Ni3AlAl2O3 composites with interpenetrating networks

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Ni,Al/A&O, COMPOSITES WITH INTERPENETRATING NETWORKS J. R(idel’, H. Prielipp2, N. Claussen2, M. Sternitzke3, K.B. Alexander4, P.F. Becher4 and J.H. Schneibe14 ‘FG Nichtmetallisch-Anorganische Werkatoffe, TH Darmstadt, 64295 Darmstadt, Germany 2Advanced Ceramics Group, TU Hamburg-Harburg, 2 1073 Hamburg, Germany 3Department of Materials, University of Oxford, Oxford OX1 3PH, U.K. 4Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 3783 1, USA (Received April 6, 1995) (Revised May 3, 1995) Introduction

Ni,Al is an ordered intermetallic with the L 1z structure and is the major precipitate phase in nickel-base superalloys. Single-phase, single-cry&line N&Alis very ductile (up to 40 % plastic elongation) at room temperature. PolycrystallineN&Alalloys, on the other hand, are very brittle at room temperature in laboratory air unless their album concentration remains below 25 at. % and they have been doped with small amounts of boron (1). Ofparticular interest is the mechanical behavior: the yield strength of Ni,Al increases with temperatum until it starts to fall off around 800 to 900 K. This feature makes this intermetallic attractive for high temperature applications. In recent years, there has been considerable interest in reinforcing Ni,Al alloys with Alz03, or toughening A&O3with N&Al.First, A&O3is thermodynamically compatible with Ni,Al. Second, A&O, additions to Ni,Al lower the density, and they may result in additional strengthening. For example, McKamey and Lee (2) extruded Ni,Al powders mixed with chopped A&O, fibers. In some cases they observed improvements in the room temperature yield strength However, since the Ni,Al grain sizes were quite small, the high temperature properties were inadequate. Schneibel et al. (3), Ringer and White (4), and Nourbakash et al. (5) incorporated continuous single crystal A&O, fibres into Ni,Al matrices. Jn the first two pieces of work, this was achieved by hot-pressing layers of fibers and foils. Nourbakash et al., on the other hand, intlltrated A&O, fibre lay-ups with liquid Ni,Al. Although Schneibel et al. found some evidence for mechanical property improvements, there appeared to be a number of problems. Nourbakash et al., in particular, found relatively weak inter-facial bond strengths ranging f!om 19 MPa in binaty N&Alto 55 MPa in Ni,Al alloyed with Cr. They found also that pmcessing decreased the strength of the A&O, fibres markedly, and that it even led to the formation of compression twins in some of the fibers. Also, minor alloying additions to the Ni,Al such as Zr or Cr resulted in reactions with the AJO,. A somewhat more promising field appears to be the toughening of A&O, with Ni,Al. Alexander et al. (6) chose a powder-metallurgical route to fabricate A&O, containing Ni,Al inclusions. Suitably processed composites approached fmcture to@nesses of 8 MPa mln, while retaining bend strengths above 500 MPa. Additional hpmwmnt in the fhcture tow might be achievable if the Ni,Al phase would be interconnected, since interfacial debonding would, in this case, not limit the achievable toughening. The desire to create such 843

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interpenetrating structures was the driving force for the present work. An alloy with the compositions Ni22Al-lZr-O.lC-O.lB (at %) was chosen, since previous work (7) showed this ahoy to wet and to bond to Al,O, reasonably well. Metal/ceramic composites with interpenetrating networks allow tailoring of fracture strength and fracture toughness, dependmg on ligament size and metal volume fraction (8). A recently developed process based on gas pressure infiltration of the liquid metal into a porous ceramic followed by pulling the composite out of the metallic melt appears to provide good versatility in processing these materials (8) and provides further technological improvement over earlier methods in producing these materials (9,lO).

Porous alumina preforms were prepared using a coarse alumina powder of grain size 4-5 urn (Alcoa CL SOOO), slip casting it into plates and sintering at 1923 K for 6 hours to yield a density of 65 vol. %. The plates of size 30 mm x 40 mm x 5 mm were attachedto a holder and embedded into chips of the intermetallic, which in turn were positioned inside an alumina crucible. The infiltration furnace was heated up in vacuum with a rate of 15 K&n to 1923 K. After a holding time of 10 min the argon gas pressure was raised with a rate of 2 MPa/min to 16 MPa with a subsequent holding time of 15 min to allow for complete infiltration. Subsequently the furnace was cooled down at a rate of 5 Wmin with a hold at 1823 K to lower the metal bath with the in-built hydraulic ram, thus pulling the composite out of the molten bath. Microstructural characterizationwas performed by optical microscopy of polished surfaces in conjunction with a quantitative determination of phase contents. Specimen were also prepared for bright field and dark field analytical transmission electron microscopy (Philips CM 20) with ion milling being performed using a cold stage specimen holder. The plates were cut and polished to a 1 urn surface finish to yield bend bars of size 25 mm x 4 mm x 3 mm to be used for both fracture strength and fracture toughness measurements. Room temperature mechanical properties were obtained either in 4-pt. bending or 3-pt. bending, the former with spans 10 mm and 20 mm, the latter with a span of 20 mm Fracture toughness at room temperature was measured using the SEPB method with starter notches about 500 pm long and 150 urn wide. A load of 50 kN in the bridge configuration yielded precracks of about 250 pm length. Then the bridged crack was cut back to have a crack about 25 % of notch length, which ensured small scale yielding conditions with mesurements providing values close to plateau toughness. Preliminary R-curves were obtained by in-situ crack extension in a SBM using a 4-pt. bend geometry.In these experiments, starter radial cracks from Vickers indents of load of either 9.81 or 18.62 N were utilized. The stress intensity factor was calculated by incorporating the Chantikul et al. (11) calibration of crack shape and residual indentation stress. Chevron notch specimens in 3-pt. bending were used to measure fracture toughness both at room temperature and at 1073 K in argon. Fracture toughness was evaluated from absorbed energy, area of the chevron and the plane strain Young’s modulus. Young’smodulus (340 GPa) was calculated horn the single phase values of Al,O, (380 GPa) and N&Al (180 GPa) using a simple linear relationship and a Poisson’s ratio of 0.25. Fractography was finally performed to study the t?acture path, the degree of interface debonding and the degree of ductile deformation of the intermetallic phase. Results

The homogeneous microstructure of the composite is shown in Fig. 1. The grain size of the alumina was about 5 pm (as in the startingpowder). The porosity was evaluated to be about 0.5 vol.-%, the alumina 79.5 vol-% and the intermetallicphase 20 vol.%. This phase distribution found by quantitative optical microscopy can be used to predict a density for the composite of 4.68 g/cm’, which was found to be in good agreement with the experimentally obtained density of the composite of 4.73 g/cm’. A TEM dark field micrograph of the composite is shown in Fig. 2. The N&Al (222) diffraction spot was chosen to image the intermetallicphase, which, in the chosen orientation, appears bright (as in Fig. 2). Several

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Figure 1. Optical micrograph of the Ni,AvA1,0, composite. Brightphase is the intermetaUic phase, dark phase is the ceramic.

Figure 2. TEM dark field micrc-graph of the Ni,AlIAl,O, composite. Bright phase is the intermetallic.

representative areas were inspected and gave an average gram size of the intermetallic phase of 20-30 pm and a ligament diameter of 1 to 5 urn. Additionally, selected area dtiaction showed that the orientation of grains of the intermetalhc phase was within 10 o over areas of about 100 t.mr.No indications for microcracking were found. Me&anicai property data of the composite are collated in Table 1. The fracture strength in 4-pt. bending was found to be slightlyhigher than 300 MPa, with the 3-pt. bending result lying somewhat higher due to the smaller volume being exposed to peak tensile stresses. The fracture strength at 1073 K was comparable to the room temperature values. Fracture toughness values as gained by SEPB and the chevron notch technique provided similar results witb values lying between 11.2 and 13.4 MPa mm. The fracture toughness at 1073 K was comparable to that at room temperature.

TABLE 1 Results of Mechanical Evaluation at Room Temperature as well as at 1073 K

Testing Method

Temperature

Fracture Strength MPa

4-pt. bending, polished and beveled tensile surface

R.T.

305 315

3-pt. bending, as ground specimens

R.T.

403 362 401 381

1073 K/Argon

Fracture Toughness MPA ml”

SEPB, 4-pt. bending

R.T.

11.5 11.3

Chevron-notch 3-pt. bending

R.T.

13.4 11.6 11.3 11.2

1073 K/Argon

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Steady state -

kJ 10

AI,O,

0

infiltrated

with

200 400 600 800 Crack Length, Frn

1000

Figure 3. R-curve results from the Ni,Al/Al,O, composite, with the triangles from the data using the 9.81 N indent, and the squares using the 18.62 N indent.

Results of the R-curve measurements are provided in Fig. 3 and compared with the average steady state room temperature toughness values of the SEPB and chevron notch technique. The toughness values are shown to rise from about 2 MPa ml’* at 400 pm crack length to 10 MPa ml’* at 900 m crack length. Microstructural observations during in-situ crack propagation showed that N&Al bridging ligaments were formed. Fracture surfaces of room temperature fracture are shown in Fig. 4. Complete debonding can be seen in regions marked by an arrow. Alumina grains are marked by ,,A“, N&Al grams are marked by ,,N“. The intermetallic phase is seen to show large ductility in the ligaments. Discussion I&ration of the porous alumina with the intermetallic phase is accompanied by further shrinkage of the porous preform during the process. The hnal composite exhibits a very homogenous microstructure (Fig. 1) with a grain size of the N&Al of 20-30 micron which is about a factor of 6 to 10 larger than the average ligament size. The ligaments are therefore predominantly in single crystalline form and deform in a completely ductile manner (Fig. 4). Though boron is also present as an alloying element in the inter-metallic phase, this ductility is therefore attributed to the single crystal deformation of N&Al. The strength of the interface between N&Al and A&O, appeared low and leads to complete debonding during crack propagation. This is in accordance with the results of Nourbakash (5) et al., who found weak interfacial bonds between N&Al and Al,O,. Complete debonding allows for unconstrained plastic deformation in the intermetallic ligaments. Weak bonds lead furthermore to interfacial fracture, as in Cu/Al,O,, as compared to transcrystallme fiacture of the ceramic, as in Al/Al,O, (12). A small inter-facial fracture resistance between the ductile reinforcement and the ceramic phase promotes unhindered crack initiation at large pores filled with the reinforcing phase and yields rather small fracture strengths. Comparable effects were found in Al/A&O, composites as compared to Cu/A&O, composites containing either large or small ligaments (12). The former had very high fracture strengths exceeding 800 MPa, while the latter had smaller fracture strengths, which were reduced even more at large ligament sizes, where microcracking at the interface and therefore preferential tYacture sites were also reported (12). The fracture strength values of the NiflA120, composite can be rationalized by the increased fracture tough-

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ness at crack instability brought about by the ductile ligament bridging, which raises the strength from about 150 MPa for the porous alumina (8) to more than 300 MPa in the composite. For metal reinforced ceramics which exhibit complete debonding the total toughening effect of the metal inclusions can be rationahzed by assuming scale invariant plastic deformation and using the stress-strain curve of macroscopic single crystal d&rmation. The energy absorbed in plastically stretching the intermetallic phase (R,,) can be written as a function of the area fraction of the reinforcement (f = 0.2) and the integral over the stresses @) taken up in the bridging ligament as a function of the crack opening displacement 2u (13), Eq. 1:

Since we cau only compare our theoretical prediction to a plateau toughness value, we compute the total energy absorbed in the ductile ligaments until i?acture with the energy taken corn the stress-strain curves measured at either room temperature (14) or 973 K for a first estimate for the high temperature behaviour (15). Since the stress-strain curves for N&Al are highly orientation dependent, only a range of obtainable absorbed energies R, can be computed. R,, furthermore scales with ligament diameter (8). Since the average ligament diameter is not accurately known this provides a further uncertainty. As a first estimate, this diameter is here taken as 3 pm and the gauge length over which deformation occurs taken as the grain size of the alumina (5 pm). This, in combination with the stress-strain curves for the literature yields a stress-crack opening displacement function p(u). Using the range of energies absorbed at room temperature (14) and at 973 K (15) for a 5 pm long ligament with diameter 3 pm aud an area traction of 20 %, we find a shielding term R, of 370400 J/m’ for room temperature and of 100-450 J/m2 for 973 K. Correlating fracture energies and fracture toughness (13) and computing the equivalent shielding term in fracture toughness notation, we find (Eq. 2):

Kp=

/ K;

+ E’R

I

-

K,

Using & = 2.0 Ml’s rnlR as crack tip toughness from the only reported crack tip toughness value in metal reinforced ceramics (16), we obtain finally plateau toughness ranges for room temperature of 11.4- 11.8 MPa mm and for 973 K of 6.2- 12.5 MPa ml’. These predictions are found to be consistent with our experimentally obtained values. Summary Homogenous microstructures of Ni,Al/A,O, composites can be produced by gas pressure infiltration of the intermetallic phase into the porous ceramic. Since the nickel ahuninide is single crystalline in the reinforcing ligaments, it exhibits ductile deformation and provides appreciable toughening for the composite. The mechanical property data of this material at room temperature and at 1073 K are comparable. AcknowledPements J.R., H.P. and N.C. acknowledge

support by the Volkswagen Foundation under contract number I 66/790.

K.B. A., PEB. and J.H.S. acknowledge support of this research by the US Department of Energy, Assistant Secretary for Energy Efhciency and Renewable Energy, O&e of Industrial Technologies, Advanced Industrial Materials (AIM) Program, and by the Division of Materials Sciences, US Department of Energy, under contract DE-AC05-84or2 1400 with Martin Marietta Energy Systems, Inc.

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C.T. Liu and KS. Kumar, JOM 45,38 (1993). C.G. McKamey and E.H. Lee., in ,,Intermetaltic Matrix Composites“, D.L: Anton, P.L. Martin, D.B. Miracle, and R McMeeking+ eds., MRS vol. 194, p. 163 (1990). J.H. Schneibel, E.P. George, C.G. McKamey, E.K. Ohriner, M.L. Sante118 and CA Carmichael, J. Mater. Res. 6,1673 (1991) D. Ringer and C.L. White, in ,,Intennetallic Matrix Composites III“, J.A Graves, RR Bowman, and J.J. Lewandowski, eds., Materials Research Sot., Vol. 350, p. 125 (1994). S. Nourbakbsb, W.H. F&e, 0. Sabin and H. Margolii Metall. Trans. A, 25A 1259 (1994). KB. Alexander, HT. L;1, J.H Sdmeibel and P.F. Becher, to be published in ,,Procesing and Fabrication of Advanced Mate&k“, V.A Ravi, T.S. Srivatsanand J.J. Moore, eds. J.H Scbneibel and KB. Alexander, in ,&ermetallic Matrix Composites III“, J.A Graves, RR Bowman and J.J. Lewandowski, eds., Materials Research Sot., Vol. 350, p. 255 (1994). H. Prielipp, M. Knechtel, N. Clausxn, S.K. Streiffer, H. Mitllejans, M. Rtihle and J. Rijdel, Mat. Sci. & Eng. A, in press. F.F. Lange, B.V. Velamakak and AG. Evans, J. Amer. Ceram. Sot. 73,388 (1990). N.A Travitzky andN. Clausen, J. Eur. Ceram. Sot. 9,61 (1992). P. Chantikul, G.R An&, B.R Lawnand D.B. Marshall, J. Amer. Ceram. Sot. 64,539 (1981). M. Knechtel, H. Prielipp, H. Mtkllejans, N. Claussen and J. R&l, Ser. Met. et Mat. 31,1085 (1994). B.R Lawn, Fracture of Brittle Solids, Cambridge University Press (1993). M.S. Kim, S. Hanada, S. Watanabe andO. Izumi, Actametall. 36,2615 (1988). M.S.Kim,S.Hanada,S. WatanabeandO.Izumi,Actametall.36,2967(1988). and N. Claussen, J. Eur. Ceram. Sot., 14,153 (1994). J. Radel, M. Sin&l, M. Dransmann, R!kinbrech