sulphidation resistance of Ti and Ti6Al4V alloy by HfN coating

,(

ELSEVIER

Materials Science and Engineering A205 (1996) 199-208

A

Enhancement of oxidation/sulphidation resistance of Ti and Ti-6A1-4V alloy by HfN coating H . L . D u a'*, P . K . D a t t a a, D . B . L e w i s b, J.S. B u r n e l l - G r a y a ~SurJace Engineering Research Group, University of Northumbria at Newcastle, Newcastle upon Tyne NE1 8ST, UK bMaterials Research Institute, Sheffield Hallam University, Sheffield $1 l WB, UK Received 28 February 1995; in revised form 31 May 1995

Abstract

An attempt has been made to enhance the oxidation/sulphidation resistance of Ti and Ti-6A1-4V alloy by the application of an HfN coating produced using Physical Vapour Deposition (PVD). The coated specimens were exposed to an atmosphere comprising high sulphur and low oxygen potentials (PS2 ~ 10- 1 Pa and pO2 ~ 10 ~s Pa) at 750 °C for periods up to 240 h. The corrosion kinetics were obtained by a discontinuous gravimetric method. Linear and linear-parabolic kinetics were recorded for Ti and for Ti-6A1-4V alloy respectively. The HfN-coated Ti and Ti-6A1 4V specimens showed enhanced corrosion resistance. The scale formed on the uncoated Ti specimen consisted of a double layer of TiO2 with an underlying TiS2 film on the substrate. The exposed Ti-6A1-4V alloy also contained a double layer of TiO2 but with A1203 precipitated in the external portion of the outer layer of the TiO2, whilst a layer of AleS3 and TiS2 with vanadium sulphide developed beneath the inner layer of TiO2. Application of the HfN coating suppressed the formation of the duplex oxide scales. However, extensive cracking developed either parallel or perpendicular to the sample surface, and severely compromised the protectiveness of the coating.

Keywords: Oxidation; Sulphidation; Titanium; Coatings

I. Introduction Ti, Ti-based alloys and metal matrix composites based on Ti and Ti-alloys are highly attractive materials in aero industries and in other advanced industrial applications. However, a major problem in the use of these high temperature materials stems from their lack of surface stability particularly in aggressive environments. These materials possess the necessary mechanical properties at elevated temperatures; but prolonged exposure brings about environmental degradation involving scaling. It has been known [1-6] that the performance of Ti and Ti-alloys at high temperatures is controlled not only by the scaling reaction, but also by the dissolution and diffusion of interstitials, especially oxygen, hydrogen and nitrogen into these materials with concomitant reduction in fatigue life. More recently, Du and Datta et al. [7,8] investigated the high temperature corrosion behaviour of Ti and T i - 6 A 1 - 4 V in air, H2/H20 and H2/H20/H2S environments and * Corresponding author. 0921-5093/96/$15.00 © 1 9 9 6 - Elsevier Science S.A. All rights reserved S S D I 0921-5093(95)09888-7

reported rapid corrosion kinetics and complex scaling processes showing significant dependence on the nature of environments. There has been much research activity seeking to develop oxidation resistant Ti-alloys. Additions of Cr, e.g. at 5 or 10 wt.% level, were observed to increase the weight gains, although interestingly both these alloys exhibited reduced scale thickness. The overall increase in weight was associated with greater amounts of dissolved oxygen [9]. Such an effect was related to the microstructure of the T i - C r alloys, Cr being a fl-stabilising element [10]. However, it is reported that at high concentrations ( > 10 wt.%) Cr has a beneficial effect on the oxidation resistance of titanium, particularly at low temperatures [11]. It is well established that additions of A1 to conventional Fe, Ni and Co-based alloys produce a stable and protective ~ - A I 2 0 3 scale. However, in Ti-A1 alloys, Ti and A1 form oxides of very similar stability [12,13], but only A1 forms a slowly growing oxide (A1203), whilst all titanium oxides have relatively high growth rates [1]. As a result, although in Ti-A1 alloy with ~ 50 at.% AI an A1203 scale does form in

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oxidising environments between 700 and 900 °C, this scale does not display expected protectivity [14]. Welsch and Kahveci [15] observed the formation of mixed oxide scales on Ti-A1 alloys when oxidised in dry oxygen (pO2 ~ 105 Pa) between 600 and 1100 °C. These scales contained increasing amounts of A1203, interspersed heterogeneously with TiO 2 and included a predominantly-rich A1203 sublayer with increasing aluminium contents. 7-TiA1 also formed alumina scales below ~ 1000 °C in oxygen but developed mixed titanium and aluminium oxide scales at higher temperatures [16]. The formation of a protective A1203 scale on A13Ti in pure oxygen in a temperature range 800-1000 °C has also been reported [17]. Datta [18] pointed out that although the new high temperature Ti alloys, IMI829 and IMI834, have been developed with temperature capabilities up to 650 °C, they lack the microstructural and thermal stability required to withstand exposure at this temperature for long durations. There is, therefore, a need to confront the question of how to protect such alloys against environmental degradation at these elevated temperatures ( > 650 °C). Here the use of coatings can offer significant benefits. A large number of ceramic materials, such as refractory borides, carbides, silicides, nitrides and oxides, having high melting points and low dissociation partial pressures [19], offer attractive possibilities for their use not only in the fabrication of high temperature structural composites but also in the development of high temperature protective coatings. The main objective of this paper is to investigate the effects of HfN coating on the durability of Ti and Ti-6AI-4V in an atmosphere of high sulphur ( P S 2 ~ 1 0 - 1 Pa) and low oxygen (PO2 ~ 10 t8 Pa) activities at 750 °C.

2. Experimental procedures The substrate materials, commercial Ti and Ti-6A14V (IM1 318), employed in this study were supplied by IMI Titanium Ltd in rod form. Coupons of 12.5 mm diameter and 1.5 mm thickness were machined from these rods. A 1 mm diameter hole was drilled near the edge of each sample to facilitate suspension in the work tube. Prior to coating, all coupons were polished using SiC papers up to 1200 grit followed by ultrasonic cleaning and degreasing in acetone for 30 min. All HfN coatings ( ~ 10 /Lm thick) used in this programme were deposited in a Multi-Arc PVD arc evaporation unit [20,21], where a vacuum arc was employed for the generation of both the coating flux and ionisation necessary for the deposition. The deposition parameters used are listed in Table 1. Fig. 1 shows a typical XRD spectrum for as-received HfN-coated Ti and clearly

reveals the presence of the HfN coating. A similar spectrum was obtained for the HfN-coated Ti-6A14V. The HfN-coated and uncoated specimens were ultrasonically cleaned again in acetone for 30 min prior to exposure. The oxidation/sulphidation kinetics of the HfNcoated and uncoated Ti and Ti-6A1-4V were determined using a discontinuous gravimetric method. A pre-mixed H2/HzO/H2S gas mixture was designed to yield an atmosphere ofpS2 "~ 10-~ Pa andpO2 ~ 10 as Pa at the reaction temperature of 750 °C. Details of the experimental rig have been described previously in Refs. [22-24]. The corrosion products generated on both HfN-coated and uncoated specimens were characterised by means of scanning electron microscopy (SEM), energy-dispersive analysis by X-ray (EDX) and X-ray diffraction (XRD).

3. Experimental results 3. I. Corrosion kinetics Weight gain data plotted against exposure time are given in Fig. 2 for the HfN-coated and uncoated pure titanium and Ti-6A1-4V alloy exposed in the test environment (pS2 ~ 10 l Pa and pO2"~ 10-18 Pa) at 750 °C for up to 240 h. It is apparent that the oxidation/sulphidation behaviour of the uncoated titanium followed a linear rate law with a linear rate constant of 5.7 x 10- 8 g c m - 2 s - 1, whilst that of the Ti-6A1-4V obeyed linear-parabolic kinetics with the corrosion rate decreasing with exposure time. The oxidation/sulphidation behaviour of the HfN-coated Ti and Ti-6AI-4V exhibited linear kinetics with linear rate constants of 1.2 x 10 -s g cm -2 s -1 and 1.0 x 10 s g cm 2 s - l , respectively. These similar values in linear rate constants for both HfN-coated Ti and Ti-6A1-4V indicate that the HfN coating remained effective even following 240 h exposure and up to this time the substrates did not substantially influence the corrosion behaviour, in terms of both the weight change and the corrosion pattern. Clearly, the HfN coating did enhance the oxidation/sulphidation resistance of the Ti and Ti6A1 4V during the 240 h exposure periods although linear rate kinetics were recorded. For pure Ti, the linear corrosion rate constant of the HfN-coated samTable 1 Deposition parameters for HfN coating on Ti and Ti 6AI-4V alloy Cathode Voltage No. of Amp Amphour Ar gas Coating material (V) evaporators (A) (Ah) pressure thickness (Pa) (/am) HfN

150

2

50

130

2

~ l0

H.L. D u e t al. / Materials Science and Engineering A205 (1996) 199-208 I

I

I

I

I

I

I

[

I

I

201

I

I

I

0 0 t~ ¢xl ¢,q

JCPDS number for HfN: 33-0592

t..)

0 0

t:5 ....

215 . . . .

310 , , ,

15 ....

410 ....

4~ ....

510 ....

515 I '

'610 . . . .

615 ' ' ,

for as-received

HfN-coated

20 (degrees) F i g . 1. X - r a y

diffraction

pattern

pie was 5-times smaller than that of the uncoated specimen. 3.2. Corrosion products 3.2.1. Uncoated 77 and T i - 6 A l - 4 V When pure titanium and Ti-6A1-4V were exposed to the test environment comprising p S 2 "-~ 10 ] Pa and pO2 ~ 10 -18 Pa at 750 °C, the corrosion products consisted of two subscale layers, which are clearly illustrated in Figs. 3 and 4 respectively. For pure Ti, XRD analysis showed that both the outer layer and the inner layer comprised TiO2 (rutile) and no compositional nor structural differences were evidenced between these two layers. It is believed that the outer layer was formed by the outward diffusion of titanium whilst the inner layer was produced by the inward diffusion of oxygen species. A film of TiS2 was found to exist beneath the rutile scale. In addition, significant quantities of titanium hydride (identified by XRD as Till2) were observed in the substrate (Fig. 3). Obviously, when Ti was exposed to the hydrogen-dominant H2/ H20/H2S environment at elevated temperature, hydrogen readily dissolved into the substrate and precipitated to give acicular hydrides. In the case of Ti-6A1-4V, the external portion of the outer layer was enriched in A1, and a mixture of AI203

Ti.

and TiO2 (recognised by XRD) was generated. However, a pure TiO 2 layer formed beneath this mixed layer. A mixture of sulphides of A1, Ti and V, as shown in Fig. 4, was located between the oxide scale and the substrate; here only A12S3 and TiS2 were recognised by XRD. Significantly, vanadium oxide was found at the interface of the outer layer and inner layer; also, a small amount of vanadium was present in the outer layer of TiO2 as a result of outward diffusion. Unfortunately, it was very difficult to identify the vanadium sulphide by XRD, due perhaps to the low amount present. No titanium hydride was evident in this alloy even after 240 h exposure. 3.2.2. HfN-coated Ti and T i - 6 A I - 4 V SEM morphological examination of the HfN-coated Ti and Ti-6A1-4V after various periods of exposure revealed no distinct differences between non-exposed specimens and samples subjected to short term exposure (e.g. < 5 h). However, HfO2 on the surface of HfN-coated Ti and Ti-6A1-4V specimens was identified by XRD at the early stages of exposure. As the exposure time increased, the number of cracks on the surface of coated specimens (after cooling) was observed to increase. It was noted that, before oxidation/ sulphidation, there were some HfO2 spots on the coating surface as determined by EDX. It is believed

H.L. Duet al. / Materials Science and Engineering A205 (1996) 199 208

202

5° I

Ti Ti-6AI-4V

40

• Ti+HfN X T i - 6 A I - 4 V + HfN

E o

3O

.~ 20 O1

10

0 0

50

1 O0

150

200

250

Time (hours) Fig. 2. Weight gain against exposure time for HfN-coated and uncoated Ti and Ti-6AI 4V exposed in an H2/H20/H2S environment at 750 °C.

that these oxides were produced during the coating deposition process and some still survived as flower-like products even after 72 h exposure (Fig. 5(b)). However, some of these HfO 2 spots were no longer apparent after further exposure, leaving holes on the coating surface,

Fig. 3. Electron image and Digimaps showing typical compositional profiles through the scale on Ti after 72 h exposure in an H2/H20/ H2S environment at 750 °C.

which subsequently became the initiation sites for TiO2, formed due to the outward diffusion of titanium to the surface, as shown in Fig. 5(a). These holes also initiated cracks. Fig. 5 (c,d) displays typical morphological features of the HfN-coated Ti and Ti-6A1-4V alloy at the end of the experiment (240 h). The HfN coatings on both substrates (after cooling to room temperature) showed substantial numbers of cracks after 240 h. XRD results indicated that two phases (TiO2 and HfO2) were present on the surface of all the samples; ,even so a great deal of HfN could still be detected. The cross-sectional analysis of the exposed specimens again displayed similar features for the coated Ti and Ti-6A1-4V. It was found that the cracks were not only present on the surface of the coating, but also existed inside the HfN coating. These cracks grew parallel to the coating surface, which implies that the internal cracks were not initiated from the surface cracks. Figs. 6 and 7 represent SEM and EDX results for the HfN-coated Ti and Ti-6AI-4V alloy following 240 h oxidation/sulphidation, respectively. In each case TiO2 formed on the sample surface in a discontinuous manner. In the HfNcoated Ti, several cracks formed parallel to the coating surface and TiO2 occupied the space within these cracks. Beneath the HfN coating on the titanium specimen, TiS2 (identified by XRD) formed a continuous film. The absence of the formation of TiO2 beneath the HfN coating indicates that there was no substantial inward diffusion of oxygen through the coating al-

H.L. D u e t al. / Materials Science and Engineering A205 (1996) 199 208

203

cracks were initiated during the heating periods and then grew during the corrosion process aided by the corrosion products. It is clear that some cracks did not contain corrosion products, which implies crack formation during the cooling periods. To elucidate the mechanisms responsible for the initiation of the cracks, two supplementary investigations were carried out by SEM. In the first series, cross-sections of HfN-coated samples were examined before exposure, whilst in the second set of tests, the HfNcoated samples were heated in the furnace until the temperature reached 750 °C, at which point the furnace was switched off. The unexposed and exposed samples were then cross-sectioned and examined. In general, the coating was dense, compact and crack-free except for a few white spots on the surface, identified as HfO2, which was generated during the deposition process; but very occasionally through-cracks were observed on the HfN-coated Ti-6A1-4V, as shown in Fig. 9(a). For the samples undergoing the heating period, the cracks developed between the coating and substrate, as revealed in Fig. 9(b,c). It appeared that the cracks on the sample undergoing heating protruded into the substrate, whilst no such protrusions were observed on the unexposed sample. Because of the rare occurrence of the throughcracks on the as-received coated samples, the mechanism described later assumes crack generation at the substrate-coating interface during the heating process. 4. Discussion Fig. 4. Electron image and Digimaps showing typical compositional profiles through the scale on Ti-6AI 4V after 72 h exposure in an H2/H20/H2S environment at 750 °C.

though some TiO2 was identified in the triangular area where the coating slightly receded upwards (Fig. 6). It was also found that a significant amount of Till 2 formed in the Ti substrate, as described earlier for the uncoated Ti. In contrast to the HfN-coated Ti, most cracks within the HfN coating on the Ti-6A1-4V formed at right angles to the coating-substrate interface, as shown in Fig. 7. However, it is also evident that some cracks parallel to coating surface did develop in the HfN coating. Like the HfN-coated Ti, a discontinuous TiO2 layer was also identified on the coating surface. Oxygen species diffused inward and produced TiO2 both in the cracks and beneath the HfN coating. AI2S3 located beneath the HfN coating was formed by the inward transport of sulphur species through the cracks. However, it is not clear whether vanadium played a part in this process. It is, however, important to note that the HfN coating on Ti-6A1-4V underwent significant deformation as shown in Fig. 8 and certain of the cracks were filled with corrosion products which consisted mainly of TiO 2. This indicates that these

Models describing the oxidation/sulphidation of pure titanium and Ti-6A1-4V at 750 °C in an H2/H20/H2S environment have been described in an early paper [7]. It was established that for both materials the outer layer of TiO 2 was formed by the outward diffusion of titanium species, whilst the inner layer of TiO 2 developed by the ingress of oxygen species. Simultaneously, a TiS2 layer on Ti and A12S3 and TiS 2 layers on Ti-6AI-4V formed together with the oxide layers. After the oxide/sulphide scale developed, concentration gradients of the reactants (O, S, Ti, A1, V) were established for both Ti and Ti-6A1-4V. At the oxide-sulphide interface, A12S 3 became unstable and decomposed to release S and A1 which became a reservoir of aluminium. The migration of aluminium from the reservoir resulted in the precipitation of A l 2 0 3 , the thermodynamically favoured product, in the external portion of the outer layer. It is believed that the dissociation of TiS2 similarly occurred, releasing Ti which subsequently migrated outwards and precipitated as TiO 2 in the external portion of the TiO2 outer layer, although this precipitated TiO 2 was not distinguishable from the previously formed TiO 2. The freed sulphur diffused through the sulphide layer towards the substrate and reacted with the alloying elements.

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Materials Science and Engineering A205 (1996) 199-208

(a)

(b)

(c)

(d)

Fig. 5. Scanning electron micrographs showing typical surface morphologiesof HfN-coated Ti and Ti-6Al-4V after exposure in an H2/H20/H2S environment at 750 °C. (a) HfN-coated Ti for 72 h, (b) HfN-coated Ti 6A1 4V for 72 h, (c) HfN-coated Ti for 240 h and (d) HfN-coated Ti-6AI-4V for 240 h. HfN has a very high melting point ( ~ 3387 °C [25]) and low dissociation partial pressure ( ~ 4 . 5 x l0 24 Pa) at 750 °C [19]. Its high stability was expected to provide a high degree of protectiveness for Ti and Ti 6AI 4V in the aggressive environment employed in this investigation. Indeed, the HfN-coated Ti and T i 6A1-4V displayed an enhanced oxidation/sulphidation resistance even up to 240 h. For example, the linear corrosion rate constant for the HfN-coated Ti (1.2 x 10 -8 g cm -2 s - j ) was about 5-times smaller than that of the uncoated Ti specimens (5.7 x 10 8 g cm -2 s-1). The HfN coating acted effectively as a protective diffusion barrier and prevented the substrate elements from migrating outwards and oxygen and sulphur species from diffusing inwards. This barrier function was particularly efficient at the early stages of oxidation/sulphidation. It is important to note that, in general, linear corrosion kinetics are considered non-protective. In the present study, due to the formation of cracks during the heating and corrosion processes, the corrosion of HfN-

coated samples followed linear rate laws. This implies that the coating only provided limited protection within a relatively short period (e.g. 240 h). However, the protectiveness of the coating may disappear after prolonged exposure. The rate constants depend on the activation energy of the reaction in the case of the linear rate kinetics. As a result, the rate constants are independent of the thickness of the HfN coating. In spite of its high stability, it is thermodynamically possible for the HfN to dissociate in the oxygen/sulphur (nil nitrogen) atmosphere employed. However, the question that arises is whether the released hafnium species would react with oxygen or sulphur in the environment. Two possible reactions would occur at the gas coating interface: 3Hf+

S2 =

H~S2

H f + O2 = HfO2

(1) (2)

The corresponding AG ° can be described: A G ~ = - R T l n a n f \ ~ s 2 / ( a n O 3 "pS 2

(3)

H.L. D u e t al. / Materials Science and Engineering A205 (1996) 199-208 A G ~ = - R T l n a H f o 2 / aHf . p 0 2

205

(4)

where AG~ = free energy of formation for reaction (1), AG~ = free energy of formation for reaction (2), pS2 = dissociation partial pressure of Hf3S2, p O 2 = d i s sociation partial pressure of HfO2, aHf = activity of Hf, aHf3S2 = activity of Hf3S2, a H f o 2 = activity of HfO2, R = gas constant, and T = reaction temperature. Using the available thermodynamic data for these two reactions, the free energies of formation at 750 °C are - 72.6 kcal mole -~ ($2) for reaction (1) [26] and - 2 2 1 . 9 kcal mole-~ (02) for reaction (2) [27]. It is reasonable to assume unity for aHr3s2 and aHfo:. Then the minimum values of the hafnium activities for the formation of Hf3S2 and HfO2 can be calculated, giving 3.2 × 10 - ~ for Hf3S2 and 3.9 × 1 0 - 4 3 for HfO2 in this specific atmosphere ( p S : ~ 10 t Pa and p O 2 ~ 10 -18 Pa) at 750 °C. Here the calculated results clearly indicate that HfO2 would be the more thermodynamically favourable reaction product. This was indeed verified experimentally by the identification of only HfO2 using XRD. It is established [28] that plastic deformation of polycrystalline solids by dislocation glide needs five inde-

Fig. 7. Electron image and EDX line-scans showing typical compositional profilesthrough the scale on HfN-coated Ti-6AI-4V after 21,0 h exposure in an H2/H20/H2S environment at 750 °C.

Fig. 6. Electron image and Digimaps showing typical compositional profiles through the scale on HfN-coated Ti after 240 h exposure in an H2/H:O/H2S environment at 750 °C.

pendent slip systems. However, most ceramics, such as oxides and nitrides, contain fewer than five independent primary slip systems and the stresses required to activate secondary slip systems are generally higher than those for the initiation of cracks below a transition temperature. Therefore, below this transition temperature only elastic deformation of the ceramics occurs before fracture, thereby leading to extremely low fracture strains. In the present study, due to severe mismatch in the thermal expansion coefficients between the HfN coating and the substrates of Ti and T i - 6 A I - 4 V , as shown in Table 2 [27,29,30], with increasing temper-

206

H.L. D u e t al. / Materials Science and Engineering A205 (1996) 199 208

atures during the heating period, a thermal strain would be generated according to the following equation: eth = (~ub - ~HfN)"AT

(5)

Where eth = thermal strain, ~.,ub= thermal expansion coefficient of the substrates, 0~HfN ~---thermal expansion coefficient of the H f N coating, and A T = difference between room temperature and reaction temperature. Therefore the substrate would expand more rapidly than the H f N coating; as a result, a compressive stress on the substrate side and a corresponding tensile stress on the H f N coating side would be produced along the substrate-coating interface. At the early stages of heating, the HfN coating-substrate system can accommodate strain by elastic deformation. However, once the elastic deformation limit is exceeded, a mechanical failure of the system may occur, Bearing in mind that ceramics such as HfN have extremely low elastic strain, then when the tensile stress approaches a critical level at which the H f N coating fractures, the cracks would be generated at the coating side, as schematically illustrated in Fig. 10(a). The cracks may be initiated at some pre-existing defects within the coating (e.g. holes, pores, and voids) or at the coating-substrate interface (e.g. concave undulation) where a high stress concentration may be present. These cracks might be extremely small and may not be observed even under the SEM. At the early stages of exposure, the cracks did not grow enough to reach the surface of the coating. It is proposed that oxygen species diffused inward into the cracks through the HfOz layer and the H f N coating where some physical defects (e.g. microcracks, dislocations, voids, etc.) were present and reacted with Ti species which migrated from the substrates and produced TiO2. The oxygen species required to form TiO2 did not originate from the dissociation of HfO2 at the H f O 2 - H f N interface since the dissociation partial pres-

Fig. 8. Scanning electron micrographs showing morphology of crosssectioned HfN-coated Ti-6A1-4V after 240 h exposure in an H~/ H20/H2S environment at 750 °C.

(a)

(b)

:c) Fig. 9. Scanning electron micrographs showing morphology of crosssectioned HfN-coated Ti-SAl 4V: (a) before exposure and (b) and (c) after undergoing heating.

H.L. D u e t al. / Materials Science and Engineering A205 (1996) 199 208

Table 2 Thermal expansion coefficientsof Ti, Ti-6AI-4V alloy, AI203, HfN and HfO2 Material

Coefficient of thermal expansion (10 ~'deg i)

T e m p e r a t u r e Reference range (°C)

Ti Ti-6Al-4V AI203 HfN HfO2

9.9 10.0 8.9-9.1 6.9 5.8

20 800 20 500 20 1400 20 1100 20- 1000

[29] [30] [29] [27] [29]

sure (3.9 x 10-43 Pa) of HfO2 is lower than that of TiO2 (1.6 x l0 34 Pa). As shown in Figs. 6 and 8, the cracks developed from the coating-substrate interface did not reach the external gas sample interface and contained TiO2. A Pilling-Bedworth ratio of 1.73 for TiO2/Ti [1] means that the formation of TiO2 further increased the stress concentration at the crack tips and made the cracks grow outward. The cracks would grow

z

f . ~

~:

z

~: ......

~. ~ % . :

.

•

::. :::::-%.%:::::::J:::::7:::::::::: "~-------~-

" - - - - - - - -

~

0 2

fN coating Cracks

207

through the whole coating after a certain exposure time. If different sizes of cracks were initiated during the heating period, it is obvious that the times for these cracks to reach the coating surface would also be different. Therefore, more cracks could be observed with lengthening exposure time. Once these cracks grew to reach the coating surface, the cracks became the diffusion paths for the reactants. Inward diffusion of oxygen led to the formation of TiO~ within the cracks and at the coating-substrate interface. Ti species migrated to the coating surface and reacted with oxygen to form TiO2. Also, sulphur species were transported inward to the coating-substrate interface and produced TiS 2 for the HfN-coated Ti and A12S3 +TiS2 for the HfN-coated T i - 6 A I - 4 V , as shown in Fig. 10(b). It is highly likely that the formation of a n H f O 2 scale on the HfN coating further constrained the plastic flow of the HfN coating itself. Indeed, the coefficient of thermal expansion for HfO2 is somewhat less than that of HfN, as reported in Table 2. Consequently, the HfN coating was kept intact at the early stages of oxidation/ sulphidation and effectively prevented the substrates from being attacked by the environment. Therefore, future work on the coating design should concentrate on improving the coating's plasticity and reducing the differences in the thermal expansion coefficients between the coating and the substrates by changing the coating constituents or by the application of multilayered graded coatings.

Substrate

(a) Formation of HfO 2 over HfN coating and occurrence of cracks

coa,.ng

Substrate

(b) Formation of TiO 2 within cracks and over HfN coating and TIS2 and/or AI2 S3 Fig. 10. Schematic model describing the degradation process of the HfN-coated Ti and Ti 6Al 4V.

5. Conclusions

(1) During exposure to an atmosphere ofpS2 ~ 10 1 Pa and p O 2 ~ 10 18 Pa at 750 °C, Ti followed a linear rate law whilst T i - 6 A 1 - 4 V obeyed a linear-parabolic rate law. (2) The applied HfN coating enhanced the oxidation/ sulphidation resistance of Ti and Ti 6A1-4V. The linear corrosion rate constant of the HfN-coated Ti was 5-times smaller than that of the uncoated samples. (3) Uncoated Ti characteristically displayed a doublelayered oxide scale of TiO2 with a TiS2 layer at the oxide-substrate interface. Uncoated T i - 6 A I - 4 V also yielded a double-layered TiO2 scale in which A120 3 was precipitated in the external portion of the outer layer of TiO2, whilst a layer of A12S3 and TiS2 with vanadium sulphide was located underneath the inner layer of TiO2. For Ti and Ti 6A1 4V the HfN coating suppressed the formation of these duplex scales. (4) Severe cracking of the HfN-coated Ti and T i 6AI-4V developed during periods of heating and isothermal corrosion deriving from the large differences in thermal expansion coefficients between the HfO2, HfN and the substrates.

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Acknowledgement Grateful acknowledgement is given to the financial funding of H.L. Du, as a Research Fellow, at the University of Northumbria at Newcastle, and D.B. Lewis, as a Senior Research Fellow, at Sheffield Hallam University, by the Science and Engineering Research Council's Rolling Grant in Surface Engineering.

[11] [12] [13] [14] [15]

[16]

[17]

References

[18]

[1] P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, Barking, 1988. [2] K. Hauffe, Oxidation of Metals, Plenum, New York, 1965. [3] N.E. Paton and J.C. Williams, Hydrogen in Metals, American Society for Metals, Metals Park, OH, 1974, p. 409. [4] B.A. Kolachev, S.A. Vigdorchick, A.V. Malkov, in K. Kimura and O. Izumi (eds.), Titanium' 80 Science and Technology, Vol. 3, Met. Soc. AIME, p. 1671. [5] P.K. Datta, K.N. Stratford and A.L. Dowson, in J. Bolton and S. Hampshire (eds.), Proc. 2nd Irish Conf. on Durability and Fracture, 1984. [6] A.L. Dowson, PhD Thesis, Newcastle upon Tyne Polytechnic, 1988. [7] H.L. Du, P.K. Datta, 13. Lewis and J.S. Burnell-Gray, High temperature corrosion of Ti and Ti-6A1 4V alloy, Oxid. Met., in press. [8] H.L. Du, P.K. Datta, B. Lewis, and J.S. Burnell-Gray, Corr. Sci., 36 (1994) 631. [9] I.A. Menzies and K.N. Stratford, Corr. ScL, 7 (1967) 23. [10] I.A. Menzies and K.N. Stratford, J. Less Common Met., 12 (1967) 85.

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