Microstructural changes and phase transformations in a plasma-sprayed zirconia-yttria-titania thermal barrier coating

8LIIITI I

TIN68

Htll#OLOT ELSEVIER

Surface and Coatings Technology 82 (1996) 284-290

Microstructural changes and phase transformations in a plasma-sprayed zirconia-yttria-titania thermal barrier coating P. Diaz, M.J.

Edirisinghe *, B. Ralph

Department of Materials Technology, Brunel University, Uxbridge, Middlesex UB8 3PH, UK Received 11 July 1995; accepted in final form 6 October 1995

Abstract

A thermal barrier coating comprising a zirconia-yttria-titania ceramic coat deposited by plasma spraying on an alloy steel substrate using an Ni-Cr A1-Co-Y intermediate bond coat has been soaked at various temperatures in the range 700-1400 °C and cooled to room temperature in static air. Some thermal treatments were also carried out in an argon atmosphere. Scanning electron microscopy (SEM) studies show that a discontinuous thickness of the bond coat is the reason for the failure of the coating due to thermal treatment. The phase transformations and microstructural changes due to thermal treatment have been investigated using X-ray diffraction and SEM. Thermal treatment did not affect the monoclinic phase of the Z r O 2 content present in the ceramic coat. In the case of thermal treatment in air and argon (below 800 °C) the cubic --*tetragonal phase transformation occurred. The appearance of the fracture surfaces of the ceramic coat suggests that interlamellar fracture in the as-sprayed condition changed to intralamellar fracture after thermal treatment below 1200 °C. At higher temperatures the fractures were intergranular in appearance. Reasons for these changes in the appearance of the fracture surfaces are suggested with reference to the microstructure. Microstructural investigations also show some changes in the area of porosity distribution of the ceramic coat. Thermal treatment heals some pores, but other voids appear (cracking of the coating).

Keywords: Microstructure; Phase transformations; Plasma spraying; Thermal barrier coating (TBC); Zirconia-yttria-titania

1. Introduction

Zirconia (ZrO2) is an attractive material for several engineering applications because, compared with other ceramics, it has superior mechanical properties such as high strength and fracture toughness combined with good wear resistance and a thermal expansion coefficient close to that of iron and iron-based alloys. However, owing to its polymorphism, it must be stabilized by the addition of other oxides [1,2]. Several binary and ternary oxide systems such as ZrOz-yttria (Y203), ZrO2-ceria (CeO2), ZrOz-calcia (CaO) and ZrO2 Y203-magnesia (MgO) have been found to be potential sources of high toughness zirconia polycrystals (TZP) [ 1-4]. In comparison, Z r O 2 systems containing titania (TiO2) have been found to be inferior [5,6]. TiO2 dissolves in Z r O 2 and decreases the transformation temperature at which the tetragonal Z r O 2 solid solution transforms into the monoclinic form, but the

* Corresponding author. 0257-8972/96/$15.00 © 1996 Elsevier Science S.A. All rights reserved S S D I 0 2 5 7 - 8 9 7 2 ( 9 5 )02721-1

tetragonal phase is not retained at room temperature. Therefore the intention to use TiO2 as a stabilizer of the tetragonal phase at room temperature has been unsuccessful [7 9]. However, investigations on the systems Z r O 2 CeO2 TiO 2 [10] and ZrO2 Y203-TiO2 [7 9] show that they are useful for the preparation of T Z P materials. Pyda et al. [8,9] found fully or nearly fully tetragonal polycrystals within a wide range of composition (Y2Oa from 0.5 to 3 m o l % and TiO 2 from 3 to 28 tool %). TiO 2 promotes grain growth of the tetragonal phase, but in compositions with a high TiOz concentration (23-28 mol %) the grain growth is inhibited [61. TiOz can also destabilize the cubic phase in TiOz~203-ZrO2 [ 11,12]. The use of ZrO2 alloyed with various stabilizing oxides such as MgO, CaO, C e O 2 and Y203 is quite well known in the preparation of thermal barrier coatings made mostly by plasma spraying [13 15]. It is widely believed that ZrO 2 ceramic coats containing 6 10 wt.% Y203 offer the most desirable properties in the service of these coatings. However, it has been suggested recently that ZrO2-Y203 TiO2 ceramic coats exhibit improved

P. Diaz et al./SurJitce and Coatings Technology82 (1996) 284-290

properties such as erosion, abrasion and corrosion resistance compared with the ZrOz-Y20 3 system [16]. A study on this system was initiated recently [ 17] establishing a comparison with the widely known ZrO2-YzO3 system in terms of the phases present. The Z r O ~ YzO3-TiO2 samples contained a higher concentration of the tetragonal phase than the ZrOz-Y20 3 samples. The stability of alloyed zirconia after various thermal treatments has been the subject of numerous investigations [18-20]. Some studies have been focused on the identification and quantification of phase transformations in ZrO2, mainly by X-ray diffraction (XRD) [21-23], and the influence of these on the properties, either in bulk-sintered solids or plasma-sprayed coatings. In the present study a plasma-sprayed ZrO2 Y203 TiO2 thermal barrier coating has been subjected to various thermal treatments. Phase changes in the ceramic coat caused by thermal treatment are identified and quantified using XRD. The microstructural variations in the ceramic due to thermal treatment are investigated by scanning electron microscopy (SEM) of the fracture surfaces.

285

ments were carried out under an argon atmosphere at temperatures up to 1200 °C. In these experiments the specimens were left to cool slowly inside the furnace and reached room temperature in about 8 h. 2.3. Characterization by X R D and S E M

Characterization of the materials used for the ceramic coat and the bond coat was carried out using XRD, SEM and energy-dispersive X-ray spectroscopy (EDS). Polished sections of the as-sprayed and thermally treated specimens were studied using SEM and EDS. An evaluation of the changes in porosity in the ceramic coat due to thermal treatment at 900'~C in air was performed using image analysis of the relevant scanning electron micrographs. Fracture surfaces of the as-sprayed and thermally treated specimens were also studied using SEM. The cubic, monoclinic and tetragonal mole fractions of ZrO2 present in each specimen before and after thermal treatment (see Section 2.2) were determined using the integrated intensities of the {111} and {400} XRD reflections as described previously [ 17,22].

2. Experimental details 3. Results and discussion 2.1. Materials and Preparation of Specimens 3.1. Ceramic and bond coat powders

An Ni-Cr 304A high alloy steel was used as the base metal. This was coated with an intermediate metallic bond coat (containing 17.5wt.% Cr, 5.5wt.% A1, 2.5 wt.% Co, 0.5 wt.% Y203 and the remainder Ni) and a top coat of a ZrO 2 alloy (containing 10 wt.% Y203 and 18 wt.% TiO2). Coated specimens, which are used as a thermal barrier material, were prepared by Plasmatec (Edo Miranda, Venezuela) and supplied by Intevep SA (Edo Miranda, Venezuela). The steel surface was sandblasted with alumina and then precoated with the bond coat by thermal spraying using an oxygen-acetylene flame (thermal spraying was used to reduce the cost). The ceramic coat was sprayed on to the bond coat using a 60 V, 600 A Hz-Ar plasma. Plasma spraying was preceded by preheating the substrate to about 110 °C and a coating 0.025 mm thick was applied per pass. Post-spray cooling was carried out in static air. Further details of the plasma-spraying process used are given in Ref. [16].

The ZrO2-Y203-TiO2 powder used in this investigation is produced by spray drying and consists of porous spherical agglomerates of varying size (less than 10 gm) as shown in Fig. 1. The XRD pattern of this powder (Fig. 2) shows that ZrO 2 is present only in the monoclinic form. Y203 (cubic) and TiO 2 (anatase) were also detected in the powder as indicated in Fig. 2. The bond coat particles (Fig. 3) correspond to a sintered and crushed powder type. Oblong particles about 100 gm in size have a higher Ni and Cr content

2.2. Thermal Treatment

The specimens (30 x 30 x 5mm 3) were soaked in static air at 700, 900, 1200 or 1400 °C for 24h and subsequently cooled to room temperature in static air in about 60 rain. In order to evaluate the influence of the oxidation of the bond coat during thermal treatment, some experi-

Fig. 1. Scanningelectronmicrographof the ZrO Y203~Ti02 powder used for plasma spraying.

P. Diaz et al./SurJace and Coatings Technology 82 (1996) 284 290

286

=2

ij, ill

15

ea

e5

~fl

,, 35

40

45

AIA 50

~

60

, ~5

?(i

7fi

8o

t15

91)

2 0 (degrees) Fig. 2. XRD pattern of the powder shown in Fig. 1. Z, Y and T denote zirconia, yttria and titania peaks respectively; m and c refer to monoclinic and cubic respectively. The zirconia peaks are of the monoclinic type, but for clarity not all these have been labelled.

Fig. 4. Scanning electron micrograph of a polished cross-section of the as-sprayed coating. C and B refer to the ceramic and bond coats respectively.

off after being exposed to 700 °C in both static air and argon. The oxidation of the bond coat has been considered widely as the primary cause of thermal spalling [-24,25]. However, it was found that specimens with a discontinuous bond coat failed even when they were subjected to thermal treatment under an argon atmosphere. Therefore the residual stress caused by the increased thermal mismatch due to the discontinuous bond coat is the reason for the thermal spalling of the coating. 3.3. Fracture of the ceramic coat

Fig. 3. Scanning electron micrograph of the Ni Cr-AI Co-Y powder used for the bond coat.

particles

of the

as detected by EDS. Smaller particles (less than l0 gin) show a higher concentration of A1, Co, Y and Mo as detected by EDS. 3.2. Thermal treatment

A continuous bond coat (uniform thickness of about 50 gm) between the substrate and the ceramic coat is present in the majority of specimens (Fig. 4). In these specimens the coating remained attached to the substrate after thermal treatment in both static air and argon. Some specimens presented a discontinuous bond coat of non-uniform thickness and this could be caused by the use of thermal spraying (Fig. 5(a)). This can result in extra residual stress on the ceramic coat during thermal treatment. Fig. 5(b) and 5(c) correspond to the X-ray mapping for Ni and Zr respectively of the crosssection shown in Fig. 5(a). It can be seen clearly that there is a discontinuity in the nickel-based bond coat present between the ZrO2 ceramic coat and the steel substrate. In these specimens the ceramic coat spalled

During the plasma-spraying process the material is deposited in discontinuous layers (lamellae) parallel to the substrate. Fracture surfaces of the as-sprayed ceramic coat show an interlamellar appearance (Fig. 6). This type of fracture has been explained in terms of weak interlamellar adhesion caused by the discontinuous solidification process inherent in plasma spraying [26]. Specimens of the as-sprayed ceramic coat cooled in liquid nitrogen were also fractured. These show intralamellar fracture (Fig. 7) and indicate that the ceramic "layers", deposited in the form of lamellae, have a fine columnar microstructure. The lamellar thickness is about 5 gm and the columns have a diameter less than 1 gm. Fracture surfaces of the ceramic coat after thermal treatment below 1200 °C in both air and argon also showed an intralamellar appearance (Fig. 8). This change in fracture behaviour of the thermally treated specimens, in contrast with the as-sprayed specimens (Fig. 6), is caused by improved interlamellar contact and adhesion, probably due to the healing of the interlamellar porosity during thermal treatment, as discussed in Section 3.4. At a higher temperature ( 1400 °C) the fracture surfaces of the specimens heat treated in air presented a clear intergranular appearance (Fig. 9). There were some

P. Diaz et aL/SurJace and Coatings" Technology 82 (1996) 284 290

287

(a) Fig. 6. Scanning electron micrograph of the fracture surface of the as-sprayed ceramic coat.

(b) Fig. 7. Scanning electron micrograph showing intralamellar fracture of the ceramic coat of a specimen cooled in liquid nitrogen before fracture.

(c) Fig. 5. (a) Scanning electron micrograph of a polished cross-section of the coating showing discontinuity in the bond coat. B indicates the few regions in which the bond coat is present. (b) X-Ray map of Ni for the area shown in (a). (c) X-Ray map of Zr for the area shown in (a). Fig. 8. Scanning electron micrograph showing intralamellar fracture of the ceramic coat of a specimen soaked at 1000 °C. m o r p h o l o g i c a l changes observed in these fracture surfaces, such as the collapse of the c o l u m n a r structure, giving rise to smaller crystals a b o u t 1 g m in size. H a r m s w o r t h a n d Stevens [26,27] explained the occurrence of the i n t e r g r a n u l a r fracture in Z r O z - Y 2 0 3 coatings by considering a weak b o n d between the ceramic grains caused by a thin i n t e r g r a n u l a r glassy phase. It is

most likely that a similar situation exists Z r O 2 - Y 2 0 3 - T i O 2 samples.

in the

3.4. Porosity

The n a t u r e of the porosity found in Z r O z - Y 2 0 3 has been described as [ 2 6 ] (i) some elongated a n d flat voids

P. Dia: et al./Surface and Coatings Technology 82 (1996) 284 -290

288

an increment in the thermal conductivity of the thermal barrier coating after treatment at elevated temperatures [-13]. 3.5. Phase transformations in the ceramic coating

Fig. 9. Scanning electron micrograph showing intergranular fracture of the ceramic coat of a specimen soaked at 1400 C in air.

produced by mismatch between lamellae, (ii) some smaller and spherical pores caused probably by entrapped gas in the material during the plasma-spraying process and (iii) larger voids associated with cracks. There were no significant differences in the pore size distribution before and after thermal treatment at 900 °C in air. However, some changes were observed in the area of porosity distribution. Fig. 10 presents a comparison between the area of porosity distribution of the as-sprayed ceramic coat and after being subjected to thermal treatment in air at 900 °C. At least three populations (or types) of porosity can be distinguished in these distributions. Some of the interlamellar voids in the ceramic coat were healed by thermal treatment and a reduction in their population (4.9% 6.0%) occurs in the area of porosity distribution (Fig. 10). At the same time, some cracks appeared owing to thermal treatment and a new population of voids (greater than 10%) is seen in the area of porosity distribution (Fig. 10). The healing of the interlamellar voids and the corresponding increase in interlamellar contact have also been associated with

IJ.. oay r

25

15

~

As sprayed

The XRD pattern of the as-sprayed ceramic coat (Fig. 11) shows that it consists of mainly tetragonal and cubic ZrO2 with about 10% of the monoclinic form (the plasma-spraying parameters used caused most of the monoclinic Z r O 2 in the powder to transform to tetragonal and cubic Z r O 2 ) . Some relatively small peaks corresponding to TiO2 (rutile) and Y 2 0 3 (cubic) can also be distinguished. The changes observed in the proportion of each ZrOz phase when heat treated in air and argon atmospheres are presented in Figs. 12(a) and 12(b) respectively. In each case the monoclinic proportion remains virtually the same after each thermal treatment. The cubic form transforms to tetragonal after thermal treatment in air (Fig. 12(a)) at all temperatures investigated. Considering that the samples were cooled rapidly to room temperature after being soaked at a high temperature, a diffusionless transformation process must be taking place in this case. A displacive and diffusionless transformation process has been described previously for quenched YzO3-alloyed ZrO2, where diffusion has been limited by the fast cooling rate E28-31 ]. A different trend was observed in specimens heated and slowly cooled in argon (Fig. 12(b)). The proportion of tetragonal ZrOz increases while the cubic ZrOz content decreases after the specimen is soaked at temperatures up to about 800 °C, which suggests that the cubic~tetragonal phase transformation is taking place. At higher temperatures the much slower cooling rate in argon helps diffusion and this is the most likely reason for the final proportion of tetragonal to cubic ZrOz being only slightly different from its initial level. The ZrO2-Y203 phase diagram has been used to explain phase transformations observed in plasma-sprayed ZrOz-Y203 thermal barrier coatings [32-34]. However, phase equilibrium information on the ZrOz-YzO3-TiO2 system is not available at present in order to understand fully the observations of the present work.

lO 4. Conclusions

0

0.5

2.7

4.9

7.1

9.3

11.5 13.7

Area of Porosity (%) Fig. 10. Area of porosity distribution in the ceramic coat in the as-sprayed condition and after thermal treatment at 900 °C in air. The area of porosity is the ratio between the area of pores present and the total area considered during image analysis.

Comparison of experimental results of thermal treatment in static air and argon shows that, compared with oxidation, the discontinuous thickness of the bond coat has a more significant influence on the failure of the ceramic coat due to thermal treatment. Thermal treatment up to 1400 °C in air did not affect the 10% by volume of monoclinic phase of zirconia present in the ceramic coat but caused the transformation

P. Diaz et al./Surface and Coatings Technology 82 (1996) 284 290

289

o t-~

fl

Ill 10

t5

,90

25

30

35

40

45

50

55

60

65

70

75

80

135

90

20 (degrees) Fig. 11. X R D pattern of the as-sprayed ZrOz Y203-TiO2 coating. Z denotes zirconia peaks; m, c and t refer to monoclinic, cubic and tetragonal respectively; T denotes a titania peak. For clarity not all the peaks have been labelled.

of cubic zirconia to tetragonal zirconia. In contrast, for thermal treatment in argon where slow cooling took place, the cubic~tetragonal transformation was evident only up to 800 °C. The ceramic coat showed a layered sequence of lamellae about 5 I-tin in thickness with a columnar microstructure. Interlamellar fracture of the ceramic coat occurred in the as-sprayed condition, but after thermal treatment up to 1200 °C intralamellar fracture was observed. At higher temperatures (1400 °C) the fi'acture surface was of an intergranular appearance where the columnar structure had broken down to small crystals about 1 gm in size. Thermal treatment also caused changes in the area of porosity distribution in the ceramic coat. Some of the interlamellar voids in the ceramic coat were healed, but another population of voids associated with cracks appeared owing to thermal treatment.

100

6o ~

Tetragonal

g 40 1 ,

Cubic

20

Monoclinic

~"~---.i _ . . ~ E

200

400

600

800

(a)

11300

1200

1400

Temperature (°C)

8O "-7".

Tetragonal ~.

60

~0

4(1

Acknowledgements

cubic

2(1

200

(b)

\

400

600

80C

j

1000

1200

Temperature (°C)

Fig. 12. Phase transformations in the Z r O 2 - - Y 2 0 3 TiO2 ceramic coating (a) after thermal treatment in static air and (b) after thermal treatment in argon.

Pedro Diaz acknowledges gratefully the financial support given by the British Council and C O N I C I T (Venezuela). Mrs. Carmen Di Rausso of Intevep SA is thanked for technical help. The helpful information supplied by Joseph Debarro of SULZER M E T C O (U.S.) Inc., New York, USA is acknowledged. Kenneth Batchelor and Mrs. Kathy Goddard are thanked for their help with the preparation of figures and the typing of the manuscript respectively.

290

P. Diaz et al.,"Surjace and Coatings" TechnoloL,y 82 (1996) 284 290

References [1] I. Nettleship and R. Stevens, lnt. J. High Technol. Ceram., 6 (1987) 1. [2] T. Gupta, J. Bechtold and R. Kvznicki, J. Mater. Sci., 12 (1977) 2421. [3] R. Garvie, J. Am. Ceram. Soc., 20 (1984) 23. [4] J. Brandon and R. Taylor, Surf. Coat. Teehnol., 39 40 (1989) 143. [5] H. Frank, H. Brown and P. Dwwez, J. Am. Ceram. Soe., 37 (1954) 129. [6] V. Pandolfelli, I. Nettleship and R. Stevens, Prot. Br. Ceram. Soc., 42 (1989) 139. [-7] R. Ingel and D. Lewis, in S. Somiya et al. (eds.), Advanees in Ceramics, Vol. 24, Science and Technology t~fl Zireonia III, American Ceramic Society, OH, 1988, p. 385. [8] W. Pyda, R. Haberko and M. Bucko, Ceram. Int., 18 (1992) 321. [9] W. Pyda, R. Haberko and M. Bucko, in S. Badwal et al. (eds.), Seienee and Technology ~[ Zirconia V, Technomic, PA, 1993, p. 137. [10] V. Pandolfelli, W. Rainforth and R. Stevens, in G. de With et al. (eds.), Euroeeramics: Properties ~f Ceramics, Vol. 2, Elsevier, London, 1989, p. 161. [ 11 ] P. Kountouros and G. Petzow, in S. Badwal et al. (eds.), Science and Technology qt'Zireonia V, Technomic, PA, 1993, p. 30. [12] C. Lin, Mater. Sci. Eng., A129 (1990) 147. [13] R. McPherson, J. Aust. Ceram. Soe., 20 (1984) 23. [14] R. Miller, Smf. Coat. Technol., 30 (1987) 1. [15] J. Sun, Su~[] Coat. Teehnol., 58 (1993) 93. [16] Tech. Bull. 143, 1984 (METCO Inc., New York). [17] P. Diaz, M.J. Edirisinghe and B. Ralph, J. Mater. Sci. Lett., 13 (1994) 1595.

[18] B. Wu and E. Chang, Thin Solid Films, 172 (1989) 185. [19] L. Lelait, S. Alperine and C. Diot. Mater. Sei. E~lg., AI21 (1989) 475. [20] R. Miller, J. Smialik and R. Garlick, in A.H. Heuer and L.W. Hobbs (eds.), Adzanees in Ceramics, Vol.3, Science and Technology ~/" Zireonia. American Ceramic Society. OH, 1981, p. 241. [21] G. Heintze and R. McPherson, in S. Somiya et al. (eds.), Advances in Ceramics, Vol. 24. Scienee and Technology ~/' Zireonia 111, American Ceramic Society, OH, 1988, p. 431. [22] N. Iwamoto, N. Umesably and S. Endo, Thin Solid Fihns, 127 (1985} 129. [23] J. Brandon and R. Taylor, Sur[~ Coat. Technol.. 46 (1991) 75. [24] B. Wu, Mater. Sei. Eng., A124 (1990) 215. [25] W. Lih, Oxid. Met., 38 (1982) 99. [26] P. Harmsworth and R. Stevens, J. Mater. Sci., 27 (1992) 616. [-27] P. Harmsworth and R. Stevens, Proc. Br. Ceram. Soc., 42 (1989) 123. [28] H. Scott, J. Mater. Sci.. I0 (1975) 1527. [29] A. Heuer, Aeta Metall., 35 (1987) 661. [30] A. Heuer, in S. Somiya et al. (eds.), Advances iH Ceranlics, Vol. 24, Science and Technology ~( Zirconia Ill, American Ceramic Society, OH, 1988, p. 241. [31] G. Heuntze and R. McPherson, in S. Somiya et al. (eds.), Adt, anees in Ceramics, Vol. 24. Science and Technology ~/" Zirconia II1, American Ceramic Society, OH, 1988, p. 43l. [32] B. Chert, J. Am. Ceram. Soe., 72 (1989) 3. [33] R. Miller, in A.H. Heuer and L.W. Hobbs (eds.), Adz~anees itl Ceramies, Vol. 3, Seienee and Technology ~f Zireonia, American Ceramic Society, OH, 1981, p. 3. [34] A. Heuer, in N. Claussen et al. (eds.), Ath,antes in Ceramics, Vol. 12, Science atut Technology ~/ Zireonia 11, American Ceramic Society, OH, 1984, p. 1.