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Rapid synthesis of Al2O3 reinforced Fe–Cr–Ni composites
Materials Science and Engineering A344 (2003) 245 /252 www.elsevier.com/locate/msea
Rapid synthesis of Al2O3 reinforced Fe Cr Ni composites /
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N. Travitzky a, , P. Kumar a, K.H. Sandhage b, R. Janssen a, N. Claussen a a
b
Advanced Ceramics Group, Technische Universita¨t Hamburg-Harburg, D-21073 Hamburg, Germany Department of Materials Science and Engineering, The Ohio State University, Columbus, OH 43210, USA Received 20 March 2002; received in revised form 3 June 2002
Abstract Short distance infiltration (sdi) was used to fabricate metal matrix composites (MMC) with very fine ceramic reinforcements. This ‘sdi-MMC’ process was used to synthesize a dense composite comprised of Fe /Cr /Ni/Al2O3. Fe2O3/Fe/Ni/Cr/Al powder compacts were rapidly heated to 900 8C and held at this temperature and a pressure of 20 MPa in air for 5 min. This thermomechanical treatment triggered the highly-exothermic thermite reaction between starting powders that, in turn, resulted in rapid in-situ formation of a composite comprised predominantly of g(fcc) Fe /Cr /Ni, along with smaller amounts of a(bcc) Cr /Fe /Ni and aalumina. SEM and TEM analyses revealed a microstructure consisting of a uniformly dispersed network of very fine (a few hundred nanometers) Al2O3 grains in a matrix of a fine-grained ( /5 microns) metallic alloy. The dense composites exhibit average bend strength and toughness values of 1100 MPa and 18 MPa (m)1/2, respectively. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Reactive syntheses; Metal matrix composites; Ceramic reinforcement
1. Introduction Ceramic-reinforced, metal-matrix composites (MMCs) with enhanced resistance to fracture, creep and oxidation are attractive for automotive, aerospace and defense applications [1 /4]. Development of advanced MMCs depends on the availability of practical processing routes. MMCs may be produced by powder metallurgical or melt casting routes, depending on process efficiency, production cost and the macrostructure (shape, surface features) and microstructure desired for the final product. High-melting MMCs with attractive microstructures (e.g. fine-scale and uniformly-dispersed phases) and with a wide range of compositions may be produced by the compaction and forming of finely-divided powder mixtures [5,6]. The widespread use of powder processing has been limited by relatively slow and expensive consolidation/forming (e.g. hot pressing, hot extrusion) and machining steps. Melt casting approaches (e.g. compocasting or melt infiltra-
Corresponding author. Tel.: /49-40-428-78-3459; fax: /49-40428-78-3427 E-mail address: [email protected] (N. Travitzky).
tion), on the other hand, are capable of rapidly producing dense composites with complex and near net shapes [2]. However, the casting of high-melting MMCs, reinforced with a fine and continuous ceramic phase, is relatively complicated and expensive. New cost-effective (low-temperature, rapid) processing routes need to be developed for the direct fabrication of dense, high-temperature MMCs. The melt casting temperature may be dramatically lowered by using a chemical reaction between a lowmelting metallic liquid and a shaped solid preform to generate a higher-melting composite. This concept has been utilized in a number of reactive melt casting techniques [7 /17]. For example, in infiltration alumina aluminide alloy (i-3A) and displasive compensation of porosity (DCP) processes, low-melting metallic liquids (e.g. pure molten Al or Mg; or low-melting alloy liquids) are allowed to infiltrate into and undergo displacement reactions with, porous preforms to obtain dense, near net-shaped composites containing higher-melting ceramic and metallic (or intermetallic) phases [10 /17]. In reactive melt penetration (RMP), a low-melting metallic liquid (e.g. molten Al) undergoes a displacement reaction with, and penetrates into, a dense oxide preform to yield a co-continuous ceramic/metallic or ceramic/inter-
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 4 1 9 - 7
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metallic composite [18 /20]. These reactive casting approaches have been used to produce a variety of dense composites (e.g. Al2O3/TiAl3, Al2O3/NbAl3, MgAl2O4/Fe /Ni /Al alloy, MgO/FeAl and ZrC/W composites) at temperatures 5/1200 8C. The composition and phase content of a composite produced by any particular reactive casting method is dependent upon and limited by: (i) the reaction-induced volume change; and (ii) the porosity of the reacting preform. With RMP, the volume fraction of the metal phase in the final composite is largely determined by the space made available as the reactant oxide is converted into a product oxide. Consequently, RMP tends to yield composites with relatively high ceramic contents of ]/65 vol.% Al2O3 [18 /20]. With i-3A and DCP, porosity present in the starting preforms can provide additional space for the metallic or intermetallic phase in the final composite. Indeed, preform porosity must be carefully tailored with these latter two processes to produce composites with desired amounts of metallic or intermetallic phases [10 /17]. Melt processing methods that can yield a wider range of compositions and phase contents, without the need for preforms of tailored porosity, should be examined. The processing time for reactive melt casting is strongly influenced by porosity and the dimensions of the preform [12,14,18 /20]. With RMP, the time required for complete reaction and penetration of the metallic liquid into the dense preform scales with the preform size, and tends to be relatively long for thick preforms (e.g. /4 h for 2.5 cm thick preforms at 1100 8C [19]). With i-3A and DCP, reactant liquids are dispersed throughout porous preforms by infiltration prior to reaction, so the subsequent reaction occurs more rapidly and uniformly throughout preforms. However, the time required for complete melt infiltration with i-3A or DCP still depends on the dimensions and porosity (i.e. the pore fraction and size) of the preform. Self-propagating high-temperature synthesis (SHS or combustion synthesis) method is among the fastest reaction-based processes. The exothermic heat of reactions rapidly converts compacted powder mixtures into products [21]. Low energy requirements due to the self generation of heat, short processing time, possibility to form metastable phases due to high thermal gradients and rapid cooling rates are the major advantages of SHS process. Thermite reactions are a sub-set of SHS processes [22]. Thermite reactions have the advantage over the reactions with elemental reagents in that they start with naturally occurring oxides that are less expensive and more readily available than elemental reactant powder. Recently, thermite-based reactions were used to fabricate dense and homogenous alumina aluminide alloys (3A) with interpenetrating networks of Al2O3 and aluminides of Ti, Fe, Nb, etc. [18,19]. The
composites were fabricated by pressureless reaction sintering in a nonoxidizing atmosphere using intensively milled metal oxide-aluminium powder compacts. The exothermic in 3A necessitate careful control of the processing parameters. Rapid temperature increases due to these reactions result in the formation of liquid phases, which may cause large pores that cannot be removed during pressureless sintering [23,24]. Since full density is of paramount significance for most structural materials, residual porosity is the main hindrance for the commercial use of SHS. In order to obtain composites with reduced porosity, several authors have conducted SHS reactions under the application of an external pressure [25,26]. For example, Holt and Munir fabricated 95% dense TiC by ignition under a mild pressure of 27.6 MPa [26]. Gutmanas and Gotman demonstrated that the application of a modest external pressure can considerably reduce the porosity in alumina/aluminide composites [25]. This method, called short distance infiltration-sdi was used to fabricate a :/99% dense composite of Al2O3 and TiAlx phases (Ti3Al, TiAl, TiAl2) by igniting a porous powder compact of Al and TiO2 (molar Al:TiO2 ratio /7:3) in a die preheated to 950 8C under a pressure of 100 MPa. The equiaxed Al2O3 grains were well-dispersed with the TiAlx phases, were of the order of a few microns in size and were present in a high concentration (note: complete reaction of an Al:TiO2 /7:3 mixture should yield a TiAl/Al2O3 composite comprised of 46.9 vol.% Al2O3 and 53.1 vol.% TiAl). In the above works, in situ processing techniques lead to composite materials with significant ceramic contents (i.e. /45 vol.%). However, fabrication of metal/ceramic composites with lower ceramic content has not been demonstrated to date. Unlike RMP, i-3A and DCP, the time required to complete such short distance infiltration and reaction should be independent of the preform dimensions and shape. Furthermore, by controlling the ratio of solid precursor and liquid in the preform, the composition and phase content of the final composite: (i) can be precisely tailored; and (ii) should not depend on the porosity of the starting preform. It should be possible to produce composites with a wide range of compositions and phase contents (including MMCs) by short distance infiltration. The purpose of this work is to demonstrate that dense MMCs with B/30 vol.% ceramic reinforcement can be designed and fabricated by a rapid in situ sdi-metal matrix composites (MMC) by short distance infiltration (sdi) process. The specific objectives are: (i) to fabricate a dense MMC comprised of a Fe0.49Cr0.32Ni0.191 with 1 Note: this Fe:Cr:Ni ratio is similar to that for ASTM HL-type heat resistant alloys [27].
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interpenetrating nano-sized Al2O3 reinforcement phase; and (ii) to evaluate the micro-chemistry, microstructure and mechanical properties of such sdi-MMC-derived composites.
2. Experimental Powders of Al (B/45 mm, 99.5% purity; EckartWerke, Fuerth, Germany), Fe2O3 (15 mm, 98% purity; Jan de Poorter BV, Geertruidenberg, The Netherlands), Fe (4 /5 mm, 99.5% purity; BASF, Ludwigshafen, Germany), Cr (B/10 mm, 99.8% purity; Johnson Matthey GmbH, Karlsruhe, Germany) and Ni (3.0 mm, 99.7% purity; Aldrich Chemical Co., Milwaukee, WI) were mixed in a molar ratio of 2.0:1.0:2.9:3.2:1.9 (as per reaction Eq. (1) in Section 4). This metal/oxide powder mixture was then attrition milled (Model PE 075 Netzsch-Feintechnik, Selb, Germany) in acetone for 1 h at a rotation rate of 500 rpm using yttria-stabilized zirconia balls (3 mm diameter, 3Y-TZP, Tosoh Co., Tokyo, Japan) as milling media. Balls (1500 g) were used to mill 50 g of powder (i.e. a ball-to-charge weight ratio of 30:1) within an alumina milling vial (10 cm diameter). After overnight evaporation of the acetone, the milled mixture was passed through a 200-micron sieve. The sieved powder was then uniaxially pressed at 25 MPa into disks that, in turn, were cold isostatically pressed (CIPped) at a peak stress of 700 MPa. The as-CIPped disks ( :/15 mm diameter and 10 mm thick) disk were placed within two tightly-fitting stainless steel cells (i.e. two flat bottom cylindrical crucibles with internal diameters of 16 and 15 mm, respectively, heights of 10 mm, wall thickness of 0.5 mm and bottom thickness of 1 mm). The walls of each cell had been sprayed with a BN coating (BN-S, SINTEC Innovation in Keramik, Buching, Germany) to prevent chemical interaction with the specimen and with the preheated rams. The steel cells, containing the green compacts, were then placed into a hot vertical box furnace located within a conventional Zwick (Type 1478, Zwick, Ulm, Germany) testing frame. The box furnace was preheated to 900 8C in air prior to insertion of the steel cells. After insertion, a uniaxial load was applied to some of the sample assemblies at a rate of 30 N s 1 until a peak pressure of 20 MPa was achieved. The pressure was released after 5 min and the samples were quickly removed from the furnace and allowed to cool in air. Other specimens were exposed to similar thermal conditions without the application of an external pressure. After cooling to room temperature, the annealed samples were stripped out of the stainless steel cells. Microstructural analyses of the synthesized samples were conducted with a scanning electron microscope (Gemini 1530 SEM, LEO, Oberkochen, Germany) and a transmission electron microscope (2000 FX II TEM,
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JEOL, Tokyo, Japan). Microchemical analyses were conducted during SEM and TEM analyses with energydispersive X-ray spectroscopy. At least five different EDS spot analyses were conducted to obtain the average composition of a given phase. The samples for SEM/ EDS analyses were ground and polished to a 1 mm diamond finish. Electron transparent specimens for TEM analyses were prepared by grinding the sample into foils with a thickness of :/70 mm followed by ionbeam milling (GATAN DuoMill, Pleasanton, USA). Selected area diffraction (SAD/TEM) and X-ray diffraction (XRD) analyses were employed for phase identification. XRD analyses (Model PW 1710 Diffractometer, Philips Inc., Eindhoven, The Netherlands) were conducted using Cu-Ka radiation at a scan rate of 0.758 min 1. Differential thermal analyses (DTA) of green compacts and annealed composites were conducted in flowing Ar at a heating rate of 20 8C min 1 to 1500 8C (STA 409 DTA/TGA Netzsch Geraetebau GmbH, Selb, Germany). The densities of green compacts and annealed composites were obtained from weight and dimension measurements. Archimede’s method, using water as the immersion medium, was used to determine the relative porosity of the annealed composites. The average values of hardness, flexural strength and fracture toughness of the annealed composites were evaluated. Hardness tests were conducted with a standard Vickers apparatus (Type 3212, Zwick) using a load of 100 N applied for 15 s. The average hardness value was determined from ten indentation measurements. The flexural strength and fracture toughness was evaluated in three-point bending with the use of a universal testing machine (Type 1478, Zwick). The span and crosshead speed were maintained at 12 and 0.5 mm min 1, respectively, for all bend tests. The average values of flexural strength and fracture toughness were determined from measurements conducted on at least four bar-shaped samples each. Bars with dimensions of 2.5 /3.5 /14 mm were cut out of the fabricated composites and ground to a 15 mm diamond finish. The tensile surfaces of the samples were then polished to a 1 mm diamond finish prior to bending. Prior to a given fracture toughness test, a Vickers indentor was pushed into the specimen surface with a load of 300 N, as per the indentation strength in bending (ISB) technique [28]. The fracture surfaces and crack paths were examined by scanning electron microscopy.
3. Results Distinct peaks for Al, Fe2O3, Fe, Cr and Ni were detected by XRD analysis of an as-CIPped green compact. The average density of the as-CIPped green
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compacts was found to be :/4.65 g cm 3, which corresponds to :/75.8% of the theoretical density for the Al/Fe2O3/Fe/Cr/Ni mixtures. Within :/1 min of achieving the peak stress of the Al/ Fe2O3/Fe/Cr/Ni powder compacts into the preheated furnace, a distinct ‘pop’ was heard to emanate from the specimens, which was presumed to be a result of combustion. The applied pressure was maintained at a constant value after the combustion had occurred. Although the processing was conducted in ambient air, visual inspection of the synthesized specimens revealed no apparent surface damage (e.g. cracking, spallation) due to oxidation. After treatment under 20 MPa peak pressure, the disk-shaped specimens possessed diameter and thickness dimensions of :/17 and 7 mm, respectively. The average bulk density after such reaction was 6.50 g cm 3, which corresponded to a relative density of 94.2% of the theoretical value for a fully-converted Fe0.49Cr0.32Ni0.19/Al2O3 composite (see reaction Eq. (1) in the Section 4). XRD analyses of specimens annealed under pressure revealed the presence of two metallic phases along with Al2O3. The predominant metallic phase was identified as an austenitic Fe-rich (g, fcc) solid solution (Fe-SS), whereas the minor metallic phase was consistent with a Cr-rich (a, bcc) solid solution (Cr-SS). A secondary electron image of a polished cross-section of this composite material is shown in Fig. 1. The specimen appeared dense and comprised of a fine-grained (5/5 mm) metallic matrix with very fine alumina grains visible as dark threads winding throughout the matrix. Brightfield TEM images of the composite are shown in Fig. 2. The Fe- and Cr-rich alloy phases can be seen in Fig. 2, along with the nano-sized (bright) Al2O3 phase. EDS analyses revealed that the average composition of the Fe-rich phase was 50.0 at.% Fe, 29.9 at.% Cr and 20.1 at.% Ni. The secondary Cr-rich metallic phase possessed an average composition of 64.4 at.% Cr, 26.8 at.% Fe
and 8.8 at.% Ni. Aluminum was not detected in either metallic phase. Selected area electron diffraction (SAD) patterns of the Fe- and Cr-rich phases were consistent with g-Fe- and a-Cr-rich solid solutions, respectively. The average lattice constants of the Fe- and Cr-rich phases obtained from SAD patterns were 0.37 and 0.30 nm, respectively. The TEM image in Fig. 2(a) reveals that the austenitic Fe-rich phase possessed a much higher dislocation density than the Cr-rich phase. As seen in Fig. 2, the a-Al2O3 grains consisted of single alumina crystals connected in series along the grain length. The alumina grain size was of the order of a few hundred nanometers. A secondary electron image of a polished crosssection of a sample annealed at 900 8C without the application of pressure is shown in Fig. 3. An XRD pattern of this sample was similar to that of the composite prepared under pressure. However, the morphology of the alumina phase was quite different. The alumina in the composite of Fig. 3 did not exhibit the fine, elongated droplet-morphology prevalent in the composites annealed under pressure. Differential thermal analyses (DTA) obtained upon heating an as-CIPped green compact to 1500 8C at 20 8C min 1 in an argon atmosphere showed strong exothermic and endothermic peaks with onset temperatures of 605 and 1405 8C, respectively. The average value of room-temperature flexural strength for composites prepared under pressure was 11009/95 MPa. Some proportion of plastic deformation at fracture was obtained in all tested samples. The average value of fracture toughness was 189/0.5 MPa m1/2. The average Vickers hardness value of the pressure-annealed composites was 4.29/0.2 GPa. A typical fracture surface and a crack in a polished cross-section (where the crack was generated upon indentation with a Vickers indentor) of a pressureannealed Fe /Ni /Cr/Al2O3 composite are shown in Fig. 4.
4. Discussion An objective of the present work was to use the following net reaction to convert a porous Al/Fe2O3/Fe/ Cr/Ni preform into a dense MMC comprised of a Fe0.49Cr0.32Ni0.19 alloy reinforced with Al2O3: 2AlFe2 O3 2:9Fe3:2Cr1:9Ni [Al2 O3 10Fe0:49 Cr0:32 Ni0:19
(1)
The product of this reaction is a composite comprised of 73.0 vol.% Fe0.49Cr0.32Ni0.19 alloy and 27.0 vol.% Al2O3 [29,30]. This net reaction can be considered to be a combination of the following two reactions: Fig. 1. Secondary electron image of a polished cross section of Fealloy/Al2O3 composite prepared by sdi-MMC method.
2Al(l)Fe2 O3 [Al2 O3 2Fe
(2)
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Fig. 2. Brightfield TEM images of Fe-alloy/Al2O3 composite fabricated by sdi-MMC method. Magnification of (a) 23 k / and (b) 75 k /.
Fig. 3. Secondary electron image of the sample treated at 900 8C without the application of a pressure.
4:9Fe3:2Cr1:9Ni [10Fe0:49 Cr0:32 Ni0:19
(3)
Reaction Eq. (2) is the well-known, exothermic thermite reaction, for which the standard enthalpy of reaction ranges from /881.4 to /879.8 kJ over the temperature range of 670 8C (just above the melting point of Al, 660.5 8C) to 1527 8C (just below the melting point of Fe, 1538 8C) [31]. Reaction Eq. (3) is an endothermic alloy formation reaction, for which the enthalpy of reaction has been measured to be a modest /36.0 kJ at
1292 8C (note: the enthalpy of formation of Fe /Cr /Ni alloys with 19 at.% Ni is endothermic for Cr contents ]/ 10 at.%) [32,33]. The DTA reveals a strong exothermic peak that commenced at 605 8C, which was below the melting point of aluminum or alloys of aluminum with iron, chromium or nickel [34 /37]. This indicated that a solidstate exothermic reaction preceded the melting of aluminum for the thermite reaction Eq. (2). A number of exothermic solid-state reactions could have occurred at particle contacts in the powder compacts of the present work. The reaction of solid Al with Fe2O3 is quite exothermic at 605 8C (i.e. DH8 (605 8C) / / 872.5 kJ mol 1 of Al2O3 formed [31]). Solid-state compound formation reactions between Al and Ni, Al and Fe and Al and Cr are also exothermic [32]. For example, the adiabatic temperature for the formation of NiAl from Ni/Al mixtures has been reported to be 1650 8C [38]. Philpot et al. reported that Ni/Al powder compacts heated in flowing argon exhibited strong exotherms (with DT values up to 800 8C) that commenced at temperatures in the range of :/510/635 8C [39]. The onset temperature for combustion increased as the Ni particle size decreased (from B/74 to 149/3 mm) and as the heating rate increased (from 1 to 5 8C min 1). Given the finer Ni particle size (3 mm) and
Fig. 4. Secondary electron image of (a) fractured surface of sdi-MMC composite and (b) polished cross section showing metallic ligaments bridging an indentation crack. The plastically deformed metal phase is indicated by arrows.
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higher heating rates (20 8C min 1 for DTA) used in the specimens in the present paper, solid-state reactions between Al and Ni (and/or Fe, Cr) may have also contributed to the relatively low onset temperature. The heat released from solid-state reactions at contacts between solid powder particles can then result in local melting and spreading of Al or Al-rich liquids2 and the onset of the thermite reaction Eq. (2). Indeed, the presence of a single strong exotherm, instead of several distinct exotherms for solid-state reaction(s) and the thermite reaction (as well as an endotherm for liquid formation), was consistent with the rapid onset of the thermite reaction shortly after the initial heat was liberated by the exothermic solid-state reaction(s). Prior work with gasless combustion reactions has shown that closely-spaced exotherms associated with a lower-temperature solid-state reaction and a higher-temperature liquid/solid reaction can merge into a single exotherm as the heating rate increases [39,40]. Although a relatively high heating rate of 20 8C min 1 was used in the DTA experiments (20 8C min 1 was the highest programmable heating rate available on the DTA instrument used in this work), a much higher rate of heating was experienced by the specimens inserted into the furnace preheated to 900 8C. Hence, the latter specimens are likely to have ignited at or before reaching a temperature of 605 8C. Within 5 min of insertion into the 900 8C furnace, the Al/Fe2O3/Fe/Cr/Ni powder preforms were fully converted into mixtures of Al2O3 and Fe /Cr /Ni alloy. Residual Fe2O3 or Al were not detected by XRD, SEM or TEM analyses. XRD and TEM/SAD analyses revealed the presence of two metallic phases: a g(fcc) Fe-rich (50.0 at.% Fe, 29.9 at.% Cr, 20.1 at.% Ni) solid solution and an a(bcc) Cr-rich (64.4 at.% Cr, 26.8 at.% Fe, 8.8 at.% N) solid solution. These fcc and bcc phases possessed lattice constants of 0.37 and 0.30 nm, respectively, which were similar to the lattice constants reported in the literature for g(fcc) Fe /Cr /Al solid solutions (0.36 nm) and a(bcc) Cr-rich solid solutions (0.29 nm) [30,41]. The formation of these g(fcc) and a(bcc) phases was consistent with the reported phase equilibria for the Fe /Ni /Cr system [42,43]. At temperatures in the range of 920 /1300 8C, an overall composition of 49 at.% Fe, 32 at.% Cr and 19 at.% Ni (i.e. the alloy product of net reaction Eq. (1)) falls within a (g/a) two-phase field, not far from the g/(g/a) phase boundary. Hence, equilibration of an alloy with a Fe0.49Cr0.32Ni0.19 composition at 920/1300 8C should yield a mixture comprised largely of the g(fcc) solid solution phase, along with a lesser amount of the a(bcc) 2 For example, an Al /Ni eutectic liquid with 3 at.% Ni can form at 640 8C [34]. An Al /Ni /Fe eutectic liquid with 0.1 at.% Fe, 3.0 at.% Ni can form at 638 8C [35].
phase. Although the a phase can congruently transform into a brittle s phase at lower temperatures, the s phase has been reported to form at a slow rate [42]. Consequently, the observed formation of composites with a metallic matrix comprised predominantly of a g(fcc) Fe /Cr /Ni solid solution, with a smaller amount of the a(bcc) Cr /Fe /Ni solid solution, was consistent with the completion of net reaction Eq. (1), equilibration at ]/ 900 8C, and then air quenching. The DTA also provided an indirect confirmation that the net reaction (Eq. (1)) could be completed during heating at 20 8C min 1. A strong endothermic peak was observed with an onset temperature of 1405 8C. This temperature falls within the reported values of the solidus and liquidus temperatures (1370 and 1420 8C, respectively) for an alloy with a composition of 49 at.% Fe, 32 at.% Cr and 19 at.% Ni [35,42]. In other words, the endotherm was consistent with the melting of a metal matrix with a composition similar to that obtained by the completion of net reaction Eq. (1). Comparison of the XRD patterns indicated that the application of an external pressure of 20 MPa during annealing at 900 8C did not significantly alter the number or compositions of the phases generated in the resulting composites. However, this external pressure did have a dramatic effect on the microstructure of the final composites. Comparison of Figs. 1 and 3 indicated that the composites produced by annealing at 20 MPa were more dense than those generated without pressure. Solidified metallic droplets were also detected on the external surfaces of the latter specimens after annealing; that is, at least some of the porosity observed in Fig. 1 was due to the loss of liquid metal from the preform during annealing. The alumina phase in the pressuretreated composites was also considerably finer (a few hundred nanometers in diameter) and possessed a more filamentary morphology than the alumina present in the specimens annealed without pressure. Higher-magnification TEM images (in Fig. 2) revealed additional features of the alumina grains in the pressure-annealed composites. The alumina grains resembled the shape of droplets joined together. Some metallic inclusions could also be seen embedded in the droplet-shaped alumina particle (Fig. 2b), further suggesting alumina to be molten upon the reaction. Such droplet-morphology of Al2O3 suggests that the amount of heat liberated upon the reaction among the green compact constituents was large enough to exceed the melting temperature above the melting point of alumina (2050 8C) [44]. The adiabatic temperature for reaction Eq. (1) was, however, calculated to be 1900 8C [33,45]. The presence of droplet-shaped alumina particles, when adiabatic temperature is lower than its melting point suggests the possibility of melting of Al2O3 locally at the reaction site. Such an event is possible if the heat generated by the reaction given by
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Eq. (2) cannot be transferred away from the reaction site fast enough. The adiabatic temperature of Eq. (2), per se, is known to well exceed the melting point of alumina [21]. These observations indicate that the reaction between Al and Fe2O3 (Eq. (2)) lead to the formation of molten Al2O3 and Fe. As most of the heat of the reaction is locally absorbed and because the adiabatic temperature for the overall reaction (Eq. (1)) is 1900 8C, it is likely that the metallic constituents (Cr, Ni, Fe) are largely in solid state after the reaction. Such a scenario can explain the nanosize thickness of Al2O3. Molten Al2O3 under the applied pressure can get squeezed into the surrounding region resulting in nano-sized droplets morphology. The molten Fe that is produced upon reaction, can go into solution with other largely solid metallic constituents. The presence of coarse alumina particles in the absence of a pressure (Fig. 3) also confirms that molten alumina at the reaction site flows under the applied pressure and crystallises into very fine droplets. As the molten alumina spreads, under an applied pressure of 20 MPa, throughout the sample and solidifies, the resulting microstructure has an interpenetrating network of alumina and metallic phases. Such a continuous network of ceramic phase with only 25 vol.% ceramic is a unique feature of this process. Furthermore, as a result of the spontaneous exothermic reaction in the present process, transformation was complete within 5 min. The alumina/TiAl composites prepared by Gutmanas et al. [25] by the pressure SHS technique yielded equiaxed alumina grains. This observation can again be explained by the fact that the adiabatic temperature for the reaction between Al and TiO2 to form Al2O3 and Ti is :/1675 8C, which is well below the melting point of Al2O3. Thus, even though a pressure of 100 MPa was applied, Al2O3 being formed in the solid state could not be squeezed to form thin ligaments. Also, the shape of the oxide phases in the composite materials formed by RMP, 3A or DCP is generally equiaxed [10 /20]. In all of these processes, oxide phase in the product forms in the solid form as a result of a reaction between a molten metal and sacrificial solid oxide. Such melt oxide reactions tend to be time consuming and can take from 1 to /10 h depending on the process and sacrificial preform. Fig. 4(a) shows the fracture surface of Fe /Cr /Ni/ Al2O3 composite. The existence of a continuous network of alumina and Fe alloy may increase the toughening effect of alumina and allow crack propagation along its network. For the Fe /Cr /Ni/Al2O3 composite with a mismatch in coefficients of thermal expansion (CTE), internal stresses are set up within and around the alumina phase. The cooling of such a system from the peak reaction temperature is expected to generate compressive microstresses in the alumina phase due its lower CTE vis-a`-vis the metal phase. The compressive
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microstresses in Al2O3 are counterbalanced by tensile stresses in the metal phase. Thus, the metal/ceramic interface is subjected to a tensile stress. In the metallic phase, however, some stress can be relaxed due to plastic deformation. The investigation of crack paths shows a tendency of the cracks to propagate through the alumina phase, indicating strong interfacial bonding between metal phase and alumina. Strong interfacial bonding leads to high geometrical constraint for the metal phase and a high degree of triaxial tension in the metal ligament, thereby increasing the uniaxial yield strength by a factor of 5 /7 [46]. This could explain the high strength obtained for the composites. As a result of strong bonding, alumina grains seem to have fractured mainly in the transcrystalline manner. The metal phase showed plastic deformation, as evident by the sharp edges in Fig. 4(a) and fractured in a completely ductile manner. Other materials systems containing nano-sized alumina reinforcement can also be designed with the knowledge and control of adiabatic temperature associated with the reactions.
5. Conclusions Metal-matrix composites reinforced with very fine ceramic phase have been fabricated at a modest applied temperature of 900 8C by a rapid, in situ, short distance infiltration (sdi-MMC) process. Dense composites comprised of a fine-grained (5/5 mm) Fe /Cr /Ni matrix reinforced with interpenetrating a /Al2O3 reinforcements (a few hundred nanometers in diameter) were produced by the insertion of porous, mechanically mixed Al/Fe2O3/Fe/Cr/Ni powder compacts into a furnace preheated to 900 8C and holding at this applied temperature for 5 min under a peak pressure of 20 MPa. This heat treatment triggered the highly-exothermic thermite reaction between liquid Al and Fe2O3 in the preform that, in turn, led to the rapid formation of a liquid Fe /Cr /Ni alloy and Al2O3. XRD, SEM and TEM analyses confirmed that the composite formation reaction was completed within this brief thermo-mechanical treatment at 900 8C (i.e. residual unreacted Al or Fe2O3 were not detected). Consistent with the Fe / Cr /Ni phase diagram, SEM/EDS and TEM/SAD analyses revealed that the metal matrix in such fullyreacted composites was comprised largely of a g(fcc) Fe /Cr /Ni solid solution along with a smaller amount of an a(bcc) Cr /Fe /Ni solid solution. The unique morphology of the alumina reinforcement phase was attributed to the pressure-assisted spreading of molten alumina throughout the specimen. The average values of fracture strength and fracture toughness of the pressureannealed Fe /Cr /Ni/Al2O3 composites obtained from
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three-point bend tests were 1100 and 18 MPa m1/2, respectively.
Acknowledgements This work was supported by the German /Israeli Foundation for Scientific Research and Development (GIF) through research Grant No. I-556-216.10/97 and the Alexander von Humboldt Foundation (for K.H. Sandhage). The authors wish to acknowledge the assistance of Dr M. Sternitzke in the performance of the TEM analysis.
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[19] M.C. Breslin, J. Ringnalda, L. Xu, M. Fuller, J.G. Seeger, G.S. Daehn, T. Otani, H.L. Fraser, Mater. Sci. Eng. A195 (1995) 113. [20] W.G. Fahrenholtz, K.G. Ewsuk, D.T. Ellerby, R.E. Loehman, J. Am. Ceram. Soc. 79 (1996) 2497. [21] J. Subrahmanyam, M. Vijaykumar, J. Mater. Sci. 27 (1992) 6249. [22] L.L. Wang, Z.A. Munir, Y.M. Maximov, J. Mater. Sci. 28 (1993) 3693. [23] S.P. Gaus, M.P. Harmer, H.M. Chan, H.S. Caram, N. Claussen, J. Am. Ceram. Soc. 83 (7) (2000) 1599. [24] S.P. Gaus, M.P. Harmer, H.M. Chan, H.S. Caram, J. Bruhn, N. Claussen, J. Am. Ceram. Soc. 83 (7) (2000) 1606. [25] E.Y. Gutmanas, I. Gotman, J. Eur. Ceram. Soc. 19 (1999) 2381. [26] J.B. Holt, Z.A. Munir, J. Mater. Sci. 21 (1986) 251. [27] Metals Handbook, Metallography and Microstructures, vol. 9, 9th ed, 1985, p. 330. [28] P. Chantikul, G.R. Anstis, B.R. Lawn, D.B. Marshall, J. Am. Ceram. Soc. 64 (1981) 539. [29] JCPDS X-Ray Diffraction Card File, Cards No. 4-787 for Al, No. 33 /664 for Fe2O3, No. 6 /696 for Fe, No. 4 /850 for Ni, No. 6 / 694 for Cr, No. 43 /1484 for Al2O3, No. 31 /619 for austenitic Fe0.966C0.034, No. 33 /397 for austenitic Fe0.70Cr0.19Ni0.11. [30] W.P. Rees, B.D. Burns, A.J. Cook, J. Iron Steel Inst. 162 (1949) 325. [31] L.B. Pankratz, J.M. Stuve, N.A. Gokcen, Thermodynamic Data for Mineral Technology, vol. 677, US Department of the Interior/ Bureau of Mines, USA, 1984, p. 262, 270. [32] R. Hultgren, P.D. Desai, D.T. Hawkins, M. Gleiser, K.K. Kelley, Selected Values of the Thermodynamic Properties of Binary Alloys, American Society for Metals, Metals Park, OH, 1973, p. 148, 161, 192, 699, 701, 711, 851. [33] O. Kubaschewski, L.E.H. Stuart, J. Chem. Eng. Data 12 (1967) 418. [34] H. Okamoto, J. Phase Equil. 14 (1993) 257. [35] P. Villars, A. Prince, H. Okamoto, Handbook of Ternary Alloy Phase Diagrams, ASM International, Metals Park, OH, 1995, p. 3532, 3533. [36] O. Kubaschewski, Iron-Binary Phase Diagrams, Springer-Verlag, Berlin, 1982, p. 5. [37] J.G. Costa Neto, S. Gama, C.A. Ribeiro, J. Alloys Comp. 182 (1992) 271. [38] Y.S. Naiborodenko, V.I. Itin, Comb. Expl. Shock Wave 11 (1975) 293. [39] K.A. Philpot, Z.A. Munir, J.B. Holt, J. Mater. Sci. 22 (1987) 159. [40] Z.A. Munir, Ceram. Bull. 67 (1988) 342. [41] N. Yukawa, M. Hida, T. Imura, M. Kawamura, Y. Mizuno, Met. Trans. 3 (1972) 887. [42] V.G. Rivlin, G.V. Raynor, Bull. Alloy Phase Diag. 2 (1981) 89. [43] M. Hasebe, T. Nishizawa, Applications of Phase Diagrams in Metallurgy and Ceramics (Special Publication SP-496), National Bureau of Standards, Gaithersburg, MD, 1978, p. 911. [44] R.S. Roth (Ed.), Phase Equilibria Diagrams, Phase Diagrams for Ceramics, Oxides, vol. XI, 1995. [45] JANAF Thermochemical Tables, 3rd ed, National Bureau of Standards, 1985. [46] M.F. Ashby, F.J. Blunt, M. Banister, Acta Metal. 37 (1989) 1847.
Rapid synthesis of Al2O3 reinforced Fe Cr Ni composites /
/
N. Travitzky a, , P. Kumar a, K.H. Sandhage b, R. Janssen a, N. Claussen a a
b
Advanced Ceramics Group, Technische Universita¨t Hamburg-Harburg, D-21073 Hamburg, Germany Department of Materials Science and Engineering, The Ohio State University, Columbus, OH 43210, USA Received 20 March 2002; received in revised form 3 June 2002
Abstract Short distance infiltration (sdi) was used to fabricate metal matrix composites (MMC) with very fine ceramic reinforcements. This ‘sdi-MMC’ process was used to synthesize a dense composite comprised of Fe /Cr /Ni/Al2O3. Fe2O3/Fe/Ni/Cr/Al powder compacts were rapidly heated to 900 8C and held at this temperature and a pressure of 20 MPa in air for 5 min. This thermomechanical treatment triggered the highly-exothermic thermite reaction between starting powders that, in turn, resulted in rapid in-situ formation of a composite comprised predominantly of g(fcc) Fe /Cr /Ni, along with smaller amounts of a(bcc) Cr /Fe /Ni and aalumina. SEM and TEM analyses revealed a microstructure consisting of a uniformly dispersed network of very fine (a few hundred nanometers) Al2O3 grains in a matrix of a fine-grained ( /5 microns) metallic alloy. The dense composites exhibit average bend strength and toughness values of 1100 MPa and 18 MPa (m)1/2, respectively. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Reactive syntheses; Metal matrix composites; Ceramic reinforcement
1. Introduction Ceramic-reinforced, metal-matrix composites (MMCs) with enhanced resistance to fracture, creep and oxidation are attractive for automotive, aerospace and defense applications [1 /4]. Development of advanced MMCs depends on the availability of practical processing routes. MMCs may be produced by powder metallurgical or melt casting routes, depending on process efficiency, production cost and the macrostructure (shape, surface features) and microstructure desired for the final product. High-melting MMCs with attractive microstructures (e.g. fine-scale and uniformly-dispersed phases) and with a wide range of compositions may be produced by the compaction and forming of finely-divided powder mixtures [5,6]. The widespread use of powder processing has been limited by relatively slow and expensive consolidation/forming (e.g. hot pressing, hot extrusion) and machining steps. Melt casting approaches (e.g. compocasting or melt infiltra-
Corresponding author. Tel.: /49-40-428-78-3459; fax: /49-40428-78-3427 E-mail address: [email protected] (N. Travitzky).
tion), on the other hand, are capable of rapidly producing dense composites with complex and near net shapes [2]. However, the casting of high-melting MMCs, reinforced with a fine and continuous ceramic phase, is relatively complicated and expensive. New cost-effective (low-temperature, rapid) processing routes need to be developed for the direct fabrication of dense, high-temperature MMCs. The melt casting temperature may be dramatically lowered by using a chemical reaction between a lowmelting metallic liquid and a shaped solid preform to generate a higher-melting composite. This concept has been utilized in a number of reactive melt casting techniques [7 /17]. For example, in infiltration alumina aluminide alloy (i-3A) and displasive compensation of porosity (DCP) processes, low-melting metallic liquids (e.g. pure molten Al or Mg; or low-melting alloy liquids) are allowed to infiltrate into and undergo displacement reactions with, porous preforms to obtain dense, near net-shaped composites containing higher-melting ceramic and metallic (or intermetallic) phases [10 /17]. In reactive melt penetration (RMP), a low-melting metallic liquid (e.g. molten Al) undergoes a displacement reaction with, and penetrates into, a dense oxide preform to yield a co-continuous ceramic/metallic or ceramic/inter-
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 4 1 9 - 7
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metallic composite [18 /20]. These reactive casting approaches have been used to produce a variety of dense composites (e.g. Al2O3/TiAl3, Al2O3/NbAl3, MgAl2O4/Fe /Ni /Al alloy, MgO/FeAl and ZrC/W composites) at temperatures 5/1200 8C. The composition and phase content of a composite produced by any particular reactive casting method is dependent upon and limited by: (i) the reaction-induced volume change; and (ii) the porosity of the reacting preform. With RMP, the volume fraction of the metal phase in the final composite is largely determined by the space made available as the reactant oxide is converted into a product oxide. Consequently, RMP tends to yield composites with relatively high ceramic contents of ]/65 vol.% Al2O3 [18 /20]. With i-3A and DCP, porosity present in the starting preforms can provide additional space for the metallic or intermetallic phase in the final composite. Indeed, preform porosity must be carefully tailored with these latter two processes to produce composites with desired amounts of metallic or intermetallic phases [10 /17]. Melt processing methods that can yield a wider range of compositions and phase contents, without the need for preforms of tailored porosity, should be examined. The processing time for reactive melt casting is strongly influenced by porosity and the dimensions of the preform [12,14,18 /20]. With RMP, the time required for complete reaction and penetration of the metallic liquid into the dense preform scales with the preform size, and tends to be relatively long for thick preforms (e.g. /4 h for 2.5 cm thick preforms at 1100 8C [19]). With i-3A and DCP, reactant liquids are dispersed throughout porous preforms by infiltration prior to reaction, so the subsequent reaction occurs more rapidly and uniformly throughout preforms. However, the time required for complete melt infiltration with i-3A or DCP still depends on the dimensions and porosity (i.e. the pore fraction and size) of the preform. Self-propagating high-temperature synthesis (SHS or combustion synthesis) method is among the fastest reaction-based processes. The exothermic heat of reactions rapidly converts compacted powder mixtures into products [21]. Low energy requirements due to the self generation of heat, short processing time, possibility to form metastable phases due to high thermal gradients and rapid cooling rates are the major advantages of SHS process. Thermite reactions are a sub-set of SHS processes [22]. Thermite reactions have the advantage over the reactions with elemental reagents in that they start with naturally occurring oxides that are less expensive and more readily available than elemental reactant powder. Recently, thermite-based reactions were used to fabricate dense and homogenous alumina aluminide alloys (3A) with interpenetrating networks of Al2O3 and aluminides of Ti, Fe, Nb, etc. [18,19]. The
composites were fabricated by pressureless reaction sintering in a nonoxidizing atmosphere using intensively milled metal oxide-aluminium powder compacts. The exothermic in 3A necessitate careful control of the processing parameters. Rapid temperature increases due to these reactions result in the formation of liquid phases, which may cause large pores that cannot be removed during pressureless sintering [23,24]. Since full density is of paramount significance for most structural materials, residual porosity is the main hindrance for the commercial use of SHS. In order to obtain composites with reduced porosity, several authors have conducted SHS reactions under the application of an external pressure [25,26]. For example, Holt and Munir fabricated 95% dense TiC by ignition under a mild pressure of 27.6 MPa [26]. Gutmanas and Gotman demonstrated that the application of a modest external pressure can considerably reduce the porosity in alumina/aluminide composites [25]. This method, called short distance infiltration-sdi was used to fabricate a :/99% dense composite of Al2O3 and TiAlx phases (Ti3Al, TiAl, TiAl2) by igniting a porous powder compact of Al and TiO2 (molar Al:TiO2 ratio /7:3) in a die preheated to 950 8C under a pressure of 100 MPa. The equiaxed Al2O3 grains were well-dispersed with the TiAlx phases, were of the order of a few microns in size and were present in a high concentration (note: complete reaction of an Al:TiO2 /7:3 mixture should yield a TiAl/Al2O3 composite comprised of 46.9 vol.% Al2O3 and 53.1 vol.% TiAl). In the above works, in situ processing techniques lead to composite materials with significant ceramic contents (i.e. /45 vol.%). However, fabrication of metal/ceramic composites with lower ceramic content has not been demonstrated to date. Unlike RMP, i-3A and DCP, the time required to complete such short distance infiltration and reaction should be independent of the preform dimensions and shape. Furthermore, by controlling the ratio of solid precursor and liquid in the preform, the composition and phase content of the final composite: (i) can be precisely tailored; and (ii) should not depend on the porosity of the starting preform. It should be possible to produce composites with a wide range of compositions and phase contents (including MMCs) by short distance infiltration. The purpose of this work is to demonstrate that dense MMCs with B/30 vol.% ceramic reinforcement can be designed and fabricated by a rapid in situ sdi-metal matrix composites (MMC) by short distance infiltration (sdi) process. The specific objectives are: (i) to fabricate a dense MMC comprised of a Fe0.49Cr0.32Ni0.191 with 1 Note: this Fe:Cr:Ni ratio is similar to that for ASTM HL-type heat resistant alloys [27].
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interpenetrating nano-sized Al2O3 reinforcement phase; and (ii) to evaluate the micro-chemistry, microstructure and mechanical properties of such sdi-MMC-derived composites.
2. Experimental Powders of Al (B/45 mm, 99.5% purity; EckartWerke, Fuerth, Germany), Fe2O3 (15 mm, 98% purity; Jan de Poorter BV, Geertruidenberg, The Netherlands), Fe (4 /5 mm, 99.5% purity; BASF, Ludwigshafen, Germany), Cr (B/10 mm, 99.8% purity; Johnson Matthey GmbH, Karlsruhe, Germany) and Ni (3.0 mm, 99.7% purity; Aldrich Chemical Co., Milwaukee, WI) were mixed in a molar ratio of 2.0:1.0:2.9:3.2:1.9 (as per reaction Eq. (1) in Section 4). This metal/oxide powder mixture was then attrition milled (Model PE 075 Netzsch-Feintechnik, Selb, Germany) in acetone for 1 h at a rotation rate of 500 rpm using yttria-stabilized zirconia balls (3 mm diameter, 3Y-TZP, Tosoh Co., Tokyo, Japan) as milling media. Balls (1500 g) were used to mill 50 g of powder (i.e. a ball-to-charge weight ratio of 30:1) within an alumina milling vial (10 cm diameter). After overnight evaporation of the acetone, the milled mixture was passed through a 200-micron sieve. The sieved powder was then uniaxially pressed at 25 MPa into disks that, in turn, were cold isostatically pressed (CIPped) at a peak stress of 700 MPa. The as-CIPped disks ( :/15 mm diameter and 10 mm thick) disk were placed within two tightly-fitting stainless steel cells (i.e. two flat bottom cylindrical crucibles with internal diameters of 16 and 15 mm, respectively, heights of 10 mm, wall thickness of 0.5 mm and bottom thickness of 1 mm). The walls of each cell had been sprayed with a BN coating (BN-S, SINTEC Innovation in Keramik, Buching, Germany) to prevent chemical interaction with the specimen and with the preheated rams. The steel cells, containing the green compacts, were then placed into a hot vertical box furnace located within a conventional Zwick (Type 1478, Zwick, Ulm, Germany) testing frame. The box furnace was preheated to 900 8C in air prior to insertion of the steel cells. After insertion, a uniaxial load was applied to some of the sample assemblies at a rate of 30 N s 1 until a peak pressure of 20 MPa was achieved. The pressure was released after 5 min and the samples were quickly removed from the furnace and allowed to cool in air. Other specimens were exposed to similar thermal conditions without the application of an external pressure. After cooling to room temperature, the annealed samples were stripped out of the stainless steel cells. Microstructural analyses of the synthesized samples were conducted with a scanning electron microscope (Gemini 1530 SEM, LEO, Oberkochen, Germany) and a transmission electron microscope (2000 FX II TEM,
247
JEOL, Tokyo, Japan). Microchemical analyses were conducted during SEM and TEM analyses with energydispersive X-ray spectroscopy. At least five different EDS spot analyses were conducted to obtain the average composition of a given phase. The samples for SEM/ EDS analyses were ground and polished to a 1 mm diamond finish. Electron transparent specimens for TEM analyses were prepared by grinding the sample into foils with a thickness of :/70 mm followed by ionbeam milling (GATAN DuoMill, Pleasanton, USA). Selected area diffraction (SAD/TEM) and X-ray diffraction (XRD) analyses were employed for phase identification. XRD analyses (Model PW 1710 Diffractometer, Philips Inc., Eindhoven, The Netherlands) were conducted using Cu-Ka radiation at a scan rate of 0.758 min 1. Differential thermal analyses (DTA) of green compacts and annealed composites were conducted in flowing Ar at a heating rate of 20 8C min 1 to 1500 8C (STA 409 DTA/TGA Netzsch Geraetebau GmbH, Selb, Germany). The densities of green compacts and annealed composites were obtained from weight and dimension measurements. Archimede’s method, using water as the immersion medium, was used to determine the relative porosity of the annealed composites. The average values of hardness, flexural strength and fracture toughness of the annealed composites were evaluated. Hardness tests were conducted with a standard Vickers apparatus (Type 3212, Zwick) using a load of 100 N applied for 15 s. The average hardness value was determined from ten indentation measurements. The flexural strength and fracture toughness was evaluated in three-point bending with the use of a universal testing machine (Type 1478, Zwick). The span and crosshead speed were maintained at 12 and 0.5 mm min 1, respectively, for all bend tests. The average values of flexural strength and fracture toughness were determined from measurements conducted on at least four bar-shaped samples each. Bars with dimensions of 2.5 /3.5 /14 mm were cut out of the fabricated composites and ground to a 15 mm diamond finish. The tensile surfaces of the samples were then polished to a 1 mm diamond finish prior to bending. Prior to a given fracture toughness test, a Vickers indentor was pushed into the specimen surface with a load of 300 N, as per the indentation strength in bending (ISB) technique [28]. The fracture surfaces and crack paths were examined by scanning electron microscopy.
3. Results Distinct peaks for Al, Fe2O3, Fe, Cr and Ni were detected by XRD analysis of an as-CIPped green compact. The average density of the as-CIPped green
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compacts was found to be :/4.65 g cm 3, which corresponds to :/75.8% of the theoretical density for the Al/Fe2O3/Fe/Cr/Ni mixtures. Within :/1 min of achieving the peak stress of the Al/ Fe2O3/Fe/Cr/Ni powder compacts into the preheated furnace, a distinct ‘pop’ was heard to emanate from the specimens, which was presumed to be a result of combustion. The applied pressure was maintained at a constant value after the combustion had occurred. Although the processing was conducted in ambient air, visual inspection of the synthesized specimens revealed no apparent surface damage (e.g. cracking, spallation) due to oxidation. After treatment under 20 MPa peak pressure, the disk-shaped specimens possessed diameter and thickness dimensions of :/17 and 7 mm, respectively. The average bulk density after such reaction was 6.50 g cm 3, which corresponded to a relative density of 94.2% of the theoretical value for a fully-converted Fe0.49Cr0.32Ni0.19/Al2O3 composite (see reaction Eq. (1) in the Section 4). XRD analyses of specimens annealed under pressure revealed the presence of two metallic phases along with Al2O3. The predominant metallic phase was identified as an austenitic Fe-rich (g, fcc) solid solution (Fe-SS), whereas the minor metallic phase was consistent with a Cr-rich (a, bcc) solid solution (Cr-SS). A secondary electron image of a polished cross-section of this composite material is shown in Fig. 1. The specimen appeared dense and comprised of a fine-grained (5/5 mm) metallic matrix with very fine alumina grains visible as dark threads winding throughout the matrix. Brightfield TEM images of the composite are shown in Fig. 2. The Fe- and Cr-rich alloy phases can be seen in Fig. 2, along with the nano-sized (bright) Al2O3 phase. EDS analyses revealed that the average composition of the Fe-rich phase was 50.0 at.% Fe, 29.9 at.% Cr and 20.1 at.% Ni. The secondary Cr-rich metallic phase possessed an average composition of 64.4 at.% Cr, 26.8 at.% Fe
and 8.8 at.% Ni. Aluminum was not detected in either metallic phase. Selected area electron diffraction (SAD) patterns of the Fe- and Cr-rich phases were consistent with g-Fe- and a-Cr-rich solid solutions, respectively. The average lattice constants of the Fe- and Cr-rich phases obtained from SAD patterns were 0.37 and 0.30 nm, respectively. The TEM image in Fig. 2(a) reveals that the austenitic Fe-rich phase possessed a much higher dislocation density than the Cr-rich phase. As seen in Fig. 2, the a-Al2O3 grains consisted of single alumina crystals connected in series along the grain length. The alumina grain size was of the order of a few hundred nanometers. A secondary electron image of a polished crosssection of a sample annealed at 900 8C without the application of pressure is shown in Fig. 3. An XRD pattern of this sample was similar to that of the composite prepared under pressure. However, the morphology of the alumina phase was quite different. The alumina in the composite of Fig. 3 did not exhibit the fine, elongated droplet-morphology prevalent in the composites annealed under pressure. Differential thermal analyses (DTA) obtained upon heating an as-CIPped green compact to 1500 8C at 20 8C min 1 in an argon atmosphere showed strong exothermic and endothermic peaks with onset temperatures of 605 and 1405 8C, respectively. The average value of room-temperature flexural strength for composites prepared under pressure was 11009/95 MPa. Some proportion of plastic deformation at fracture was obtained in all tested samples. The average value of fracture toughness was 189/0.5 MPa m1/2. The average Vickers hardness value of the pressure-annealed composites was 4.29/0.2 GPa. A typical fracture surface and a crack in a polished cross-section (where the crack was generated upon indentation with a Vickers indentor) of a pressureannealed Fe /Ni /Cr/Al2O3 composite are shown in Fig. 4.
4. Discussion An objective of the present work was to use the following net reaction to convert a porous Al/Fe2O3/Fe/ Cr/Ni preform into a dense MMC comprised of a Fe0.49Cr0.32Ni0.19 alloy reinforced with Al2O3: 2AlFe2 O3 2:9Fe3:2Cr1:9Ni [Al2 O3 10Fe0:49 Cr0:32 Ni0:19
(1)
The product of this reaction is a composite comprised of 73.0 vol.% Fe0.49Cr0.32Ni0.19 alloy and 27.0 vol.% Al2O3 [29,30]. This net reaction can be considered to be a combination of the following two reactions: Fig. 1. Secondary electron image of a polished cross section of Fealloy/Al2O3 composite prepared by sdi-MMC method.
2Al(l)Fe2 O3 [Al2 O3 2Fe
(2)
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Fig. 2. Brightfield TEM images of Fe-alloy/Al2O3 composite fabricated by sdi-MMC method. Magnification of (a) 23 k / and (b) 75 k /.
Fig. 3. Secondary electron image of the sample treated at 900 8C without the application of a pressure.
4:9Fe3:2Cr1:9Ni [10Fe0:49 Cr0:32 Ni0:19
(3)
Reaction Eq. (2) is the well-known, exothermic thermite reaction, for which the standard enthalpy of reaction ranges from /881.4 to /879.8 kJ over the temperature range of 670 8C (just above the melting point of Al, 660.5 8C) to 1527 8C (just below the melting point of Fe, 1538 8C) [31]. Reaction Eq. (3) is an endothermic alloy formation reaction, for which the enthalpy of reaction has been measured to be a modest /36.0 kJ at
1292 8C (note: the enthalpy of formation of Fe /Cr /Ni alloys with 19 at.% Ni is endothermic for Cr contents ]/ 10 at.%) [32,33]. The DTA reveals a strong exothermic peak that commenced at 605 8C, which was below the melting point of aluminum or alloys of aluminum with iron, chromium or nickel [34 /37]. This indicated that a solidstate exothermic reaction preceded the melting of aluminum for the thermite reaction Eq. (2). A number of exothermic solid-state reactions could have occurred at particle contacts in the powder compacts of the present work. The reaction of solid Al with Fe2O3 is quite exothermic at 605 8C (i.e. DH8 (605 8C) / / 872.5 kJ mol 1 of Al2O3 formed [31]). Solid-state compound formation reactions between Al and Ni, Al and Fe and Al and Cr are also exothermic [32]. For example, the adiabatic temperature for the formation of NiAl from Ni/Al mixtures has been reported to be 1650 8C [38]. Philpot et al. reported that Ni/Al powder compacts heated in flowing argon exhibited strong exotherms (with DT values up to 800 8C) that commenced at temperatures in the range of :/510/635 8C [39]. The onset temperature for combustion increased as the Ni particle size decreased (from B/74 to 149/3 mm) and as the heating rate increased (from 1 to 5 8C min 1). Given the finer Ni particle size (3 mm) and
Fig. 4. Secondary electron image of (a) fractured surface of sdi-MMC composite and (b) polished cross section showing metallic ligaments bridging an indentation crack. The plastically deformed metal phase is indicated by arrows.
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higher heating rates (20 8C min 1 for DTA) used in the specimens in the present paper, solid-state reactions between Al and Ni (and/or Fe, Cr) may have also contributed to the relatively low onset temperature. The heat released from solid-state reactions at contacts between solid powder particles can then result in local melting and spreading of Al or Al-rich liquids2 and the onset of the thermite reaction Eq. (2). Indeed, the presence of a single strong exotherm, instead of several distinct exotherms for solid-state reaction(s) and the thermite reaction (as well as an endotherm for liquid formation), was consistent with the rapid onset of the thermite reaction shortly after the initial heat was liberated by the exothermic solid-state reaction(s). Prior work with gasless combustion reactions has shown that closely-spaced exotherms associated with a lower-temperature solid-state reaction and a higher-temperature liquid/solid reaction can merge into a single exotherm as the heating rate increases [39,40]. Although a relatively high heating rate of 20 8C min 1 was used in the DTA experiments (20 8C min 1 was the highest programmable heating rate available on the DTA instrument used in this work), a much higher rate of heating was experienced by the specimens inserted into the furnace preheated to 900 8C. Hence, the latter specimens are likely to have ignited at or before reaching a temperature of 605 8C. Within 5 min of insertion into the 900 8C furnace, the Al/Fe2O3/Fe/Cr/Ni powder preforms were fully converted into mixtures of Al2O3 and Fe /Cr /Ni alloy. Residual Fe2O3 or Al were not detected by XRD, SEM or TEM analyses. XRD and TEM/SAD analyses revealed the presence of two metallic phases: a g(fcc) Fe-rich (50.0 at.% Fe, 29.9 at.% Cr, 20.1 at.% Ni) solid solution and an a(bcc) Cr-rich (64.4 at.% Cr, 26.8 at.% Fe, 8.8 at.% N) solid solution. These fcc and bcc phases possessed lattice constants of 0.37 and 0.30 nm, respectively, which were similar to the lattice constants reported in the literature for g(fcc) Fe /Cr /Al solid solutions (0.36 nm) and a(bcc) Cr-rich solid solutions (0.29 nm) [30,41]. The formation of these g(fcc) and a(bcc) phases was consistent with the reported phase equilibria for the Fe /Ni /Cr system [42,43]. At temperatures in the range of 920 /1300 8C, an overall composition of 49 at.% Fe, 32 at.% Cr and 19 at.% Ni (i.e. the alloy product of net reaction Eq. (1)) falls within a (g/a) two-phase field, not far from the g/(g/a) phase boundary. Hence, equilibration of an alloy with a Fe0.49Cr0.32Ni0.19 composition at 920/1300 8C should yield a mixture comprised largely of the g(fcc) solid solution phase, along with a lesser amount of the a(bcc) 2 For example, an Al /Ni eutectic liquid with 3 at.% Ni can form at 640 8C [34]. An Al /Ni /Fe eutectic liquid with 0.1 at.% Fe, 3.0 at.% Ni can form at 638 8C [35].
phase. Although the a phase can congruently transform into a brittle s phase at lower temperatures, the s phase has been reported to form at a slow rate [42]. Consequently, the observed formation of composites with a metallic matrix comprised predominantly of a g(fcc) Fe /Cr /Ni solid solution, with a smaller amount of the a(bcc) Cr /Fe /Ni solid solution, was consistent with the completion of net reaction Eq. (1), equilibration at ]/ 900 8C, and then air quenching. The DTA also provided an indirect confirmation that the net reaction (Eq. (1)) could be completed during heating at 20 8C min 1. A strong endothermic peak was observed with an onset temperature of 1405 8C. This temperature falls within the reported values of the solidus and liquidus temperatures (1370 and 1420 8C, respectively) for an alloy with a composition of 49 at.% Fe, 32 at.% Cr and 19 at.% Ni [35,42]. In other words, the endotherm was consistent with the melting of a metal matrix with a composition similar to that obtained by the completion of net reaction Eq. (1). Comparison of the XRD patterns indicated that the application of an external pressure of 20 MPa during annealing at 900 8C did not significantly alter the number or compositions of the phases generated in the resulting composites. However, this external pressure did have a dramatic effect on the microstructure of the final composites. Comparison of Figs. 1 and 3 indicated that the composites produced by annealing at 20 MPa were more dense than those generated without pressure. Solidified metallic droplets were also detected on the external surfaces of the latter specimens after annealing; that is, at least some of the porosity observed in Fig. 1 was due to the loss of liquid metal from the preform during annealing. The alumina phase in the pressuretreated composites was also considerably finer (a few hundred nanometers in diameter) and possessed a more filamentary morphology than the alumina present in the specimens annealed without pressure. Higher-magnification TEM images (in Fig. 2) revealed additional features of the alumina grains in the pressure-annealed composites. The alumina grains resembled the shape of droplets joined together. Some metallic inclusions could also be seen embedded in the droplet-shaped alumina particle (Fig. 2b), further suggesting alumina to be molten upon the reaction. Such droplet-morphology of Al2O3 suggests that the amount of heat liberated upon the reaction among the green compact constituents was large enough to exceed the melting temperature above the melting point of alumina (2050 8C) [44]. The adiabatic temperature for reaction Eq. (1) was, however, calculated to be 1900 8C [33,45]. The presence of droplet-shaped alumina particles, when adiabatic temperature is lower than its melting point suggests the possibility of melting of Al2O3 locally at the reaction site. Such an event is possible if the heat generated by the reaction given by
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Eq. (2) cannot be transferred away from the reaction site fast enough. The adiabatic temperature of Eq. (2), per se, is known to well exceed the melting point of alumina [21]. These observations indicate that the reaction between Al and Fe2O3 (Eq. (2)) lead to the formation of molten Al2O3 and Fe. As most of the heat of the reaction is locally absorbed and because the adiabatic temperature for the overall reaction (Eq. (1)) is 1900 8C, it is likely that the metallic constituents (Cr, Ni, Fe) are largely in solid state after the reaction. Such a scenario can explain the nanosize thickness of Al2O3. Molten Al2O3 under the applied pressure can get squeezed into the surrounding region resulting in nano-sized droplets morphology. The molten Fe that is produced upon reaction, can go into solution with other largely solid metallic constituents. The presence of coarse alumina particles in the absence of a pressure (Fig. 3) also confirms that molten alumina at the reaction site flows under the applied pressure and crystallises into very fine droplets. As the molten alumina spreads, under an applied pressure of 20 MPa, throughout the sample and solidifies, the resulting microstructure has an interpenetrating network of alumina and metallic phases. Such a continuous network of ceramic phase with only 25 vol.% ceramic is a unique feature of this process. Furthermore, as a result of the spontaneous exothermic reaction in the present process, transformation was complete within 5 min. The alumina/TiAl composites prepared by Gutmanas et al. [25] by the pressure SHS technique yielded equiaxed alumina grains. This observation can again be explained by the fact that the adiabatic temperature for the reaction between Al and TiO2 to form Al2O3 and Ti is :/1675 8C, which is well below the melting point of Al2O3. Thus, even though a pressure of 100 MPa was applied, Al2O3 being formed in the solid state could not be squeezed to form thin ligaments. Also, the shape of the oxide phases in the composite materials formed by RMP, 3A or DCP is generally equiaxed [10 /20]. In all of these processes, oxide phase in the product forms in the solid form as a result of a reaction between a molten metal and sacrificial solid oxide. Such melt oxide reactions tend to be time consuming and can take from 1 to /10 h depending on the process and sacrificial preform. Fig. 4(a) shows the fracture surface of Fe /Cr /Ni/ Al2O3 composite. The existence of a continuous network of alumina and Fe alloy may increase the toughening effect of alumina and allow crack propagation along its network. For the Fe /Cr /Ni/Al2O3 composite with a mismatch in coefficients of thermal expansion (CTE), internal stresses are set up within and around the alumina phase. The cooling of such a system from the peak reaction temperature is expected to generate compressive microstresses in the alumina phase due its lower CTE vis-a`-vis the metal phase. The compressive
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microstresses in Al2O3 are counterbalanced by tensile stresses in the metal phase. Thus, the metal/ceramic interface is subjected to a tensile stress. In the metallic phase, however, some stress can be relaxed due to plastic deformation. The investigation of crack paths shows a tendency of the cracks to propagate through the alumina phase, indicating strong interfacial bonding between metal phase and alumina. Strong interfacial bonding leads to high geometrical constraint for the metal phase and a high degree of triaxial tension in the metal ligament, thereby increasing the uniaxial yield strength by a factor of 5 /7 [46]. This could explain the high strength obtained for the composites. As a result of strong bonding, alumina grains seem to have fractured mainly in the transcrystalline manner. The metal phase showed plastic deformation, as evident by the sharp edges in Fig. 4(a) and fractured in a completely ductile manner. Other materials systems containing nano-sized alumina reinforcement can also be designed with the knowledge and control of adiabatic temperature associated with the reactions.
5. Conclusions Metal-matrix composites reinforced with very fine ceramic phase have been fabricated at a modest applied temperature of 900 8C by a rapid, in situ, short distance infiltration (sdi-MMC) process. Dense composites comprised of a fine-grained (5/5 mm) Fe /Cr /Ni matrix reinforced with interpenetrating a /Al2O3 reinforcements (a few hundred nanometers in diameter) were produced by the insertion of porous, mechanically mixed Al/Fe2O3/Fe/Cr/Ni powder compacts into a furnace preheated to 900 8C and holding at this applied temperature for 5 min under a peak pressure of 20 MPa. This heat treatment triggered the highly-exothermic thermite reaction between liquid Al and Fe2O3 in the preform that, in turn, led to the rapid formation of a liquid Fe /Cr /Ni alloy and Al2O3. XRD, SEM and TEM analyses confirmed that the composite formation reaction was completed within this brief thermo-mechanical treatment at 900 8C (i.e. residual unreacted Al or Fe2O3 were not detected). Consistent with the Fe / Cr /Ni phase diagram, SEM/EDS and TEM/SAD analyses revealed that the metal matrix in such fullyreacted composites was comprised largely of a g(fcc) Fe /Cr /Ni solid solution along with a smaller amount of an a(bcc) Cr /Fe /Ni solid solution. The unique morphology of the alumina reinforcement phase was attributed to the pressure-assisted spreading of molten alumina throughout the specimen. The average values of fracture strength and fracture toughness of the pressureannealed Fe /Cr /Ni/Al2O3 composites obtained from
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three-point bend tests were 1100 and 18 MPa m1/2, respectively.
Acknowledgements This work was supported by the German /Israeli Foundation for Scientific Research and Development (GIF) through research Grant No. I-556-216.10/97 and the Alexander von Humboldt Foundation (for K.H. Sandhage). The authors wish to acknowledge the assistance of Dr M. Sternitzke in the performance of the TEM analysis.
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