Journal Archive

Johnson Matthey Technol. Rev., 2017, 61, (2), 93
doi: 10.1595/205651317X695064

Iridium Coating: Processes, Properties and Application. Part II

Analysis of the coatings against high-temperature oxidation and corrosion

  • Wang-ping Wu*
  • School of Mechanical Engineering, Institute of Energy Chemical Equipment and Jiangsu Key Laboratory of Materials Surface Science and Technology, Changzhou University, Changzhou 213164, P.R. China
  • Zhao-feng Chen
  • International Laboratory for Insulation and Energy Efficiency Materials, College of Material Science and Technology, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, P.R. China
  • *Email: wwp3.14@163.com

Article Synopsis

Iridium as a barrier coating is an important area of high-temperature application. In Part I, the introduction was presented and the different deposition processes were reviewed (1). This paper, Part II, describes the texture and structure evolution, mechanical properties, growth mechanisms and applications of Ir coatings. The mechanisms of micropore formation after high-temperature treatment are also investigated in some detail.

1. Texture, Structure Evolution and Mechanical Properties of Iridium Coatings

1.1 Texture Evolution

The crystal orientation of the coating can influence its structure and further affect the properties and hence the feasibility of the coating for use in extreme environments (2, 3). For fcc metals, films formed by MOCVD, PVD and electrodeposition present a <111> orientation. In general, the orientation of Ir coatings formed by DCMS, RFMS or MOCVD is the (111) crystal plane. The (111) orientation is most likely favoured because it is the most closely packed arrangement with the lowest surface energy. The orientation of Ir films created by RFMS deposited on a yttria stabilised zirconia film depended on the deposition rate (4, 5). When the deposition rate was low, an epitaxial (100) Ir film was formed. A highly (111)-orientated Ir film was grown at higher deposition rates.

Murakami, Yano and Sodeoka (6) suggested that an Ir coating formed by magnetron sputtering was strongly orientated to the (220) direction whilst a coating formed by electron beam PVD did not exhibit such tendencies. The orientation of the coating depended on the substrate temperature, the partial pressure ratio of argon to the entire total deposition atmosphere gas, the deposition rate and the thickness of the coating (7, 8). Wessling et al. (9, 10) observed that the crystallographic texture of the Ir film changed from a <220> to a <111> orientation with increasing incident energy. The <220>-orientated films were rougher, with lower densities and stronger void formation. The coating with <111> texture was smoother and denser than that with <110> texture regardless of the grain size. With increasing current density, the surface roughness of the electrodeposited Ir coating increased while the grain size decreased and the preferred orientations changed in the order <111>→<110>→<311> (11). With increasing deposition temperature, the surface roughness decreased while the grain size increased, and the preferred orientation changed in the order <111>→<110>→<111>. The substrate was composed of polycrystalline Nb (Figure 1(a)). (110), (200), (211) and (220) Nb diffraction peaks appeared. The preferential growth orientation could be determined by a texture coefficient (TC(hkl)), which was calculated using Equation (i) (13):


(i)

TC(hkl) is the texture coefficient of the (hkl) plane, I(hkl) is the measured intensity of (hkl) plane, I0(hkl ) is the corresponding recorded intensity in the Joint Committee on Powder Diffraction Standards (JCPDS) data file and n is the number of preferred growth directions of Nb or Ir. According to Equation (i) and the JCPDS Card (No. 35-0789), TC(110), TC(200), TC(211) and TC(220) of Nb were 0.346, 2.703, 0.463 and 0.519, respectively. It was indicated that the substrate had a preferred (200) orientation. The Ir coating had a dominant (220) orientation in the XRD pattern (see Figure 1(b)). The diffraction intensities of the other peaks were weak. The epitaxial growth of the (220)-orientated Ir coating did not occur on the (200)-orientated Nb substrate. The (220)-textured coating was due to the deposition of initial nuclei with preferred growth on the substrate surface. The crystal orientation of the DGP Ir coating was further investigated by an electron backscatter diffraction (EBSD) technique (3).

Fig. 1.

X-Ray diffraction (XRD) patterns of: (a) Nb substrate; and (b) Ir coating formed by a DGP process (Reproduced with permission of Elsevier (12))

In Figure 2, the red, blue and green colours represent Ir (001), Ir (111) and Ir (101) crystal faces, respectively. The orientation distribution map shows that the vast majority of colour on the coating surface was green, indicating that the DGP Ir coating had a preferred (101) orientation. The coating was mainly composed of submicrometre-sized grains of 0.2–0.3 μm (Figure 3).

Fig. 2.

Orientation distribution map for Ir grains according to the inverse pole figure (Reproduced with permission of Elsevier (12))

Fig. 3.

Grain size distribution map of the surface of a DGP Ir coating (Reproduced with permission of Elsevier (12))

Figure 4 shows the maximum value distributed in the centre of the (101) pole figure. This indicates that the preferred orientation of the coating was a (101) texture. This result was in agreement with that of the XRD pattern (Figure 1).

Fig. 4.

Pole figures of {001}, {101} and {111} textures for the surface of a DGP Ir coating (Reproduced with permission of Elsevier (12))

The <110>-orientated Ir coating formed by DGP was an interesting phenomenon, which can be related to the density of broken metallic bonds and the relative surface energy of the dominant crystal faces (14). Because the atom density of each crystal face was different, the density of broken metallic bonds varied when different crystal faces were split. The (111) crystal face has the highest atom density and interplanar spacing compared with other crystal faces in the fcc crystal structure (Table I), meaning that the lowest energy is needed to pull the (111) crystal face apart. The density of broken metallic bonds in the (111), (100) and (110) crystal faces were 3.464 a–2, 4 a–2 and 4.242 a–2 (where a denotes the lattice parameter), respectively. If the (111) crystal face energy was assumed to be 1, then the (100) crystal face energy and (110) crystal face energy would be 1.154 and 1.223, respectively.

Table I

Atomic Structure (Hard Sphere Model) and Surface Energy (Shkl) and (hkl) Plane of Iridium with Monoatomic fcc Structure

{hkl } {111} {100} {110}
{hkl } crystal planes

Surface density 2.3/a2 2/a2 1.4/a2
Surface energy Shkl, J m–2 2.59 2.95 3.19

According to the principle of crystal growth, the preferred growth direction usually has the fastest growing speed. From surface energy considerations, the (111) orientation was expected to be preferred in the fcc structure. The surface energies S111, S100 and S110 for fcc Ir were 2.59 J m–2, 2.95 J m–2 and 3.19 J m–2 (15), respectively. The <111> orientation is typical for the growth of fcc metals, since it presents the lowest surface energy and the highest atomic density. Thompson and Carel (16) suggested that energetic constraints such as surface and interface energy minimisation and strain energy minimisation could lead to texture selection. The <111>-textured grain was probably the result of surface and interface energy minimisation during coating formation and growth. The <110>-textured grain was the result of strain energy minimisation during grain growth. The <111>-orientated Ir coating exhibited excellent diffusion barrier properties against oxygen (17). On high-energy face <111> Ir, the adatoms were likely to diffuse quickly to a cluster edge and expand the cluster laterally. On low energy face <110> Ir, diffusion was limited and growth proceeded in the vertical direction. Due to developing a height advantage and the resulting shadowing, more particle flux was intercepted, which led to faster growth of <110> Ir grains (9). At low adatom mobilities, <110>-orientated grains could outgrow <111>-orientated grains. The <110>-orientated Ir coating was rough, with low densities and void formation. Vigorous bombardment from high-energy ions prevents the formation of low energy (100) and (111) crystal faces. Only high-energy (110) crystal faces are formed and grow under these conditions, resulting in the preferred (110) orientation. The glow discharge and the vigorous bombardment therefore make the (100) and (111) diffraction peaks weak. Due to their lower atomic density along the growth direction, (110)-textured grains grew faster than (111)-textured grains with the same atomic deposition rate and finally won the competition (9, 10).

When the deposition of the Ir coating began, the substrate temperature was relatively low, resulting in a low surface mobility of the deposited atoms. <110> grains grew faster than <111> grains when the deposited atoms had a low surface mobility. For fcc metals, incident energetic particles favour the development of <110> over the <111> texture, due to the effect of channeling and thereby less localised radiation damage in the <110> grains. For the DGP process, the Ir layer formed suffered from the vigorous high-energy ion bombardment and sputtered ions. So, a <110>-textured Ir coating was easily obtained by a DGP process.

1.2 Structure Evolution

Polycrystalline coatings are usually formed through the nucleation of isolated crystals on the surface of a substrate. Moreover, the microstructures and properties of the coatings depend on grain nucleation, growth, coalescence and film thickness growth (18–21). Figure 5 shows the scanning electron microscope (SEM) micrographs of a DGP Ir coating on a Nb substrate. The surface of the coating consisted of many small aggregates. There were no microcracks or pores on the coating surface (Figure 5(a)). The chemical composition of the coating consisted only of Ir, as demonstrated by the energy-dispersive X-ray spectroscopy (EDS) pattern. The fracture surface of the coating consisted of columnar grains (Figure 5(b)). The columnar grains extended through the entire coating thickness. The coating thickness was 6–7 μm.

Fig. 5.

SEM micrographs of: (a) surface; and (b) fracture surface of an Ir coating on a Nb substrate (Reproduced with permission of Elsevier (12))

The growth and coalescence processes are part of the coating formation. The islands growth occurred until they touched each other to form a continuous coating. The coalescence process was important for the design of coatings with special properties. After solid-like coalescence of two islands, there might be a grain boundary between them or they might form a boundary-free island. Surface energies and super saturation controlled the coalescence and growth process through material transport by surface and bulk diffusions. From Figures 6(a) to 6(c), the rough surface for the thin coating evolved into a smooth surface for the thick coating. The relatively rough surface consisted of some large aggregated particles (Figure 6(a)). The coating consisted of some hillock-like particles, but the surface was relatively smooth (Figure 6(b)). Compared with Figures 6(a) and 6(b), the structure evolution of the coating was due to the low substrate temperature and the shadowing effect, which resulted in low mobility of the adatoms and the limitation of adatom diffusivity. The high concentration of Ir atoms in the boundary layer led to their agglomeration. The agglomerated clusters and the sputtered Ir particles were deposited to form hillock-like particles (Figure 6(b)). For Figures 6(a) and 6(b), the coating belonged to the growth and coalescence processes. Figure 6(c) shows that the surface of the coating was uniform and dense and the surface was composed of small fine grains. The surface morphology was typical for the resputtering of the coating (Figure 6(d)). The surface was homogeneous and dense. The adsorption and diffusion of the Ir atoms were accelerated due to the high substrate temperature with increasing bias voltage. When the bias voltage (Us) was changed from –100 V to –400 V, the coating became thick. The bias voltage not only affected the substrate temperature, but also the deposition rate of the coating (23). The microstructure of the coating was influenced by its increasing thickness.

Fig. 6.

SEM images of the surface of a DGP Ir coating on Mo substrates: (a) Us = –100 V; (b) Us = –200 V; (c) Us = –300 V; (d) Us = –400 V (Reproduced with permission of Elsevier (22))

The microstructure of the coating could affect its properties and hence its possible applications. The homologous temperature TS/TM, defined as the ratio of the substrate temperature TS to the melting point of thin coating material TM, is one of the main parameters. The substrate temperature plays a key role in determining the adatom surface mobility and the bulk diffusion rates. The structure zone models (21, 24) classified the microstructure of metallic coatings into three different zones as a function of TS/TM:

  • Zone I (TS/TM < 0.3) is the columnar structure with lots of pores and open columnar boundaries. This structure is promoted by a high gas pressure or a low homologous temperature. Shadowing dominates

  • Zone II (0.3 < TS/TM < 0.5) consists of columnar grains separated by distinct and intercrystalline boundaries. The grain sizes increase with TS/TM and may extend through the coating thickness at high TS/TM. Surface diffusion is significant

  • Zone III (0.5 < TS/TM <1) consists of equiaxed grains with a bright surface. The grain diameters increase with TS/TM and the bulk diffusion of adatoms is predominant.

Generally, RFMS and DCMS produce a columnar grained Ir coating. However, a DCMS Ir coating deposited at 25°C had a rough surface, contained irregularities and had a porous columnar structure due to the effect of shadowing (25). A RFMS Ir coating deposited at 800°C had a dense columnar structure. RFMS Ir coatings deposited at less than 100°C show an ‘orange-peel-like’ surface topography which consists of densely packed sub-100 nm grains (26). A MOCVD Ir thin coating presents compact columnar and coarse grains. An electrodeposited Ir coating had large columnar grains as well as large equiaxed grains. PLD Ir films deposited at 300°C and 500°C have a polycrystalline structure with nanometre grain size. A DGP Ir coating deposited at 800–1000°C showed a columnar grain with sub-micrometre sized grains. For the DGP process, the columnar Ir coating should belong to the Zone II structure. The microstructure of Zone II was in agreement with that of the coating. The columnar grain resulted from the deposited atoms having sufficient surface mobility to diffuse and to increase the grain size. However, the Ir coating had many nanovoids at the interface ascribed to low mobility of the deposited atoms and the high deposition rate of the coating.

1.3 Mechanical Properties

The deposition temperature could enhance the grain growth of the Ir coating. However, the grain size influences the mechanical properties of the coating such as its hardness and electrical resistivity. The microhardness of Ir products fabricated by different processes is shown in Table II. According to the Hall-Petch relationship (29, 30), polycrystalline material displays an increase in hardness with decreasing grain size. Vickers hardness (Hv) can be related to the grain size (D) by Equation (ii):


(ii)
Table II

Hardness of Iridium Fabricated by Different Processing Methods (26–28)

Process method Magnetron Sputtering DGP Electrodeposition CVD Casting
Hardness, GPa 21 9.5 4.42 2.98 2.10-2.4
Grain size Nanometre Sub-micrometre Micrometre Micrometre Millimetre

where H0 and kH are constants. The highest hardness was found in the magnetron sputtered Ir films, owing to their fine crystal grain size. Kuppusami et al. (31) found that the hardness of an Ir coating formed by DCMS was ~22.5 GPa. Hagen et al. (32) reported the mechanical properties of Ir coatings with a titanium sub-layer deposited by RFMS. The obtained values agreed well with data from Kuppusami, Murakami and Ohmura (31). The heat-treated Ir coatings after indentation showed low hardness values compared with the as-deposited Ir coating (33, 34).

The electrical resistivity of the coatings mainly depends on the grain size (35). El Khakani, Chaker and Drogoff (36) pointed out that the resistivity of an Ir film was highly influenced by the energetic particle bombardment conditions during film growth. The resistivity of Ir films measured by four-point resistivity tests decreased with increasing substrate temperature (37), which was attributed to the grain size. When the RFMS deposition rate of an Ir film was increased, its crystalline quality decreased while the resistivity increased (38).

For different deposition processes and conditions, the deposited Ir coating has different growth modes and textures, leading to differences in the microstructure. The microstructure and growth modes of coatings play an important role in their physical and mechanical properties. The adhesive force of the coating on the substrate is a key value and affects the application of the coating in extreme environments. A general rule of thumb says (39) that a critical load of 30 N in scratch testing with a Rockwell C diamond tip is generally sufficient for tooling applications. Chen et al. (40) reported that the adhesive force of DGP Ir coatings on refractory metals were measured by a scratch tester at about 50–60 N. Thus the DGP Ir coating adhered well to the refractory metals, which was due to the strong metallurgical bond and the formation of a transition layer at the interface. The adhesion forces of a RFMS Ir coating (26) on stainless steel, silica (SiO2) and tungsten carbide (WC) substrates were 100 mN, 125 mN and 165 mN, respectively. The adhesive forces of MOCVD Ir coatings (35) on TiO2/SiO2/Si, SiO2/Si and Si substrates were 15 N, 9 N and 5 N, respectively, probably due to the fact that the lattice constant of Ir is closer to that of TiO2 than to those of SiO2 and Si. Good adhesion between an electrodeposited Ir coating and a Re substrate was due to the excellent match between their thermal expansions and good solid solubility between the two elements (27).

2. Growth Mechanism of Iridium Coatings

For a PVD system, if the species changes from vapour to solid at a pressure p, a free enthalpy change is involved (41, 42), Equation (iii):


(iii)

where p0 is the equilibrium vapour pressure and η is the degree of species’ saturation. The formation of the coating by a vapour deposition process is characterised by the formation and growth of nuclei. Depending on the interaction energies of the substrate atoms and the sputtered atoms, the following three growth modes can occur (43, 44):

  • (a) Layer by layer or Frank–van der Merwe (FM)

  • (b) Island or Volmer–Weber (VW)

  • (c) Layer plus island or Stranski–Krastanov (SK).

The growth modes of the coating were in terms of surface energies (γ). Wetting angle (θ) of a nucleus on a substrate surface was described by Young’s equation (45), Equation (iv):


(iv)

where γS is the surface energy of the substrate, γF is the surface energy of the coating and γS/F was the interface energy. VW growth (θ > 0) required that γS < γS/F + γF, whereas FM growth (θ = 0) required that γS >γS/F + γF. SK growth occurred because the interface energy increased with film thickness. Gerfin et al. (46) suggested that the growth mechanism of an Ir film by MOCVD followed the VW mode. The growth mode of a DGP Ir coating on a graphite substrate presented a conical grain structure formed via the SK mode due to poor wetting between Ir and carbon (47). The interface bond between the Ir coating and the graphite substrate was the only mechanical locking effect. However, the interface bond between an Ir coating and a metallic substrate is an ionic bond, which could help to improve the adhesion of the coating to the substrate. In Figure 7, the growth mode of a DGP Ir coating was VW, resulting in a columnar grain structure. Usually continued growth of the layer, after initial island (VW) growth, occurred by columnar growth. Therefore, the coating growth probably began with island clusters which grew three dimensionally.

Fig. 7.

Thin film growth process

The growth mode of the coating was controlled not only by interface energies but also by saturation. The degree of species’ saturation and the growth kinetics of the coating were responsible for the difference in the coating structure. The degree of species’ saturation played a decisive role in nucleation on the substrate, which can be explained as follows (42). For three-dimensional (3D) nucleation, a number N of atoms formed the nucleus, Equation (v):


(v)

where Y is the contribution of the interface energy. 3D nucleation for the coating required supersaturation. The faster growth rate of the coating deposited at a stable substrate temperature, resulting in supersaturation conditions, might cause a SK growth mode to be followed (42). If the deposition continues, the nucleation centres become more numerous and then coalescence of two-dimensional (2D) islands occurs. As the size of the 2D islands increases, the coverage of the substrate surface is enlarged. The growth process of the film was controlled by migration, re-evaporation, nucleation, coalescence and thickness growth (see Figure 7).

2.1 Metal-Organic Chemical Vapour Deposition of Iridium Coating

The growth mechanism of Ir coatings formed using MOCVD depends not only on the deposition parameters and the substrates, but precursor chemistry also plays a key role. Sun et al. (48) investigated the effects of both oxygen and substrate on Ir film growth by MOCVD. Without oxygen, the Ir film contained carbon. Oxygen removed carbon and also prevented carbon incorporation from other reactive gas components. Oxygen could also control the deposition rate of the film and had a significant impact on the morphology of the film. Garcia and Goto (49), and Goto, Vargas and Hirai (50) investigated the MOCVD of Ir coating using Ir(acac)3 complex with and without the addition of oxygen, and found that the Ir coating consisted of Ir clusters with diameter 1–4 nm surrounded by amorphous carbon. The carbon composition was about 20 wt%. The addition of oxygen was effective to suppress carbon incorporation into the Ir coating. A highly pure Ir coating containing less than 1 wt% carbon and further oxygen could be obtained by controlling the oxygen flow rate to avoid the formation of IrO2. A change of deposition rate at a turning point around 600°C was observed, corresponding to the change of activation energy from 160 kJ mol–1 to 24 kJ mol–1. Such changes are normally linked to the change of the CVD rate controlling step from a surface chemical reaction to a gaseous diffusion process with increasing temperature.

Semyannikov and coworkers (51) conducted a set of studies on the deposition of Ir using Ir(CO)2(acac). The compound decomposed in vacuum and in hydrogen with the formation of the following major products: acetylacetone (H(acac)), carbon monoxide, carbon dioxide and acetyl (C2H3O). Thermal decomposition of Ir(CO)2(acac) in the presence of oxygen was accompanied by the formation of water vapour. The appearance of columnar structures in the Ir coating was observed on substrates of different natures (metals and oxides) (52). At the initial stages of growth, a coating with compact structure was deposited, then the growth mechanism changed and a layer with a columnar structure was formed. Furthermore, a thin layer with high carbon concentration was formed on the growing metal surface, which led to structural changes as the process continued and resulted in the formation of coatings with a non-uniform layered structure. The growth rate of an Ir film on a titanium carbon nitride (TiCN) substrate was significantly higher than that on a SiO2 substrate. The Ir film on SiO2 was rough due to an initial 3D growth mode on isolated islands.

Vargas et al. (53) showed that an Ir film could be epitaxially grown by MOCVD on sapphire single crystal substrates using Ir-acetylacetonate precursors. Gerfin et al. (46) reported that the growth rate of a MOCVD Ir film depended on the substrate temperature. Two growth regimes could be distinguished, i.e. the kinetically controlled (Tsub<350°C) and the mass-flow-controlled region. The Ir film consisted of islands, thus the growth mechanism of the Ir film followed the VW mode. Gelfond and coworkers (54) studied an Ir film deposited with Ir(acac)3 on SiO2 and found a transition layer of IrSiO. The Ir films consisted of randomly orientated, closely spaced particles of different shape and size.

2.2 Atomic Layer Deposition of Iridium Coating

Song et al. (55) found that a hybrid ALD Ir film followed a 3D island growth pattern. Hybrid ALD could increase the rate of formation of the Ir seed layer on the substrate surface. Knapas and Ritala (56) reported the reaction mechanism in two ALD processes using Ir(acac)3 as a precursor: Ir(acac)3-O2 and Ir(acac)3-O3-H2 processes at 300°C and 195°C, respectively. For the first process, the reaction of Ir(acac)3 occurred upon its adsorption, as the adsorbed oxygen atoms combusted part of the ligands during the Ir(acac)3 pulse. For the second process, the adsorption of Ir(acac)3 was molecular on the film surface. The formation mechanism of the ALD Ir film was self-limiting (57). Growth proceeded in a FM mode in one monolayer or a fraction of a monolayer. The Ir film was deposited during one growth cycle. When the precursor dose was high enough to saturate the precursor adsorption and possibly the surface reactions, the film growth became self-limiting. The growth rate was constant over the whole surface. Due to the self-limiting growth mechanism, the ALD Ir films had conformality and uniformity over a large area. The growth and mechanism of an ALD Ir film from Ir(acac)3 followed two steps, a nucleation step and steady-state growth. The formation of Ir nuclei on the substrate surface was crucial for the growth of the Ir film. The common oxygen-based Ir ALD processes rely on the Ir surface to catalytically dissociate molecular oxygen to form reactive atomic oxygen for film growth. The first metallic nuclei were most likely formed by some minor decomposition of the Ir complex precursor. Then these first nuclei catalysed the growth of the Ir film.

Christensen and Elam (58) studied the mechanism of ALD Ir film formation using an Ir(acac)3-O2 mixture and proposed a two-step mechanism for Ir ALD. The first step was the reaction of an Ir(acac)3 precursor with adsorbed oxygen species on the ALD Ir film surface, releasing one or two of the acetylacetonate ligands through ligand exchange and ~0.1 ligand through combustion. The second step was that the remaining acetylacetonate ligands were released by combustion during the subsequent oxygen exposure and the film surface was repopulated with oxygen species. Knapas and Ritala (56) presented the growth mechanism of an ALD Ir film from an Ir(acac)3 precursor. In contrast with previous work (58), three chemical systems with the corresponding chemical reactions based on the Ir(acac)3 precursor were synthesised, Equations (vi) to (viii):

(a) Ir(acac)3-O2 system at 225°C (Ir deposition)


(vi)

(b) Ir(acac)3-O3-H2 system at 195°C (Ir deposition)


(vii)

(c) Ir(acac)3-O3 system at 195°C (IrO2 deposition)


(viii)

2.3 Physical Vapour Deposition of Iridium Coating

Chang, Wei and Chen (59) used the embedded atom method (EAM) potential to study the structures, adsorption energies, binding energies, migration paths and energy barriers of Ir adatoms and small clusters on Ir (100), (110) and (111) surfaces and found that the barrier for single adatom diffusion was lowest on the (111) surface, higher on the (110) surface, and highest on the (100) surface. The exchange mechanisms of adatom diffusion on (100) and (110) surfaces were energetically favoured. On all three Ir surfaces, Ir2 dimers with nearest neighbour spacing were the most stable. Hörmann et al. (5) studied the growth and structural properties of Ir films deposited by electron beam evaporation on the substrate surface. At 950°C, the growth of the Ir film proceeded via 3D nucleation, coalescence of the isolated islands and subsequent layer-by-layer growth, finally resulting in flat Ir films. Wessling et al. (9, 10) studied the effect of process parameters such as power bias voltage, pressure and target-to-substrate distance on the structure of the Ir film. Growth of porous Ir films with high specific surface area delivered low surface mobility of the deposited atoms.

2.4 Electrodeposition of Iridium Coatings

The structure evolution of the Ir coating with increasing current density is related to crystallisation kinetics and mass transfer during the electrodeposition process. Toenshoff et al. (60) electrodeposited an Ir coating from a chloride electrolyte at 570°C and a deposition rate of ~20 μm h–1. The electrodeposited Ir coating had a columnar structure and high ductility. Zhu et al. (11) studied the effects of current density and temperature on the structure of an Ir coating. When the deposition temperature was increased, the nucleation rate decreased but the growth rate increased, thus the macroscopic compactness of the coating decreased with increasing grain size. Huang et al. (34) reported the laminar structure of an Ir coating which resulted from using high and low current densities alternately, and suggested a growth mechanism for this coating. A two-segment cathodic current was used, and coarse Ir grains were found to be deposited on the top side under the lower current density, while ultra-fine Ir grains were applied on the bottom side by the higher current density segment. Higher current density would result in larger over-potential in the electrode, promoting nucleation of Ir grains in the coating. Conversely, a small over-potential in the electrode promoted growth of the Ir grains, which resulted in a thick and columnar grain structure in the Ir coating. Therefore, a laminar columnar structure of Ir coating was obtained by the repeated two-segment cathodic current. The thickness of each layer was determined by the quantum of electrons transferring at the cathode in each two-segment current cycle.

2.5 Double Glow Plasma Deposition of Iridium Coatings

For the DGP process, the bias voltage not only affected the deposition rate but also the substrate temperature. Under the same deposition conditions, the Ir coating exhibited (220) orientation regardless of the substrate. The substrates affected the growth rate of the Ir coating as well as the surface topography. The structure of the Ir coating was influenced by bias voltage, gas pressure and the effect of the substrate. An Ir coating formed by DGP on a Ti substrate was composed of irregular compacted columnar grains with many nanovoids at the interface (see Figure 8). The DGP Ir coating exhibited a <110> texture on different substrates due to the formation of initial nuclei with preferred growth on the substrate surface and a channelling effect during the deposition process.

Fig. 8.

Field emission scanning electron microscope (FESEM) image of the interface between an Ir coating and a Ti substrate (Reproduced with permission (61))

When only the substrate bias voltage was applied, the whole surface of the substrate was sputtered out (Figure 9(a)). In this case, the glow discharge caused the substrate to be bombarded with positive ions of the inert atmosphere and such ion bombardment heated the surface of the substrate to an elevated temperature. Another effect of the glow discharge was to clean and activate the substrate surface. When the target bias voltage was applied, the substrate and the target were co-sputtered out (Figure 9(b)). This resulted in the formation of a transition zone between the coating and the substrate (Figure 9(c)). Because of the large difference between the target and the substrate bias voltages, the sputtered substrate and coating atoms were formed on the substrate surface with time and temperature. But the target bias voltage was higher than substrate bias voltage, resulting in a high sputtering rate for the target. The substrate atoms were suppressed by the intense coating atoms and formed a mixed boundary layer on the substrate surface. A pure Ir coating was obtained (Figure 9(d)).

Fig. 9.

Overview of grain structure evolution during deposition of a polycrystalline Ir coating by DGP: (a) sputtered substrate atoms; (b) co-sputtered coating and substrate atoms, nucleation process; (c) growth and coalescence to form a continuous film; (d) thickness growth (Reproduced with permission (61))

The growth mechanism of the Ir coating by DGP depended on a kinetic adsorption and diffusion process with nucleation, coalescence and thickness growth. At the beginning of the deposition process, the growth of the coating was mainly controlled by the nucleation rate (61). Due to the low substrate temperature resulting in low mobility of the deposited atoms, some nanovoids were present at the interface. During the deposition process, the substrate temperature was increased and then kept steady. After this, the growth of the coating was governed by the growth rate. A higher substrate temperature provided enough energy for surface mobility of adatoms.

2.6 Pulsed Laser Deposition of Iridium Coating

Aziz (62) reviewed the fundamentals of growth morphology evolution in PLD in two prototypical growth modes: metal-on-insulator island growth and semiconductor homoepitaxy. In metal-on-insulator film growth, the same morphology progression occurred: equiaxed islands, elongated islands, percolating metal film and hole filling. The kinetic freezing model, involving the competition between island-island coalescence and deposition driven island-island impingement, explained the morphological transitions in both thermal deposition and PLD. For low temperatures, the high island density of PLD dominated the morphology evolution and PLD films reached percolation sooner than thermally deposited films. As the temperature was increased, the PLD percolation thickness approached and then exceeded the thermal percolation thickness. It indicated the increasing importance of an energetic effect. This effect appeared to be kinetic energy induced adatom-vacancy pair creation, which had the net effect of moving atoms upward, resulting in a vertical shape transition of the islands. Taller islands coalesced more rapidly, thereby delaying the point of the elongation transition, where coalescence was overwhelmed by impingement.

Chen et al. (42, 63) studied the epitaxial growth of Pt and Ir films by PLD grown on MgO-buffered Si(100) substrate. It was found that both Pt and Ir films featured remarkable atomic-scale smooth surfaces and had the same epitaxial relationship with substrates. Different from the non-compact surface morphology of Pt film, the morphology of Ir film offered a rectangular grain shape and the grains were arrayed regularly and compactly. The difference in the surface morphology of both electrode films was due to the degree of species’ saturation. Hence, the fast growth rate of Pt film resulted in a supersaturation condition, which might cause a SK growth mode rather than a FM mode due to its 3D nucleation. In contrast, the slow growth rate of the Ir film resulted in a subsaturation condition, which led to the FM mode because of its 2D nucleation.

3. Micropore Formation Mechanism in Iridium Coatings after High-Temperature Treatment

In our previous publications (40, 64–72), the thermal stability and oxidation resistance of DGP Ir coatings on refractory materials were studied. After high-temperature treatment at 1400°C under an inert atmosphere, micropores and microbubbles appeared on the coating surface (68). After high-temperature treatment at 2000°C under an oxyacetylene flame, micropores were still formed on the Ir coating surface and these could provide pathways for oxygen to attack the substrates (65, 67, 70, 71, 73). To date, although Reed, Biaglow and Schneider (74) claimed that CVD is the only established process for the fabrication of Ir-coated Re combustion chambers operated at 1800~2200°C, some publications (12, 27, 68, 75, 76) suggest that micropores are created on the surface after high-temperature treatment.

Micropores diffused outward in a MOCVD Ir coating after heat treatment (77), and appeared in an electrodeposited Ir coating on a Re substrate at the interface of the coating and substrate after thermal resistance testing (27). Huang et al. (78) ascertained that micropores in an electrodeposited Ir coating on a Re substrate were found at the grain boundaries and were concentrated to penetrating holes with the growth of Ir grains, which resulted in disastrous failure of the coating. The micropores in the Ir coating moved outwards after thermal cycling due to the underdense as-deposited coating (75, 79). Hu (80) suggested that micropores appeared on the surface of an Ir coating after high-temperature oxidation. Micropores were formed in Ir coatings formed by various deposition processes after high-temperature treatment in oxygen or inert environments.

A polycrystalline Ir coating consisted of a large number of individual grains separated by boundaries. A grain boundary is the interface between two grains in a polycrystalline material. Corrosion and oxidation occurred preferentially at the grain boundaries. The mean misorientation angles of grains on the surface and cross-section of the coating were 38.6° and 45.6°, respectively (Figure 10). The defect density in the coating was shown to be higher because a high-angle boundary was imperfectly packed compared to the normal lattice.

Fig. 10.

Misorientation angle distribution images of grains on: (a) surface; (b) cross-section of the as-deposited Ir coating

After heat treatment, equiaxed grains with indistinct grain boundaries were formed. Some large grains and many pores appeared on the coating as seen in Figure 11(a). There are two main reasons for this pore formation. On the one hand, it may be attributed to the recrystallisation of the Ir crystals (72). After heat treatment, equiaxed grains replace the columnar grains in the coating. The removal of Ir atoms during heat treatment results in the creation of vacancies, which accumulate to form micropores. A moving grain boundary might also be expected to be a suitable sink for vacancies to form micropores. On the other hand, the pores could result from the aggregation of inner micropores between grain boundaries. In Figure 11(b), some micropores and microbubbles appear at the surface and at the interface. The size and quantity of micropores are distributed in a gradient along the cross-section of the coating. This may be because the grain size became larger after heat treatment, the micropores in the grain boundaries diffused outward and became larger in the coating. The micropore formation at the interface between the coating and the substrate was attributed to the Kirkendall effect (80), owing to different diffusion rates for Ir and Mo at the interface between the coating and the substrate. In Figure 12(a), the as-prepared coating was composed only of Ir. No new phase was formed. A new Ir21.5Mo8.5 phase was formed in the Ir coating after heat treatment (Figure 12(b)), which may be due to a reaction at the Mo-Ir interface. The diffraction peaks of the Ir coating became sharper and narrower, indicating recrystallisation and grain growth in the coating. According to the Mo-Ir phase diagram (81), four intermetallic phases (i.e. IrMo, Ir3Mo, ɛ, IrMo3) can exist at 1200–1475°C, with Ir able to dissolve over 20 at% Mo at 1400°C. The intermetallic compound Ir21.5Mo8.5 (71.7 at% Ir) belonged to the Ir3Mo phase with hexagonal D019 structure type. A solid solution of Mo in Ir at 1400°C is realisable. Mo reacted with Ir alloy to produce two distinct intermetallic layers, IrMo and Ir3Mo, at high temperature (82, 83). The Ir3Mo phase was found at high temperature, as reported by Chen et al. (84). In Figure 12(b), the IrMo phase was not formed, possibly because the heat treatment temperature and time were insufficient.

Fig. 11.

SEM images of: (a) surface; and (b) cross-section of the Ir coating after heat treatment at 1400°C for 1.5 h in an argon atmosphere

Fig. 12.

XRD patterns of the Ir coating on a Mo substrate: (a) before; (b) after heat treatment at 1400°C for 1.5 h in an argon atmosphere

The DGP monolayer Ir coatings were ablated by an oxyacetylene torch with a flame temperature of ~2000°C for 35 s. The surface of the as-ablated coating was composed of equiaxed grains with distinct grain boundaries (Figure 13(a)). Some micropores appear on the coating surface. In Figure 13(b), some micropores appear as fractures on the external surface of the coating. These micropores could provide tunnels for oxygen diffusion, resulting in the oxidation consumption of the substrate, which would weaken the Ir-substrate bond and eventually cause the failure of the coating. Under inert gas protection, formation of micropores could occur for the following reasons. First, the micropores may partly result from grain growth of the Ir at high temperature and represent pre-existing intercolumnar defects arising from the original as-deposited growth habit of the Ir coating. Second, they may result from the aggregation of inner defects such as vacancies, voids or grain boundaries. Subsequently, these inner defects in the as-deposited coating aggregate together resulting in the formation of micropores. However, different conditions apply for ablated Ir coatings in an ablating or oxidising environment. Ablation processing is an erosive phenomenon, which can be interpreted as the combined effect of thermochemical ablation with thermophysical and thermomechanical attacks from high temperature, pressure and velocity of the combustion flame.

Fig. 13.

SEM images of the as-ablated Ir coating on Mo substrate: (a) coating surface; (b) fracture surface

Since Ir did not form a condensed oxide, micropores appeared on the surface under severe conditions. The surface of the ablated monolayer coating presented imperfections including pores, bulges and cracks after exposure to the flame; however, the Ir coating kept sufficient adhesion to limit the weight loss of the Ir-coated Mo specimen. The deposited monolayer Ir coating had a preferred (220) orientation. A previous publication (85) showed that a <110>-textured Ir coating was easily obtained by DGP. The <110>-textured Ir coating was not dense, however a <111>-textured Ir coating was dense (101). Some pinholes were present in the monolayer Ir coating; however, the multilayer Ir coating could support high-temperature ablation. After ablation, some micropores appeared on the surface (See Figure 14(a)). The surface porosities of multilayer and monolayer after ablation were about 0.55% and 2%, respectively. Therefore, the porosity on the surface of the multilayer coating was less than that of the monolayer. In Figure 14(b), the three layers appear on the fracture surface of the coating. The multilayer Ir coating was debonded from the substrate by the released gas. The grains became large aggregates, then grain recrystallisation and growth occurred. The interface between adjacent layers became relatively unclear due to the interdiffusion effect. The diameter of micropores was about 0.1–0.6 μm. Many pinholes in the Ir coating were aggregated to form micropores and diffused outward, but were reduced after heat treatment (75).

Fig. 14.

SEM images of: (a) surface; (b) fracture surface of the multilayer Ir coating on Nb substrate after ablation at 2000°C for 40 s

In Figure 15 the multilayer Ir coating may be thick enough to protect the substrate from oxidation at high temperature. The multilayer Ir coating could inhibit the formation of cracks or micropores extended through the thickness of the coating (Figure 15(b)). Although the surface of the deposited Ir coating was dense, Ir was not expected to form a protective oxide coating because of the formation of the volatile IrO3. The evaporation of the Ir and its oxide resulted in the formation of micropores. These pinholes or micropores on the surface of the coating acted as pathways for oxygen to penetrate the coating. Once oxygen reached a reaction site on the Nb substrate, it quickly reacted with Nb to form an oxide. Tuffias, Melden and Harding (86) reported that the outward diffusion of Nb was the major factor contributing to interdiffusion. This diffusion resulted in void formation which compromised the structural integrity of the coating. Because the Nb substrate ignited with oxygen in the quickly flowing flame, masses of gas were released from the surface of the as-coated specimen. The released gases severely influenced the thermal stability of the multilayer Ir coating (Figure 15(c)). An intermediate layer of Re or W could be employed to prevent interdiffusion of the coating and the substrate (86). The underdense multilayer Ir coating and the interdiffusion between the coating and the substrate resulted in the failure of the multilayer Ir coating.

Fig. 15.

Schematic diagram of failure of multilayer Ir coating during the ablation process: (a) coating structure; (b) the duration of the process; (c) coating failure (Reproduced with permission of Springer (70))

At about 1100°C, Ir does not form a condensed oxide due to the formation of the volatile oxide (IrO3), Figure 16. The as-deposited coating was composed of columnar grains with high-angle boundaries. Numerous grain boundaries gave preferred sites for the onset of oxidation during high-temperature treatment. Therefore, the Ir grain boundaries in the coating could easily react with oxygen and formed gaseous IrO3. Although the quantity of volatile Ir oxide was low during the temperature rise and cooling periods, the grain boundaries provided tunnels for oxygen diffusion and oxidation of the Ir grains then occurred at the interface. The gaseous IrO3 was formed at the grain boundaries, resulting in the formation of micropores. For the high-angle boundary of the columnar grains (Figure 10), the coating had a higher energy state due to a combination of high interfacial and surface energies. A columnar-to-equiaxed grain transition occurred, resulting in a more stable atomic arrangement, which minimised the total energy in the coating. The coating configuration must have sufficient energy to overcome an activation energy barrier. During high-temperature treatment, recrystallisation occurred. There was an energetic driving force for grain boundary formation, which eliminated the energies of the free surfaces of the two contacting columnar grains in exchange for the lower energy of a new grain boundary. After high-temperature treatment, the micropores in the coating by rapid surface or bulk diffusion mechanisms reduced the high surface energy associated with the coating configuration (79). The micropores in the coating diffused outward after high-temperature treatment due to the extrusion effect under grain growth. Although the surface of a pure Ir coating produced some pinholes and micropores after high-temperature treatment, ceramic coatings may protect the Ir coating to some extent. An as-coated specimen was laid on an Al2O3 firebrick. Because the ablation temperature was higher than the melting point of Al2O3, one small portion of the firebrick located at the surface of the Ir coated specimen was melted. The melted Al2O3 flowed on the surface of the coating which was then partially covered by melted Al2O3 (see Figure 17). The surface of the Ir coating appeared as many volcano-like micropores. The coating seemed to be melted and many micropores were found. However, the remaining surface was dense, because the melted Al2O3 filled in the pores on the surface of the coating. Al2O3 could therefore be used as an outer layer for oxidation protection because of its attractive properties such as low vapour pressure, low oxygen diffusion coefficient, good corrosion resistance and oxidation stability (87, 88). Therefore, Al2O3 as a top layer could seal micropores on the surface of the coating and inhibit the evaporation of the Ir. Even if the ceramic oxide layer cracked during rapid cooling, the evaporation of Ir would be minimised and orders-of-magnitude less oxygen would be permitted to attack the coating.

Fig. 16.

Micropore formation mechanism in an Ir coating after high-temperature treatment in an oxidising environment (Reproduced with permission of Elsevier (72))

Fig. 17.

SEM images of the surface of multilayer DGP Ir coating after oxidising ablation at 2200°C for 90 s (Reproduced with permission of Elsevier (71))

4. Conclusion and Outlook

Re-Ir material is currently of particular interest for engineering applications. In China, a short nozzle of Re-Ir has been tested and the temperature of combustion chamber wall reached 2090°C for 300 s without failure (89). There are a number of processes available to produce thin or thick Ir coatings on various substrates (as shown in Part I (1)), with varying results as to the quality of the as-deposited coatings. The texture of the coating affects its microstructure and mechanical properties. The texture evolution of the coating is controlled by the deposition parameters and conditions. A <111>-textured Ir coating is dense, which could provide an effective diffusion barrier against oxygen. The growth mechanism of Ir coatings produced using different deposition processes is different, and the process should therefore be selected according to the requirements of the intended application. For vapour and solution deposition processes, the formation of the coatings is controlled by migration, nucleation, coalescence, growth and thickness growth. After high-temperature treatment, micropores were formed on the surface and the fracture surface of the Ir coating. Micropores may supply tunnels for oxygen to attack the substrate. Grain boundaries may act as sources and sinks for vacancies in a high-temperature environment. Inner defects such as grain boundaries and vacancies, accumulated in the coating during high-temperature treatment, result in the formation of micropores. Evaporation of volatile Ir oxide occurred preferentially at the grain boundaries, again resulting in micropore formation. The columnar-to-equiaxed grain transition resulted in a more stable microcrystalline arrangement. The micropores in the Ir coating diffused outward after high-temperature treatment due to the extrusion effect under grain growth. Therefore, a dense and thick Ir coating could improve the life of engineering components in very high-temperature environments. Furthermore, a ceramic oxide layer on the surface of an Ir coating may minimise the evaporation and oxidation of the coating.

Acknowledgements

This work has been supported by the National Natural Science Foundation of China (Grant Number: 50872055/E020703) and the Natural Science Foundation of Jiangsu Province (Grant Number: BK20150260). The authors wish to thank the referees for their helpful suggestions, and Editors Ms Sara Coles and Ming Chung for the editing.

References

  1. 1.
    W.-P. Wu and Z.-F. Chen, Johnson Matthey Technol. Rev., 2017, 61, (1), 16 LINK https://www.technology.matthey.com/article/61/1/16-28/
  2. 2.
    J.-P. Zhao, X. Wang, Z.-Y. Chen, S.-Q. Yang, T.-S. Shi and X.-H. Liu, J. Phys. D: Appl. Phys., 1997, 30, (1), 5 LINK http://dx.doi.org/10.1088/0022-3727/30/1/002
  3. 3.
    S. Eroglu and B. Gallois, J. Phys. IV France, 1993, 3, C3-177 LINK http://dx.doi.org/10.1051/jp4:1993322
  4. 4.
    H. Osamura, ‘Development of Long Life and High Ignitability Iridium Spark Plug’, Seoul 2000 FISITA World Automotive Congress, Seoul, South Korea, 12th–15th June, 2000, F2000A144 LINK http://210.101.116.115/fisita/pdf/A144.pdf
  5. 5.
    F. Hörmann, H.-Y. Peng, Th. Bauer, Q. Li, M. Schreck, Y. Lifshitz, S.-T. Lee and B. Stritzker, Surf. Sci., 2002, 513, (3), 525 LINK http://dx.doi.org/10.1016/S0039-6028(02)01852-6
  6. 6.
    H. Murakami, T. Yano and S. Sodeoka, Mater. Trans., 2004, 45, (9), 2886 LINK http://dx.doi.org/10.2320/matertrans.45.2886
  7. 7.
    M. Hecq and A. Hecq, J. Vac. Sci. Technol., 1981, 18, (2), 219 LINK http://dx.doi.org/10.1116/1.570727
  8. 8.
    D.-Y. Park, D.-S. Lee, H.-J. Woo, D.-I. Chun and E.-J. Yoon, Tong Yang Cement Corp, US Patent 6,025,205; 2000
  9. 9.
    B. Wessling, W. Mokwa and U. Schnakenberg, J. Electrochem. Soc., 2008, 155, (5), F61 LINK http://dx.doi.org/10.1149/1.2844805
  10. 10.
    B. Wessling, D. Lüsebrink, W. Mokwa and U. Schnakenberg, J. Electrochem. Soc., 2008, 155, (5), F66 LINK http://dx.doi.org/10.1149/1.2844818
  11. 11.
    L. Zhu, S. Bai, H. Zhang and Y. Ye, Appl. Surf. Sci., 2013, 265, 537 LINK http://dx.doi.org/10.1016/j.apsusc.2012.11.041
  12. 12.
    W.-P. Wu, Z.-F. Chen, X.-W. Cheng and Y.-W. Wang, Nucl. Instr. Meth. Phys. Res. Sect. B: Beam Int. Mater. Atoms, 2013, 307, 315 LINK http://dx.doi.org/10.1016/j.nimb.2012.12.069
  13. 13.
    L.-B. Wang, Z.-F. Chen, P.-Z. Zhang, W.-P. Wu and Y. Zhang, J. Coat. Technol. Res., 2009, 6, (4), 517 LINK http://dx.doi.org/10.1007/s11998-008-9123-7
  14. 14.
    Z. Zhao, C. Wang, M. Li, L. Wang and L. Kong, Appl. Surf. Sci., 2006, 252, (12), 4257 LINK http://dx.doi.org/10.1016/j.apsusc.2005.07.005
  15. 15.
    V. G. Beshenkov, L. A. Fomin, D. V. Irzhak, V. A. Marchenko, V. I. Nikolaichik and A. G. Znamenskii, Thin Solid Films, 2012, 520, (23), 6888 LINK http://dx.doi.org/10.1016/j.tsf.2012.07.033
  16. 16.
    C. V. Thompson and R. Carel, Mater. Sci. Eng. B, 1995, 32, (3), 211 LINK http://dx.doi.org/10.1016/0921-5107(95)03011-5
  17. 17.
    J. Goswami, C.-G. Wang, P. Majhi, Y.-W. Shin and S. K. Dey, J. Mater. Res., 2001, 16, (8), 2192 LINK http://dx.doi.org/10.1557/JMR.2001.0300
  18. 18.
    C. V. Thompson, Ann. Rev. Mater. Sci., 2000, 30, 159 LINK http://dx.doi.org/10.1146/annurev.matsci.30.1.159
  19. 19.
    C. V. Thompson, Ann. Rev. Mater. Sci., 1990, 20, 245 LINK http://dx.doi.org/10.1146/annurev.ms.20.080190.001333
  20. 20.
    I. V. Markov, “Crystal Growth for Beginners: Fundamentals of Nucleation, Crystal Growth and Epitaxy”, 2nd Edn., World Scientific Publishing Co Pte Ltd, Singapore, 2003
  21. 21.
    C. R. M. Grovenor, H. T. G. Hentzell and D. A. Smith, Acta Metall., 1984, 32, (5), 773 LINK http://dx.doi.org/10.1016/0001-6160(84)90150-0
  22. 22.
    W.-P. Wu, Z.-F. Chen, X. Lin, B.-B. Li and X.-N. Cong, Vacuum, 2011, 86, (4), 429 LINK http://dx.doi.org/10.1016/j.vacuum.2011.09.003
  23. 23.
    K. Mumtaz, J. Echigoya, T. Hirai and Y. Shindo, Mater. Sci. Eng.: A, 1993, 167, (1–2), 187 LINK http://dx.doi.org/10.1016/0921-5093(93)90353-G
  24. 24.
    H. Levinstein, J. Appl. Phys., 1949, 20, (4), 306 LINK http://dx.doi.org/10.1063/1.1698362
  25. 25.
    S. M. Sabol, B. T. Randall, J. D. Edington, C. J. Larkin and B. J. Close, “Barrier Coatings for Refractory Metals and Superalloys”, B-MT-(SPME)-25, TRN: US0603658, Bettis Atomic Power Laboratory (BAPL), Pennsylvania, USA, 2006, pp. 1–28 LINK http://dx.doi.org/10.2172/884669
  26. 26.
    H.-U. Kim, D.-H. Cha, H.-J. Kim and J.-H. Kim, Int. J. Prec. Eng. Manuf., 2009, 10, (3), 19 LINK http://dx.doi.org/10.1007/s12541-009-0042-z
  27. 27.
    L.-A. Zhu, S.-X. Bai and H. Zhang, Surf. Coat. Technol., 2011, 206, (6), 1351 LINK http://dx.doi.org/10.1016/j.surfcoat.2011.08.058
  28. 28.
    W.-P. Wu, X. Lin, Z.-F. Chen, Z. Chen, X.-N. Cong, T.-Z. Xu and J.-L. Qiu, Plasma Chem. Plasma Proc., 2011, 31, (3), 465 LINK http://dx.doi.org/10.1007/s11090-011-9293-4
  29. 29.
    S.-C. Tjong and H. Chen, Mater. Sci. Eng.: R: Reports, 2004, 45, (1–2), 1 LINK http://dx.doi.org/10.1016/j.mser.2004.07.001
  30. 30.
    T. Hanamura and H. Qiu, ‘Ultra-Fine Grained Steel: Relationship Between Grain Size and Tensile Properties’, in “Analysis of Fracture Toughness Mechanism in Ultra-Fine-Grained Steels”, NIMS Monographs, Springer, Japan, 2014, pp. 9–25 LINK http://dx.doi.org/10.1007/978-4-431-54499-9_2
  31. 31.
    P. Kuppusami, H. Murakami and T. Ohmura, J. Vac. Sci. Technol. A, 2004, 22, (4), 1208 LINK http://dx.doi.org/10.1116/1.1763913
  32. 32.
    J. Hagen, F. Burmeister, A. Fromm, P. Manns and G. Kleer, Plasma Process. Polym., 2009, 6, (S1), 678 LINK http://dx.doi.org/10.1002/ppap.200931701
  33. 33.
    H.-J. Li, H. Xue, Q.-G. Fu, Y.-L. Zhang, X.-H. Shi and K.-Z. Li, J. Inorg. Mater., 2010, 25, 337 LINK http://dx.doi.org/10.3724/SP.J.1077.2010.00337
  34. 34.
    Y.-L. Huang, S.-X. Bai, H. Zhang and Y.-C. Ye, Int. J. Refract. Metals Hard Mater., 2015, 50, 204 LINK http://dx.doi.org/10.1016/j.ijrmhm.2015.01.009
  35. 35.
    Y. Ritterhaus, T. Hur’yeva, M. Lisker and E. P. Burte, Chem. Vap. Deposition, 2007, 13, (12), 698 LINK http://dx.doi.org/10.1002/cvde.200706630
  36. 36.
    M. A. El Khakani, M. Chaker and B. Le Drogoff, J. Vac. Sci. Technol. A, 1998, 16, (2), 885 LINK http://dx.doi.org/10.1116/1.581029
  37. 37.
    Y.-S. Gong, C.-B. Wang, Q. Shen and L.-M. Zhang, Vacuum, 2008, 82, (6), 594 LINK http://dx.doi.org/10.1016/j.vacuum.2007.09.003
  38. 38.
    T. D. Khoa, S. Horii and S. Horita, Thin Solid Films, 2002, 419, (1–2), 88 LINK http://dx.doi.org/10.1016/S0040-6090(02)00761-7
  39. 39.
    S. Hogmark, S. Jacobson and M. Larsson, Wear, 2000, 246, (1–2), 20 LINK http://dx.doi.org/10.1016/S0043-1648(00)00505-6
  40. 40.
    Z.-F. Chen, W.-P. Wu, X.-N. Cong and L.-B. Wang, Adv. Mater. Res., 2011, 314–316, 214 LINK http://dx.doi.org/10.4028/www.scientific.net/AMR.314-316.214
  41. 41.
    J. A. Venables, “Introduction to Surface and Thin Film Processes”, Cambridge University Press, Cambridge, UK, 2000 LINK http://dx.doi.org/10.1017/CBO9780511755651
  42. 42.
    T.-L. Chen, X.-M. Li, S. Zhang and X. Zhang, Appl. Phys. A, 2005, 80, (1), 73 LINK http://dx.doi.org/10.1007/s00339-004-2978-2
  43. 43.
    E. Bauer and J. H. van der Merwe, Phys. Rev. B, 1986, 33, (6), 3657 LINK http://dx.doi.org/10.1103/PhysRevB.33.3657
  44. 44.
    J. A. Venables, G. D. T. Spiller and M. Hanbucken, Rep. Prog. Phys., 1984, 47, (4), 399 LINK http://dx.doi.org/10.1088/0034-4885/47/4/002
  45. 45.
    N. Kaiser, Appl. Optics, 2002, 41, (16), 3053 LINK http://dx.doi.org/10.1364/AO.41.003053
  46. 46.
    T. Gerfin, W. J. Hälg, F. Atamny and K.-H. Dahmen, Thin Solid Films, 1994, 241, (1–2), 352 LINK http://dx.doi.org/10.1016/0040-6090(94)90456-1
  47. 47.
    Z.-F. Chen, W.-P. Wu, L.-B. Wang and Y. Zhang, Surf. Eng., 2011, 27, (4), 242 LINK http://www.tandfonline.com/doi/full/10.1179/174329409X397787
  48. 48.
    Y.-M. Sun, J. P. Endle, K. Smith, S. Whaley, R. Mahaffy, J. G. Ekerdt, J. M. White and R. L. Hance, Thin Solid Films, 1999, 346, (1–2), 100 LINK http://dx.doi.org/10.1016/S0040-6090(98)01458-8
  49. 49.
    J. R. V. Garcia and T. Goto, Mater. Trans., 2003, 44, (9), 1717 LINK http://dx.doi.org/10.2320/matertrans.44.1717
  50. 50.
    T. Goto, R. Vargas and T. Hirai, J. Phys. IV France, 1993, 3, (C3), 297 LINK http://dx.doi.org/10.1051/jp4:1993341
  51. 51.
    P. P. Semyannikov, N. B. Morozova, K. V. Zherikova, S. V. Trubin, I. K. Igumenov and N. V. Gelfond, ECS Trans., 2009, 25, (8), 887 LINK http://dx.doi.org/10.1149/1.3207682
  52. 52.
    V. Yu. Vasilyev, N. B. Morozova, T. V. Basova, I. K. Igumenov and A. Hassan, RSC Adv., 2015, 5, (41), 32034 LINK http://dx.doi.org/10.1039/C5RA03566J
  53. 53.
    R. Vargas, T. Goto, W. Zhang and T. Hirai, Appl. Phys. Lett., 1994, 65, (9), 1094 LINK http://dx.doi.org/10.1063/1.112108
  54. 54.
    N. V. Gelfond, I. K. Igumenov, A. I. Boronin, V. I. Bukhtiyarov, M. Yu. Smirnov, I. P. Prosvirin and R. I. Kwon, Surf. Sci., 1992, 275, (3), 323 LINK http://dx.doi.org/10.1016/0039-6028(92)90804-F
  55. 55.
    S.-I. Song, J.-H. Lee, B.-H. Choi, H.-K. Lee, D.-C. Shin and J.-W. Lee, Surf. Coat. Technol., 2012, 211, 14 LINK http://dx.doi.org/10.1016/j.surfcoat.2011.09.074
  56. 56.
    K. Knapas and M. Ritala, Chem. Mater., 2011, 23, (11), 2766 LINK http://dx.doi.org/10.1021/cm103490v
  57. 57.
    T. Aaltonen, M. Ritala, V. Sammelselg and M. Leskelä, J. Electrochem. Soc., 2004, 151, (8), G489 LINK http://dx.doi.org/10.1149/1.1761011
  58. 58.
    S. T. Christensen and J. W. Elam, Chem. Mater., 2010, 22, (8), 2517 LINK http://dx.doi.org/10.1021/cm9031978
  59. 59.
    C.-M. Chang, C.-M. Wei and S.-P. Chen, Phys. Rev. B, 1996, 54, (23), 17083 LINK http://dx.doi.org/10.1103/PhysRevB.54.17083
  60. 60.
    D. A. Toenshoff, R. D. Lanam, J. Ragaini, A. Shchetkovskiy and A. Smirnov, ‘Iridium Coated Rhenium Rocket Chambers Produced by Electroforming’, 36th AIAA/ASME/SAE/ASEE Joint Propulsion Conference and Exhibit, Las Vegas, USA, 24th–28th July, 2000 LINK http://dx.doi.org/10.2514/6.2000-3166
  61. 61.
    W.-P. Wu and Z.-F. Chen, Acta Metall. Sin. (Engl. Lett.), 2012, 25, (6), 469 LINK http://dx.doi.org/10.11890/1006-7191-126-469
  62. 62.
    M. J. Aziz, Appl. Phys. A, 2008, 93, (3), 579 LINK http://dx.doi.org/10.1007/s00339-008-4696-7
  63. 63.
    T. Chen, X.-M. Li and X. Zhang, J. Cryst. Growth, 2004, 267, (1–2), 80 LINK http://dx.doi.org/10.1016/j.jcrysgro.2004.03.005
  64. 64.
    Z.-F. Chen, W.-P. Wu and X.-N. Cong, J. Mater. Sci. Technol., 2014, 30, (3), 268 LINK http://dx.doi.org/10.1016/j.jmst.2013.06.002
  65. 65.
    W.-P. Wu, Z.-F. Chen, X. N. Cong and L.-B. Wang, Rare Metal Mater. Eng., 2013, 42, (2), 435 (In Chinese) LINK http://caod.oriprobe.com/articles/31751798/Review_on_High_Temperature_Oxidation_Resistant_Iridium_Coating_for_Ref.htm
  66. 66.
    Z.-W. Zhang, Z.-H. Xu, J.-M. Wang, W.-P. Wu and Z.-F. Chen, J. Mater. Eng. Perf., 2012, 21, (10), 2085 LINK http://dx.doi.org/10.1007/s11665-012-0133-3
  67. 67.
    W.-P. Wu and Z.-F. Chen, J. Wuhan Univ. Technol.-Mater. Sci. Ed., 2012, 27, (4), 652 LINK http://dx.doi.org/10.1007/s11595-012-0522-3
  68. 68.
    W.-P. Wu, Z.-F. Chen and X. Lin, Adv. Mater. Res., 2011, 189–193, 688 LINK http://dx.doi.org/10.4028/www.scientific.net/AMR.189-193.688
  69. 69.
    W.-P. Wu, Z.-F. Chen, H. Cheng, L.-B. Wang and Y. Zhang, Appl. Surf. Sci., 2011, 257, (16), 7295 LINK http://dx.doi.org/10.1016/j.apsusc.2011.03.108
  70. 70.
    W.-P. Wu, Z.-F. Chen and L.-B. Wang, Prot. Met. Phys. Chem. Surf., 2015, 51, (4), 607 LINK http://dx.doi.org/10.1134/S2070205115040358
  71. 71.
    W.-P. Wu, J.-J. Jiang and Z.-F. Chen, Acta Astronaut., 2016, 123, 1 LINK http://dx.doi.org/10.1016/j.actaastro.2016.03.004
  72. 72.
    W.-P. Wu and Z.-F. Chen, Surf. Interface Anal., 2016, 48, (6), 353 LINK http://dx.doi.org/10.1002/sia.5986
  73. 73.
    J.-M. Wang, Z.-W. Zhang, Z.-H. Xu, X. Lin, W.-P. Wu and Z.-F. Chen, Corros. Eng. Sci. Technol., 2011, 46, (6), 732 LINK http://dx.doi.org/10.1179/1743278210Y.0000000023
  74. 74.
    B. D. Reed, J. A. Biaglow and S. J. Schneider, Mater. Manuf. Proc., 1998, 13, (5), 757 LINK http://dx.doi.org/10.1080/10426919808935297
  75. 75.
    K. Mumtaz, J. Echigoya, H. Enoki, T. Hirai and Y. Shindo, J. Mater. Sci., 1995, 30, (2), 465 LINK http://dx.doi.org/10.1007/BF00354413
  76. 76.
    D. H. Lowndes, D. B. Geohegan, A. A. Puretzky, D. P. Norton and C. M. Rouleau, Science, 1996, 273, (5277), 898 LINK http://dx.doi.org/10.1126/science.273.5277.898
  77. 77.
    W.-B. Yang, L.-T. Zhang, Y.-F. Hua and L.-F. Cheng, Int. J. Refract. Metals Hard Mater., 2009, 27, (1), 33 LINK http://dx.doi.org/10.1016/j.ijrmhm.2008.02.001
  78. 78.
    Y.-L. Huang, S.-X. Bai, H. Zhang and Y.-C. Ye, Appl. Surf. Sci., 2015, 328, 436 LINK http://dx.doi.org/10.1016/j.apsusc.2014.12.063
  79. 79.
    K. Mumtaz, J. Echigoya and M. Taya, J. Mater. Sci., 1993, 28, (20), 5521 LINK http://dx.doi.org/10.1007/BF00367824
  80. 80.
    C.-Y. Hu, ‘Study on the CVD Iridium Coated Rhenium Composite’, PhD thesis, Central South University, Changsha, China, Dissertations from CNKI, 2002 (in Chinese)
  81. 81.
    L. Brewer and R. H. Lamoreaux, ‘Molybdenum: Physico-Chemical Properties of its Compounds and Alloy’, Atomic Energy Review, International Atomic Energy Agency, Vienna, Austria, 1980, 7, 263
  82. 82.
    E. K. Ohriner and E. P. George, J. Alloys Compd., 1991, 177, (2), 219 LINK http://dx.doi.org/10.1016/0925-8388(91)90075-7
  83. 83.
    J. C. Hamilton, N. Y. C. Yang, W. M. Clift, D. R. Boehme, K. F. McCarty and J. E. Franklin, Metall. Trans. A, 1992, 23, (3), 851 LINK http://dx.doi.org/10.1007/BF02675562
  84. 84.
    S. Chen, C.-Y. Hu, J.-M. Guo and J.-M. Yang, Rare Metal Mater. Eng., 2005, 34, (6), 916 LINK http://caod.oriprobe.com/articles/9362013/Study_on_Interdiffusion_of_Iridium_Molybdenum.htm
  85. 85.
    J. Hämäläinen, E. Puukilainen, M. Kemell, L. Costelle, M. Ritala and M. Leskelä, Chem. Mater., 2009, 21, (20), 4868 LINK http://dx.doi.org/10.1021/cm901687w
  86. 86.
    R. H. Tuffias, G. J. Melden and J. T. Harding, ‘High Temperature Oxidation-Resistant Thruster Materials, Phase II’, NASA Contractor Report 187205, 1991 LINK http://www.dtic.mil/dtic/tr/fulltext/u2/a293027.pdf
  87. 87.
    K. Mumtaz, J. Echigoya, H. Enoki, T. Hirai and Y. Shindo, J. Alloys Compd., 1994, 209, (1–2), 279 LINK http://dx.doi.org/10.1016/0925-8388(94)91114-2
  88. 88.
    K. Mumtaz, J. Echigoya, H. Enoki, T. Hirai and Y. Shindo, J. Mater. Sci., 1996, 31, (19), 5247 LINK http://dx.doi.org/10.1007/BF00355932
  89. 89.
    X.-H. Zhang, F.-T. Xu, Z.-H. Jia, H. Q. Li, K.-M. He and D. Chen, Mater. China, 2013, 32, (4), 203 (in Chinese) LINK http://en.oversea.cnki.net/kcms/detail/detail.aspx?QueryID=25&CurRec=2&dbCode=CJFD&filename=XJKB201304003&dbname=CJFD2013
 

The Authors


Wang-ping Wu received his doctorate in Materials Processing Engineering at Nanjing University of Aeronautics and Astronautics, China, in 2013 and held a Pikovsky Valazzi Scholarship at Tel Aviv University, Israel, where he was a Postdoctoral Fellow. He is now a Senior Lecturer at the School of Mechanical Engineering in Changzhou University, China. His research interests are mainly directed towards the synthesis and characterisation of films and coatings of the noble metals and their alloys.


Zhao-feng Chen is Professor of Materials Science and Director of the International Laboratory for Insulation and Energy Efficiency Materials at the College of Material Science and Technology at Nanjing University of Aeronautics and Astronautics. His research interests include advanced insulation composite materials and the application of coatings of noble metals.

Read more of this issue

Related articles

On the Criticality of Palladium in Organic Synthesis: A Perspective

View
Download

On the Sustainability of Palladium in Organic Synthesis: A Perspective

View
Download
MORE ARTICLES

Find an article

ArticleSearch