Interaction between titanium and carbon at moderate temperatures

Journal of Alloys and Compounds 368 (2004) 116–122

Interaction between titanium and carbon at moderate temperatures C. Arvieu a,b,∗ , J.P. Manaud b , J.M. Quenisset a,b a

Institut de Chimie de la Matière Condensée de Bordeaux, CNRS 87, Avenue du Docteur Schweitzer, 33608 Pessac Cedex, France b Laboratoire de Génie Mécanique, IUT, Université Bordeaux 1, 33405 Talence Cedex, France Received 28 February 2003; received in revised form 4 August 2003; accepted 7 August 2003

Abstract The examination of numerous studies devoted to titanium–carbon interactions shows that most domains of investigation are related to high temperatures (above 800 ◦ C), for which the diffusion of carbon within titanium is mainly controlled by carbon diffusion through titanium carbide. In order to understand the phenomena of Ti–C interaction at moderate temperatures, Ti-coated carbon disks prepared by PVD were heat treated under vacuum at temperatures in the range of 400–750 ◦ C, and analysed mainly by RBS. At low temperatures and for short isothermal exposure, the formation of a titanium carbide interphase is discontinuous and controlled by the diffusivity of C in Ti. As soon as a continuous titanium carbide layer is formed between C and Ti, the interaction is controlled by the kinetics of diffusion of C in titanium carbide. © 2003 Elsevier B.V. All rights reserved. Keywords: Titanium–carbon interaction; Diffusion kinetics; Interface; Titanium carbide

1. Introduction During the last five decades, the interaction between titanium and carbon has given rise to many studies with various objectives. For instance, the determination of the diffusion rates of carbon in both alpha and beta titanium over a wide range of temperatures was required in order to control metallurgical processes of heat treatment, homogenisation and aging of titanium alloys [1–3]. The kinetics of titanium carbide formation was studied through the assessment of the chemical diffusivity of carbon in TiC [4–9]. Other studies were concerned with the production of high-density substoichiometric titanium carbide by liquid sintering or hot isostatic pressing of titanium carbide and titanium powders [10,11]. Also, impeding the C–Ti interaction during processing and high-temperature use of titanium matrix composites reinforced by carbon fibres or carbon-coated SiC filaments has been of great importance [12–14]. All these studies have led to significant differences in results, particularly for diffusivity and related thermal activation energy. However, significant differences in experimental conditions and investigation methods are able to explain the apparent discrepancy in the obtained diffusion param∗

Corresponding author. E-mail address: [email protected] (C. Arvieu).

0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.08.051

eter values. For instance, the coupled materials are either amorphous, polycrystalline or single crystals, and the compositions of titanium carbide are significantly different. The domains of study concern high (600–1400 ◦ C) or very high temperatures (1500–2700 ◦ C). The methods of investigation are based on the use of tracers or on the determination of interphase growth rates, and lead to different kinds of characteristics. As a consequence, it is difficult to extrapolate the published diffusion parameters for C–Ti interaction conditions, which are significantly different from those related to the previously reported results. This is particularly the case in the present contribution, which aims at evaluating the significance of the Ti–C interaction in a domain of moderate temperatures ranging between 400 and 700 ◦ C. As a matter of fact, carbon-reinforced titanium matrix composites (TMC) are able to withstand such temperature during their elaboration by powder metallurgy (650–700 ◦ C), during their use at 450 ◦ C or higher temperatures up to 700 ◦ C for a short time [15]. Plane Ti–C couples were prepared by physical vapour deposition (PVD), heat treated at various temperatures and analysed mainly by Rutherford back scattering (RBS), in order to design the interfacial zone required for preventing excessive formation of brittle interfacial layers in thin carbon fibre reinforced TMC. Experimental conditions were

C. Arvieu et al. / Journal of Alloys and Compounds 368 (2004) 116–122

chosen in order to investigate an interfacial zone some hundreds of nanometres thick. Given the discussion and comparison of the derived diffusion parameters, microscopic observations and other results obtained on composites and reported elsewhere, a fibre protection may be proposed and a maximum temperature for composite use may be estimated [16].

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In the chamber, the samples were placed in a titanium crucible located at the very centre of an RF heating furnace. Before heating at temperatures of isothermal exposure, the chamber was evacuated below 5 × 10−5 Pa to remove gas from it and from the samples heated at 250 ◦ C for 20 min. Next, samples were submitted to isothermal exposure under an absolute pressure of 4 × 104 Pa of purified Argon. The titanium crucible surrounding the samples and used as a susceptor of RF heating was expected to play the role of getter for residual impurities, since the susceptor temperature is higher than that of the sample and both of them are not in equilibrium with the small oxygen partial pressure in the chamber. Isothermal exposure time was fixed at 30 min, but some samples were heat treated at 500 ◦ C for 30, 90 and 180 min in order to determine the growth kinetics of the interfacial zone versus time.

2. Experimental 2.1. Sample preparation Ti–C couples were prepared by depositing very thin titanium coatings (715 nm thick) on planar vitreous carbon substrates from Carbone Lorraine Company (14 mm in diameter and 2 mm thick, impurity ratio less than 20 ppm). After polishing one face (Ra = 150 Å), carbon substrates were coated by RF magnetron sputtering commercial system in a chamber backed below 3 × 10−5 Pa, using a turbomolecular pump and a liquid nitrogen trap. Most impurities of the argon used as glow discharge gas were eliminated through a purifier, and the titanium target purity from CERAC matched 99.995% with a residual carbon content of 12 ppm. Prior to each deposition, the polished surface of the carbon substrate was etched for 5 min under a power density discharge of 0.25 W cm−2 . During the depositions, substrates were maintained at room temperature and other coating parameters were as follow: 2 W cm−2 power density, 0.5 Pa of argon, 70 mm distance between target and substrate and deposition rate of 11.1 nm s−1 . The sputtering rate was determined through measurement of a step obtained by a ‘lift-off’ method. Then, resulting Ti-coated carbon substrates were annealed in a cold-wall chamber at various temperatures: 400, 450, 500, 550, 600 and 750 ◦ C.

2.2. Analyses Microstructures of titanium deposits were characterized before and after isothermal exposure by X-ray diffraction (XRD) at a low incident angle (1◦ < θ < 3◦ ). The deposits consisted of small crystallites of titanium. After heat treatment, the characteristic peaks of titanium carbide appear on spectra as illustrated in Fig. 1. Considering the Ti–C interface, X-ray photoelectron spectroscopy (XPS) analyses (Al monochromatized source, spot size 150 ␮m) were performed on some samples before isothermal exposure in order to check the absence of carbide at the Ti–C interface. As shown in Fig. 2, XPS spectra reveal only a small quantity of titanium carbide (1 at.% estimated), confirming that neither PVD deposition nor XPS sputtering induces any significant carbide component at the interface. After these preliminary analyses, the chemical composition of the as-deposited and annealed films was determined

Ti TiC C

counts

30' at 750˚C

30' at 600˚C

30' at 500˚C

20˚C

15

20

25

30

35

40

45

50

55

60

65

70

75

80

2θ (˚) Fig. 1. XRD spectra obtained at low angle of incidence on titanium coating.

85

90

118

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Fig. 2. XPS depth profile and C1S spectra obtained at different stages of sputtering through the Ti film and the Ti–C interfacial zone.

40

Ti peak

Ti/vit.C 25C Ti/vit.C 400C 30min Ti/vit.C 550C 30min Ti/Vit.C 750C 30min

Normalized Yield

30

20

wr for 550˚C wi 10

wt C substrate

0 100

200

300

400

500

600

700

800

Channel Fig. 3. RBS spectra obtained through the Ti-coated carbon substrate before and after annealing at various temperatures for 30 min.

C. Arvieu et al. / Journal of Alloys and Compounds 368 (2004) 116–122

by 2 MeV ␣ particles RBS (sample tilt angle: 0◦ ; exit angle: 20◦ ) (Fig. 3). The spectrum related to the as-deposited sample shows a titanium peak alone (continuous line). For annealed samples, the titanium peak appears to be altered. In addition, the spectra show a carbon peak due to the carbide formation. For the sample annealed at 550 ◦ C (cross dashed line), the titanium peak can be broken up into a front one, corresponding to a pure titanium layer, and a rear lower one, corresponding to the titanium in TiC growing interphase. The deviation in density related to the presence of carbon gives rise to peak broadening. A RUMP simulation [17] shows this peak fits a computed peak related to a pure titanium carbide 715 nm thick layer, the density of which is 4.97 g/cm3 . That means the Ti to titanium carbide transformation is paired with a modification of density but there is no modification in thickness. RUMP simulation consists in computing a spectrum of a virtual layer, knowing the scattering cross section, the stopping power of its constitutive atoms [18] and its physical density. Under these conditions, the thickness set for this spectrum is the physical thickness [17]. When the theoretical spectrum fits the data we can assume the virtual sample is imaging the real one.

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to 87.5% ([MTi /ρTi ]/[(MTi + MC )/ρTiC ], with M the molar mass and ρ the density, and considering the titanium carbide as quasi-stoichiometric, which remains questionable [19]). The thickness x was also obtained by comparing the spectra width (wi − wr ) related to the thickness of Ti deposits, partly transformed into titanium carbide, with wt of a Ti deposit completely transformed at 750 ◦ C for half an hour into a TiC of thickness yt = 720 nm measured by profilometer. Thus, x = yt (wi − wr )/wt . The mean values x of titanium carbide thickness are considered as apparent values, since it is not certain that a continuous titanium carbide layer forms as soon as the Ti–C interaction begins. In general, two measurements were performed with each method. The values of titanium carbide thickness derived from the samples heat treated at 500 ◦ C for different durations are reported in Fig. 4. As shown in Fig. 4, the average thickness x deduced from the previously defined procedure for various durations of isothermal exposure is not easily approached by the following relation, which is commonly adopted for depicting the growth of a continuous interphase by diffusion of interactive elements: x = [kt]1/2

3. Exploitation and discussion

(1)

where k is the chemical interaction coefficient and t the duration of isothermal exposure. The linear correlation of the experimental results corresponding to the parabolic relation (1) is not obvious for short isothermal exposure. This noticeable discrepancy can be explained by three remarks: (i) The uncertainty related to the measurement of TiC thickness is particularly significant for the weakest Ti–C interactions. (ii) At the beginning of this interaction at moderate temperature, the mechanism of diffusion might be different from that involved in the growth of a continuous titanium carbide interphase. (iii) The titanium carbide composition is presumably non-homogeneous

3.1. Determination of the diffusion parameters The thickness x of the titanium carbide interphase was derived from the RBS spectra (Fig. 3) by comparing the spectra width wr , related to the residual thickness yr , of Ti with wi of the initial Ti deposit, thickness yi = 715 nm measured by profilometer. Thus, yr = yi wr /wi and x = η(yi –yr ) with η = VTiC /VTi = 1.14, where VTiC is the volume of carbide used by transformation of a Ti volume VTi . Indeed, it is considered that the ratio between the transformed Ti thickness and the resulting titanium carbide thickness is equal

TiC thickness (nm)

250

200

150

100

50

0 0

20

40

60

t

1/2

80

100

(s)

Fig. 4. Average titanium carbide layer thickness vs. time at 500 ◦ C.

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Table 1 Titanium carbide thickness and related chemical interaction coefficients k for 1800 s of isothermal exposures at various temperatures Temperature (◦ C)

TiC thickness (nm)

k (10−14 cm2 s−1 )

400 400 400 400 450 450 450 450 500 500 500 500 550 550 550 550 600 600 600 600

30 15 35 16 47 30 53 30 123 100 130 112 261 200 260 219 500 431 472 425

0.5 0.125 0.68 0.14 1.23 0.5 1.56 0.5 8.4 5.56 9.39 6.97 37.8 22.2 37.6 26.6 139 103 124 100

through its thickness and the rate of non-stoichiometry might be dependent on the duration of Ti–C interaction. Consequently, an assessment of coefficient k was obtained from relation (1) by considering that each mean value of apparent TiC thickness (related to a particular temperature and a duration of heat treatment) corresponds to one predominant interaction mechanism, which can be depicted by relation (1). The resulting experimental values of k, related to 30 min heat treatment at various temperatures, are reported in Table 1. Reporting these values (Fig. 5) with semi-logarithmic coordinates allows the kinetics of the titanium carbide forma-25 -26 -27

Lnk (cm2.s-1)

-28 -29

tion to be depicted by the following relation:
 Q k = k0 exp − RT

(2)

where the frequency coefficient k0 is 1.48 × 10−3 cm2 s−1 and the thermal activation energy is Q = 153 ± 18 kJ mol−1 between 400 and 600 ◦ C. Although domains of temperature investigation are quite different, the comparison between the present experimental results (Fig. 6) and those previously reported by other investigators allows the following remarks to be deduced: • The thermal activation energy related to the interaction between carbon and titanium at rather low temperature is significantly lower than the values obtained for the diffusion of carbon in titanium carbide (250–400 kJ mol−1 ) [4–14]. Otherwise, the Q value is close to those reported for the diffusion of carbon in titanium (130–180 kJ mol−1 ) [1–3]. • An attempt of extrapolation for lower temperatures of the previously reported results (‘our results’ in Fig. 6) shows that the chemical interaction coefficients obtained in the present study are lower than those related to carbon diffusion in titanium, but higher than those related to the diffusion of carbon in titanium carbide. • An increase in temperature of isothermal exposure between 400 and 600 ◦ C tends to draw the related chemical interaction coefficients closer to the those determined in the carbide. • A transition in the mechanism of carbon–titanium interaction at moderate temperatures from a diffusion of carbon in titanium towards a diffusion of carbon in titanium carbide is expected to induce an apparently lower thermal activation energy, compared with that of carbon diffusion in titanium. This may justify the noticeable difference between the thermal activation energy derived from this study (153 kJ mol−1 ) and those reported elsewhere for the diffusion of carbon in titanium (182 kJ mol−1 ) [1]. It also explains the diference between the experimental results obtained for short heat treatments and the linear correlation associated with a parabolic growth of interphase. 3.2. Interaction mechanisms

-30 -31 -32 -33 -34 -35 1,10E-03

1,20E-03

1,30E-03

1,40E-03

1,50E-03

-1

1/T(K ) Fig. 5. Arrhenius plots of Ti–C chemical interaction coefficients obtained at temperatures ranging between 400 and 600 ◦ C.

The previous comments concerning the obtained experimental results allow a model of Ti–C interaction mechanism at moderate temperature to be proposed. As schematically illustrated in Fig. 7a, during the first minutes of interaction, the carbon of the substrate preferably diffuses along titanium grain boundaries with a presumable high kinetic of diffusion compared with that of carbon within titanium grains. As a matter of fact the related values of the apparent TiC thickness is much higher than that expected from the classical Eq. (1) (Fig. 4). Consequently, a net of C forms in the vicinity of the Ti–C interface, within a Ti layer of about 10 nm thickness.

C. Arvieu et al. / Journal of Alloys and Compounds 368 (2004) 116–122

121

0

-5

-10

2

LnD (cm /s)

[3 ]

in Ti

[1 ]

[6 ]

-15 [9 ] [7 ]

-20

[3 ]

[9 ]

[11]

our results

[4 ]

[7 ]

[5 ] [13]

-25

-30

in TiC -35 3, E -04

4, E -04

5, E -04

6, E -04

7, E -04

8, E -04

9, E -04

1, E -03

1, E -03

1/ T ( /K ) Fig. 6. Arrhenius plots of carbon diffusivities in titanium and in titanium carbide as obtained by various investigators.

Then, as the carbon concentration increases at Ti grain boundaries, titanium carbide crystallites grow into the intergranular spaces at the expense of the adjacent Ti grains. At this stage of the Ti–C interaction, there is no interposition of a continuous interphase at the Ti–C interface (Fig. 7b). The transfer of carbon from the C substrate towards the Ti film preferably occurs at the titanium carbide crystallite–Ti grain boundaries and in the titanium adjacent to the C–Ti interface rather than through the titanium carbide crystallites. After this initial stage, which will be particularly illustrated in the next cross section (Fig. 8), the titanium carbide crystallites are able to grow in directions parallel to the C–Ti interface up to the formation of a continuous titanium carbide interphase (Fig. 7c and d). Finally, carbon diffuses from the substrate through the

(a)

titanium carbide interphase, towards the residual titanium film. The previously described mechanism leads to consider that, during the first period of an isothermal exposure at moderate temperature, the C–Ti interaction occurs by diffusion of C within Ti not with the formation of any titanium carbide continuous layer but with the formation of titanium carbide crystallites. Next, the C–Ti interaction mechanism is controlled by the diffusion of C through the titanium carbide interphase. Thus, it is not surprising to obtain diffusion characteristics close to those which can be expected from the diffusion of carbon within titanium for the lowest temperatures and durations of isothermal exposure and close to those related to the diffusion of C through

(b)

C

(c)

Ti

(d)

Titanium carbide

The arrows show the diffusion way of carbon Fig. 7. Schematic representation of the C–Ti mechanism of interaction during isothermal exposure at moderate temperatures.

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C. Arvieu et al. / Journal of Alloys and Compounds 368 (2004) 116–122

aries. After growing perpendicularly to the Ti–C interface, the carbide spreads parallel to the interface until it forms a continuous layer through which the carbon has to diffuse later in order to extend the titanium carbide interphase layer. These different periods in the mechanism of Ti–C interaction explain that (i) the lowest isothermal exposure temperatures preferentially reveal diffusion characteristics related to those reported elsewhere for C diffusion within Ti and (ii) the highest temperatures correspond to diffusion characteristics related to C diffusion within titanium carbide. Acknowledgements Fig. 8. Electron micrograph of titanium carbide crystallites formed at the fibre–matrix interface in carbon-fibre-reinforced titanium matrix composites [15].

titanium carbide for higher temperatures and annealing times (Fig. 6). 3.3. Design of fibre–matrix interfacial zone in carbon-fibre-reinforced titanium matrix composites The C–Ti interaction mechanism presented in the previous section leads to propose the interposition of an artificial continuous titanium carbide layer some tens of nanometres thick between the C substrate and Ti film. Such interfacial tailoring is expected to limit C transfer into Ti as soon as the Ti–C couple is thermally exposed and thus, to prevent the formation of mechanically unfavourable titanium carbide crystallites as illustrated in Fig. 8 [15]. Indeed, the size of the crystallites leading to an apparent titanium carbide layer some tens of nanometres thick (Table 1) is about 500 nm long, which is critical from the mechanical point of view. As a consequence, in case of elaboration at high temperature (600–700 ◦ C) for a short time (30 min) and use of carbon-fibre-reinforced titanium matrix composites (TMC) at moderate temperatures, coating the C fibres with a very thin titanium carbide layer prior to TMC processing prevents the initial part of the Ti–C interaction mechanism from taking place, that is, it impedes the uncontrolled fibre–matrix (FM) interaction leading to titanium carbide crystallites shown in Fig. 8.

4. Conclusion RBS analyses and observations which were made on Ti-coated carbon disks prepared by PVD and heat treated at moderate temperatures have shown that the Ti–C interaction begins by a C diffusion within titanium, giving rise to the formation of titanium carbide crystallites at Ti grain bound-

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