toughness of the composites generally showed a nearly type-I loaddisplacement record. The values of the P,,,,JPo_ratio, where Pm,~is the maximum load, for all compacts tested were nearly 1.10 or slightly larger. Hence, Ko may be considered as Kzc or near K~cvalues. Kzc of the compacts shows a decrease with the increase in BM duration. This trend appears to be closely related to the results of a corresponding increase in microhardness.
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Prolong BM duration caused microhardness to increase due to work hardening. Owing to the increase in hardness, densification of the hardened powders by cold compact becomes more difficult resulting in a decease in density. According to FIG. 2, although 5.4% liquid phase can be obtained, it is not enough to fill the pores to enhance densification. Rice (1719) has demonstrated that the porosity dependence of tensile strength, young's modulus, fracture energy and hence fracture toughness should usually be similar. The ratio of a property at some volume fraction porosity P to the property at P=0 can generally be represented as e bp where e is the Napierian logarithmic base and b is the slope from a semilogarithmic plot of the property versus P. The value of b is constant for a given type of porosity and property. Hence, sintering temperature was raised to 898 K where the liquid phase is about 21.6%. Klc value of the compact without BM sintered at 868 K was found to be 9.2 MPam 1/2 when the sintering temperature was increased to 898 K, Kic value was 11 MPam -1/2. A dramatic increase has likewise been observed for the ball milled compacts. At a sintering temperature of 868 K, Klc was only 4.2 MPam "1/2because of low compact density (about 84% of its theoretical density). However, when the sintering temperature was increased to 898 K, Kic value was increased to 6.2 MPam 1/2. This increment is about 50%. Fracture Surface and Microstrueture Analysis. FIG. 4 shows the fracture surfaces of the compacts sintered at 868 K. Agglomeration of SiC particulates in the matrix can be easily observed in the blended powders (FIG. 4 (a)). Large clusters of SiC particulates are more commonly seen on the fracture surface at shorter BM durations than those at longer BM durations. Because of the lack of liquid phase during sintering, the pores and the gaps between SiC particulates cannot be filled by molten A1 alloy (FIG. 5). Fracture surface analysis suggests that fracture is due mainly to the segregation of SiC particulates and/or debonding at the matrix-particle interfaces in conventional blending and at shorter BM duration where a uniform distribution of SiC particulates has not been achieved. Another reason for the agglomeration of SiC particulates are due to too little liquid phase formed during sintering. Liquid phase with high viscosity cannot penetrate into the pores between SiC particulates under capillary force. The agglomerations of SiC particulates in the matrix and the interfacial bond strength may, however, be improved with the increase in BM duration. Although almost no agglomeration of SiC particulates can be observed at longer BM duration, the longer duration shows some evidence of poor bonding between the matrix composite particles as indicated in FIG. 4 (b) and (d). FIG. 4 (b) and (c) show an interesting finding in that some particles manifested very smooth surfaces which indicate the formation of liquid phase during the sintering process. Since there existed only little amount of liquid phase, it was insufficient to fill the gaps between different particles and thus to produce diffusion bonding. This result suggests that a sintering temperature of 868 K is too low to have sufficient liquid phase for LPS of MMCs. FIG. 6 shows the fracture surfaces of the compacts sintered at 898 K. In comparison with FIG. 4, it is evident that bonding between particles has been improved by increasing the sintering temperature. FIG. 7 shows microstructure of a compact ball milled for 7.2 ks and sintered at higher temperature of 898 K, showing less porosity. Prolonged sintering, however, was observed to have little effect on densification. German (7) reported that porosity and pore size decease at all temperatures prior to liquid formation. For long sintering time, densification is accomplished by an increase in pore size due to coarsening
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459
I0 lam
10 I.tm
10 ktm
i
i
i
i
FIG. 4 Fracture surface of the compacts sintered at 868 K: (a) blended; at 863 K: (b) 36ks of BM and (c) 54ks of BM; at 868 K: (d) 72ks of BM.
FIG. 5 Microstructure of polished surface of a compact ball milled for 7.2 ks and sintered at 868 K shows large amount of porosity due to incomplete sintering.
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FIG. 6 Fracture surface of the compacts sintered at 898 K: (a) blended and (b) 72 ks of BM.
effect. Therefore, the process of sintering can be divided into several stages: densification is initially associated with pore shrinkage, followed by grain growth and pore coarseningin the later stage of densification. All these changes take place prior to the formation of the first liquid. One of the advantage ofLPS of MMC is that it enables liquid to penetrate into clusters of SiC particulates and hence improve their bond strength. Janowski (20) found that the LPS process results in unique particulate and matrix interfaces due to the locally high concentrations of alloying elements that is present at the interface during sintering. The reaction between SiC and AI alloy that leads to the faceted SiC particulates improves the properties of MMCs. Liquid-Phase Sintering of MMCs. In the present study, the maximum liquid phase during the sintering of matrix materials of composition A1-4.5wt.% was about 15%. Beyond that, the sintered compacts had been found to slump. However, the maximum liquid phase in the sintering of MMCs, in general, can be raised to 26% which is much higher than that of traditional LPS of metallic materials. The increase in the liquid phase can be understood from the following simple model as shown in FIG. 8.
FIG. 7 Microstructure of polished surface of a compact ball milled for 7.2 ks and sintered at 898 K shows less porosity.
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Matrixparticle
(a)
~
(b)
Ccr,'mficp~ticlcs Liquidphase (c)
FIG. 8 Liquid phase sintering of three different modes: (a) metallic powders; Co) metallic and ceramic powders; and (c) metallic and ceramic powders after ball milling. For LPS of a metallic material, the amount of liquid is to form the thin films surrounding the panicles and the grain boundaries. If the liquid film is too thick, the viscosity of the sintered compact will be very low resulting consequently in a slump of the compact. For sintering of MMCs, the melting temperature of the reinforcement can be much higher than the matrix materials. Capillary force leads to the movement of liquid whereas SiC particulates essentially remain at their original position. Because of the existence of SiC particulates, the viscosity is very low even though thick film has been known to form. To obtain good bonding, SiC particulates have to be surrounded by a liquid phase. Hence, an extra liquid is required. The amount of the extra liquid phase can be assumed by two extreme cases. In the first case, SiC particulates are agglomerated together. Since SiC particulates cannot be deformed, the amount of the liquid phase required is equivalent to
TABLE 1 Data Used for the Calculation of the Amount of Extra Liquid Phase Symbol Symbol 8 F~ Fi G P£
7Lv Ssic psic C
Description grain boundar7 film thickness = 1 I~m (6) Critical fractionalcoverage for metallic material --0.89 46) Ratio of liquid in grain interior = 0 Grain edge length = 10 Ilm (6) Green density =87 % Liquid-vaporsurface energy = 1 J/m2 (6) Particle size of SiC = 5 ~m Density of SiC = 3.2 ~/cm3 Particle size = 50 pm
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the amount needed to fill into a block of loose SiC particulates. Namely, i f a volume of SiC particulates Vt is agglomerated, it will occupy a space of volume Vs. The difference between Vs and Vt is the amount o f liquid phase needed in LPS.
-
-
/(w,
p, ao)
(7)
where do is the size o f the particulate, Ad, the thickness of the liquid film and 4L 2 /(w,
p, a o) =
6L+3
,/5 V-L2 + 0.268L + 0.025 2
and
where L is the cluster factor, W, p are the weight and density of reinforcement. If the reinforcement particulates are fully agglomerated, there is no liquid film between the reinforcement particulate. Therefore, Eq.[7] can be rewritten as:
V/ V
"g 1
-
- - f ( W ,
24 s
P, do)
(7)
In the second case, particulates are well distributed such that no clusters can be found, for example, after ball milling. If the particulates are considered to be surrounded by thin film (FIG. 8 (c)), the amount of extra liquid can be calculated as:
In general, the amount of extra liquid phase lies somewhere between these two extreme cases. If the data in Table 1 are used, the liquid phase and sintered density will respectively be 11% and 89% of its theoretical density for metallic body sintering. However, for sintering MMC with 15wt.% SiC, the liquid phase can be increased to 20.5% if the SiC particulates are considered to be surrounded by a film of 1 ~tm thickness. With the increase in size o f the reinforcement, at the same thickness of liquid film, the volume fraction of liquid phase will decrease.
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CONCLUSIONS 1.
2.
3.
4.
Fracture toughness values of ball milled compacts were found to decrease with increasing BM duration when they were sintered at 868 K. The main reason for the decrease is insufficient and poor diffusion bonding between the ball milled particles. Density o f compacts was increased if the compacts were sintered at temperature as high as 893 K. According to the AI-Cu phase diagram, liquid phase is approximately 33% which is much higher than that for LPS of metallic materials. Larger amount of liquid phase promotes interdiffusion between particles. In addition, the liquid phase may react with the SiC particulates to form better interfacial zone between matrix particle and the SiC particulates. Consequently, fracture toughness of the compacts can be increased. Fracture analysis shows three types of failure, namely, (a) debonding between matrix particles and the SiC particulates at short BM duration, (b) debonding between the composite powder particles due to insufficient diffusion at long BM duration but low sintering temperature, and (c) more ductile fracture when high sintering temperature is used.0 By using a simple model, volume fraction of the liquid phase in LPS can be determined. ACKNOWLEDGMENT
This project is supported by a grant from the National University of Singapore under RP3940667. REFERENCES
1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13.
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