Production and characterization of in situ Al4C3 reinforced aluminum-based composite produced by mechanical alloying technique

Materials and Design 25(2004) 31–40

0261-3069/04/$ - see front matter� 2003 Elsevier Ltd. All rights reserved.doi:10.1016/S0261-3069(03)00163-8

Production and characterization of in situ Al C reinforced aluminum-4 3

based composite produced by mechanical alloying technique

Halil Arik

Department of Metallurgy, Faculty of Technical Education, Gazi University, Ankara 06500, Turkey

Received 27 January 2003; accepted 25 July 2003

Abstract

In this study, aluminum powders were mixed with carbon black and mechanically alloyed in a high-energy ball mill for up to30 h under argon atmosphere. XRD results showed that no Al C formed during the mechanical alloying(MA) process.4 3

Mechanically alloyed powders were then compacted at 650 MPa pressure to produce transverse rupture blanks. These blankswere then sintered at 600 and 6508C for 5, 10 and 20 h under argon atmosphere. X-Ray diffraction analysis of the sinteredblocks showed the presence of Al C phase. Increasing MA processing time, sintering temperature and duration, resulted in an4 3

increase in the amount of Al C phase. XRD and SEM analysis showed that the most Al C transformation occurred in the blank,4 3 4 3

which was mechanically alloyed for 30 h and sintered for 20 h at 6508C. The highest hardness value(347 Hv) was obtained inthis blank. Transverse rupture strengths of the sintered blanks were between 6.32 and 234.16 MPa. The highest value was obtainedfor the blank, which was mechanically alloyed for 15 h and sintered for 20 h at 6508C.� 2003 Elsevier Ltd. All rights reserved.

Keywords: Composite; Mechanical alloying; Powder metallurgy; Mechanical properties; Scanning electron microscopy; X-Ray analysis

1. Introduction

There are vast reserves of aluminum ores in the worldand it is the second most used metal after ironw1x. Dueto ease of formability and lightweight, it is finding manyapplications. Since aluminum and its alloys have poorelevated temperature properties, much research workhad been carried out to strengthen them. This is gener-ally done by dispersion of high temperature resistantfine particles of oxides like Al O , ThO or Y O and2 3 2 2 3

carbides like TiC or SiC, in the aluminum structurew2–4x. Dispersion of fine particles is either carried out inthe solid or liquid-state. In the liquid methods, particlesare added to liquid aluminum by stirring before casting,but the resulting distribution is generally inhomoge-neous. The solid-state route is powder metallurgy proc-essing of the blended powders of each constituentw4–10x. Another way of producing the fine and homogenousdistribution of these hard particles is their production insitu, generally by reaction milling and annealing pro-cesses. Mechanical alloying is a simple and usefultechnique to synthesize both equilibrium and non-equi-

E-mail address: [email protected](H. Arik).

librium phases of commercially useful materials startingfrom elemental powders. This method was developedduring the late 1960s to produce high temperaturematerialsw11–15x. MA produces a homogeneous distri-bution of inert, fine particles within the matrix andavoids many problems associated with melting andsolidification. It has been recognized that MA can beused to induce chemical reactions in powder mixturesat room temperature or at least at much lower tempera-tures than normally requiredw9x. Raw materials used forMA are pure or alloy powders that have particle sizesfrom 1 to 200mm. A process control agent(PCA) isadded to the powder mixture during milling, especiallywhen the powder mixture involves a substantial fractionof a ductile component. The PCAs are mostly organiccompounds, which act as surface-active agents. PCAsminimize cold welding and inhibit agglomeration. Com-mon PCAs are stearic acid(zinc stearate), hexane andoxalic acid, which are used at levels of 1–4 wt.% ofthe total powder charge. The actual process of MA startswith mixing powders in the right proportion and loadingthe powder into the mill along with the grinding media(generally steel balls). This mix is then milled for the

32 H. Arik / Materials and Design 25 (2004) 31–40

Table 1Specifications of carbon black

Reflection(%, 325 mesh–sieve Moisture Density Sulfurwith toluene) oversize(%) (%) (gy1) (%)

Minimum 80 – – 320 –Maximum – 0.1 25 380 1

Table 2MA conditions for the powder mixture

Vessel volume(cm )3 750 Rotor speed(rev. min )y1 450Mass of aluminum powder(g) 48.5 MA atmosphere ArgonMass of carbon black(g) 1.5 Cooling WaterCharge ratio(mass of grinding 6:1 Milling time(h) 0.5–30balls: mass of powder in mill)Steel ball diameter(mm) 10 PCA(%) 2–3

Fig. 1. Mechanical alloying set up:(1) argon cylinder;(2) manometer;(3) tube furnace;(4) attritor.

desired time until a steady state is reached. During high-energy milling, the powder particles are repeatedlyflattened, welded, fractured and re-welded. In the earlystages of milling, the particles are soft and their tendencyto weld together is highw16x. A broad range of particlesize develops, with some as large as three times biggerthan the starting particles. The composite particles atthis stage have a characteristic layered structure consist-ing of various combinations of starting constituents.With continued deformation, the particles become workhardened and fracture by a fatigue failure mechanismandyor by the fragmentation of fragile flakes. Fragmentsgenerated by this mechanism can continue to reduce insize in the absence of strong agglomeration forcesw17–21x. The aim of this study was to produce Al C4 3

reinforced aluminum composite powders by reactionmilling through MA of carbon black and aluminumpowder mixtures. These composite powders were thencompacted and sintered at certain temperatures for cer-tain durations under argon gas. Finally, the mechanicalbehavior of these blocks of composite materials wasinvestigated.

2. Experimental procedure

2.1. Material

Gas atomized Al powders were produced by GaziUniversity PM Lab. A maximum particle size of 150mm (y100 mesh) powders was produced from 99%pure Al ingots, which were supplied from ETIBANK(Turkish Aluminum Producing Co.).Carbon black of 99% purity and mean powder size

(agglomerate size) of 2.4 mm were obtained fromYARPET (Turkish Petrochemical Trade Co.). The prop-erties of carbon black are given in Table 1.

2.2. Mechanical alloying

Aluminum powders of 48.5 g, 1.5 g of carbon blackand 300 g of steel balls, with diameter 10 mm, wereplaced into a 750 ml capacity tank of a high energyattritor. Furthermore, 1 g of stearic acid was added tothe mixture to eliminate sticking of Al powders to theballs and to the walls of the milling tank. In order to

33H. Arik / Materials and Design 25 (2004) 31–40

Fig. 2. XRD results for the samples(a) before MA, (b) after MA for 20 h.

eliminate oxidation of aluminum powders during MA,the process was conducted under argon an environment.Ar gas was purified of residual oxygen by passing itthrough Cu chips, which were heated to 6008C. The

tank of the attritor was water cooled during MA and themixtures were milled at various milling time(1, 2.5, 5,7.5, 10, 15, 20, 25 and 30 h) (Fig. 1). The parametersfor mechanical alloying are given in Table 2.


Fig. 3. Effect of MA processing time on particle size.

Fig. 4. Morphology of samples mechanically alloyed for(a) 1 h, (b) 5 h, (c) 10 h, (d) 20 h and(e) 30 h.


Fig. 5. Effect of MA time on the sample density.

Fig. 6. Deformation of the powder particles depending on pressing direction(mechanically alloyed for 7.5 h sintered at 6508C for 20 h).

2.3. Characterization of the alloyed powders

In order to determine whether any carbide formationor any other structural changes had taken place duringMA, powders were analyzed by using a Rigaku–Geig-erflex X-ray diffractometer(XRD) after milling and thepeaks in the XRD traces were compared with those ofthe peaks obtained from the elemental powders. Powdersizes were measured by using a Malvern Master SizerE version 1.2b laser scattering machine and their shapeswere analyzed by using a Joel JSM 6400-Noran Instru-ments Series II scanning electron microscopy(SEM)before and after MA processing.

2.4. Compaction and sintering

Mechanical alloyed powders were compacted at 650MPa pressure to produce blocks 6.35=12.70=31.70

mm in size, according to ASTM B 312, by using asingle action press. Block samples were put to graphiteboats which were placed in an atmosphere controlledtube furnace and heated to test temperature in a flowingargon atmosphere for predetermined times. The heatingrate of the furnace to the desired temperature wasapproximately 58C min . The furnace was held at thaty1

temperature with an accuracy of"5 8C and then cooledto room temperature at 58C min . Sintering wasy1

performed at 600 and 6508C for 5, 10 and 20 h.

2.5. Compact characterization

The surfaces of all sintered blanks were lightly groundand polished to remove any irregularity or debris. Thesamples were then subjected to transverse rupture testsin a special device, designed and manufactured accord-ing to MPIF Standard 41 at Gazi PyM Lab. The loadwas applied by a lever mechanism with balls of 20 g(minimum load reading). Hardnesses of all sampleswere measured by the Vickers hardness method andmean of at least five readings was taken. In order todetermine the effect of sintering temperature and thetime the formation of Al C particles, all compacts were4 3

analyzed using XRD and SEM microscopy.

3. Results and discussion

XRD results for mechanically alloyed powdersshowed the characteristic peaks of pure aluminum,which indicated no chemical interaction between Alpowders and carbon black during MA processing. How-ever, XRD peaks taken from alloyed powders showedbroadening compared to those for elemental powders.


Fig. 7. Cross-sectional part of the fractured surface of sample mechanically alloyed for 20 h and sintered for 20 h at 6508C.

This indicates that mechanical alloying causing highdeformation of powders, which results in the formationof amorphous structure(Fig. 2). A similar effect wasreported by other investigators, especially when usinghigh ballypowder ratio and it was found that Al hadturned from a crystalline structure to amorphous struc-ture w21–24x.The measurement of powder size was conducted

before and after mechanical alloying and the meanparticle size of un-processed and 30 h mechanicalalloyed Al–C powders were found to be 157.86mmand 13.30mm, respectively(Fig. 3). Due to the carbonblack and stearic acid, added to the powders as alubricant, welding and agglomeration of the powderswere hindered considerably. In the early stages ofmechanical alloying shown in Fig. 4a. The soft Alpowders deformed intensively and formed flaky andneedle-like structures, which resulted in an increase inthe particle size. As milling precede, cold welding andfragmentation occurred throughout the process andresulted in the formation of finer particles. SEM micro-graphs show a considerable change in particle size andmorphology during mechanical alloying(Fig. 4a–e).Mechanical alloyed powders were compacted at 650

MPa and block samples were produced. In the earlystages of MA, the green density of powders increasedwith alloying time. However, a decrease in green densitywas observed in samples mechanical alloyed for 15 hor more (Fig. 5). This is due to the size and hardnessof the mixed powders. In the early stages of mechanicalalloying (up to 15 h) the density of the blocks increasedwith decreasing the particle size. In this stage, powdersare still soft and can be deformed easily during compac-tion. Optic pictures taken from the fracture surface of

blocks clearly show in Fig. 6 the deformation behaviorof the powders during compaction. The cross-sectionalpart of the fractured surface of the sample, takenperpendicular to the pressing direction, shows a flake-like structure(Fig. 6). However, powders alloyed for20 h or more exhibited work hardening. Thus, compac-tion of these powders was difficult and consequentlylow green densities were achieved. Almost no defor-mation was observed in samples produced from powdersalloyed for 20 h or more(Fig. 7).XRD analysis of the sintered blanks revealed no

change in phases of the powders mechanical alloyed for7.5 h or less. Powders mechanically alloyed more than7.5 h, however, exhibited the Al C phase and the4 3

number and intensity of the peaks increased withincreasing mechanical alloying time(Fig. 8). At lowerMA durations, the powder particle size was not smallenough and contained a surface oxide layer, whichretarded the reaction between Al and C and consequent-ly, the Al C phase did not form. Longer MA durations4 3

resulted in fracturing of surface oxide layer on the Alparticles and production of clean particle surfaces. Theseparticles cold welded together, fractured and carbonblack diffused into Al particles. For this reason, Al C4 3

particles formed in the bulk of the powders rather thenon the surface of them during sintering. Furthermore,longer MA durations resulted in higher deformation ofAl particles and caused a high internal energy. Thisenergy facilitated the formation of the Al C phase4 3

during sintering.The lowest transverse rupture strength of 6.32 MPa

was obtained in the blank, which was MA for 1 h andsintered at 6008C for 5 h. The most important factorfor this low transverse rupture strength is low MA


Fig. 9. SEM micrograph of the sample, mechanically alloyed for 2.5h and sintered at 6008C for 1 h, showing layers of Al particles andC black, which hindered the sintering.

Fig. 8. X-Ray diffraction patterns for samples mechanically alloyed for 7.5–30 h and sintered for 20 h at 6508C.

duration where carbon black did not have enough timeto diffuse into the Al particles. Carbon black coveredthe surface of the Al particles and adversely affectedthe sintering behavior of the block(Fig. 9). Test resultsindicated that the transverse rupture strength of theblanks, obtained from powders MA’ed for shorter times,increased with both increasing temperature and sinteringtime (Fig. 10). This is most probably due to the highgreen or sintered or both densities of blocks togetherwith the decreasing negative effect of the carbon black,which was dedected on the surface of Al particles.Furthermore, the higher deformation rate, obtained atlonger MA durations, increased the internal energy,which caused better sintering and resulted in highertransverse rupture strengths. Powders alloyed for 15 hor more exhibited lower transverse rupture strengthvalues. This is due to the decreasing green density ofthe blocks and the formation of high amounts of brittleAl C phases in the matrix. The highest transverse4 3

rupture strength of 234.16 MPa was obtained on the


Fig. 10. Effect of MA duration on the transverse rupture strength ofsamples sintered at various temperatures and durations.

Fig. 11. SEM micrograph of the sample mechanically alloyed for 15h and sintered at 6508C for 20 h showing brittle fracture behavior.

Fig. 12. Effect of MA duration on the hardness of samples sinteredat various temperatures and durations.

blank, which was mechanically alloyed for 15 h andsintered at 6508C for 20 h. The highest transverserupture strength for this sample is due to the high greenand sintering densities.The fractured surfaces of the blanks indicated that the

Al particles still showed ductile fracture behavior. Withincreasing the MA duration, the amount of deformation,due to compaction, decreased and after MA for 20 halmost no deformation was seen. Fracture of thesesamples occurred at the contact points of the particles,with lower strengths and samples exhibited brittle frac-ture behavior(Fig. 11).The results indicated that the most important factor

for the hardness was MA duration. However, highsintering temperature and duration also effected thehardness. Among the blanks, the highest hardness valueof 347 Hv was obtained for the sample MA for 30 hand sintered at 6508C for 20 h. A sharp increase inhardness was observed for the sample MA’ed for 10 h.This is most probably due to formation of Al C , which4 3


Fig. 13. Dispersion of Al C particles in the Al matrix,(a) Al C4 3 4 3

particles in the Al powders,(b) Al C particles formed at the interface4 3

between Al powders.

formed after 10 h MA(Fig. 12). XRD results alsoconfirm this postulate. It can be concluded that longerMA duration results in the formation of larger amountsof Al C , which increases the hardness value of the4 3

sample.SEM studies of sintered blocks showed cubic Al C4 3

particles in the matrix with a size of approximately 1–3 mm. These synthesized carbide particles mostlyformed in the Al particles rather than particle surfaces(Fig. 13a–b). This situation is different from the Albased composite produced by different method such aspowder metallurgy and the bonding at the particle–matrix interface looks much better and may give bettermechanical properties. Powder alloyed for less than 7.5h did not show any Al C phase(Fig. 8) whereas4 3

powders alloyed more than 7.5 h exhibited Al C parti-4 3

cles, which was proportional the MA time. The highestconcentration of Al C was seen in the sample, which4 3

exhibited highest hardness value.

4. Conclusion

Al–C powders were mechanical alloyed and blocksamples were produced by using PM routes and thefollowing results were obtained:

1. Up to 30 h of mechanical alloying, no formation ofAl C phase was observed.4 3

2. With increasing MA duration, the particle size of themilled powders decreased and the morphologychanged from globular to flake-like structure.

3. During the MA process, work hardening of thepowders occurred which resulted in a reductor in thepowder compressibility and resulted in lower densitiesand lower mechanical properties of blocks.

4. The high deformation of Al particles, which occurredduring MA, increased the internal energy and actedas a driving force during sintering and enhanced theformation of Al C particles.4 3

5. Residual carbon and the surface oxide layer on theAl powders after MA adversely affected the sinteringbehavior of powders, which resulted in lower mechan-ical properties.

References

w1x Tulgar HE. Metallurgy of non-ferrous metals.Istanbul TechnicalUniversity Publication, 1987. p. 352–402.

w2x Arik H, Turker M, Saritas S. Investigation of the mechanicalproperties of in situ Al C reinforced aluminum based com-4 3

posites by mechanical alloying technique. Powder MetallurgyWorld Congress and Exhibition, Kyoto, 2000, pp. 543–549.

w3x Ibrahim IA, Mohamed FA, Lavemia EJ. Particulate reinforcedmetal matrix composite—a review. J Mater Sci 1991;26:1137–1156.

w4x Lu YX, Meng XM, Lee CS. Microstructure and mechanicalbehavior of a SiC particles reinforced Al–5Cu composite underdynamic loading. J Mater Processing Technol. 1999;94:175–178.

w5x Benjamin JS. Mechanical alloying—a perspective. In: Arzt E,Schultz C, editors. New materials by mechanical alloyingtechniques.Colw-Hirson(FRG), 1988. p. 3–19.

w6x Benjamin JS, Volin TE. The mechanism of mechanical alloy-ing. Metall Trans 1974;5:1929–1933.

w7x Fischer JJ, Weber JH. Mechanical alloying. Adv Mater Pro-cesses 1990;10:43–50.

w8x German RM. Powder metallurgy science. New Jersey: MetalPowder Industries Federation, 1984.

w9x Tweed JH. Manufacture of 2014 aluminum reinforced with SiCparticulate by vacuum hot pressing. Mater Sci Eng1991;A135:73–76.

w10x Jangg G, Slesar M, Besterci M. Influence of heat treatmentduring manufacturing of Al–Al C materials on microstructure4 3

and properties. Pmi 1989;21:25–30.w11x Singer RF, Oliver WC, Nix WD. Identification of dispersoid

phases created in aluminum during mechanical alloying. MetallTrans A 1980;11A:1895–1901.

w12x Bronsveld PM, Bruinsma P. Microstructural analysis of hotisostatically pressed Al–SiC. J. Th. De. Hosson: Mater SciEng 1991;A135:77–81.


w13x Rodriguez JA, Gallardo JM, Herrera EJ. Structure and proper-ties of attrition-milled aluminium powder. J Mater Sci1997;32:3535–3539.

w14x Suryanarayana C. Bibliography on mechanical alloying andmilling. Cambridge International Science Publishing, 1995.

w15x Gilman PS, Nix WD. The structure and properties of aluminumalloys produced by mechanical alloying: powder processingand resultant powder structures. Metall Trans A 1981;12A-:813–823.

w16x Wu NQ, Wu JM, Wang GX. Amorphization in the Al–Csystem by mechanical alloying. J Alloys Compd1997;260:121–126.

w17x Lawley A, Koczak MJ, Orban RL. Fabrication of Nanocrystal-line NiAl Compacts via Combined Mechanical Alloying andSelf Propagating Synthesis. European Conference on AdvancedPM Materials, Birmingham, 23–25 1995;2: pp. 88–95.

w18x Besterci M, Kohutek I, Sulleiova K. Analysis of spatialarrangement of particles in thin foil of Al–Al C material. J4 3

Mater Sci 1999;34:1055–1060.

w19x Besterci M, Slesar M, Kovac L. Influence of strain rate onfracture of dispersion strengthened Al–Al C systems. Scr4 3

Materialia 1997;37:1077–1080.w20x Zhu LH, Ma XM, Zhao L. Study on phase transformation of

Fe–Ni powders during mechanical alloying. J Mater Sci2001;36:5571–5574.

w21x Kim HS, Kum D, Hanada S. Structural evolution duringmechanical alloying and annealing of a Nb-25 at.%A1 alloy. JMater Sci 2000;33:235–239.

w22x Zhu M, Che XZ, Li ZX. Mechanical alloying of immisciblePb–Al binary system by high energy ball milling. J Mater Sci1998;33:5873–5881.

w23x Liu ZG, Guo J-T, Zhou LZ. Mechanical alloying synthesis andstructural characterisation of ternary Ni–Al–Fe alloys. J MaterSci 1997;32:4857–4864.

w24x Wen CE, Kobayashi K, Sugiyama A. Synthesis of nonocrys-tallite by mechanical alloying and in situ observation of theircombustion phase transformation in Al C . J Mater Sci4 3

2000;35:2099–2105.

Documents

Production and characterization of in situ Al4C3 reinforced aluminum-based composite produced by mechanical alloying technique