Materials Science and Engineering A239–240 (1997) 665–671

Ni3Al intermetallic compound as second phase in Al2O3

ceramic composites

V.M. Sglavo a,*, F. Marino b, B.R. Zhang b, S. Gialanella a

a Uni6ersita di Trento, Dipartimento di Ingegneria dei Materiali, Via Mesiano 77, 38050 Trento, Italyb Dipartimento di Scienza dei Materiali e Ingegneria Chimica, Politecnico di Torino, C.so Duca degli Abruzzi 24, 10129 Torino, Italy

Abstract

Several metals have been proposed as second phases in ceramic matrix composites in order to improve their fracture toughness.Unfortunately, the use of metals is limited by low melting temperature, as for Al and Ag, poor oxidation resistance, as for Ni,Mo and W, and decrease of mechanical strength as temperature increases. In these respects, high temperature structuralintermetallics show better properties. This work presents the preparation and the characterization of a Ni3Al reinforced-alumina.A ceramic composite containing 10 vol% Ni3Al powder was prepared by hot-pressing at 1350°C for 1.5 h green compacts of themixture of ceramic and intermetallic powder. Microstuctural features were investigated by scanning electron microscopy (SEM).Elastic modulus, flexural strength and fracture toughness were measured at room and high temperatures and correlated to themicrostructural characteristics of the material. A toughening mechanism due to plastic deformation of the intermetallic particlesduring crack propagation was seen to operate both at room and at high temperature. © 1997 Elsevier Science S.A.

Keywords: Composite; Ni3Al; Alumina; Strength; Fracture toughness; High-temperature

1. Introduction

One of the most promising approaches for improvingthe toughness of brittle ceramics involves the incorpora-tion of a ductile phase. The toughening mechanismsactually operating will depend upon the ductile phasemorphology. In the case of fibres, foils, etc., crackbridging will most likely occur and the fracture energyis dissipated by the plastic deformation of the ductilephase. Such mechanisms are particularly effective whenlimited debonding between the brittle matrix and theductile second phase occurs [1].

For particulate systems the situation is more complexowing to the larger influence of microstructural aspects.Crack bridging can still be an important tougheningmechanism. However, partial debonding can be difficultto control and, in general, a strong particle-matrixadhesion is recommended in order to achieve a suffi-ciently high level of plastic deformation of the ductiledispersion.

Several metallic phases have been considered so farin the attempt of improving toughness of alumina, oneof the most popular ceramic material whose widespreaduse as structural material is limited by unreliable valuesof fracture toughness and strength. Nickel [2–4], aNi–Ti alloy [4], iron [5], silver [6], etc., have been usedas second phases and have usually provided interestingresults in improving room temperature fracture tough-ness. On the other hand the selection of a metal maypresent some drawbacks related to a comparativelylower melting temperature, as compared to the ceramiccomponent, poor mechanical properties at intermediateand high temperatures and limited oxidation resistance.

Therefore it is interesting to explore the capabilitiesoffered in this field by suitable intermetallic phases,with better mechanical and surface stability properties.A NiAl powder was incorporated inside an aluminamatrix to produce composite compacts having excellentmechanical properties [7]. Another nickel aluminide, i.e.Ni3Al, is currently under investigation (unpublished). Inthe present study some microstructural and mechanicalaspects of alumina composites with a Ni3Al dispersionare discussed.* Corresponding author.

0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved.

PII S0921 -5093 (97 )00650 -3

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2. Experimental

2.1. Sample preparation

The starting powders were a commercial aluminapowder and an intermetallic atomised powder, kindlymade available by Wright and Knibloe (Idaho Nat.Eng. Lab., USA), based on the Ni3Al phase, having acomposition (at%) 73.12Ni–18.82Al–8.06Cr–0.019Mo–0.1B plus traces of S, N and O as impurities.The presence of chromium is mainly meant to improvethe mechanical strength, via solid solution strengthen-ing [8], but certainly has also a beneficial effect on theoxidation resistance of the alloy. Boron is a ductilizer,as it segregates along grain boundaries, improves theircohesion and reduces the intergranular fracture propa-gation mode.

Powder mixtures containing 10 vol% Ni3Al-base par-ticles were prepared. These were introduced after apreliminary ball-milling process, which deformed theinitially spherical grains in order to improve the sinter-ing behaviour of the mixture. Near full density sampleswere consolidated by hot pressing the powder mixturesfor 90 min at 1350°C under a reducing hydrogen atmo-sphere and with a pressure of 25 MPa.

2.2. Microstructural and mechanical techniques

Specimen bars were obtained from the original hot-pressed disk for the mechanical characterization. Bothelastic modulus, E, and strength, sf, were measured ona universal mechanical testing machine equipped with ahigh temperature furnace. Tests were performed in airat temperature up to 1000°C using silicon carbidefixtures. The heating rate was controlled at :160°Cmin−1 and sample was tested :10 min after theselected temperature was reached.

The elastic modulus was measured by four-pointbending tests with inner and outer span equal to 20 and40 mm, respectively. Bars with sizes equal to 45 mm×4mm×3 mm were used in these tests. The elastic mod-ulus was evaluated from the deflection of the barmeasured by an extensometer. At least four measure-ments were perfomed at each temperature.

The bending strength was evaluated from three-pointbending tests. A span equal to 20 mm was used.Specimen sizes were around 25 mm×3 mm×2 mm.The prospective tensile face was previously polishedand the edges were chamfered using diamond paste upto 3 mm. Four measurements were performed at eachtemperature.

Fracture toughness, Kc, was measured at room tem-perature by an indentation technique. Vickers indenta-tions were produced on polished surfaces of thecomposite using loads of 98 and 294 N. The length ofthe radial cracks developed from the corner of the

indentation site were measured and was calculated onthe basis of the formula proposed by Anstis et al. [9]:

KC=0.016�E

H�0.5 P

c1.5 (1)

where P is the indentation load, c, the radial cracklength and H, the Vickers hardness.

Fracture toughness was also measured at higher tem-peratures by flexure of notched beams. Through-thick-ness sharp cracks were introduced in 25 mm×3mm×2 mm bars by using the technique presented byPancheri et al. (unpublished). Crack lengths in therange 1 mm–2 mm were obtained. Notched spacimenswere then subjected to three-point bending tests and Kc

was evaluated from the maximum load and the cracklength by the arguments proposed by Nose and Fujii[10]. At least two measurements were performed at eachtemperatures.

3. Results and discussion

The elastic modulus as a function of testing tempera-ture is shown in Fig. 1. The elastic modulus is constantfor temperatures up to 800°C. It abruptly decreases athigher temperatures. This decrease can be related to thestiffness variation of the intermetallic dispersion withtemperature [8]. The experimental value measured atroom temperature can be compared with the valuewhich can be calculated on the basis of the elasticmoduli of pure polycrystalline Al2O3 and Ni3Al. Valuesequal to 397 and 179 GPa can be taken for the ceramicand the intermetallic compound, respectively. The cal-culation can be accomplished by using Voigt and Reussmodels which correspond to the upper and lower limitfor elastic modulus [11]. The Voigt model (rule ofmixtures) furnishes a value of equal to 375 GPa whileby the Reuss model a value of 354 GPa can be calcu-lated. The experimental value (33793 GPa) is slightlylower than the theoretical values. This disagreementcan be related to the residual porosity (:2%) presentin the hot pressed material. Such porosity results froma minor mismatch in the thermal expansion coefficients(TEC) of the intermetallic and ceramic phases. As amatter of fact, the intermetallic compound has a TECslightly larger than alumina (:10.10−6°C−1 for Ni3Alas compared to :8.10−6°C−1 for Al2O3 [12]). There-fore, tensile stresses are created upon cooling after thehot pressing process. These stresses can promote theformation of microcracks, as shown in the SEM mi-crograph reported in Fig. 2, and decrease the strengthof the interface.

The evolution of the bending strength, sf, with tem-perature is shown in Fig. 3. A different behaviour, ascompared to the elastic modulus, is shown by sf, whichin fact continuously decreases as the testing tempera-ture increases.

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Fig. 1. Evolution of the elastic modulus, E, as a function of the testing temperature.

Eq. (1) was used for the determination of fracturetoughness at room temperature by indentation. Vickershardness was measured to be equal to 15.091.2 GPa.This value is lower than the hardness of pure polycrys-talline alumina (:16–18 GPa) and can be related tothe presence of 10 vol% of Ni3Al whose hardness canbe estimated to be equal to 1.8–2.0 GPa as the yieldstress of Ni3Al at room temperature is around 600 MPa[8]. Values of fracture toughness equal to 2.890.4 MPam and 3.290.3 MPam were measured at98 and 294 N, respectively. It must be reminded thatfracture toughness measured by indentation is usuallyconsidered as the fracture toughness associated to smalldefects. The artificially high crack opening displacement(COD) values as found in Vickers indentation crackssometimes prevent the development of the tougheningmechanisms [13,14] related to bridging effects between

crack faces [15]. In this respect Kc measured by indenta-tion should not be greatly affected by any tougheningmechanism [13]. In fact, fracture toughness for purealumina is around 3 MPam [17]. Conversely, Kc val-ues measured by flexure of precracked beams allow topoint out the presence of toughening effects. The frac-ture toughness measured on precracked specimens atvarious temperatures is shown in Fig. 4. Similar trend isshown by sf and Kc as function of testing temperatureat least up to 800°C (Figs. 3 and 4).

At room temperature a fracture toughness value of:7 MPam are obtained. This value is sensibly largerthan the results obtained by indentation and more thandouble with respect to typical Kc values for pure alu-mina [16,17]. These results reveal the presence of anR-curve effect, i.e. of a strong toughening mechanismwhich can be associated to the plastic deformation ofthe intermetallic particles.

The fracture surfaces obtained at room temperature,600 and 1000°C from the specimens used for Kc and sf

measurement are shown in Fig. 5. Ni3Al particles arealigned along a prefential direction, perpendicular tothe symmetry axis of the original disk obtained fromthe hot pressing procedure. Different features are ob-served at various temperatures. First of all, at increas-ing testing temperatures a larger number of Ni3Alparticles are pulled-out resulting in a larger number ofcavities.

A closer look at the intermetallic second phase showssevere plastic deformation at room temperature (Fig.6(a)). Ni3Al particles are stretched to failure by neckingto a line. These particles act as plastic bridges [15]during crack propagation and account for the measuredrelatively high fracture toughness of the composite. The

Fig. 2. Secondary electron SEM micrograph showing microcracks atthe Al2O3 matrix (darker)–Ni3Al particle (brighter) interface.

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Fig. 3. Evolution of the bending strength, sf, as a function of the testing temperature.

observation of fracture surfaces allows point out thatthis plastic deformation is always accompanied by aclear debonding from the alumina matrix. Plastic defor-mation is more evident for irregularly shaped particles.When the shape of the intermetallic particles is smooth,fracture becomes brittle and no debonding is observed(Fig. 6(b)). These observations reveal that the mechani-cal interlocking of the Ni3Al particles within the alu-mina matrix represents a stronger effect than thechemical bond between the two phases. Owing to thehighly reducing hot-pressing atmosphere, no significantreactions are expected to occur at the metal-ceramicinterface, as observed for similar systems, e.g. Al2O3-Ni[18], when treated in oxidizing atmospheres. Thestrength of the interface is also decreased by thermalstresses and microcracks previously discussed (Fig. 2)caused by the difference in the thermal expansion co-efficients of the two phases.

At intermediate temperatures, e.g. 500–600°C, thefractured intermetallic particles are considerably de-formed. Moreover debonding, determined by the weak-ening of the mechanical interlocking between particlesand ceramic matrix, is present (Fig. 7).

At higher temperature a lower amount of plasticdeformation is observed in Ni3Al particles. Moreover,debonding at the interface can be observed on simplypulled-out particles. Fig. 8 shows detail of the fracturesurface of a sample tested at 1000°C. The Ni3Al particlehas kept its original shape and does not appear signifi-cantly deformed. This is apparently in contrast with thetypical behaviour of Ni3Al compound at high tempera-ture. This intermetallic compound has a yield strengthwhich is constant up to 700°C and then decreases for

higher temperatures. Conversely, the total tensile elon-gation initially decreases up to 800°C and then in-creases sensibly for higher temperatures [8]. The resultsof this work may suggest that the load transfer at themetal-oxide interface decreases as the temperature isincreased. Therefore, the toughening effect which isactive at room temperature is partially lost at hightemperature and this is reflected by the decrease of Kc

and the failure stress.The effect of testing temperature is also important on

the fracture morphology of the alumina matrix. Atroom temperature fracture is mainly transgranularwhile, at higher temperatures, larger amount of inter-granular failure can be detected. Apparently, the addi-tion of Ni3Al has a significant effect on the aluminagrain boundary strength at various temperatures. Thestrength and fracture toughness of pure polycrystallinealumina is substantially constant for temperatures up to1000°C and no transition fromtransgranular to inter-granular fracture is observed [16]. Conversely, the oc-currence of intergranular fracture at high temperaturein the composite used in this work is an indication ofthe reduction of alumina grain boundary strength andfracture energy. This effect contributes to the decreas-ing of strength shown in Fig. 3.

The toughening effect of the intermetallic phase canbe quantified on the basis of the results of previoustheoretical analysis [15,17]. When a ductile phase isdispersed within a brittle matrix the toughening incre-ment is proportional to the yield stress sp

y, the rupturestrain, op, the volume fraction, Vp and the radius, rp, ofthe dispersed particles through the relation [17]:

DKc=aEspyopVprp (2)

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Fig. 4. Fracture toughness, Kc, as a function of the testing temperature.

where E is the elastic modulus of the composite and a

is a numerical coefficient (:1) depending on the modeof particle detachment from the matrix, the work hard-ening rate, etc. [15,17]. Substituting the values corre-sponding to the composite considered in this work inEq. (2) (sp

y:600 MPa [8], op:0.4 [8], Vp:0.1, rp:20 mm, E:340 GPa) the toughening increment at roomtemperature is equal to :13 MPam. This value islarger than the measured fracture toughness. Neverthe-less, from SEM observation it can be estimated thatabout one third only of the intermetallic particles un-dergo plastic deformation and, therefore, participate tothe toughening process (Fig. 5(a)). Therefore, the valuefor Vp should be corrected accordingly. The tougheningincrement, DKc, becomes :6 MPam and the totalfracture toughness is :9 MPam which is in goodagreement with measured Kc values. The presence ofresidual stresses, the irregular shape of the Ni3Al parti-cles and the imperfect load transfer between the parti-cles and the matrix can account for the discrepancybetween calculated and measured fracture toughnessvalues.

At higher temperatures the toughening effect dependsstrongly on the yielding and rupture behaviour of theNi3Al particles. The decrease of the fracture toughness,as well as that of failure stress, is smooth up to 600°C.Then, both Kc and sf show a marked drop between 600and 800°C. This behaviour can be correlated to themarked decrease of yield stress shown by Ni3Al in thistemperature range [8]. sp

y is equal to :600 MPa and:400 MPa at 600 and 800°C, respectively, while noevident variations are detected for op. In addition, thedecrease of Kc can be also influenced by the fracturebehaviour of the alumina matrix as already pointed

out. At 1000°C the rupture strain of the intermetallicphase sensibly increases (:1) [8] and this can accountfor the slight increase of Kc between 800 and 1000°C.Correspondingly, the bending strength remains almostconstant in this temperature range (Fig. 3).

These observation allow to point out that a strongtoughening effect has been achieved by adding theNi3Al particles to the alumina matrix. Though a chem-ical bond has not been fully developed between the twophases, the interlocking effect of the intermetallic parti-cles correlated to their irregular shape is sufficient forthe development of toughening mechanisms by plasticdeformation of the second phase. Due to the puremechanical nature of the interaction between Ni3Alparticles and alumina, the toughening effect decreasesat 800°C, even though the experimental Kc value stillremains larger than typical values for fine grained-alu-mina. The possibility of introducing an intermediateconnecting phase between the matrix and the particlescertainly would allow to obtain larger Kc values both atroom and at high temperature.

4. Conclusions

An alumina matrix composite containing 10 vol%Ni3Al base alloy powder was prepared by hot-pressing.Elastic modulus, flexural strength and fracture tough-ness were measured over a temperature interval rangingfrom 25 to 1000°C. Flexural strength is only slightlyaffected by the temperature up to 600°C as a conse-quence of the toughening effect of the metallic particu-late. A good mechanical interlocking between thematrix and the intermetallic phase, which allows load

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transfer and plastic deformation of the metal, is the

Fig. 6. Secondary electron SEM micrographs showing details of thefracture surface of the specimen tested at room temperature. (a)Irregurarly shaped particle which undergoes heavy plastic deforma-tion. An empty site is also visible. (b) Debonded particle featuring asmooth surface on which the traces of the matrix grains are visible.

Fig. 5. Secondary electron SEM micrographs showing fracture sur-faces of the samples tested at: (a) room temperature; (b) 600°C; and(c) 1000°C.

dominant mechanism at this stage. At higher tempera-tures a softening of the metallic component occurs and

Fig. 7. Secondary electron SEM micrograph showing the fracturesurface of the sample tested at 500°C. The intermetallic particle isconsistently plastically deformed to fracture.

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Fig. 8. Secondary electron SEM micrograph showing the fracturesurface of the sample tested at 1000°C.

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this reduces the adhesion strength between matrix andparticles. Correspondingly a reduction in the fracturetoughness is also observed. However Kc still remains athigher values than those typically measured for purealumina.

Acknowledgements

We wish to thank Dr E. Degasperi for his collabora-tion during mechanical testing.

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