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TiC composites
Intermetallics 7 (1999) 947±955
In-situ production and microstructures of iron aluminide/TiC composites S.H. Ko*, S. Hanada Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 16 September 1998; received in revised form 24 December 1998; accepted 28 December 1998
Abstract In this study we have tried to produce the titanium carbide reinforced iron aluminide composites by in-situ reaction between titanium and carbon in liquid iron±aluminum alloy doped with titanium and carbon. A homogeneous distribution of titanium carbide particles in the iron aluminide matrix up to about 16 vol% of titanium carbide was intended without agglomeration. The composition of TiC formed during in-situ reaction was investigated by ICP analysis and the Combustion-Infrared Absorption method after chemical dissolution of the iron aluminide matrix. It is found that the composition of titanium carbide formed during melt processing is an average of Ti±48.4 mol% C. In addition, titanium carbide has very low solubility of Fe and Al. The microstructure of composites consists of three dierent regions; primary large TiC particles of 5±40 mm, matrix with small dendritic TiC particles of about 1 mm and particle-free regions around primary large TiC particles. The formation of this complex microstructure can be explained by assuming the Fe3Al±TiC pseudo-binary system containing the eutectic reaction. Particle-free regions are halos of iron aluminide phase and the formation of halos is explained by coupled zone concept. Subsequent heat treatment at 1373 K for 48 h induces spheroidization and/or coarsening of small TiC particles, while microstructure after heat treatment at 973 K for 48 h exhibits the additional formation of small TiC precipitates. Though excess 1 mol% Ti addition over the Ti content for TiC formation is soluble to Fe±28 mol% Al, excess 1 mol% C addition forms the secondary Fe3AlC phase during melt processing. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Intermetallic, miscellaneous; A. Iron aluminides, based on Fe3Al; C. Heat treatment; C. Melting; D. Microstructure
1. Introduction Because of good resistance to high-temperature sul®dizing and oxidizing environments, iron aluminides have high potential for structural applications at high temperatures under corrosive environments [1±3]. In addition, they have relatively low density compared to stainless steel, which leads to a better strength-to-weight ratio [1]. Recently, a few studies have been carried out to utilize carbides or borides as reinforcements in iron aluminide via a powder metallurgy process [4,5]. Among them titanium carbide with high hardness, low density and low chemical reactivity was found to be a suitable reinforcement for wear and high temperature applications [4]. Although many processes have been developed to incorporate reinforcing particles into iron aluminide * Corresponding author now at present address: Tohoku National Industrial Research Institute, Agency of Industrial Science and Technology, MITI, Sendai 983-8551, Japan.
matrix, an in-situ technique in which a thermodynamically stable reinforcing phase is produced during processing can be the most economical on synthesis of composites. A typical example will be the conventional melt processing based on melting and casting, by which in-situ composites can be prepared at low production costs and high eciency. Furthermore, this processing has an advantage that surfaces of reinforcements are not contaminated with oxidation ®lms or other detrimental surface reactions [6±8]. Over the past several years a number of studies have been made on in-situ TiC synthesis by melt processing [7±12]. However, only a little is known about the reactions between titanium and carbon, nucleation and growth of the TiC particles, and particle distributions in the liquid solution, which are necessary to understand and control the process [9]. Furthermore, only few attempts have so far been made on study for intermetallic compound alloys, and no study has ever tried to fabricate in-situ composites based on iron aluminides.
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966 -9 795(99)00002 -3
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On the other hand, some other studies [13,14] about Fe3Al based alloys containing Fe3AlC precipitates have been reported, where these alloys exhibited higher yield stress than binary Fe3Al alloy from room temperature up to 873 K maintaining some room temperature ductility (above 3.5%). In addition, these alloys show reduced susceptibility to environmental embrittlement compared to binary alloy, which was explained by the trapping eects of hydrogen due to Fe3AlC precipitates [14]. In this study TiC reinforced Fe±28 mol% Al based composites are synthesized by in-situ reaction between titanium and carbon in liquid iron±aluminum alloy and in¯uence of contents of titanium and carbon on microstructures is investigated. In addition, we investigate the eect of heat treatment at various temperatures and melt processing condition on microstructures. The composition of TiC formed during in-situ melt processing is also examined. 2. Experimental procedure Three types of samples were prepared to study in-situ formation of titanium carbide in Fe±28 mol% Al alloys (unless otherwise noted, all compositions will be given in mol%) through liquid phase reaction. As for raw materials 99.9 mass% electrolytic iron, 99.9 mass% aluminum, 99.9 mass% titanium and Fe±3.9 mass% C master alloy were used. A vacuum induction furnace was used for preparing of Fe±3.9 mass% C master alloy. The samples were obtained by arc melting. Each sample was remelted four times to ensure composition homogeneity. In order to investigate the forming composition of TiC during melt processing, ICP analysis and the Combustion±Infrared Absorption method were performed after preliminary chemical dissolution of the iron aluminide matrix using a HCl solution. Since iron aluminide matrix is dissolved in a HCl+H2O (1:1) solution, TiC particles can be chemically extracted. Extracted particles were dissolved in a KHSO4 solution to examine the composition of titanium using ICP analysis. The CombustionInfrared Absorption method was performed to investigated the composition of carbon in the particles. The cast samples were cut into 1.5 mm thickness plates and were given heat treatment at 973 and 1373K for 48 h in a mue furnace after encapsulating samples in argon-®lled silica tubes. After desired period of heat treatment samples were water quenched from the respective temperatures. Samples were mechanically polished using diamond slurry and a chemical solution. All the samples were examined metallographically before and after etching, and by X-ray diractometry for phase identi®cation. In addition, the matrix of each sample was partially etched away to observe the true particle shape by a scanning electron microscope. The
samples were put in a HCl+H2O (1:1) solution at room temperature for 2 min to dissolve the surface of matrix. The eect of melt processing time on microstructure was investigated under a given melt processing condition. One sample was arc-melted for 5 min, and the other sample was arc-melted for 20 min at one time. Both the samples were remelted four times. 3. Results From the binary phase diagram of Ti±C [15], it can be found that titanium carbide (TiC) has the composition range from Ti±32% C to Ti±48.5% C at 1923 K. Preliminary samples were prepared and their compositions were analyzed by ICP analysis and the CombustionInfrared Absorption method to investigate the composition of TiC formed during melt processing. The nominal composition of each preliminary sample was Ti±43% C, Ti±45% C and Ti±47% C. Titanium carbide particles were found to be formed at an average composition of Ti±48.4% C in all samples. Scatter of compositions of formed titanium carbide particles was negligibly small from sample to sample. Table 1 shows nominal compositions of samples in present work. Type A samples were alloyed with carbon such that the average mol percent ratio of titanium to carbon coincides with that of TiC (hereafter, titanium carbide with composition of Ti±48.4% C will be simply called TiC) formed by the liquid reaction. The type B and C samples were alloyed so as to have 1% titaniumrich and 1% carbon-rich compositions, respectively. Compositions of particles and matrices of some samples are summarized in Table 2. It is found that TiC particles contain only minimal amounts of Fe and Al, indicating that TiC has very low solubility of Fe and Al. Correspondingly, the compositions of matrices were close to the nominal composition, Fe±28% Al, although titanium contents show some variation ranging from 0.04 to 0.20%. Optical micrographs of unetched, as-cast (Fe±28Al)± 3.2Ti±3C and (Fe±28Al)±9.6Ti-9C and their schematic microstructure are shown in Fig. 1. The characteristic distribution of large particles in the Fe±28%Al matrix Table 1 Nominal compositions of alloys (compositions are given in mol%) Samples
Type Fe
Al
Fe±28A])±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±4.2Ti±3C (Fe±28A])±10.6Ti±9C (Fe±28A])±3.2Ti±4C (Fe±28A])±9.6Ti±10C (Fe±28A1)±0.53Ti±0.SC (Fe±28A1)±1.6Ti±1.5C
A
26.27 3.20 3.00 22.50 9.59 9.00 25.99 4.20 3.00 22.52 10.59 9.00 25.99 3.20 4.00 22.52 9.59 10.00 27.71 0.53 0.50 27.13 1.60 1.50
B C A
67.54 58.62 66.82 57.90 66.82 57.90 71.26 69.80
Ti
C
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can be achieved by the liquid reaction without agglomeration. As depicted in Fig. 1(c), the microstructures of as-cast samples consist of two types of TiC particles concerning particle size, shape and distribution, and Table 2 Compositions of particles and matrices Samples
Particles Fe-mol% Al-mol% Ti-mol% C-mol%a
(Fe±28A1)±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±0.53Ti±0.5C (Fe±28A1)±1.6Ti±1.5C
0.50 0.11 0.41 0.51
Samples (mol%)
Matrix
0.23 0.11 0.24 0.23
51.6 51.7 50.5 50.4
47.7 48.1 48.9 48.9
Fe-mol% Al-mol% Ti-mol% (Fe±28A])±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±0.53Ti±0.5C (Fe±28A1)±1.6Ti±1.5C
71.7 72.2 72.1 71.8
28.2 27.6 27.8 28.0
0.04 0.16 0.09 0.20
a These values were obtained by the Combustion±Infrared Absorption method. The others were obtained by ICP±OES analysis.
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particle-free regions, which are dendrites of iron aluminide phase growing around large TiC particles. (Fe±28Al)±4.2Ti3C and (Fe±28Al)±10.6Ti±9C showed microstructures similar to (Fe±28Al)±3.2Ti±3C and (Fe±28Al)±9.6Ti±9C, respectively. However, carbon-rich samples, (Fe±28Al)± 3.2Ti±4C and (Fe±28Al)±9.6Ti±10C, show the additional formation of precipitates which have two dierent shapes, i.e. spherical and needle-like shapes, as shown in Fig. 2. XRD patterns of the as-cast samples are presented in Fig. 3. One can see Fe3Al and TiC peaks in all samples, indicating the formation of TiC particles during melting and casting, while Fe3AlC peaks can be seen in (Fe± 28Al)±3.2Ti±4C and (Fe±28Al)±9.6Ti±10C. The matrix of (Fe-28Al)-3.2Ti-3C sample was partially removed by chemical deep etching to observe the two types of TiC particles in detail. A SEM micrograph is shown in Fig. 4. Large TiC particles with sizes of about 5±40 mm are formed as faceted crystals, while small ones of about 1 mm have a dendritic structure. Discontinuous distribution in size and dierent shapes of the TiC particles imply that these particles are produced through dierent formation paths. From Figs. 1 and 4, microstructures can be divided into three regions
Fig. 1. Micrographs of as-cast samples of (a) (Fe±28Al)±3.2Ti±3C, (b) (Fe±28Al)±9.6Ti±9C and (c) schematic illustration of microstructure.
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Fig. 2. Micrographs of as-cast samples of (a) (Fe±28Al)±3.2Ti±4C and (b) (Fe±28Al)±9.6Ti±10C.
Fig. 4. SEM micrograph of as-cast (Fe±28Al)±3.2Ti±3C sample after chemical removal of Fe3Al matrix. Fig. 3. XRD results of as-cast samples.
consisting of large particles, matrix with small particles and particle-free region around large particles. Fig. 5 shows optical micrographs of etched, as-cast (Fe±28Al)± 3.2Ti±3C and (Fe±28Al)±9.6Ti±9C. Almost all boundaries of grains delineated by dierent contrasts in (Fe±28Al)± 3.2Ti±3C correspond to boundaries between particles-free matrices and small particle regions, while grain boundaries in (Fe±28Al)±9.6Ti±9C coincide partially with boundaries of small particle regions. Fig. 6 shows optical micrographs of (Fe±28Al)± 3.2Ti±3C heat treated at 973 and 1373 K for 48 h followed by water-quenching. The heat treatment at 973 K causes additional formation of ®ne precipitates, whereas the heat treatment at 1373 K leads to spheroidization and/or coarsening of already existing small particles. XRD analysis indicated that the additional precipitates at 973 K are identi®ed as TiC particles, and no other phase is formed at 1373 K except Fe3Al and TiC. Fig. 7 shows a SEM micrograph of (Fe±28Al)±3.2Ti±3C heat treated at 1373 K for 48 h followed by water-quenching,
where the matrix was partially removed by chemical etching. It is evident from comparison with Fig. 4 that morphology of small particles is changed from dendritic needles to ellipsoids. The compositions of large and small particles were examined using as-cast and heat treated samples. In this experiment large and small particles were distinguished from the dierence in sedimentation speed of particles in acetone. Compositional analysis indicates that particles have almost the same composition independent of particle size and heat treatment (Table 3). Fig. 8 shows the microstructures of unetched, as-cast (Fe±28Al)±0.53Ti±0.5C and (Fe± 28Al)±1.6Ti±1.5C. Decrease in titanium and carbon contents results in the disappearance or the reduced volume fraction of large particles. Large particles greater than 5 mm were not observed in (Fe±28Al)± 0.53Ti±0.5C, but observed with a very low density in (Fe±28Al)±1.6Ti±1.5C. Fig. 9 shows optical micrographs of unetched, as-cast (Fe±28Al)±3.2Ti±3C with dierent holding times in melting. Prolonged holding in the molten state appears to result in coarsening of large particles.
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Fig. 5. Micrographs of as-cast samples after chemical etching of (a) (Fe±28Al)±3.2Ti±3C and (b) (Fe±28Al)±9.6Al±9C.
Fig. 6. Micrographs of (Fe±28Al)±3.2Ti±3C sample heat-treated at (a) 973 K and (b) 1173 K for 48 h followed by water-quenching.
Fig. 7. SEM micrograph of (Fe±28Al)±3.2Ti±3C sample after heat treatment at 1373 K for 48 h followed by chemical removal of Fe3Al matrix.
4. Discussion TiC can be formed in the iron aluminide by a reaction between titanium and carbon, as mentioned before. Provided that a TiC particle has theoretical density, 4.93 g/cm3,
the volume fraction of TiC particles can be calculated. After complete dissolution of the iron aluminide matrix in a HCl+H2O solution, the weight of extracted TiC particles was measured. The mass percent of extracted TiC particles in (Fe±28Al)±3.2Ti±3C and (Fe±28Al)±9.6Ti± 9C were 4.01, and 13.18 mass%, respectively. From the densities of 6.52 g/cm3 for Fe±28%Al and 4.93 g/cm3 for TiC, the volume fraction of TiC particles was evaluated to be 5.2% in (Fe±28Al)±3.2Ti±3C and 16.7% in (Fe±28Al)±9.6Ti±9C. Assuming that alloyed titanium and carbon are entirely consumed to form TiC, one can calculate a theoretical volume fraction of TiC. As shown in Table 4, volume fractions measured from mass percent of extracted TiC particles are in reasonable agreement with those calculated from nominal mass percent of titanium and carbon. In the Fe±Al±Ti±C system, at 1723 K the liquid phase containing only 4.9% C and 64 mol ppm Ti are in equilibrium with Ti0.512C0.488 [4]. Therefore, most TiC particles may precipitate in liquid iron-aluminum alloy at least at 1723 K. Particle-free regions, i.e. dendrites of the iron aluminide phase, are formed around large TiC particles and then interdendritic eutectic grows as shown in Figs. 1 and 4. Based on these observations the
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sequence of TiC formation can be explained by assuming the Fe3Al±TiC pseudo-binary system containing the eutectic reaction between Fe3Al and TiC phases as shown in Fig. 10. Since TiC has a much higher melting point than Fe3Al, the eutectic composition may lie much closer to Fe3Al than to TiC: Fe3Al±TiC system is an anomalous eutectic structure [16]. In such a case the coupled zone, a range on the phase diagram de®ning compositions and temperatures at which the two eutectic Table 3 Compositions of large and small particles in (Fe±28Al)±3.2Ti±3C insitu composite
As-cast Heat-treated at 1373 K for 48 h a
Kind of particles
Elements
Large particles Small particles Large particles Small particles
0.48 0.45 0.49 0.14
Fe± Al± Ti± C± mol% mol% mol% mol%a 0.24 0.20 0.23 0.14
52.0 52.4 52.2 52.0
47.3 47.0 47.1 47.7
These values were obtained by the Combustion±Infrared Absorption method. The others were obtained by ICP analysis.
phases can grow with similar velocities, will incline strongly to TiC side as shown in Fig. 11. [16±18]. It has also shown that the coupled zone for faceted-nonfaceted eutectic systems lies at or near the extended non-faceted liquidus and faceting is a signi®cant feature of anormalous eutectic structure [18±20]. Since the primary large TiC particles and dendrites of small TiC particles show faceted structure as shown in Figs. 4 and 12, this assumption about the shape of coupled zone is proper. At ®rst, it is reasonable to consider that large particles are formed in the melt because the liquid iron-aluminum alloy during melt processing is in the L+TiC two phase region. During cooling the temperature of melt containing primary TiC particles attains the eutectic temperature of Fe3Al and TiC phases. If primary large TiC is a poor nucleant of the Fe3Al phase, the composition of melt including primary large TiC particles will follow the extension of the liquidus line to a temperature which is below the eutectic temperature [17,18,21]. Since the coupled zone in hypothetical Fe3Al-TiC pseudo-binary phase diagram is skewed to TiC side as shown in Fig. 11, the composition of melt lies outside the coupled zone and the melt is supersaturated in Fe3Al. Therefore, dendritic growth of Fe3Al phase forming the particlefree region, halo, will proceed around the primary large
Fig. 8. Micrographs of as-cast samples of (a) (Fe±28Al)±0.53Ti±0.5C and (b) (Fe±28Al)±1.6Al±1.5C.
Fig. 9. Eects of melt processing time on microstructure of (Fe±28Al)±3.2Ti±3C sample. (a) (for 5 min)4 times and (b) (for 20 min)4 times.
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TiC particles before eutectic grows. It brings the composition of melt back to the coupled zone, so that the eutectic of dendritic TiC and Fe3Al phases can grow [17,18,21]. The size of the halo is proportional to the undercooling required for second phase nucleation, and halos exist only when the primary phase is not a good nucleant of the second phase [17]. Since (Fe±28Al)± 0.53Ti±0.5C, that is, the alloy containing small amounts of titanium and carbon is on very near-eutectic composition, it does not show the occurrence of large particle above 5 mm in diameter. We can, however, con®rm the Table 4 Comparison of measured volume fraction, Vm (%), and theoretical volume fraction, Vt (%), of TiC particles in the composites Samples
Vm (%)
Vt (%)
(Fe±28Al)±3.2Ti±3C (Fe±28A1)±9.6Ti±9C
5.2 16.7
5.3 16.3
Fig. 10. Hypothetical pseudo-binary phase diagram of Fe3Al±TiC.
Fig. 11. Schematic drawing for coupled zone.
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growth of halos as shown in Fig. 8. It represents that the composition of (Fe±28Al)±0.53Ti±0.5C is on the TiC side of the eutectic composition because eutectic alloys exhibit non-reciprocal nucleating characteristic which means that one primary phase is a good nucleant of the other phase, but not vice versa [17,22±24]. In addition, in order to investigate the emerging mechanism of large TiC particles, two types of experiments were carried out; (1) TiC powder was mixed with aluminum powder, cold compacted and then arc-melted with electrolytic iron, (2) titanium, graphite and aluminum powders were mixed, cold compacted and then arcmelted with electrolytic iron. Though the mean size of TiC and graphite powders as raw materials was 1.36 mm, the microstructures of the samples were found to be identical to those of samples made by Fe±C master alloy, showing primary large particles, dendrites of iron aluminide and interdendritic eutectic. When the molten iron aluminide contacts TiC powder, substantial dissolution of TiC will occur. It is obvious, therefore, that the dissolution±reprecipitation reaction occurs and the primary large TiC particles precipitate during melting because the liquid iron±aluminum alloy containing large TiC particles is in the L+TiC two phase region. In this study, the large TiC particles showed cuboidal shape. Other study also showed that cuboidal TiC particles with (100) facets precipitate in iron melt by in-situ processing with other types of TiC particles, spherical, starlike and dendritic crystals [9]. However, 12-faced polyhedron TiC particles were reported to form in aluminum melt by insitu processing [25]. Formation mechanism of the cuboidal shape TiC is not clear at present. Increasing titanium and carbon additions decreased the volume fraction of small particles as shown in Fig. 1. Provided that melting conditions and the solubility of titanium and carbon in liquid iron-aluminum alloy at a given temperature are the same, the volume fraction of
Fig. 12. SEM micrograph for dendrites of small TiC particles extracted from as-cast (Fe±28Al)±3.2Ti±3C sample.
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small particles will depend on solidi®cation process: that is, decreasing temperature during solidi®cation will induce the decreasing solubility of titanium and carbon in liquid iron-aluminum alloy, and the precipitation process of small particles will be aected by diusivity of titanium and carbon. As a result, titanium and carbon atoms around primary large TiC particles will transfer to the surface of the large TiC particles via a path with the closest distance, and grow the large particles. Therefore, evolved microstructure will be dependent on density of primary large TiC particles. Since increasing titanium and carbon additions cause an increase in density of primary large TiC particles, the volume fraction of small particles decreases. It is also found that increased titanium and carbon contents do not increase the size of large particles (Fig. 1). Although the reason for this is not clear, one possible explanation is that this phenomenon may result from the dierence in the interfacial free energy between dierent phases [9]. If surface energy between TiC particles with dierent orientations is larger than 2 times that between liquid iron-aluminum alloy and TiC particle, TiC particles will not agglomerate [9]. On the other hand, prolonged holding in melt is inclined to coarsen the large particles as shown in Fig. 9. SEM observation after extraction of TiC from samples shows that there are a lot of large particles connected each other with parallel or 45 twisted facet orientation. This indicates that large TiC particles can grow by clustering of particles with the same crystal orientation. It has already been shown that at 1273 K iron aluminides are compatible with a wide variety of ceramics such as carbides, borides and nitrides [26]. XRD results revealed no formation of the secondary phase in type A and B samples of present work, and this is consistent with results by Subramanian et al. [4]. Recent work by them showed that certain carbides (TiC and WC) and borides (TiB2 and ZrB2) are compatible with liquid iron-aluminum alloy at 1723 K. These results may be interpreted in terms of the thermodynamic compatibility between liquid iron-aluminum alloy and solid titanium carbide. Furthermore, Table 3 shows that Fe and Al contents in primary large TiC particles, which may be formed at higher temperatures than the eutectic temperature, are more or less the same as those in small TiC particles, which may be formed at the eutectic temperature. In addition, the contents are very small. This means that solubility limits of Fe and Al to TiC are very small and they look unchanged above the eutectic temperature. On the other hand, as shown in Figs. 2 and 3, excess carbon remained after the reaction between Ti and C formed Fe3AlC phase in (Fe±28Al)± 3.2Ti±4C and (Fe±28Al)±9.6Ti±10C. The Fe3AlC phase remains after heat treatment at 973 and 1373 K for 48 h, although it shows coarsening. Evidently, excess 1% carbon addition forms the secondary phase, Fe3AlC, while excess 1% titanium is soluble to Fe±28%Al in
these composites, which is similar to the case of Fe± 28%Al binary alloy [13,14]. Heat treatment at 973 K for 48 h resulted in additional formation of small precipitates. It may be due to supersaturation of titanium and carbon during fast cooling. Similar observation has been reported in Fe± TiC in-situ composites: TiC particles can precipitate in Fe matrix by heat treatment [12]. On the other hand, the spheroidization of small TiC particles, which can be explained by Ostwald ripening, occurred in the samples heat-treated at 1373 K for 48 h. Although TiC has little solubility in the matrix, supersaturated matrix due to fast cooling and dendritic structure with narrow interspacings can induce Gibbs±Thomson eect. As shown in Fig. 5, small TiC particles change their shapes from dendritic needles to ellipsoids and their size of small particles increase by dissolution and growth. However, there was little change on large TiC particles. 5. Conclusions In-situ iron aluminide based TiC composites containing up to 16.7 vol% of TiC can be fabricated without agglomeration by arc-melting. The in-situ carbide formed have been identi®ed as TiC by X-ray diraction analysis, ICP analysis and the Combustion-Infrared Absorption method. The method to determine the composition of the titanium carbide formed in iron aluminides has been established. TiC has an average composition of Ti±48.4 mol% C with very small scatter for all samples examined. The microstructure of composites containing Ti and C above 3.2 and 3%, respectively, consists of primary large TiC particles of 5±40 mm, particle-free dendrites of iron aluminide and matrix with small TiC particles of about 1 mm. Large TiC particles are in cuboidal shape, while small ones are in dendritic shape. This complex microstructure is caused by eutectic reaction between Fe3Al and TiC phases. Particle-free dendrite of iron aluminides grows as a halo because primary large TiC is a poor nucleant of the Fe3Al phase. In addition, excess 1% carbon addition in these insitu composites forms the secondary spherical or needlelike Fe3AlC phase, although excess 1% titanium addition is soluble in Fe±28%Al matrix. Subsequent heat treatment at 973 K for 48 h produces very small TiC precipitates due to the supersaturated titanium and carbon in the matrix, and heat treatment at 1373 K for 48 h induces the spheroidization of small TiC particles due to Ostwald ripening. Longer melting time is inclined to coarsen the large particles by growth due to dissolution± reprecipitation and/or by clustering of particles with the same crystal orientation. Although we have not examined the reaction temperature and cooling rate in this study, these factors may also signi®cantly in¯uence the size and distribution
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of carbides in the matrix. Since the distribution and size of TiC particles aect signi®cantly mechanical properties of the composites, further study to investigate the formation mechanism of TiC and thereby to control the microstructure of composites is needed to improve mechanical properties of composites. Acknowledgements The authors gratefully acknowledge the experimental assistance provided by Mr. S. Ishikuro in analysis of TiC composition. References [1] Mckamey CG, Devan JH, Tortorelli PF, Sikka VK. J Mater Res 1991;6:1779. [2] Prakash U, Buckley RA, Jones H, Sellars CM. ISIJ Int 1991;31:1113. [3] Liu CT. In: Baker I, Darolia R, Whittenberger JD, Yoo MH, editors. High-temperature ordered intermetallic alloys V, MRS Symp Proc, Vol. 288. Pittsburgh (PA): MRS, 1993. 3±19. [4] Subramanian R, Schneibel JH. JOM 1997;49:50. [5] Schneibel JH, Carmichael CA, Specht ED, Subramanian R. Intermetallics 1997;5:61.
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[6] Aikin Jr. RM. JOM 1997;49:35. [7] Popov AA, Gasik MM. Scripta Metall 1996;35:629. [8] Raghunath C, Bhat MS, Rohatgi PK. Scripta Metall 1995;32:577. [9] Liu Z, Fredriksson H. Metall Mater Trans A 1997;28A:471. [10] Liu Z, Fredriksson H. Metall Mater Trans A 1996;27A:407. [11] Nukami T, Flemings MC. Metall Mater Trans A 1995;26A:1877. [12] Terry BS, Chinyamakobvu OS. J Mat Sci 1992;27:5666. [13] Baligidad RG, Prakash U, Radhakrishna A. Mater Sci Eng 1997;A230:188. [14] Baligidad RG, Prakash U, Radhakrishna A, Ramakrishna Rao V, Rao PK, Ballal NB. Scripta Metall 1997;36:667. [15] Massalski TB, editor. Binary alloy phase diagrams, vol. 2. Metals Park (OH): ASM, 1990. 888±91. [16] Elliott R. Eutectic solidi®cation processing. Butterworths monographs in materials. London: Butterworths 1983. 55±6. [17] Bluni ST, Notis MR, Marder AR. Acta Metall Mater 1995;43:1775. [18] Gilgliotti MFX, Colligan GA, Powell GLF. Metall Trans 1970;1:891. [19] Hunt JD, Jackson KA. Trans AIME 1967;239:864. [20] Gilgliotti MFX, Colligan GA, Powell GLF. Metall Trans 1970;1:1038. [21] Barclay RS, Niessen P, Kerr HW. J Cryst growth 1973;20:175. [22] Holloman JH, Turnbull D. Trans AIME 1951;9:803. [23] Sundquist BE, Mondolfo LF. Trans AIME 1961;221:157. [24] Lemaignan C. Acta Metall 1981;29:1379. [25] Morimoto H, Nomura M, Ashida Y. J Japan Inst Metals 1995;59:429 [26] Misra AK. Metall Mater Trans A 1990;21A:441.
In-situ production and microstructures of iron aluminide/TiC composites S.H. Ko*, S. Hanada Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 16 September 1998; received in revised form 24 December 1998; accepted 28 December 1998
Abstract In this study we have tried to produce the titanium carbide reinforced iron aluminide composites by in-situ reaction between titanium and carbon in liquid iron±aluminum alloy doped with titanium and carbon. A homogeneous distribution of titanium carbide particles in the iron aluminide matrix up to about 16 vol% of titanium carbide was intended without agglomeration. The composition of TiC formed during in-situ reaction was investigated by ICP analysis and the Combustion-Infrared Absorption method after chemical dissolution of the iron aluminide matrix. It is found that the composition of titanium carbide formed during melt processing is an average of Ti±48.4 mol% C. In addition, titanium carbide has very low solubility of Fe and Al. The microstructure of composites consists of three dierent regions; primary large TiC particles of 5±40 mm, matrix with small dendritic TiC particles of about 1 mm and particle-free regions around primary large TiC particles. The formation of this complex microstructure can be explained by assuming the Fe3Al±TiC pseudo-binary system containing the eutectic reaction. Particle-free regions are halos of iron aluminide phase and the formation of halos is explained by coupled zone concept. Subsequent heat treatment at 1373 K for 48 h induces spheroidization and/or coarsening of small TiC particles, while microstructure after heat treatment at 973 K for 48 h exhibits the additional formation of small TiC precipitates. Though excess 1 mol% Ti addition over the Ti content for TiC formation is soluble to Fe±28 mol% Al, excess 1 mol% C addition forms the secondary Fe3AlC phase during melt processing. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Intermetallic, miscellaneous; A. Iron aluminides, based on Fe3Al; C. Heat treatment; C. Melting; D. Microstructure
1. Introduction Because of good resistance to high-temperature sul®dizing and oxidizing environments, iron aluminides have high potential for structural applications at high temperatures under corrosive environments [1±3]. In addition, they have relatively low density compared to stainless steel, which leads to a better strength-to-weight ratio [1]. Recently, a few studies have been carried out to utilize carbides or borides as reinforcements in iron aluminide via a powder metallurgy process [4,5]. Among them titanium carbide with high hardness, low density and low chemical reactivity was found to be a suitable reinforcement for wear and high temperature applications [4]. Although many processes have been developed to incorporate reinforcing particles into iron aluminide * Corresponding author now at present address: Tohoku National Industrial Research Institute, Agency of Industrial Science and Technology, MITI, Sendai 983-8551, Japan.
matrix, an in-situ technique in which a thermodynamically stable reinforcing phase is produced during processing can be the most economical on synthesis of composites. A typical example will be the conventional melt processing based on melting and casting, by which in-situ composites can be prepared at low production costs and high eciency. Furthermore, this processing has an advantage that surfaces of reinforcements are not contaminated with oxidation ®lms or other detrimental surface reactions [6±8]. Over the past several years a number of studies have been made on in-situ TiC synthesis by melt processing [7±12]. However, only a little is known about the reactions between titanium and carbon, nucleation and growth of the TiC particles, and particle distributions in the liquid solution, which are necessary to understand and control the process [9]. Furthermore, only few attempts have so far been made on study for intermetallic compound alloys, and no study has ever tried to fabricate in-situ composites based on iron aluminides.
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966 -9 795(99)00002 -3
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On the other hand, some other studies [13,14] about Fe3Al based alloys containing Fe3AlC precipitates have been reported, where these alloys exhibited higher yield stress than binary Fe3Al alloy from room temperature up to 873 K maintaining some room temperature ductility (above 3.5%). In addition, these alloys show reduced susceptibility to environmental embrittlement compared to binary alloy, which was explained by the trapping eects of hydrogen due to Fe3AlC precipitates [14]. In this study TiC reinforced Fe±28 mol% Al based composites are synthesized by in-situ reaction between titanium and carbon in liquid iron±aluminum alloy and in¯uence of contents of titanium and carbon on microstructures is investigated. In addition, we investigate the eect of heat treatment at various temperatures and melt processing condition on microstructures. The composition of TiC formed during in-situ melt processing is also examined. 2. Experimental procedure Three types of samples were prepared to study in-situ formation of titanium carbide in Fe±28 mol% Al alloys (unless otherwise noted, all compositions will be given in mol%) through liquid phase reaction. As for raw materials 99.9 mass% electrolytic iron, 99.9 mass% aluminum, 99.9 mass% titanium and Fe±3.9 mass% C master alloy were used. A vacuum induction furnace was used for preparing of Fe±3.9 mass% C master alloy. The samples were obtained by arc melting. Each sample was remelted four times to ensure composition homogeneity. In order to investigate the forming composition of TiC during melt processing, ICP analysis and the Combustion±Infrared Absorption method were performed after preliminary chemical dissolution of the iron aluminide matrix using a HCl solution. Since iron aluminide matrix is dissolved in a HCl+H2O (1:1) solution, TiC particles can be chemically extracted. Extracted particles were dissolved in a KHSO4 solution to examine the composition of titanium using ICP analysis. The CombustionInfrared Absorption method was performed to investigated the composition of carbon in the particles. The cast samples were cut into 1.5 mm thickness plates and were given heat treatment at 973 and 1373K for 48 h in a mue furnace after encapsulating samples in argon-®lled silica tubes. After desired period of heat treatment samples were water quenched from the respective temperatures. Samples were mechanically polished using diamond slurry and a chemical solution. All the samples were examined metallographically before and after etching, and by X-ray diractometry for phase identi®cation. In addition, the matrix of each sample was partially etched away to observe the true particle shape by a scanning electron microscope. The
samples were put in a HCl+H2O (1:1) solution at room temperature for 2 min to dissolve the surface of matrix. The eect of melt processing time on microstructure was investigated under a given melt processing condition. One sample was arc-melted for 5 min, and the other sample was arc-melted for 20 min at one time. Both the samples were remelted four times. 3. Results From the binary phase diagram of Ti±C [15], it can be found that titanium carbide (TiC) has the composition range from Ti±32% C to Ti±48.5% C at 1923 K. Preliminary samples were prepared and their compositions were analyzed by ICP analysis and the CombustionInfrared Absorption method to investigate the composition of TiC formed during melt processing. The nominal composition of each preliminary sample was Ti±43% C, Ti±45% C and Ti±47% C. Titanium carbide particles were found to be formed at an average composition of Ti±48.4% C in all samples. Scatter of compositions of formed titanium carbide particles was negligibly small from sample to sample. Table 1 shows nominal compositions of samples in present work. Type A samples were alloyed with carbon such that the average mol percent ratio of titanium to carbon coincides with that of TiC (hereafter, titanium carbide with composition of Ti±48.4% C will be simply called TiC) formed by the liquid reaction. The type B and C samples were alloyed so as to have 1% titaniumrich and 1% carbon-rich compositions, respectively. Compositions of particles and matrices of some samples are summarized in Table 2. It is found that TiC particles contain only minimal amounts of Fe and Al, indicating that TiC has very low solubility of Fe and Al. Correspondingly, the compositions of matrices were close to the nominal composition, Fe±28% Al, although titanium contents show some variation ranging from 0.04 to 0.20%. Optical micrographs of unetched, as-cast (Fe±28Al)± 3.2Ti±3C and (Fe±28Al)±9.6Ti-9C and their schematic microstructure are shown in Fig. 1. The characteristic distribution of large particles in the Fe±28%Al matrix Table 1 Nominal compositions of alloys (compositions are given in mol%) Samples
Type Fe
Al
Fe±28A])±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±4.2Ti±3C (Fe±28A])±10.6Ti±9C (Fe±28A])±3.2Ti±4C (Fe±28A])±9.6Ti±10C (Fe±28A1)±0.53Ti±0.SC (Fe±28A1)±1.6Ti±1.5C
A
26.27 3.20 3.00 22.50 9.59 9.00 25.99 4.20 3.00 22.52 10.59 9.00 25.99 3.20 4.00 22.52 9.59 10.00 27.71 0.53 0.50 27.13 1.60 1.50
B C A
67.54 58.62 66.82 57.90 66.82 57.90 71.26 69.80
Ti
C
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can be achieved by the liquid reaction without agglomeration. As depicted in Fig. 1(c), the microstructures of as-cast samples consist of two types of TiC particles concerning particle size, shape and distribution, and Table 2 Compositions of particles and matrices Samples
Particles Fe-mol% Al-mol% Ti-mol% C-mol%a
(Fe±28A1)±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±0.53Ti±0.5C (Fe±28A1)±1.6Ti±1.5C
0.50 0.11 0.41 0.51
Samples (mol%)
Matrix
0.23 0.11 0.24 0.23
51.6 51.7 50.5 50.4
47.7 48.1 48.9 48.9
Fe-mol% Al-mol% Ti-mol% (Fe±28A])±3.2Ti±3C (Fe±28A1)±9.6Ti±9C (Fe±28A])±0.53Ti±0.5C (Fe±28A1)±1.6Ti±1.5C
71.7 72.2 72.1 71.8
28.2 27.6 27.8 28.0
0.04 0.16 0.09 0.20
a These values were obtained by the Combustion±Infrared Absorption method. The others were obtained by ICP±OES analysis.
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particle-free regions, which are dendrites of iron aluminide phase growing around large TiC particles. (Fe±28Al)±4.2Ti3C and (Fe±28Al)±10.6Ti±9C showed microstructures similar to (Fe±28Al)±3.2Ti±3C and (Fe±28Al)±9.6Ti±9C, respectively. However, carbon-rich samples, (Fe±28Al)± 3.2Ti±4C and (Fe±28Al)±9.6Ti±10C, show the additional formation of precipitates which have two dierent shapes, i.e. spherical and needle-like shapes, as shown in Fig. 2. XRD patterns of the as-cast samples are presented in Fig. 3. One can see Fe3Al and TiC peaks in all samples, indicating the formation of TiC particles during melting and casting, while Fe3AlC peaks can be seen in (Fe± 28Al)±3.2Ti±4C and (Fe±28Al)±9.6Ti±10C. The matrix of (Fe-28Al)-3.2Ti-3C sample was partially removed by chemical deep etching to observe the two types of TiC particles in detail. A SEM micrograph is shown in Fig. 4. Large TiC particles with sizes of about 5±40 mm are formed as faceted crystals, while small ones of about 1 mm have a dendritic structure. Discontinuous distribution in size and dierent shapes of the TiC particles imply that these particles are produced through dierent formation paths. From Figs. 1 and 4, microstructures can be divided into three regions
Fig. 1. Micrographs of as-cast samples of (a) (Fe±28Al)±3.2Ti±3C, (b) (Fe±28Al)±9.6Ti±9C and (c) schematic illustration of microstructure.
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Fig. 2. Micrographs of as-cast samples of (a) (Fe±28Al)±3.2Ti±4C and (b) (Fe±28Al)±9.6Ti±10C.
Fig. 4. SEM micrograph of as-cast (Fe±28Al)±3.2Ti±3C sample after chemical removal of Fe3Al matrix. Fig. 3. XRD results of as-cast samples.
consisting of large particles, matrix with small particles and particle-free region around large particles. Fig. 5 shows optical micrographs of etched, as-cast (Fe±28Al)± 3.2Ti±3C and (Fe±28Al)±9.6Ti±9C. Almost all boundaries of grains delineated by dierent contrasts in (Fe±28Al)± 3.2Ti±3C correspond to boundaries between particles-free matrices and small particle regions, while grain boundaries in (Fe±28Al)±9.6Ti±9C coincide partially with boundaries of small particle regions. Fig. 6 shows optical micrographs of (Fe±28Al)± 3.2Ti±3C heat treated at 973 and 1373 K for 48 h followed by water-quenching. The heat treatment at 973 K causes additional formation of ®ne precipitates, whereas the heat treatment at 1373 K leads to spheroidization and/or coarsening of already existing small particles. XRD analysis indicated that the additional precipitates at 973 K are identi®ed as TiC particles, and no other phase is formed at 1373 K except Fe3Al and TiC. Fig. 7 shows a SEM micrograph of (Fe±28Al)±3.2Ti±3C heat treated at 1373 K for 48 h followed by water-quenching,
where the matrix was partially removed by chemical etching. It is evident from comparison with Fig. 4 that morphology of small particles is changed from dendritic needles to ellipsoids. The compositions of large and small particles were examined using as-cast and heat treated samples. In this experiment large and small particles were distinguished from the dierence in sedimentation speed of particles in acetone. Compositional analysis indicates that particles have almost the same composition independent of particle size and heat treatment (Table 3). Fig. 8 shows the microstructures of unetched, as-cast (Fe±28Al)±0.53Ti±0.5C and (Fe± 28Al)±1.6Ti±1.5C. Decrease in titanium and carbon contents results in the disappearance or the reduced volume fraction of large particles. Large particles greater than 5 mm were not observed in (Fe±28Al)± 0.53Ti±0.5C, but observed with a very low density in (Fe±28Al)±1.6Ti±1.5C. Fig. 9 shows optical micrographs of unetched, as-cast (Fe±28Al)±3.2Ti±3C with dierent holding times in melting. Prolonged holding in the molten state appears to result in coarsening of large particles.
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Fig. 5. Micrographs of as-cast samples after chemical etching of (a) (Fe±28Al)±3.2Ti±3C and (b) (Fe±28Al)±9.6Al±9C.
Fig. 6. Micrographs of (Fe±28Al)±3.2Ti±3C sample heat-treated at (a) 973 K and (b) 1173 K for 48 h followed by water-quenching.
Fig. 7. SEM micrograph of (Fe±28Al)±3.2Ti±3C sample after heat treatment at 1373 K for 48 h followed by chemical removal of Fe3Al matrix.
4. Discussion TiC can be formed in the iron aluminide by a reaction between titanium and carbon, as mentioned before. Provided that a TiC particle has theoretical density, 4.93 g/cm3,
the volume fraction of TiC particles can be calculated. After complete dissolution of the iron aluminide matrix in a HCl+H2O solution, the weight of extracted TiC particles was measured. The mass percent of extracted TiC particles in (Fe±28Al)±3.2Ti±3C and (Fe±28Al)±9.6Ti± 9C were 4.01, and 13.18 mass%, respectively. From the densities of 6.52 g/cm3 for Fe±28%Al and 4.93 g/cm3 for TiC, the volume fraction of TiC particles was evaluated to be 5.2% in (Fe±28Al)±3.2Ti±3C and 16.7% in (Fe±28Al)±9.6Ti±9C. Assuming that alloyed titanium and carbon are entirely consumed to form TiC, one can calculate a theoretical volume fraction of TiC. As shown in Table 4, volume fractions measured from mass percent of extracted TiC particles are in reasonable agreement with those calculated from nominal mass percent of titanium and carbon. In the Fe±Al±Ti±C system, at 1723 K the liquid phase containing only 4.9% C and 64 mol ppm Ti are in equilibrium with Ti0.512C0.488 [4]. Therefore, most TiC particles may precipitate in liquid iron-aluminum alloy at least at 1723 K. Particle-free regions, i.e. dendrites of the iron aluminide phase, are formed around large TiC particles and then interdendritic eutectic grows as shown in Figs. 1 and 4. Based on these observations the
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sequence of TiC formation can be explained by assuming the Fe3Al±TiC pseudo-binary system containing the eutectic reaction between Fe3Al and TiC phases as shown in Fig. 10. Since TiC has a much higher melting point than Fe3Al, the eutectic composition may lie much closer to Fe3Al than to TiC: Fe3Al±TiC system is an anomalous eutectic structure [16]. In such a case the coupled zone, a range on the phase diagram de®ning compositions and temperatures at which the two eutectic Table 3 Compositions of large and small particles in (Fe±28Al)±3.2Ti±3C insitu composite
As-cast Heat-treated at 1373 K for 48 h a
Kind of particles
Elements
Large particles Small particles Large particles Small particles
0.48 0.45 0.49 0.14
Fe± Al± Ti± C± mol% mol% mol% mol%a 0.24 0.20 0.23 0.14
52.0 52.4 52.2 52.0
47.3 47.0 47.1 47.7
These values were obtained by the Combustion±Infrared Absorption method. The others were obtained by ICP analysis.
phases can grow with similar velocities, will incline strongly to TiC side as shown in Fig. 11. [16±18]. It has also shown that the coupled zone for faceted-nonfaceted eutectic systems lies at or near the extended non-faceted liquidus and faceting is a signi®cant feature of anormalous eutectic structure [18±20]. Since the primary large TiC particles and dendrites of small TiC particles show faceted structure as shown in Figs. 4 and 12, this assumption about the shape of coupled zone is proper. At ®rst, it is reasonable to consider that large particles are formed in the melt because the liquid iron-aluminum alloy during melt processing is in the L+TiC two phase region. During cooling the temperature of melt containing primary TiC particles attains the eutectic temperature of Fe3Al and TiC phases. If primary large TiC is a poor nucleant of the Fe3Al phase, the composition of melt including primary large TiC particles will follow the extension of the liquidus line to a temperature which is below the eutectic temperature [17,18,21]. Since the coupled zone in hypothetical Fe3Al-TiC pseudo-binary phase diagram is skewed to TiC side as shown in Fig. 11, the composition of melt lies outside the coupled zone and the melt is supersaturated in Fe3Al. Therefore, dendritic growth of Fe3Al phase forming the particlefree region, halo, will proceed around the primary large
Fig. 8. Micrographs of as-cast samples of (a) (Fe±28Al)±0.53Ti±0.5C and (b) (Fe±28Al)±1.6Al±1.5C.
Fig. 9. Eects of melt processing time on microstructure of (Fe±28Al)±3.2Ti±3C sample. (a) (for 5 min)4 times and (b) (for 20 min)4 times.
S.H. Ko, S. Hanada / Intermetallics 7 (1999) 947±955
TiC particles before eutectic grows. It brings the composition of melt back to the coupled zone, so that the eutectic of dendritic TiC and Fe3Al phases can grow [17,18,21]. The size of the halo is proportional to the undercooling required for second phase nucleation, and halos exist only when the primary phase is not a good nucleant of the second phase [17]. Since (Fe±28Al)± 0.53Ti±0.5C, that is, the alloy containing small amounts of titanium and carbon is on very near-eutectic composition, it does not show the occurrence of large particle above 5 mm in diameter. We can, however, con®rm the Table 4 Comparison of measured volume fraction, Vm (%), and theoretical volume fraction, Vt (%), of TiC particles in the composites Samples
Vm (%)
Vt (%)
(Fe±28Al)±3.2Ti±3C (Fe±28A1)±9.6Ti±9C
5.2 16.7
5.3 16.3
Fig. 10. Hypothetical pseudo-binary phase diagram of Fe3Al±TiC.
Fig. 11. Schematic drawing for coupled zone.
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growth of halos as shown in Fig. 8. It represents that the composition of (Fe±28Al)±0.53Ti±0.5C is on the TiC side of the eutectic composition because eutectic alloys exhibit non-reciprocal nucleating characteristic which means that one primary phase is a good nucleant of the other phase, but not vice versa [17,22±24]. In addition, in order to investigate the emerging mechanism of large TiC particles, two types of experiments were carried out; (1) TiC powder was mixed with aluminum powder, cold compacted and then arc-melted with electrolytic iron, (2) titanium, graphite and aluminum powders were mixed, cold compacted and then arcmelted with electrolytic iron. Though the mean size of TiC and graphite powders as raw materials was 1.36 mm, the microstructures of the samples were found to be identical to those of samples made by Fe±C master alloy, showing primary large particles, dendrites of iron aluminide and interdendritic eutectic. When the molten iron aluminide contacts TiC powder, substantial dissolution of TiC will occur. It is obvious, therefore, that the dissolution±reprecipitation reaction occurs and the primary large TiC particles precipitate during melting because the liquid iron±aluminum alloy containing large TiC particles is in the L+TiC two phase region. In this study, the large TiC particles showed cuboidal shape. Other study also showed that cuboidal TiC particles with (100) facets precipitate in iron melt by in-situ processing with other types of TiC particles, spherical, starlike and dendritic crystals [9]. However, 12-faced polyhedron TiC particles were reported to form in aluminum melt by insitu processing [25]. Formation mechanism of the cuboidal shape TiC is not clear at present. Increasing titanium and carbon additions decreased the volume fraction of small particles as shown in Fig. 1. Provided that melting conditions and the solubility of titanium and carbon in liquid iron-aluminum alloy at a given temperature are the same, the volume fraction of
Fig. 12. SEM micrograph for dendrites of small TiC particles extracted from as-cast (Fe±28Al)±3.2Ti±3C sample.
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small particles will depend on solidi®cation process: that is, decreasing temperature during solidi®cation will induce the decreasing solubility of titanium and carbon in liquid iron-aluminum alloy, and the precipitation process of small particles will be aected by diusivity of titanium and carbon. As a result, titanium and carbon atoms around primary large TiC particles will transfer to the surface of the large TiC particles via a path with the closest distance, and grow the large particles. Therefore, evolved microstructure will be dependent on density of primary large TiC particles. Since increasing titanium and carbon additions cause an increase in density of primary large TiC particles, the volume fraction of small particles decreases. It is also found that increased titanium and carbon contents do not increase the size of large particles (Fig. 1). Although the reason for this is not clear, one possible explanation is that this phenomenon may result from the dierence in the interfacial free energy between dierent phases [9]. If surface energy between TiC particles with dierent orientations is larger than 2 times that between liquid iron-aluminum alloy and TiC particle, TiC particles will not agglomerate [9]. On the other hand, prolonged holding in melt is inclined to coarsen the large particles as shown in Fig. 9. SEM observation after extraction of TiC from samples shows that there are a lot of large particles connected each other with parallel or 45 twisted facet orientation. This indicates that large TiC particles can grow by clustering of particles with the same crystal orientation. It has already been shown that at 1273 K iron aluminides are compatible with a wide variety of ceramics such as carbides, borides and nitrides [26]. XRD results revealed no formation of the secondary phase in type A and B samples of present work, and this is consistent with results by Subramanian et al. [4]. Recent work by them showed that certain carbides (TiC and WC) and borides (TiB2 and ZrB2) are compatible with liquid iron-aluminum alloy at 1723 K. These results may be interpreted in terms of the thermodynamic compatibility between liquid iron-aluminum alloy and solid titanium carbide. Furthermore, Table 3 shows that Fe and Al contents in primary large TiC particles, which may be formed at higher temperatures than the eutectic temperature, are more or less the same as those in small TiC particles, which may be formed at the eutectic temperature. In addition, the contents are very small. This means that solubility limits of Fe and Al to TiC are very small and they look unchanged above the eutectic temperature. On the other hand, as shown in Figs. 2 and 3, excess carbon remained after the reaction between Ti and C formed Fe3AlC phase in (Fe±28Al)± 3.2Ti±4C and (Fe±28Al)±9.6Ti±10C. The Fe3AlC phase remains after heat treatment at 973 and 1373 K for 48 h, although it shows coarsening. Evidently, excess 1% carbon addition forms the secondary phase, Fe3AlC, while excess 1% titanium is soluble to Fe±28%Al in
these composites, which is similar to the case of Fe± 28%Al binary alloy [13,14]. Heat treatment at 973 K for 48 h resulted in additional formation of small precipitates. It may be due to supersaturation of titanium and carbon during fast cooling. Similar observation has been reported in Fe± TiC in-situ composites: TiC particles can precipitate in Fe matrix by heat treatment [12]. On the other hand, the spheroidization of small TiC particles, which can be explained by Ostwald ripening, occurred in the samples heat-treated at 1373 K for 48 h. Although TiC has little solubility in the matrix, supersaturated matrix due to fast cooling and dendritic structure with narrow interspacings can induce Gibbs±Thomson eect. As shown in Fig. 5, small TiC particles change their shapes from dendritic needles to ellipsoids and their size of small particles increase by dissolution and growth. However, there was little change on large TiC particles. 5. Conclusions In-situ iron aluminide based TiC composites containing up to 16.7 vol% of TiC can be fabricated without agglomeration by arc-melting. The in-situ carbide formed have been identi®ed as TiC by X-ray diraction analysis, ICP analysis and the Combustion-Infrared Absorption method. The method to determine the composition of the titanium carbide formed in iron aluminides has been established. TiC has an average composition of Ti±48.4 mol% C with very small scatter for all samples examined. The microstructure of composites containing Ti and C above 3.2 and 3%, respectively, consists of primary large TiC particles of 5±40 mm, particle-free dendrites of iron aluminide and matrix with small TiC particles of about 1 mm. Large TiC particles are in cuboidal shape, while small ones are in dendritic shape. This complex microstructure is caused by eutectic reaction between Fe3Al and TiC phases. Particle-free dendrite of iron aluminides grows as a halo because primary large TiC is a poor nucleant of the Fe3Al phase. In addition, excess 1% carbon addition in these insitu composites forms the secondary spherical or needlelike Fe3AlC phase, although excess 1% titanium addition is soluble in Fe±28%Al matrix. Subsequent heat treatment at 973 K for 48 h produces very small TiC precipitates due to the supersaturated titanium and carbon in the matrix, and heat treatment at 1373 K for 48 h induces the spheroidization of small TiC particles due to Ostwald ripening. Longer melting time is inclined to coarsen the large particles by growth due to dissolution± reprecipitation and/or by clustering of particles with the same crystal orientation. Although we have not examined the reaction temperature and cooling rate in this study, these factors may also signi®cantly in¯uence the size and distribution
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of carbides in the matrix. Since the distribution and size of TiC particles aect signi®cantly mechanical properties of the composites, further study to investigate the formation mechanism of TiC and thereby to control the microstructure of composites is needed to improve mechanical properties of composites. Acknowledgements The authors gratefully acknowledge the experimental assistance provided by Mr. S. Ishikuro in analysis of TiC composition. References [1] Mckamey CG, Devan JH, Tortorelli PF, Sikka VK. J Mater Res 1991;6:1779. [2] Prakash U, Buckley RA, Jones H, Sellars CM. ISIJ Int 1991;31:1113. [3] Liu CT. In: Baker I, Darolia R, Whittenberger JD, Yoo MH, editors. High-temperature ordered intermetallic alloys V, MRS Symp Proc, Vol. 288. Pittsburgh (PA): MRS, 1993. 3±19. [4] Subramanian R, Schneibel JH. JOM 1997;49:50. [5] Schneibel JH, Carmichael CA, Specht ED, Subramanian R. Intermetallics 1997;5:61.
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