**1. Introduction**

In recent years, to improve the service temperature of the hot section of an aircraft engine, besides cooling gas film, thermal barrier coatings (TBCs) consisting of an oxidation-resistant metallic bond coat and a thermally insulating ceramic topcoat of yttria-stabilized zirconia (YSZ) have been applied due to their low thermal conductivity and high thermal expansion coe fficient [1–4]. At present, air plasma spraying (APS) and electron beam–physical vapor deposition (EB-PVD) are the two main technologies that are used for TBC deposition [5]. However, both of the above technologies exhibit limitations. YSZ coatings prepared by the APS process have a typical layered microstructure with a large number of interfaces, micro-cracks and pores, which contribute to low thermal conductivity and poor thermal shock resistance [6]. Di fferent from the APS case, TBCs with super strain tolerance and improved spallation lifetime with high thermal conductivity can be obtained through the EB–PVD process. This is because of the columnar microstructure with an inter-columnar gap that is perpendicular to the top-coat/bond-coat interface [7,8].

To meet the increasing temperature demands of engines, coatings with better thermal insulation and high strain tolerance are preferred. Recently, many e fforts have been devoted to investigating plasma spray–physical vapor deposition (PS–PVD) TBCs with various microstructures, including dense coatings, PVD-like columnar coatings, nano-sized solid clusters columnar coatings (quasi-PVD) and mixed microstructure coatings [9–13]. As we know, PS–PVD combines the advantages of APS (high deposition rate and cost e fficiency) and EB–PVD (the ability to produce columnar structured coatings). TBCs with the initially favorable columnar microstructure can be prepared by PS–PVD, which combines the high thermal insulating property of the APS coatings and the high strain tolerance of the EB–PVD coatings [9–11]. Therefore, many existing results show that the PS–PVD process has the promising potential to fabricate durable YSZ coatings.

It is noted that YSZ TBCs must be exposed to high temperature during harsh service and their thermal stability is a key factor for evaluating the performance of the coatings. As a consequence, the ceramic top-coat may be exposed to high temperatures for a long time, which has grea<sup>t</sup> influence on the phase composition, microstructure and mechanical properties. It has been reported that thermal insulation ability and strain tolerance of the as-sprayed TBCs dramatically decreased upon annealing at high temperature because of the micro-cracks and pores healing during the sintering process [14–16]. The initial metastable tetragonal (t-ZrO2) phase is the main phase of the as-sprayed TBCs, which is believed to be a direct consequence of both the slow di ffusion rate of Y3+ ions and a small driving force [17]. The t-ZrO2 phase could toughen the ceramic by re-orienting its *c*-axis in crystal cells and absorbing fracture energy to increase the resistance to cracking [18]. The t-ZrO2 phase is not the equilibrium phase, and transforms to equilibrium tetragonal (t-ZrO2) and cubic (c-ZrO2) phase at high temperature. Upon further cooling, the newly precipitated t-ZrO2 will transform to monoclinic (m-ZrO2) phase. Moreover, the T/M phase transformation will cause cracking due to the huge volume change (approximately 5%) caused by the large density di fference between T and M phases [19]. For these reasons, it is necessary to qualitatively analyze the phase and microstructural stability at high temperature and their relation to mechanical properties. However, research on the relationship between phase degradation and properties in PS–PVD YSZ TBCs at high temperature is rarely reported. Therefore, in the current work, with the aim of studying the durability of the PS–PVD coatings, the e ffect of exposure time and temperature on the high-temperature stability of PS–PVD YSZ coatings will be evaluated. Thermal stability is investigated through the evolution of phase composition, microstructure, and mechanical properties under various temperatures and times. This basic research might provide some useful insight for beneficial adjustments and further improvements to the performance of PS–PVD YSZ TBCs.

## **2. Materials and Methods**

## *2.1. Preparation and Heat Treatment of PS–PVD Coating*

The as-deposited YSZ coatings with a thickness of about 310 μm were fabricated on polished graphite substrates (50 mm in length, 10 mm in width, and 10 mm in thickness) using a PS–PVD system (Oerlikon Metco, Wohlen, Switzerland) using an O3CP-type plasma gun with a 60CD powder feeder. An agglomerated ZrO2–(6–8 wt.%) Y2O3 powder (YSZ, Metco 6700) was used as feedstock; the particle size range was 5–22 μm. Prior to depositing the TBC, the polished graphite substrates were preheated by plasma flame flow. The detailed spraying parameters are listed in Table 1.

The free-standing samples used for thermal aging treatment were stripped from graphite substrates through the mismatch between the expansion coe fficients of the substrate and coating. Free-standing YSZ coatings were isothermally heat-treated in a high-temperature chamber furnace (HTK 20/17, ThermConcept, Bremen, Germany) at 1200–1600 ◦C for 24 h and at 1550 ◦C for 20–100 h, respectively. Free-standing YSZ coating samples were put in a furnace and heated at about 10 ◦C/min to the target temperature, followed by holding for the selected time. Then, the samples were cooled naturally in the furnace to room temperature. All of the free-standing coating samples were placed in vacuum with epoxy, cut to supply the observation of the cross-sections, and finely polished by routine metallographic methods.


**Table 1.** Spray parameters for YSZ coatings by PS–PVD.

## *2.2. Coating Characterization*

X-ray di ffraction (XRD) and Raman spectroscopy (RS) were used to characterize the phase composition of the as-deposited and isothermally heated coating samples. XRD measurements were performed on the X-ray di ffractometer (D/Max 2550 V, Rigaku, Tokyo, Japan) operating in the reflection mode with Cu-K α radiation (40 kV, 100 mA), using a step scan mode with the step of 0.02◦ (2θ) and 1 s per step. A Raman microscope system (Invia, Renishaw, Glouchestershire, UK) was used to record Raman spectra at the room temperature with a 532 nm laser excitation. Microstructures and grain size of the heat-treated samples together with the as-deposited coating samples were characterized by a scanning electron microscope (SEM, MIRA3 LM, TESCAN, Brno, Czech Republic) in backscattered electron image mode. A rectangular intercept method on secondary electron images which recorded the fracture surfaces of the coatings was used for measuring the grain size of the coatings. The average grain size, *D*, was then given by:

$$D = \sqrt{\frac{4A}{\pi (n\_i + \frac{n\_o}{2})}} \tag{1}$$

where *A* is the rectangle area, *ni* and *no* are the grain numbers in the rectangle and on the rectangular boundary, respectively [20]. The number and magnification of the as-sprayed coating micrographs employed for grain size measurements were 10 and 100,000 times, those of the coatings which were heat-treated in air were 10 and 10,000 times. Hardness ( *H*) and Young's Modulus (*E*) of all coatings were measured on the cross-section before and after thermal treatment through a G-200 nanoindenter (Agilent Technologies, Oak Ridge, TN, USA) equipped with a Berkovich indenter. Specific test schemes are listed in the literature [21].

## **3. Results and Discussion**
