Formation of TiC in in situ processed composites via solid-gas, solid-liquid and liquid-gas reaction...

Materials Science and Engineering, A162 (1993) 153-162 153

Formation of TiC in in situ processed composites via solid-gas, solid-liquid and liquid-gas reaction in molten A1-Ti

Subhash Khatri and Michael Koczak Department of Materials Engineering, Drexel University, Philadelphia, PA 19104 (USA)

Abstract

A novel technique to generate fine single crystal TiC platelets in an aluminum based matrix has been developed (M. J. Koczak and K. S. Kumar, US Patent 4,808,372 (1989)). The process involves decomposition of a carbonaceous gas (CH4) and reaction of nascent carbon with a strong carbide former such as Ti in an aluminum matrix at a relatively high temperature (1200-1400 °C). The highly exothermic process is moderated by means of a carrier gas and leads to a fine distribution of carbides of size 0.1-3/~m. A nucleation and growth study was carried out to understand the decomposition of the methane, distribution and subsequent reaction with an aluminum-titanium alloy to form the titanium carbide. It was observed that formation of TiC occurs in stages. Following the CH 4 decomposition, the solid carbon particles are distributed and trapped in the alloy. The reaction to form TiC is probably limited by diffusion of titanium to carbon and thereafter the carbide. After inoculating carbon in the alloy, the reaction can be completed in solid or liquid state. Trans- mission electron microscopy studies confirmed the presence of 40-50 nm amorphous carbon particulates in the alloy. It is also postulated that, in liquid state, once the reaction proceeds, the first phase to form is aluminum carbide or an aluminum-titanium carbide of the form. Given sufficient time for completion, the reaction proceeds to form the most stable carbide, i.e. TiC.

1. Introduction

During the past decade, specific material property requirements for advanced aerospace structures have escalated where conventional alloy systems are no longer suitable. Therefore attempts have been made to enhance the material properties via a reinforcing ceramic second phase of higher strength and stiffness. Various materials such as oxides, carbides and nitrides have been investigated for improving the properties of the monolithic metals. The reinforcing phase and the metal matrix are combined by various processing techniques such as powder metallurgy, preform infiltration, spray deposition, casting technologies, i.e. squeeze casting, rheocasting and compocasting [1-4]. These artificial composites have resulted in the realization of strength and modulus goals, however rarely have the composites been economically viable. Additionally, these processes also have problems such as interface reactions leading to undesirable products, thermodynamic and mechanical incompatibility, internal stresses owing to the mismatch of coefficients of thermal expansion of the various phases.

In the last decade, new in situ fabrication technologies have been developed for processing metal and ceramic matrix composites. Simply stated, in situ processes involve a chemical reaction resulting in the formation of a thermodynamically stable reinforcing

ceramic phase. Some of these technologies include DIMOX TM, XD TM, SHS and reactive gas infiltration [1, 5,6].

One of the promising in situ techniques to fabricate metal matrix composites has been recently patented by Koczak and Kumar [6]. The composite is processed by infiltration of a reactive gas, i.e. CH4, NH3, N 2 in an inert carrier gas such as argon in a molten alloy consisting of the matrix constituent and a transition metal solute which is a strong carbide and/or a nitride former. The process has been found suitable to form composites with a fine reinforcement size (0.1-3/zm) at low to moderate volume fractions ( < 15%). Owing to the in situ formation of the nascent carbide, the interface is very clean and continuous. Additionally, the process promises to be an economically feasible, near net shape fabrication technology, e.g. squeeze casting, centrifugal casting. The reaction is typically accomplished at a high temperature (1200-1400 °C) and involves decomposition of the gas followed by a reaction to form the carbide and/or the nitride. Decomposition of the carbonaceous gas can result in amorphous carbon in the form of fine particulate carbon black [7], or in spherical and fibrous graphitic form depending on the temperature and the substrate conditions (currently being researched at Drexel University). Under favorable conditions, the decomposition can also result in formation of diamond-like

0921-5093/93/$6.00 © 1993 - Elsevier Sequoia. All rights reserved

154 S. Khatri, M. Koczak / Formation of TiC composites in molten AI-Ti

particles [8]. The thermodynamic requirement is that the solute derived reinforcing phase should be very stable compared with the carbide/nitride of the matrix material, e.g. TiC vis-d-vis AI4C 3. Typically, the reaction sequence can be represented for an alloy M-X as

C H 4 = C + 2H 2

M - X + C = M - X + X C

2NH 3 = 2N + 3H 2

2M - X + N 2 = 2M + 2XN

In addition to methane, other carbonaceous gases such as acetylene, CzH 4 and carbon tetrachloride, CCI 4 can be utilized to conduct the reaction. The free energy for dissociation for these gases are [9]

CH 4 = C + 2H 2 AG° = 69 218 - 51.33 T log T+ 65.57 T J mol-

CC14 = C + 2CI 2 A G ° = 109 988 + 21.62 T log T+ 206.7 T J mol- 1

A point to note is that decomposition of different gases varies with temperature and the processing temperature must account for the variation. The fraction of the gas decomposed can be calculated as follows [10]

m C,Xm = nC + ~- X 2

where X can be either hydrogen or chlorine and A G ° is the standard free energy change for the reaction. If the activity of carbon is one, from thermodynamics [11]

AG °= - R T I n (pxm/Z/Pc,x,)

where Px2 m/z and Pc,x, are the partial pressures of hydrogen/chlorine and the carbonaceous gas respectively. Thus from knowledge of the thermodynamic quantities, the equilibrium ratio of the partial pressure of the gases, and hence fraction decomposition of the carbonaceous gas can be calculated. The decomposition of the gases, C H 4 and CC14, as a function of temperature is depicted in Fig. 1. It is clear that CC14

decomposes at lower temperatures as compared with CH4. However among hydrocarbon gases, higher molecular weight, e.g. C2H6, would result in higher decomposition temperatures. The gases with higher molecular weights also, however, have higher C / H 2

ratio, implying that lower volumes of gases would be required to complete the carburization reaction.

Upon gas decomposition, carbon is released in an Ar/H 2 bubble, and carbon-metal reaction can occur at the gas-metal interface [12]. The liquid-solid reaction is limited by diffusion of carbon inside the gas bubble and that of titanium in the liquid metal to the vapor-liquid interface. The gas decomposition reaction is also limited by the bubble residence time in the high temperature melt. In addition to aluminum, the most common matrix material, copper, nickel and intermetallic compounds, e.g., A13Ti, NiTi, etc. show potential for being matrix material by this process. The reinforcing phases mainly include carbides and nitrides of transition metals. In addition to the transition metal carbides and nitrides, SiC and A1N have also been formed via reactive gas infiltration processing. Although the A1/TiC system has been most extensively studied, A1/AIN, AI/TiN, AI-Si/SiC, Cu/TiC, Ni/TiC and HfC, TaC and NbC in aluminum matrix have also been processed successfully [13].

Depending on the activity of the carbon, alloy chemistry and the partial pressure of oxygen in the gas bubble, a series of products such as Ti, TiC, A14C3,

TiO2 and AI3Ti are possible [14]. The phase stability of these components is depicted in Fig. 2 at a temperature of 1150 °C. All the calculations are carried out for an alloy of composition AI-5wt.%Ti. It can be seen that a carbon activity of at least 2.12 x ] 0- 5 has to be main- tained and the oxygen partial pressure should not exceed 9.96 x ]0 -24, SO that TiC is stable as compared with AI3Ti and T i O 2. Similar two dimensional and three dimensional phase stability diagrams can be constructed to understand conditions for stability of TiC vs. AI4C 3 as a function of temperature, Po2, ac, aAl and avi.

1 . 0 - --

0.8'

Q . ~ 0.6'

a

.~ 0.4'

u _

0.2'

0 . 0 - - - ~ ' ~ I Carbon Tetrachlodde

300 500 700 '900 11 O0 1300 1500

Temperature (°K) Fig. 1. Fraction of gas decomposed as a function of temperature.

0 -

-2 -

..-1

- 6 -

- 8 -

AI + AI3Ti

-10 - 2 6 - 2 4

AI + TiC

AI +rio2

- 2 2 - 2 0

Log po2

Fig. 2. Al-Ti-C phase stability diagram at 1150 °C.

In addition to the thermodynamic stability, the kinetic of the reactions are important as the reaction is diffusion controlled. In this study, a series of in situ experiments were conducted by bubbling CH 4 through a dilute A1-Ti alloy at 1150 °C to form an AI-TiC composite. The experiments were conducted for very short times to understand the decomposition of methane, carbon formation, and initial nucleation and growth of the reinforcing carbide.

The key features in the process optimization would involve improving the efficiency of the reaction in terms of consumption of the available carbon as well as the production of very fine and homogeneous carbide size distribution, preferably in the sub-micron range. In addition, complete reaction of titanium to TiC should be achieved, or formation of coarse AI3Ti needles results in a brittle structural composite [9]. The control of the reaction chemistry and the subsequent reinforcing phase size can be better achieved if the gas decomposition kinetics and nucleation and growth of the reinforcing phase is understood.

2. Experimental procedure

The carbonaceous gas infiltration reactions were carried out in a Vacuum Industries System 7 furnace induction heated by a Inductotherm generator. The experimental control and processing system is shown

RS-232C TO TEMPERATURE CONTROLLER

in Fig. 3. The melt temperatures and gas flow rates are monitored via computer controlled electronic gas flowmeters from Porter Instruments. Al- lwt.%Ti charges of approximately 85 g were heated in an 99.9% crystallized alumina crucible placed in a graphite susceptor. The titanium was added in the form of Al-10wt.%Ti master alloy. The alloy was diluted by addition of commercially pure aluminum to final composition. The furnace was evacuated and heated to 300 °C for 30 min to remove moisture. Thereafter the furnace was filled with an argon atmosphere and heated to 1150 °C at a rate of about 10 °C min-~. The alloy was held at 1150 °C for 30 min to allow for complete AI3Ti dissolution. At 1150 °C, an alumina tube of 3-mm diameter was inserted in the alloy and a mixture of argon and methane was passed through the melt. The various flow rates and time periods for which the gas was bubbled through the melt are tabulated in Table 1.

This study consisted of four types of experiments, i.e. reactions which are denoted as A, B, C, and D. These experiments have been outlined in Fig. 4. For sample A, methane was bubbled through the melt for 2.5 min at 0.5 1 min-1. The amount of carbon infiltrated in the melt was about 300% of the required amount for full conversion of Ti to TiC. Thereafter the alumina tube was taken out of the furnace and the sample was cooled to 25 °C. In addition, sample A was subjected to solid state solutionizing treatment at a

[ I-'11~ I-q

o 0 o T/C

FLOW CONTROLLERS

S. Khatri, M. Koczak / Formation of TiCcomposites in moltenAl-Ti 155

Fig. 3. Schematic of the experimental apparatus with various control features.


TABLE 1. Summary of the solid and liquid state reaction for AI-TiC-C composites synthesis via in situ processing

Reaction A B C D

Processing 2.5 min C H 4 2,5 min C H 4 2.5 min CH 4 (0.5 lmin-I)at 1150 °C (0.5 I min-t) plus 18 h (0.51 min- l)plus 120

at 640 °C min argon (0.51 min- 1) at 1150 °C

Microstructure TiC ( < l/tm) AI3Ti (50-500/~m)

TiC ( < 1 ktm) AI3Ti (50-500 #m) and carbides at AI3Ti-AI interface

TiC ( < 10 #m) AI3Ti (30-50/~m)

2.5 min C H 4 (0.51 min- 1) plus 120 rain C H 4 (0.1 1 min- ~) at 1150 °C

TiC ( < 10/~m)

Reaction A CH4 Infiltration for

2.5 minutes

• Reaction B / Solid State t Homogenization for 1

hrs. at 640°C

Reaction C Liquid State

Argon (0.5 I/min) Infiltration for 120

minutes

Reaction D Liquid State

CH4 (0.1 I/min) Infiltration for 120

minutes

I / / Analysis

(SEM, TEM, Optical and XRD)

Fig. 4. Flowchart outlining the various reaction sequences.

temperature of 640 °C for 18 h, Reaction B (Table 1). Further reactions to the particles were studied by bubbling pure argon and a mixture of argon and methane through the melts as in experiments C and D respectively.

The in situ processed composites were sectioned and polished for microscopic examination. The composites were examined optically using a Zeiss microscope and via scanning electron microscope in JEOL Model 35 CE In addition thin foils were electro- polished in a perchloric acid solution at - 2 0 °C for transmission electron microscopy studies. The specimens were observed in a transmission electron microscope, JEOL 100 kV. The specimens were examined for crystal structure and chemical composition via a Siemens X-ray diffractometer. In addition, an aluminum matrix of the composite was dissolved in a 20% NaOH solution and the undissolved powder residue was analyzed by X-ray diffraction.

3. R e s u l t s

3.1. Carbon inoculation (reaction A) Figures 5(a) and 5(b) show the X-ray diffraction

patterns of the as-processed composite and the powder

i

J

L

t t

[

2,

A 1

AI3 "

AI3Ti

' 1

AI3Ti J

(a) x 2 % h e t a _u : 6 0 1 , L i n e a P 88,i~)

i l TiCZ']~ C~Kal +2 ~i

i

H

C C

i TiC

, C: AI3Ti i

( Z~.80~ × : 2t}~eta 9 : 1679, LineaP 88,1~)

(b)

Fig. 5. X-ray diffraction patterns of samples after bubbling CI-'I 4 for 2.5 min (reaction A): (a) composite; (b) powder,

residue remaining behind after dissolution of aluminum in a NaOH solution (sample A). The diffraction pattern shows the presence of aluminum and A13Ti. There are also several unmarked peaks which have been identified as various ferrous aluminides, e.g. FeA13. Iron was present in the commercially pure aluminum as an impurity. Although the presence of TiC is not clear from Fig. 5(a), XRD of the extracted reinforcing powder ~shows the presence of TiC, Fig. 5(b). It should be mentioned that three times the required amount of CH4 for complete conversion of titanium to TiC was infiltrated in the melt. The presence of unreacted titanium was observed optically

S. Khatri, M. Koczak / Formation of TiC composites in molten Al-Ti 157

(Fig. 6(a)) in the form of AI3Ti sharp precipitate needles of up to 50 ~m in thickness and 500/~m long. Figure 6(b) shows 1-10/~m iron aluminide platelets. In addition, the presence of fine submicron carbide phase is also seen. Figure 7 reveals the presence of a large amount of very fine (10-50 nm) circular particulates, many of them present in clusters. The particles are amorphous in nature and are identified as carbon. These particles are of an average size of 43.3 nm with a standard deviation of 8.8 nm.

In summary, short time methane inoculation results in methane decomposition producing 40 nm spherical or circular amorphous carbon particles leaving the coarse AI3Ti unreacted with a minor amount of interdendritic iron aluminide as an impurity.

3.2. Solid state homogen&ation reaction (reactions A and B)

Following short time carbon inoculation (e.g. reaction A), the C/AI/A13Ti composite sample was subjected to solid solution treatment at 640 °C for 18 h. Figures 8(a) and 8(b) are the scanning and transmission electron micrographs of the specimen after the solid state diffusional homogenization anneal. The scanning

Fig. 6. Typical micrograph after bubbling C H 4 for 2.5 min (reaction A): (a) optical; (b) SEM.

electron micrograph shows initiation of a solid state homogenization process, where the titanium from the A13Ti needles dissolves and diffuses through the aluminum matrix and forms fine particles near the needles. A majority of these particles are AI3Ti. Some of the particles as marked on the micrograph show a brighter contrast and thus indicate a greater presence of a higher atomic species, i.e. Ti. The transmission electron micrograph, Fig. 8(b) shows the presence of carbon particulates, as in Fig. 7 in the liquid state reaction (reaction A). The average diameter of the carbon particles is 56.6 nm (standard deviation, 9.3 nm) indicating coarsening during the solid state homogenization, Fig. 8(b); a SAD pattern shows the presence of diffusive rings, indicating the presence of amorphous carbon in the material. In summary, carbon inoculation coupled with a solid state homogenization at 640 °C, produced limited carbon coarsening and an interfacial reaction on the TiAI 3 intermetallics.

3.3. Liquid state homogenization reaction Following the carbon inoculation (reaction A), argon

was bubbled through the melt at 1150 °C for 120 min (reaction C). Figures 9(a) and 9(b) are the XRD patterns for sample C, i.e. after bubbling C H 4 for 2.5 min and argon for 120 min. Although not conclusive, there is some evidence of the presence of Ti3AIC. The powders also clearly show conversion of C to TiC. The coarse AI3Ti needles are seen in the scanning electron and transmission electron micrographs, Figs. 10(a) and 10(b) and are 10-20/~m in size. The AI3Ti needles are not as coarse as seen after reactions A and B. The scanning electron micrographs also revealed presence of TiC particles of less than 10 ~m. The transmission electron micrograph, Fig. 10(b), shows the presence of the amorphous carbon particulates present in clusters.

Fig. 7. TEM showing carbon particles after bubbling C H 4 for 2.5 rain (reaction A).


i!:i! i ii ̧̧ i

A1 A1

C ~ a l ÷ 2

I

A1 A1

A1

I B: Ti3AIC

28.8~ x : 2theta g : 604. Linea~ g8,108>

(a)

l l- ! I

2 8 , 8 ~ x : 2 t i a r a g : 6 1 9 , h i n e a ~ 8 8 . 1 0 0 >

(b) Fig. 9. X-ray diffraction patterns of samples after bubbling C H 4

for 2.5 min and argon for 120 min (reaction C): (a) composite; (b) powder.

10 #m in size, with the majority of them being in the size range 0.5 to 3 pm.

Fig. 8. Typical micrographs after bubbling C H 4 for 2.5 min and homogenization at 640 °C for 18 h (reaction B): (a) A13Ti-AI interface; (b) carbon particulates, and (c) TiC particles.

The carbon particles were of an average size of 43.3 nm with a standard deviation of 10.3 nm.

Figure 11 shows the XRD patterns of a TiC/AI composite and reinforcing phase powder for sample D, after the additional bubbling of methane through it for over 2 h. The patterns show the absence of A13Ti, and TiC presence is clearly established. Similarly, the optical and scanning electron micrographs, Figs. 1 l(a) and 1 l(b) do not show any aluminide needles. The TiC particles are present in the interdendritic region along with FeAI 3. The TiC particles are in general less than

4. D i scus s ion

4.1. Carbon inoculation The experimental liquid and solid state reactions

results provide a significant insight into the formation of carbon and the nucleation and growth of TiC particulates in aluminum melt. From the XRD and electron microscopy results of reaction A, i.e. carbon inoculation, it is clear that all the titanium was not converted to TiC even through 300% of the required amount of methane was infiltrated in the specimen. Carbon appeared as 40 nm amorphous particles. The presence of A13Ti is confirmed by the XRD pattern, Figs. 5(a) and 5(b).

From earlier modeling studies, it is clear that at the methane decomposition reaction temperature, L e. 1150 °C, over 98% of the CH 4 dissociates into carbon and hydrogen. The transmission electron micrographs reveal the presence of carbon particles of the order of 20-50 nm. The above two observations suggest that formation of TiC is not instantaneous even at these

S. Khatri, M. Koczak Formation of TiC composites in molten AI-Ti 159

li N i Alli I ~ 1 ~11 A1

! ~ , Ji I I

i i [ :

I, i ~ !! ! ~ J l ' I L ~ I ' i'

; I: i li II ! i ~ ! , , i ! ! :

It / 4 / 1 1 :,, ~ I i I ! ' ± < i t c~ II i • I t

]

L . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

(a)

2 . 0 ~ Cul,:al +2

Fig. 10. Electron micrographs after bubbling C H 4 for 2.5 mln and argon for 120 min (reaction C): (a) TiC particles; (b) carbon particles.

I: i

(b)

TiC

TiC

~l, TiC

11 i/ c ! L I I [ , Ti i

Fig. 11. X-ray diffraction patterns of samples after bubbling C H 4

for 2.5 min at 0.5 1 min -I and at 0.1 1 min -t for 120 min (reaction D): (a) composite; (b) powder.

elevated temperatures, i.e. 1150 °C. Formation of TiC is determined by a time dependent diffusional process of titanium in the melt to carbon and/or an incubation period which depends on the nucleation of TiC particulates. Thus, it is suggested that upon gas injection, the methane dissociates into hydrogen and carbon immediately. A small part of the carbon goes into the solution in the alloy. The solubility of carbon in molten aluminum, however, is very low (< 0.1 wt.%). There- fore most of the carbon remains dispersed in the alloy as fine amorphous particulates. The presence of amorphous carbon particulates is confirmed via direct TEM observations.

4.2. Liquid and solid state reactions (reactions B and C) The initial short time inoculation of methane in the

melt results in the formation of carbon particles, with no detectable reaction with titanium to form TiC. The dispersed carbon in the melt can subsequently react with either aluminum or titanium to form their respec- tive carbides. This is confirmed by the fact that carbides are seen to form if the reaction is further allowed to occur by bubbling only argon through the melt. Although not as rapid as in the liquid state, the

carbide formation reaction could also be carried out in the solid state, via solute diffusion. The titanium from the AI3Ti needles diffuses into the aluminum matrix and reacts with carbon to form carbides. The carbide nucleation in Fig. 8(a) are identified as brighter particles and are indicated by the arrows. Figure 8(b) shows a transmission electron micrograph of sample B after solutionizing treatment.

It can be seen that the carbon particles coarsen upon solutionizing without noticeable carbide formation and A13Ti dissolution. The coarsening of carbon particles occurs by solid state diffusion. Since solid state diffusion data of carbon in aluminum is not available, as an estimate, the diffusivity of carbon was taken as self- diffusion of aluminum. The self-diffusion coefficient for aluminum at 640 °C is about 10 -12 s I [15]. This can serve as a very approximate upper limit for the diffusion of carbon in aluminum, i.e. if carbon is present as a substitutional element in the lattice. If carbon is present at the interstitial sites, at these very high homologous temperatures, the diffusivity would be expected to be lower, Le. at temperatures close to the melting point of aluminum, the high vacancy con- centration results in a high diffusivity of the atoms at the substitutional sites. Assuming the diffusion of


carbon to be a rate controlling step, the coarsening rates can be calculated from the LSW theory. The corresponding equation is [16, 17]

(OoVmCa( °° )] dt3 - d°3 = ~ ~ 2RT ]

where D, a, V m and Ca(~o) are the diffusion of the controlling species (assumed to be C), the interfacial energy between the precipitate and the matrix, the molar volume of carbon and the equilibrium solubility of the alloying element, respectively. The interfacial energy between carbon and the molten alloy was assumed to be 0.5 J m -2. The equilibrium solubility of C in aluminum at 640 °C was taken as 0.0005mo1.% [18]. Based on these assumptions, the coarsening rate of carbon in aluminum at 640 °C was calculated to be 3.65 x 10 -8 m s -~ which is two orders of magnitude faster than the experimentally observed coarsening rate of 3.31 x 10- 10 m s-1. The higher theoretical coarsening rates can be attributed to unavailability of a variety of data, e.g. accurate diffusion coefficients of carbon in aluminum, solubility of carbon in aluminum and the interfacial energy between carbon particulates and aluminum matrix. In addition, the carbon particles react with titanium to form TiC. This can lead to two results: (a) two competing processes, where carbon particles grow via coarsening but are shrinking due to conversion to TiC, and (b) a TiC coating is formed on the carbon particle surface, which acts as a diffusion barrier resulting in a slower coarsening. Nevertheless, a significantly slower coarsening is observed compared with the best available calculated value.

4.3. Growth and nucleation of carbon and TiC Observation of the AI-Ti-C phase diagram

provides significant insight into the eventual formation of TiC and the metastable precursor nucleating phases. Several reaction paths are possible in the liquid and solid state. The existence of a third phase of the type Ti3AIC also suggests that before the formation of TiC, various intermediate phases like AI4C 3 and Ti3AIC may be formed, Fig. 12 [19]. The phase diagram outlines two reaction sequences for the solid and the liquid state reactions, marked B and C respectively. In the solid state reaction path B, it is postulated that the AI3Ti needles react with carbon to form Ti3A1C. If sufficient time and carbon is available for the reaction to proceed, the Ti3AIC takes part in further reaction to form the thermodynamically stable TiC. In the liquid state reaction C, owing to the absence of the AI3Ti needles, a different reaction path is proposed. Although TiC is thermodynamically the most favorable product at the reaction temperatures, it must be remembered that the activity of aluminum in the

melt is about 0.99 as compared with about 0.01 for titanium (assuming that activity is proportional to the mole fraction of the component in the solution). Hence, kineticaUy it may be more favorable for the formation of aluminum and/or aluminum-titanium carbide intermediate products before the eventual formation of titanium carbide. A reaction sequence which may be possible is suggested below

4AI + 3C = A14C 3

ml4C 3 + 12Ti + C = 4Ti3AIC

Ti3A1C + 2C = 3TiC + AI

Figure 13 shows the gradual conversion of an A1-Ti alloy and carbon to an A1-TiC composite in the form of a diffusion couple. The fraction of titanium and carbon in the reaction product increases from about 1% in the alloy to 50 at.% in the most stable thermodynamic product. However, the fraction of aluminum decreases from 99% to 0% in the reaction product.

Based on the above observations, a flow chart has been constructed to outline various possible routes for reactive synthesis of metal matrix composites, Fig. 14. The most well known route is the vapor-solid-liquid

C

H = TI2AIC

oo/ k

20 TIAI3 40 TIAI 60 TI3AI 80 a-Tlss I~'Ti~s

at%Tl Fig. 12. A I - T i - C temary isotherm showing two possible reaction paths for formation of TiC in solid and liquid state reactions.

AI-Ti AI4C3 Ti3AIC TiC C

Fig. 13. A diffusion couple showing conversion of titanium and carbon in an AI -T i -C composite alloy to TiC.

S. Khatri, M. Koczak / Formation of TiC composites in molten Al- Ti 161

Uz-]

t>~// ~ > 6o rain.

{ SolidState I [ LiquidState I Reaction Reaction

Fig. 14. Flowchart outlining various reactions.

process in which the gaseous phase, i.e. with sufficient time at temperature, C H 4 is allowed to decompose in the molten alloy, and the transition element is allowed to react with nascent carbon to form the titanium carbide. It is also possible to let the gas decompose in the liquid and solidify the alloy before the carbon has reacted; thus entrapping the carbon particle. There- after the alloy could either undergo a solid state reaction or liquid state reaction to form carbides. A variation on the solid state reaction could involve rapid solidification of the alloy with dispersed carbon, followed by solid state heat treatment to obtain strengthening via a fine dispersion of intermetallics and carbides during the partial homogenization process. The possibility also exists for graded graphite/carbide structures generated by partial graphite conversion. The final processing route involves a reaction between fine particles of carbon injected externally and liquid A1-Ti alloy (liquid-solid process). This process normally would result in carbides in the range of 1-3 ktm.

5. Conclusions

( 1 ) The liquid-solid reaction between an aluminum- titanium alloy and carbon to form TiC is not instantaneous. This is borne out by the observation of 40 nm spherical, amorphous carbon particles that after a short period TiC is not formed even though excess carbon is available in the alloy following short term elevated temperature exposure.

(2) The carbon particles formed in the liquid alloy via methane decomposition can be trapped in the alloy upon solidification. This carbon containing composite alloy can be further treated in the solid or liquid state to partially or completely form carbides. The existence of amorphous carbon particles was confirmed via TEM observations and the particles were observed to coarsen very slowly upon solid state heat treatment,

possibly due to the formation of a fine carbide reaction coating.

(3) In the solid state and liquid state homogenization process, it is proposed that the reaction does not proceed rapidly to its thermodynamic equilibrium mixture of TiC and aluminum. Before thermodynamic equilibrium is achieved, metastable carbides such as Ti3AIC and A14C 3 may form due to a higher activity of aluminum in the alloy and rapid kinetics of formation. The evidence for these intermetallics is not conclusive, nevertheless solid state models suggest intermediate phase formation during low temperature reactions.

(4) Observations of carbon conversion in the solid state reactions suggests the possibilities of forming carbide coated aluminides and possibly carbide coated carbon particles, while liquid state reactions result in conversion of AI3Ti to TiC. As a result, the gas-liquid reaction process creates the possibility of the formation of graphite, carbide and mixed reinforcements.

Acknowledgments

The authors gratefully acknowledge the Office of Naval Research and Alcoa supporting their metal matrix composite synthesis work at the Department of Materials Engineering, Drexel University.

References

1 M. J. Koczak et al., Metal Matrix Composites for Ground Vehicle, Aerospace and Industrial Applications, Butter-

worths, 1992, in press. 2 T.W. Clyne and M. G. Bader, in W. C. Harrigan, J. Strife and

A. K. Dhingra (eds.), Proc. ICCM V, TMS-AIME, 1985, p. 755.

3 C. G. Levi, C. J. Abbaschian and R. Mehrabian, Metall. Trans. A, 9A (1978) 697.

4 A. P. Divecha, S. G. Fishman and S. D. Karmarker, J. Met., 33(1981) 12.

5 M.S. Newkirk et al., J. Mater. Sci., 1 ( 1 ) ( 1986) 81-89. 6 M.J. Koczak and K. S. Kumar, In-situ process for producing

a composite containing refractory material, US Patent 4,808,372(1989).

7 E. Dannenberg, in Martin Grayson (ed.), Encyclopedia of Composite Materials and Components, Wiley, New York, 1983, p. 230.

8 P. Bachmann, D. Leers and H. Lydtin, Diamond and Related Materials, 1 (1991) 1.

9 0 . Kubaschewski and E. Evans, Metallurgical Thermo- chemistry, Metal Physics and Physical Metallurgy, Vol. 1, Pergamon, Oxford, 1955.

10 P. Sahoo and M. J. Koczak, Mater. Sci. Eng., A131 (1991) 69. 11 R. Swalin, Thermodynamics of Solids, Wiley, New York,

1972. 12 P. Sahoo and M. J. Koczak, Mater. Sci. Eng., A 144 ( 1991 ) 37. 13 P. Sahoo and M. J. Koczak, Advanced metal and ceramic

matrix composites, in TMS Conf Proc., Feb. 1990.

162 S. Khatri, M. Koczak / Formation of TiC composites in molten Al-Ti

14 D. Gaskell, Introduction to Metallurgical Thermodynamics, McGraw Hill, 1980.

15 J. E. Hatch, Aluminum: Properties and Physical Metallurgy, American Society of Metals, 1984, p. 140.

16 R. Doherty and J. Martins, Stability of Microstructures in Metallic Systems, Cambridge Solid State Science Series, Cambridge, 1980.

17 P. Haasen, Physical Metallurgy, Cambridge University Press, Cambridge, 1986.

18 L. E Mondolfo, Aluminum Alloys: Structure and Properties, Butterworths, Guildford, 1976.

19 J. Schuster, H. Nowotny and C. Vaccaro, Ternary phases in the system Ti-A1-C, J. SolidState Chem., 32 (1980) 213.

Formation of TiC in in situ processed composites via solid-gas, solid-liquid and liquid-gas reaction...

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Transcript of Formation of TiC in in situ processed composites via solid-gas, solid-liquid and liquid-gas reaction...