interpenetrating microstructure and properties of si3n4/al–mg composites fabricated by...

ORIGINAL PAPER

Interpenetrating Microstructure and Propertiesof Si3N4/Al–Mg Composites Fabricatedby Pressureless Infiltration

Wang Shouren & Geng Haoran &

Zhang Jingchun & Wang Yingzi

# Springer Science + Business Media, Inc. 2006

Abstract Si3N4/Al–Mg composites based on Al–Mg alloy reinforced by ceramicinterpenetrating network structure were fabricated via pressureless infiltrationtechnology. Infiltration temperature and infiltration time are the key parametersdistinctly effecting on infiltration processes. Moreover, the increasing of Mg content(2–8 wt.%) resulted in an increased amount of infiltration. Microstructuralcharacterization of the composites reveals a special topology of skeleton and goodintegrity of metal/ceramic interface. The presence of second reinforced phase resultsin a significant increase in 0.2% offset yield and ultimate tensile strength ofcomposites materials. However, when the volume fraction of reinforcement is largethan 6%, there are a distinctly reduction of strength. The presence of additionalsecondary brittle phase in matrix results in the reduction in ductility and increase inhardness of 3-DNRMMCs. The failure features as cracking and void in reinforce-ment, interface cracking and interface debonding as well as matrix damage result inthe decreases of fracture toughness. With the increases of volume fraction ofreinforcement, 3-DRMMC exhibits excellent wear-resistance property.

Key words silicone nitride skeleton . metal matrix composites . pressurelessinfiltration . mechanical properties

1. Introduction

The traditional approaches to fabricate metal matrix composites is to combine thecontinuous phase (matrix) with one or more discrete reinforcement phases, shaped

Appl Compos Mater (2006) 13: 115–126DOI: 10.1007/s10443-006-9015-x

W. Shouren (*) :Z. JingchunSchool of Mechanical Engineering, Jinan University,Jiwei road 106#, Jinan, Shandong 250022,People’s Republic of Chinae-mail: [email protected]

G. Haoran :W. YingziSchool of Material Science, Jinan University,Jinan 250022, People’s Republic of China

as long fibers [1], short fibers [2], whiskers [3] or particles [4], etc. Recently, a novelfabricated technology, light weight metal matrix can be infiltrated into the porousand continuous preform at moderate infiltration temperatures and pressures, hasbeen proposed [5–8]. These kinds of composites have such special microstructurecharacteristic as: Both the matrix and reinforcement phase are continuous and tri-dimensional, inter-twisting and interpenetrating. These metal matrix compositesreinforced by three dimensional (3-D) network structure (3-DNRMMCs) are veryinteresting due to both microstructure characteristic and special mechanicalproperties and performance. Owing to each phase is a continuous networkpenetrated by similar networks of the other constituents, 3-DNRMMCs exhibitrather high specific strength and stiffness, wear-resistance, excellent thermal andelectrical conductivities and can be an attractive candidate for structural andfunctional materials [9].

In the time spanning the last two decades, aluminum matrix composites arerecognized as an increasing applied engineering materials in the fields of aerospace,automobile and chemical industry. Selection of reinforcement for monolithicaluminum alloy is a key factor for its performance. The effect of reinforcement type,size, distribution and interfacial reaction on composites strength and performance areissues of academic and practical interest. Most research studies have investigatedSi3N4 as a viable reinforcement in aluminum magnesium alloy due to its highstrength and high temperature refractory properties. However, up until 2006 there islittle documentation in the published literature of effort made at using Si3N4

reticulated structure reinforcement to improve the properties of aluminum alloys.The objective of the present study was to fabricate the reticulated porous Si3N4

ceramic skeleton by precursor replica process and to fabricate 3-DNRMMCs bypressureless infiltration technology. An emphasis is placed on the microstructureand mechanical properties of composites.

2. Experimental Procedure

2.1. Composites Fabrication

High-purity b-Si3N4 powder was used as starting material for fabrication of porousceramic skeleton. The sintering additives including 2 wt.% Al2O3, 5 wt.% ZrO2 and5 wt.% Al were mixed with the starting material and ball-milled for 4 h using Al2O3

balls. A reticulated polymeric sponge as preceramic materials with interconnectedpores was chosen to prepare the porous preform by the replica process. The clearedsponge was cut into standard form of 760 � 100 mm and immersed into thehomogeneous slurry for about 10 min. the samples were placed into the drying stoveand dried for 20 h at 160 -C to remove most of the moisture, and then, sinteringprocess was carried through at 1,400 -C for 2 h in oxidizing atmosphere. Then, thereticulated porous ceramics (RPCs) with different the porosity as 94%, 91%, 88%,85%, 80%, respectively, were fabricated.

Commercial purity Al ingot (99.7% in wt.) and Mg ingot (99.95% in wt.) wereused in preparing the Al–Mg alloys by melting in a clay-graphite crucible. ElementMg was wrapped with Al foil in order to prevent evaporation losses. The chemicalcompositions of infiltrated metal used in this experiment are shown in Table I. The

116 Appl Compos Mater (2006) 13: 115–126

Si3N4 RPCs skeleton is heated up in a furnace under nitrogen atmosphere togetherwith the infiltration die. After achieving the appointed temperature, the liquid metalwas infiltrated into the infiltration die. Then, 3-DNRMMCs reinforced by differentvolume fraction as 6%, 9%, 12%, 15%, 20%, respectively, were fabricated andpapered to carry out properties tests.

2.2. Characterization

The density and porosity of RPCs were measured using Archimedes measurements(ASTM C373). The morphologies and fracture surface characteristic of RPCs and3-DNRMMCs were characterized by a scanning electron microscopy (SEM,Hitachi, No. S-2500). Chemical element distributions were examined by the energyspectrum analyses (ESA, OXFOED INCA). The wetting angle of Al molten metalabout Si3N4 substrate was measured through sessile drop technique.

The specimens with different volume fraction reinforcement were carried outtensile strength, fracture toughness, wear resistance test and hardness test. All thetests were carried out in the room temperature (25 -C) and laboratory air envi-ronment (relative humidity a 55%).

The tensile specimens were machined in accordance with procedures outlined inASTM E8M-96. There were no visible machining marks on the specimens and noevidence of fracture nucleation emanating from the specimen surface. The speci-mens were deformed to failure on a universal testing machine (Instron 5569) withinitial strain rate of 0.005 minj1 and the strain was monitored by an extensometerwith a gage length of 25 mm.

Three-point (3-P) strength was measured over a span of 30 mm using an universaltesting machine (Instron 5569) with a cross-head speed of 0.5 mmminj1. The notchwas cut through electro-discharge machining which has a geometry characteristic ofroot radius (r) of 200 mm, normal width (c) of 2 mm and a normal length (b) ofapproximately 3–4 mm.

The hardness was then measured using Rockwell hardness testing machine withscale B. In order to measure the hardness, each sample were cut to a height of40 mm and turned to 25 mm diameter. A coolant was used during turning toreduce the rise in temperature, which may affect the precision of the hardnessmeasurement.

Mg Zn Mn Fe Si Be Cu Al

2õ8 0.70 0.15 0.06 0.30 0.04 0.20 Balance

Table I Chemical composition ofAl–Mg alloy (wt.%).

Figure 1 Schematic diagram of wearapparatus

Appl Compos Mater (2006) 13: 115–126 117

The wear experimental apparatus is illustrated schematically in Figure 1, whoserotating En43 steel disk heat-treated to HRC 50õ55 simulates the axle in a journalbearing unit. Test samples were machined to 7 � 8 � 18.5 mm3. The wear specimenswere tested under dry conditions and the block and the disk are degreased andcleaned with an alcohol solution. The applied loads ( p) are 50, 100, 200 N, re-spectively, and the rotating velocity of the steel disk (v) is 0.64 m/s. The wear rate ofthe 3-DNRMMCs was monitored by worn track width (mm).

3. Results and Discussion

3.1. Microstructure Characteristic

A model of 3-DNRMMCs was proposed in present work (Figure 2). It is indicatedthat the kind of materials have anisotropic, interpenetrating and intertwistingmicrostructures. The metal matrix and ceramic reinforcement formed intercon-nected network structure. The precursor replicating technique offer a higher inter-connectivity of the pores and a higher isotropy of the pore arrangement comparedto the conventional RPCs fabricated techniques. Molt infiltration technique offerinterconnection and interpenetration of two phases. The representative image of themacroscopic structure of RPCs is presented in Figure 3 a (porosity as 91%). Due tothe low density, RPCs exhibit unique properties such as a high stiffness at lowweight, a localized strain and fracture capability and a low thermal and electricalconductivity [10, 11]. The microstructure characterization of 9% Si3N4–Al–Mgcomposites with 3-D interpenetrating network is performed on a scanning electronmicroscopy which presented in Figure 3b. The region marked 1 is Si3N4 ceramicskeleton, the region marked 2 is metal matrix phase. Chemical element distributionsof two regions of metal/ceramic interface examined by the energy spectrum analyseswere shown in Figure 4. From Figure 4, the infiltration and diffusion of molt Al andMg element to Si3N4 skeleton were observed. It is indicated that matrix phase andreinforcement phase were integrated organically.

Figure 2 Schematic diagram ofinterpenetrating composites


3.2. Infiltration Kinetics

Wettability of metal and ceramic can be recognized as a significant problem infabricating 3-DNRMMCs. It is defined as the ability of a liquid to spread on a solidsurface. The problem of the wetting of the ceramic by molten metal is one of surfacechemistry and surface tension. The chemistry of the particle surface, including anycontamination, or oxidation, the melt surface and oxide layer must be considered[12]. There are some basic means used to improve wetting. One is to increase thesurface energies of the solid, second is to decrease the surface tension of the liquidmatrix alloy, and third is to decrease the solid–liquid interfacial energy at the metal–ceramic interface [12, 13]. Wettability can be described by contact angle (q) betweenmetal and ceramic. If q = 0-, it is recognized as perfect wettability, if q = 180-, it isrecognized as poor wettability and if 0- < q < 180-, partial wetting occurs. q = 90-was taken as the critical wetting angle for pressureless infiltration. If there are large

Figure 4 Chemical element distributions of two regions of metal/ceramic interface

Figure 3 The morphology of RPCs and 3-DNRMMCs. (a) Ceramic skeleton. (b) Compositereinforced by ceramic skeleton


extents of contact, the wetting angle must be smaller than 90-. Therefore, someparticular means must be adopted to reduce the wetting angle. The addition ofalloying elements such as magnesium to the molten matrix alloy, improvement oftemperature and adequate infiltration time can promote the wetting of metal toceramic. In present paper, the contents of 2% Mg, 5% Mg, 8% Mg, respectively,were proposed and the weight gain of samples were examined by thermo-analyticalbalance. The relations of infiltration temperature and weight gain at different Mgcontent were shown in Figure 5. From Figure 5, the weight gain increased with theincreasing of Mg content. The infiltration temperature was divided into threeregions as initial infiltration temperature region (I), middle infiltration temperatureregion (II) and high infiltration temperature region (III). In I region, the samplesgained weight rapidly up to the set point temperature for Al – 2% Mg, Al – 5% Mg,Al – 8% Mg alloy, the infiltration rate get rapidly showing a distinct slopes. But pureAl had no volume of gained weight which is demonstrated that no infiltrationoccurred. The initiation infiltration temperature descended with the increasingcontent of Mg. It well known that the content of Mg can promote the infiltration.Low content of Mg play a small role because of its low vapor pressure andprematurely lost from the system, the same conclusion was investigated in literature[14]. The content of Mg beyond 10 wt.% has little further effect [15]. In II region,the weight gain curve flattened off rather abruptly at the temperature of 850 -Cwhich indicated that the infiltration processing had been achieved. In III region, theweight gained curve turned gradual flattening and little precipitous with theincreasing of the temperature but the curve of pure Al has no ascent. The reasonwhy the curve has a little precipitous is that Mg reduces the viscosity and surfacetension of aluminum.

It is well known that infiltration time affect the infiltration results. The curve ofweight gain vs infiltration time for commercial purity Al, Al – 2% Mg, Al – 5% Mgand Al – 8% Mg alloys at 1,000 -C are described in Figure 6. The infiltrationprocesses can be divided into three stages as infiltration initiation (0–15 min),infiltration continuation (15–45 min) and infiltration termination (45–60 min). In

Figure 5 Relation of infiltration temperature and weight gain


initiation of the infiltration, the sample initially gained weight rapidly not only pureAl but also Al–Mg alloys. Experiment revealed that self-gravity promotes the startof the infiltration owing to the pores of preform exhibit a suitable dimension. Incontinuation of the infiltration, the curve of the pure Al in this stage has alreadybecome even because of the poor wetting of pure Al with Si3N4 preform. So apremature termination stage of infiltration is observed. Therefore, the curve of Al–Mg alloys have a continue ascent after breaking through a flat roof. The plainascend gradually with the increasing of the content of Mg, so the curve of Al –8 wt.% Mg hold the highest plateau. The rate of weight gain drastically decreasedowing to the block of capillarity force. With the continuation of infiltration, one partof some pores of the preform were walled up and blocked or dwindled, As a result,the capillarity force turn into one kind of infiltration impetus. In the same time, Mgenhanced the wetting of ceramics by molten metal and lowered the viscosity andsurface tension of aluminum. So the content of Mg also became the impetus ofinfiltration. In termination of infiltration, the curve of all alloys became even linewhich indicated that the infiltration processes had been achieved, so appropriateinfiltration time is 50 min, prolonging the infiltration time continuously has nosignificance.

With the increases of volume fractions of reinforcement, the strength (0.2% YSand UTS) of 3-DNRMMCs exhibits a significant improvement (Table II). When thevolume fraction of reinforcement is 6%, the 0.2% YS is 212 MPa and the UTS is278 MPa. This increase is attributed to the following: (a) the special topology ofreinforcement, (b) reinforcement exhibiting good interfacial integrity with mono-lithic alloy, (c) the infiltration of molten metal and diffusion of metal element toskeleton and (d) the effective transfer of load from metal matrix to reinforcingphase. The strength decreases with the continuous increases of volume fractionreinforcement (>6%). This decrease is attributed to two side of influence. One is theinfluence of non-homogeneous and porous struts of skeleton. Another is theinfluence of division of metal matrix by overfull reinforcement. The lower ductilityof 3-DNRMMCs (Table II) is attributed to the presence of additional secondarybrittle phase in matrix. The cracks in skeleton result in the reduction in ductility.

The effect of volume fraction of reinforcement on the Rockwell hardness numberis also shown in Table II. This figure indicates that the hardness of 3-DNRMMCs

Figure 6 Relation between in-filtration time and weight gain(1,000 -C)


increased with increasing the volume fraction of reinforcement. It is worthy tomention here that Si3N4 is considered to be an industrial ceramic and has a highhardness value (micro-hardness as 18 GPa). Due to the plastic deformation underthe compressive load, plastic flow was hindered by the reinforcement skeleton andlead to a remarkable reduction. It is indicated that the restrain of plastic flow lead toan increase in the hardness of 3-DNRMMCs.

3.3. Fracture Behavior

Figure 7 shows the relation between the fracture toughness (KIC) and volumefractions of reinforcement. It is indicated that the volume fraction of reinforcementplays a significant role in controlling fracture toughness from both positive andnegative side. The positive side is that, owing to the exists of a small quantity ofbrittle phases, crack propagation was blocked off and fracture stresses were releasedin the interface of metal/ceramic, the value of fracture toughness increases. Forexample, 3-DRMMC with volume fraction of reinforcement as 6% exhibit KIC of25.5 MPa m1/2. The negative side is that, the failure features as cracking and void inreinforcement, interface cracking and interface debonding as well as matrix damageresult in the decreases of fracture toughness. Owing to a large amount of brittlephase interpenetrating in ductile phase and dissevering metal matrix, there is nopossibility for fracture toughness of reinforced composites to be in excess of that ofmatrix metal. Figure 8 shows the fractographs of Al–Mg alloy and 20% Si3N4/Al–Mg composites. Owing to the adequate ductility of monolithic alloy, much and

Materials 0.2% YS

(MPa)

UTS

(MPa)

Ductility

(%)

Hardness

(HRB)

Al–Mg–0% Si3N4 198 265 7.5 65

Al–Mg–6% Si3N4 212 278 7.2 68

Al–Mg–9% Si3N4 196 264 4.0 70

Al–Mg–12% Si3N4 187 233 3.2 71

Al–Mg–15% Si3N4 164 210 2.3 75

Al–Mg–20% Si3N4 147 189 1.8 82

Table II Summary ofmechanical properties.

Figure 7 Relation betweenfracture toughness and volumefraction of reinforcement


smaller dimples and crater are observed in SEM fractographs (Figure 8a) and thereare big craters and dimples with a bimodal and multimodal distribution character-istic. The crater coalescence is the predominant failure mode of metal matrix. Whilethe second brittle phases were introduced to the metal matrix, it turns into mixed-mode fracture mechanisms (Figure 8b). The extent of brittle fracture region andlocally ductile failure were found in Figure 8b. A large number of ceramic phaseswere found to be engulfed in crater enclosed by surrounded matrix.

3.4. Wear Behavior

The changes of friction coefficient at different volume fraction of reinforcementunder a 50 N load and 0.64 m/s sliding speed are shown in Figure 9, it is seen that thefriction coefficients of all samples are lower at first, after a short time, all of thefriction coefficients rise due to wear debris on the friction surface. The valueof friction coefficient decreases greatly with the increases of volume fraction ofreinforcement: The non-reinforced alloys have a large value of friction coefficientwhile the 20% Si3N4/Al–Mg composites has a smallest value. Figure 10 shows thechange of wear track width with the increases of volume fraction of reinforcement atdifferent applied loads. It is indicated that the wear rate decreases greatly with the

Figure 8 Fractographs of Al–Mg alloy and 15% Si3N4/Al–Mg composites

Figure 9 Friction coefficient vsdry slide time at differentvolume fraction of reinforce-ment ( p = 50 N, v = 0.64 m/s)


increases of volume fraction of reinforcement. However, the wear rate increaseswith the increases of applied load.

The morphology of the worn surface changed gradually from fine scratches todistinct grooves and to even severe deformation with the increase of the loads.Figure 11a and b show the microstructure of worn surfaces of Al–Mg alloy and 15%

Figure 11 The microstructure of the wear surfaces of (a) Al–Mg alloy under 50 N; (b) the 15 wt.%Si3N4–Al–Mg composite under 50 N; (c) Al–Mg alloy under 200 N; and (d) 15 wt.% Si3N4–Al–Mgcomposite under 200 N

Figure 10 The change of weartrack width with the increasesof volume fraction of rein-forcement at different appliedload


Si3N4/Al–Mg composite at the lower load of 50 N. It can be indicated that a mixedabrasion-plastic deformation process was observed on the worn surface. At higherloads, all samples developed severe plastic deformation because the worn surfacetemperature of the samples is high enough to low the shear strength of the sub-surface layer. The addition of second reinforcement improves the seizure resistanceand reduces the wear by minimizing the friction-induced heating [16]. Figure 11cand d showed the microstructure of the worn surfaces of alloy and its composites atthe higher applied load of 200 N. The surface of Al–Mg alloy exhibits more severeplastic reformation. However, these features are not observed in the Al–Mg–Si3N4

composites, confirming the beneficial effect of the Si3N4 network structure. It isindicated the surface of 15% Si3N4/Al–Mg composite is steadier than that of matrixalloy.

4. Conclusions

1) Three dimensional network structure reinforced aluminum magnesium matrixcomposites had been fabricated via pressureless infiltration technology. Infiltra-tion temperature and infiltration time are the important variables in fabrication.The presence of Mg content resulted in an increased amount of infiltration.

2) When the volume fraction of reinforcement is 6%, the 0.2% YS is 212 MPa andthe UTS is 278 MPa. Then with the increases of volume fraction of rein-forcement, the strength (0.2% YS and UTS) decrease. The cracks in skeletonresult in the reduction in ductility. The presence of second brittle phase resultsin the increases in hardness of composites.

3) 3-DRMMC with volume fraction of reinforcement as 6% exhibit KIC of25.5 MPa m1/2. The failure features as cracking and void in reinforcement,interface cracking and interface debonding as well as matrix damage result inthe decreases of fracture toughness.

4) With the increases of volume fraction of reinforcement, 3-DRMMC exhibitsexcellent wear-resistance property.

Acknowledgement Foundation item: Project (50371047) supported by the National Natural ScienceFoundation of China.

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