Surface & Coatings Technology 299 (2016) 153–161
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Lanthanum hydride doped tungsten-based coating fabricated by supersonic atmospheric plasma spraying Qing Yu Hou a,b,⁎, Lai Ma Luo a, Zhen Yi Huang b, Ping Wang b, Ting Ting Ding b, Yu Cheng Wu a,⁎⁎ a b
Laboratory of Nonferrous Metal Materials and Processing Engineering of Anhui Province, Hefei University of Technology, Hefei, Anhui 230009, China Anhui Key Laboratory of Metal Materials and Processing, Maanshan, Anhui 243002, China
a r t i c l e
i n f o
Article history: Received 18 December 2015 Revised 5 May 2016 Accepted in revised form 5 May 2016 Available online 07 May 2016 Keywords: Supersonic atmospheric plasma spraying W/La2O3 composite coating Rare earth hydrides Elastic failure Fracture toughness
a b s t r a c t In this work, as a kind of rare earth hydrides, 1.5 wt.% LaH2 powder was introduced into a tungsten powder and then deposited by a supersonic atmospheric plasma spraying (SAPS) machine. The purpose was to form La2O3 in the LaH2-doped tungsten-based coating. The microstructure and properties of the undoped and LaH2-doped tungsten-based coatings were characterized using scanning electron microscopy (SEM) with energy dispersive spectroscopy (EDX), X-ray photoelectron spectroscopy (XPS), X-ray diffractometer (XRD), nano-indentation tester and Vickers hardness tester. The results showed that SAPS coatings exhibited typical lamellar microstructure. Pores located mainly at lamellar gaps in association with oxidation were also observed. The type of tungsten oxide was mainly WO3. During SAPS process, La2O3 was formed in the LaH2-doped tungsten-based coating by the transformation of LaH2. The formed La2O3 was distributed mainly at lamellar gaps, having a joining effect on the adjacent lamellae. The oxygen content, porosity and thermal conductivity of the LaH2-doped tungsten-based coating were lower than those of the undoped tungsten coating. The introduction of 1.5 wt.% LaH2 in the tungsten coating improved its ability to resist elastic fracture and improved its fracture toughness. © 2016 Elsevier B.V. All rights reserved.
1. Introduction It is confirmed that tungsten has a high melting point and threshold for physical sputtering, low deuterium etching and vapor pressure, and excellent thermal conductivity. Therefore, it is thought as a promising candidate armor material for plasma-facing components (PFCs) of nuclear fusion devices [1,2]. However, its use in PFCs is limited because of its serious brittleness in several situations and a sharping decrease in high-temperature mechanical properties [3,4]. These brittleness and high-temperature mechanical properties are microstructure-sensitive. They can be alleviated and improved by forming or adding a small amount of dispersed ceramic particles in tungsten to form ceramic particles doped tungsten-based composite materials [4,5]. Among these ceramic particles, rare earth oxides were believed to have some special chemical characteristics, contributing to improving the brittleness and high-temperature mechanical properties of tungsten material [6,7]. It is well known that oxygen impurity can cause tungsten materials embrittlement [8]. The introduction of rare earth (RE) elements in the
⁎ Correspondence to: Q.Y. Hou, Laboratory of Nonferrous Metal Materials and Processing Engineering of Anhui Province, Hefei University of Technology, Hefei, Anhui 230009, China. ⁎⁎ Corresponding author. E-mail address:
[email protected] (Q.Y. Hou).
http://dx.doi.org/10.1016/j.surfcoat.2016.05.012 0257-8972/© 2016 Elsevier B.V. All rights reserved.
metallic state instead of the oxidation state was believed to be better for fabrication of high performance tungsten-based materials because of the high affinity of RE elements with oxygen [8]. However, RE elements in the metallic state would be easily oxidized before they were introduced into a tungsten-based material because their activity was high enough in air [9], limiting their positive effect in improving the properties of the tungsten-based materials. Therefore, one effective way to introduce RE elements into a tungsten-based material should be explored. From the latest investigation conducted by B. Wang et al. [9], one way to introduce RE elements into a tungsten-based material to improve its properties might be considered. They found when LaH2 powder was added into a Ti-Fe-Mo alloy powder then sintered under vacuum atmosphere at about 1573 K, hydrogen (H) would release from LaH2 then reacted with the oxygen (O) presented in the alloy to form La2O3, improving the properties of the Ti-Fe-Mo alloy. The choice to use lanthanum (La) in the hydrogenated state instead of the metallic state was because the activity of La in air is greater than that of LaH2 [9]. Thus, La would be more easily oxidized before it was introduced into the Ti-Fe-Mo alloy than LaH2 did, decreasing the beneficial effects of La on the properties of the sintering Ti-Fe-Mo alloy. Based on the reported result [9], it was reasonable to predict that La would be obtained when LaH2 was introduced into a tungsten-based material under a certain thermal effect. The obtained La would react with tungsten
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oxides or oxygen presented in the tungsten-based material to form La2O3. In addition, the introduced LaH2 might react with oxygen to form La2O3. Based on the above result and consideration, therefore, it can be predicted that the properties of the tungsten-based materials should be improved after the introduction of LaH2, which is attributed to the decreasing of tungsten oxides and the formation of La2O3 in the tungsten-based materials. Other disadvantages of tungsten for using in PFCs were its poor machinability and weldability [10]. In order to overcome these disadvantages, a possible solution to employ tungsten at a plasma-facing surface is the coating of the copper-based heat sink or steel-based support structure with a tungsten layer [10–12]. Among the potential coating techniques, atmospheric plasma spraying (APS) looks attractive because of its ability to cover large areas with thick coatings and its relative simplicity and cost-effectiveness [13,14]. However, the tungsten coatings fabricated by APS technique commonly have a high porosity and oxygen content, which has a negative effect on the physical and mechanical properties of the coatings [15,16]. Although vacuum plasma spraying (VPS), low-pressure plasma spraying (LPPS), and inert gas atmosphere plasma spraying (IPS) techniques have been used to fabricate the tungsten coatings whose properties were commonly greater than those fabricated by APS technique [17–19], the complexity and low cost-effectiveness might limit their application [15,20]. Compared with the aforementioned plasma spraying techniques, supersonic atmospheric plasma spraying (SAPS) technique was thought as possessing the advantages of reducing oxidation of powder and sprayed layers, providing a convenient and cost-effective way to fabricate tungsten coatings using in PFCs [20]. However, there are almost no attention having been focused on fabrication LaH2-doped tungsten-based composite coating by plasma spraying techniques, especially by SAPS technique. It will be interesting and important to investigate the effects of the introduced LaH2 on the microstructure and properties of the tungsten-based composite coating fabricated by SAPS technique in order to evaluate its potential for using in PFCs. This was the motivation of this work. In this work, as part of a proof-of-concept for fabricating La2O3 modified tungsten-based composite coating prototypes, LaH2 powder was introduced into a tungsten powder to form W/LaH2 composite powder and then was sprayed by a SAPS machine to form a coating. The purpose was to improve the properties of the tungsten coating by forming La2O3 in the LaH2-doped tungsten-based coating. Although it has been reported by most of the investigations [6,21,22] that the tungsten-based composite materials with 1 wt.% La2O3 particles commonly have excellent properties as compared with the pure tungsten materials, the introduced LaH2 powder in the present work was set at 1.5 wt.%. The reason is that, supposing all the introduced 1.5 wt.% LaH2 was transformed into La2O3, the theoretical content of the La2O3 particles formed in the LaH2-doped tungsten-based composite coating should be about 1.7 wt.%. However, the fact would not be so because the loss of LaH2 powder from the W/LaH2 composite powder during flight in the plasma jet, which should be similar to that of the HfC particles loss from a W/HfC composite powder [23], should not be ignored. Because it was not easy to evaluate the ratio of LaH2 powder loss from the W/LaH2 composite powder during flight in the plasma jet, it was hard to predict how much LaH2 powder should be introduced in order to form about 1 wt.% La2O3 in the LaH2-doped tungsten-based coating fabricated by SAPS. Therefore, from a perspective of conservative, referring to the content of LaH2 powder in a sintering Ti-Fe-Mo alloy (0.6 wt.%–3.0 wt.% [9]), the content of LaH2 powder in the W/LaH2 composite powder was set at 1.5 wt.%, expecting to form about 1.0 wt.% La2O3 in the LaH2-doped tungsten-based coating to evaluate its potential for using in PFCs. The microstructure and properties of the tungsten-based coatings before and after LaH2 addition were examined and compared. The possible explanations of the difference of the microstructure and the properties between the SAPS-W and SAPS-W/LaH2 coatings were discussed.
2. Experimental procedures 2.1. Powders characteristics and coatings fabrication Commercially available tungsten powder (purity N 99.9%) and LaH2 powder (purity N 99.5%) were used as feedstocks. The two types of powder were blended in a glove box by a nominal composition of 98.5 wt.% W and 1.5 wt.% LaH2 (W/LaH2) to form LaH2-doped tungstenbased composite powder. The W/LaH2 composite powder was then charged into a vessel made of WC together with WC balls and then installed in a planetary ball mill for 24 h with a rotation speed of 50 rpm for homogenizing purpose. The inner atmosphere of the vessel for blending was a purified argon gas (purity 99.999%). The coating fabrication was carried out by a supersonic atmospheric plasma spraying (SAPS, HEPJet-II) system whose jet temperature might be 10,000–20,000 K [10]. Argon and hydrogen were used as plasma forming gases with flow rates of 58 L/min and 6 L/min, respectively. Reduced activation steel substrates were grit-blasted and ultrasonically cleaned before SAPS treatments. During SAPS, the backsides of the substrates were kept at a low temperature by air jet cooling to avoid the detachment of the coatings from the substrates. Coatings about 0.3 mm thick were deposited on the substrates. 2.2. Microstructure characterization Microstructure and element composition of the W/LaH2 composite powder before and after SAPS treatments were examined by scanning electron microscope (SEM, SU-8020 and JSM-6490, Japan) with energy dispersive X-ray spectrometer (EDX, Oxford, Japan). The sizes of the blending powders were estimated from SEM micrographs taken from several locations of a specimen by image analyzing software of ImageJ; ten SEM images with a 2000× magnification were used to build a statistical analysis. The fabricated coatings with polished surfaces were analyzed by X-ray photoelectron spectrometer (XPS, Escalab250Xi, Thermo Scientific, USA) to detect the type of tungsten oxide formed in the coatings. The Al Kα was used for X-ray source. Both chambers were evacuated to less than 10− 7 Pa in order to remove the gaseous impurities, such as H2O, O2, and H2. Phase characteristic of the W/LaH2 composite powder and LaH2 powder before and after SAPS treatment was analyzed by X-ray diffractometer (XRD, D/MAX-2500V, Japan) operating with Cu Kα (λ = 1.5406 Å) radiation. XRD profiles were subsequently analyzed using the Rietveld refinement method program RIETAN-FP [24]. The goodness-of-fit of the refinement results was evaluated by minimizing weighted-profile R-factors (Rwp) and Rietveld standard deviations of the structural parameter χ2 (S) quantities [25,26]. 2.3. Property characterization The density and open porosity were measured according to the Archimedes' principle [27]. The oxygen content was detected utilizing the Nitrogen/Oxygen analyser (TC600, Leco, USA). Three measurements were taken to determine the averaged values of density, porosity and oxygen content for each coating. The standard deviation for each measurement was given to show the measurement error. The specific heat capacity (Cp) was measured on free-standing coatings of Φ 5 mm × 200 μm using a synchronous TG-DSC thermal analyser (STA449C, Germany) at 300 K. Three specimens for each free-standing coating were applied to build a statistical analysis. The thermal diffusivity (α) was conducted on free-standing coatings of Φ 12.7 mm × 200 μm using a laser flash diffusivity system (LFA447, Germany) at 300 K. The thermal conductivity (λ) was calculated using the measured Cp, α and density (ρ) of the coatings, as follows [28]: λ ¼ α Cp ρ
ð1Þ
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The hardness (H) and Young's moduli (E) were measured using a MTS nano-indentation machine (NanoXP™, USA). Nano-indentation test was conducted on the polished and etched cross section of each specimen. The specimens were aligned such that the test surface was perpendicular to the force. At least twenty indentations, which depended on the thickness of the obtained coatings, were performed on the different locations along the coating thickness of the specimens to gain the result having statistic characteristic. The capacity of controlling loading was set at 100 mN and unloading rate was set at 500 μN·min−1. The nano-indentation effective Young's modulus E* = E/(1 − ν2) and hardness H were calculated based on the loading/unloading curves which were measured with a Berkovich indenter using the Oliver-Pharr method [29], where E and ν are the Young's modulus and Poisson's ratio of the coatings. Vickers indentations were conducted on the polished coating surfaces using a microhardness tester to evaluate the ability of the coatings to resist elastic fracture and to evaluate the fracture toughness of the coatings. The indentation load was set at 4.96 N. The morphologies of the indented coating surfaces were investigated by SEM. The crack lengths, which were produced after Vickers indentation tests and taken photos by SEM, were measured using image analyzing software of ImageJ. Ten Vickers indentations in each SAPS coating were used to build a statistical analysis. 3. Results and discussion 3.1. Characteristics of the blending powders SEM morphologies of W, LaH2, and W/LaH2 powders are presented in Fig. 1a–c. It can be seen from Fig. 1a and b that W powder exhibit a more uniform morphology than LaH2 powder does and their sizes are
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about 4–20 μm and 0.2–11 μm, respectively. When the two types of powder were blended by a planetary ball mill, the W/LaH2 composite powder consisted of W powder with a large size and LaH2 powder with a small size (Fig. 1c). EDX result present in Fig. 1d gives evidence that the big particles shown in Fig. 1c are tungsten; and the small particles are LaH2. The existence of carbon (C) in the W/LaH2 composite powder was mainly attributed to the contamination of the surroundings. The existence of oxygen (O) might be attributed to the formation of oxide or hydroxide in the W/LaH2 composite powder. The morphology and size of the W powder changed a little, but that of the LaH2 powder changed much when the two types of powder were blended by a planetary ball mill. The LaH2 powder with a large and sharp edge was milled into a small and round blunt one, attributing to the higher brittleness of RE hydride [30]. Fig. 2a shows the XRD Rietveld refinement result for the W/LaH2 composite powder, in which W (JCPDS No. 04-806 [31]), LaH2 (JCPDS No. 89-4196 [31]) and La(OH)3 (JCPDS No. 36-1481 [31]) can be indexed. The formation of La(OH)3 might be ascribed to the introduction of H2O in the LaH2 powder during the blending process of the W/LaH2 composite powder. Comparing the strongest peak for LaH2 and that for La(OH)3, it can be obtained that the relative intensity of the former is higher than that of the latter, indicating the relative content of LaH2 in the W/LaH2 composite powder was greater than that of La(OH)3. Nevertheless, it should be clarified that La(OH)3 mainly formed during the blending process of the W/LaH2 composite powder; and only a small amount of La(OH)3 formed. The XRD Rietveld refinement result of the LaH2 powder shown in Fig. 2b and the quantitative analysis results shown in Table 1 can be used to clarify such result. It can be seen from Fig. 2b that only LaH2 phase can be indexed in the LaH2 powder, indicating that La(OH)3 phase in the W/LaH2 composite powder was formed during
Fig. 1. SEM micrograph of the powder and the corresponding EDX results: (a) W powder; (b) LaH2 powder; (c) W/LaH2 composite powder; (d) EDX results of the phases shown in Panel c.
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3.2. Microstructure and composition characteristics of the coatings
Fig. 2. Observed (+), calculated (line) and residual (bottom line) X-ray diffraction patterns of the powders: (a) W/LaH2 composite powder; (b) LaH2 powder.
the blending process. From the refinement result shown in Table 1, it can be seen that the relative contents of W, LaH2 and La(OH)3 phases in the W/LaH2 composite powder are about 98.1 wt.%, 1.2 wt.% and 0.7 wt.%, respectively. Although the W/LaH2 composite powder was blended in a planetary ball mill for homogenizing purpose, a complete homogenization was not easily obtained. Therefore, the relative content of W in the used sample to conduct XRD measurement might be different from the nominal composition. In addition, the introduction of H2O in the LaH2 powder decreased the relative content of W in the W/LaH2 composite powder. Finally, the calculation error, which should also be considered, was unavoidable during the refinement. Considering the three factors, the calculated result agrees well with the given nominal compositions of the W/LaH2 composite powder. It can also be proposed from Table 1 that only a small amount La(OH)3 formed in the W/LaH2 composite powder.
Table 1 Phase constitution, structural parameters, and phase abundance of W/LaH2 composite powder as determined by XRD. Materials Powders Rwp = 8.45% S = 6.52
Phases α(W) LaH2 La(OH)3
Space group
Lattice parameters (nm)
Im3m Fm3m P63/m
0.31671(4) 0.56303(8) 0.65378(9)
a
b
c
Abundance (wt.%)
0.38668(8)
98.1 1.2 0.7
Fig. 3a and b show a typical cross-sectional morphology of the SAPS-W coating. Generally, lamellar structures and splat boundaries of thermal spray coatings are observed. The well-melted tungsten particles forming the lamellar layers are called splat; and the dark lines between splats are called splat boundaries [32]. Pores can also be found mainly at lamellar gaps. Plasma spraying is a two-stage process where the first in-flight stage, in which the powder is injected into the plasma jet, melted and accelerated; is followed by the second stage (flattening, solidification and cooling down). During SAPS process, the as-injected powder was melted and propelled as molten/semi-molten droplets or un-melted particles toward the substrate. Because SAPS process was conducted in an ambient air environment, the oxidation would occur because of the entrainment of ambient air. Oxidation of the coatings commonly occurred in the first stage of spraying process [32]. Therefore, oxides on the droplets would inhibit their flattening and contacting between splats, resulting in lamellar gaps and pores. In addition, another typical microstructure of plasma spraying coatings, which was made up of un-melted/semi-melted particles and splashed debris from the impacting tungsten splats, can be seen in the SAPS-W coating (Fig. 3a and b). These particles and debris have irregular morphologies, which was the result of the random packing of unmelted/semi-melted tungsten particles and splashed debris. EDX results shown in Table 2 give evidence that the matrix microstructure (marked by A in Fig. 3b) of the SAPS-W coating was tungsten. Compositions of the lamellae gap (marked by B in Fig. 3b) and the pore (marked by C in Fig. 3b) consisted of W and O, indicating the formation of tungsten oxide. Because it is not accurate to detect light elements using EDX technique, such as oxygen, what kinds of tungsten oxide formed in the SAPS-W coating could not be obtained from EDX result. As pointed out by S.C. Cifuentes et al. [33] that oxidation kinetics of tungsten was quite complex, resulting in forming a wide variety of oxides during oxidation. In the present work, further information about the type of tungsten oxide was investigated using XPS technique, as shown in Fig. 3c and d. The peaks of the W4f and the O1s at binding energy of about 35.38 eV and 530.58 eV can be observed, respectively. The two binding energies can be assigned to those values in WO3 configuration [34]. In other words, it can be concluded from Fig. 3c and d that WO3 formed in the SAPS-W coating. It should be pointed out that the smoothness of the two XPS curves is relatively low because the polished coating to conduct XPS measurement was not removed from the steel substrates, which might lead to the magnetization of the tested samples. Additionally, columnar crystals within the individual splats were formed because of the rapid nucleation of the sprayed and flattened droplets and the temperature gradient from the substrate to the coating. In comparison, the typical lamellar microstructure can also be seen in the SAPS-W/LaH2 coating, though some dark gray phases are also observed (Fig. 4a and b). EDX results presented in Table 2 show that the matrix (marked by A in Fig. 4b) of the SAPS-W/LaH2 coating is tungsten. Tungsten oxide (marked by B1 in Fig. 4b) and dark gray phase (marked by B2 in Fig. 4b) can be detected at lamellar gap. Additionally, tungsten oxide can also be detected in the pores (marked by C in Fig. 4b). Similarity to the SAPS-W coating, the oxidization of the SAPS-W/ LaH2 coating to form tungsten oxide was also inevitable during SAPS process. EDX results shown in Table 2 give evidence that tungsten oxide was formed in the SAPS-W/LaH2 coating. The evidence to show the type of tungsten oxide in the SAPS-W/LaH2 coating was provided by XPS measurement, as shown in Fig. 4c and d. The binding energy for W4f and O1s can be assigned to those values in WO3 configuration [34]. The result shows that the type of tungsten oxide in the SAPS-W/ LaH2 coating is same to that in the SAPS-W coating. It is logical to propose from the EDX result shown in Table 2 that the dark gray phase in Fig. 4a and b is La2O3 phase. However, further explanation should be provided. As reported by Wang et al. [9] that La2O3 was formed in a vacuum sintered LaH2-doped Ti-Fe-Mo alloy.
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Fig. 3. SEM micrograph and XPS spectra of the SAPS-W coating: (a, b) SEM micrograph; (c) W4f XPS spectra of WO3 in the polished coating; (d) O1s XPS spectra of WO3 in the polished coating.
They proposed that LaH2 would be dehydrogenated at high temperature, leading to forming La and reacting with O presented in the alloy to form La2O3. In the present work, the temperature of the plasma jet of the SAPS machine is far greater than that of the sintering temperature used in the reference [9], thus the dehydrogenation of LaH2 was inevitable, leading to forming La and reacting with tungsten oxides (WO3) or oxygen to form La2O3 in the SAPS-W/LaH2 coating. Therefore, it can be proposed that the formation of La2O3 in the SAPS-W/LaH2 coating can be briefly summarized as following routes: 3 1 1 WðsÞ þ O2 ðgÞ ¼ WO3 ðsÞor WðsÞ þ H2 OðgÞ ¼ WO3 ðsÞ þ H2 ðgÞ 2 3 3
ð2Þ
LaH2 ðsÞ ¼ La ðsÞ þ H2 ðgÞ
ð3Þ
WO3 ðsÞ þ 2La ðsÞ ¼ W ðsÞ þ La2 O3 ðsÞ or 4La ðsÞ þ 3O2 ðgÞ ¼ 2La2 O3 ðsÞ ð4Þ
2H2 ðgÞ þ O2 ðgÞ ¼ 2H2 O ðgÞ
ð5Þ
Table 2 EDX results of different regions in SAPS coatings (at.%). Coatings
Region
W
La
O
H
W coating
Matrix (A) Splat boundary (B) Pore (C) Matrix (A) Splat boundary (B1) Splat boundary (B2) Pore (C)
100 28.23 36.58 100 35.08 0.87 23.71
– – – – – 29.62 20.61
– 71.77 63.42 – 64.92 69.51 55.68
– – – – – – –
W/LaH2 coating
The formed La2O3 would locate in the SAPS-W/LaH2 coating. However, the obtained H2 would get out of the coating or reacted with O2 to form H2O in gaseous phase and get out of the coating. It was logical to show the formation of La2O3 in the SAPS-W/LaH2 coating by the above routes. However, it should not be ignored that La2O3 might be formed by a reaction between LaH2 and O2 during SAPS-W/LaH2 process, as shown in Eq. (6). 5 1 LaH2 ðsÞ þ O ðgÞ ¼ La O3 ðsÞ þ H2 O ðgÞ 2 2 2 2
ð6Þ
In the present work, the formation of La2O3 in the SAPS-W/LaH2 coating might be the result of the above-mentioned ways. Additionally, it can also be seen from Fig. 4b (marked by squares) that dark gray La2O3 at lamellar gaps has a joining effect on the adjacent lamellae, contributing to enhancing the bonding strength of the adjacent lamellae. The existence of La2O3 in the pores between the lamellae had a filling effect, which was similar to that copper in tungsten [17,35], contributing to decreasing the open porosity of the SAPS-W/LaH2 coating (about 8.03%) compared to that of the SAPS-W coating (about 9.33%), as shown in Table 3, where oxygen content and thermal conductivity are also summarized. It can be seen from Table 3 that oxygen content of the SAPS-W/LaH2 coating is lower than that of the SAPS-W coating. From the above discussion about the formation of La2O3 in the SAPS-W/LaH2 coating, it can be determined that H2O in gaseous phase could be formed. The formed water vapor would move out of the SAPS-W/LaH2 coating, decreasing the oxygen content of this coating when compared to that of the SAPS-W coating. It is worth to note from Table 3 that it seems impossible that adding 1.5 wt.% LaH2 in the tungsten coating accounts for 2.4% decrease in its density. However, that is the case. As proposed in the ‘Introduction’
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Fig. 4. SEM micrograph and XPS spectra of the SAPS-W/LaH2 coating: (a, b) SEM micrograph; (c) W4f XPS spectra of WO3 in the polished coating; (d) O1s XPS spectra of WO3 in the polished coating.
Section, the theoretical content of La2O3 phase formed in the SAPS-W/ LaH2 coating should be about 1.7 wt.%. Therefore, the ideal density of the SAPS-W/LaH2 coatings is about 18.7 g/cm3. It can be obtained that the density of the SAPS-W/LaH2 coating is lower than that of the SAPS-W coating (its ideal density is about 19.3 g/cm3) by about 3.23% calculated from the theoretical point of view. In reality, the loss of LaH2, the formation of WO3 and the pores were inevitable. Thus it is logical to obtain that the density of the SAPS-W/LaH2 coating was lower than that of the SAPS-W coating by about 2.4%, the lower porosity of the SAPS-W/LaH2 coating compared to that of the SAPS-W coating might be the main reason. Although a lower oxygen content and porosity were commonly beneficial to increase the thermal conductivity of the tungsten-based materials [17,32,36], the thermal conductivity of the SAPS-W coating with a lower oxygen content and porosity was greater than that of the SAPS-W/LaH2 coating with a larger oxygen content and porosity (Table 3). As proposed above, La2O3 phase, whose thermal conductivity was far lower than that of W phase, was formed in the SAPS-W/LaH2 coating. Therefore, the conductivity of the SAPS-W/LaH2 coating with La2O3 formation might be lower than that of the SAPS-W coating without La2O3 formation, though the oxygen content and porosity of the SAPS-W/LaH2 coating were lower than those of the SAPS-W coating.
Table 3 Some basic properties of SAPS coatings. Coatings
W coating
W/LaH2 coating
Oxygen content (wt.%) Porosity (%) Density (g/cm3) Thermal conductivity (W/m K)
0.60 ± 0.021 9.33 ± 0.75 17.45 ± 0.39 98.4 ± 0.3
0.54 ± 0.007 8.03 ± 0.62 17.04 ± 0.50 87.1 ± 0.2
Fig. 5 shows the Rietveld refinement results of the SAPS-W and SAPS-W/LaH2 coatings. Only W phase can be indexed in the two coatings, which does not agree well with the above SEM + EDX analysis (Figs. 3 and 4, Table 2). As obtained in Figs. 3 and 4, WO3 was formed in the two coatings during the SAPS process. However, it cannot be indexed from the XRD patterns shown in Fig. 5. Similarly, no peaks for La2O3 phase can be indexed in the XRD pattern of the SAPS-W/LaH2 coating, though the formation of La2O3 in this coating was logical (Eqs. (2)–(6)) and could be verified by EDX analysis (Table 2). It can be seen from Figs. 3, 4 and Table 2 that oxides existed mainly at lamellar gaps. Because the distribution of the lamellae in the SAPS-W and SAPS-W/LaH2 coatings was perpendicular to their coating surfaces, most of the formed oxides would be masked by the lamellae, decreasing the opportunity of the formed oxides to be detected by XRD tests that were conducted on the coatings surface. Although the relative content of La2O3 in the LaH2-doped tungsten-based coating might as high as about 1.7 wt.% supposing all the introduced 1.5 wt.% LaH2 was transformed into La2O3, it might be not the case. This is due to the fact that the loss of LaH2 from the W/LaH2 was inevitable during flight in the plasma jet, which was logical to be predicted from the loss of an additive HfC particles from an HfC-doped tungsten-based coating during flight in the plasma jet [30]. Therefore, the relative content of La2O3 transformed from LaH2 in the LaH2 -doped tungsten-based coating decreased, decreasing the opportunity of La2O3 to be detected by XRD test. It seems reasonable to explain the reasons that WO3 and La2O3 could not be indexed in the XRD patterns using the above-mentioned theory. However, it should also be clarified whether WO3 and La2O3 had an amorphous form instead of crystal form presenting in the tungstenbased coatings because the XRD peaks of an amorphous were also unobvious. Because SAPS is a layer by layer process, although the rapid cooling rate of the SAPS process might be favorable to form
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Fig. 5. Observed (+), calculated (line) and residual (bottom line) X-ray diffraction patterns of the SAPS coatings: (a) W coating; (b) W/LaH2 coating.
amorphous, the reheating effect of the subsequent deposition process on the pre-deposited layer would lead to the crystallization of the pre-formed amorphous. In other words, WO3 and La2O3 were more likely to be present in the tungsten-based coatings in the form of crystal form rather than in the amorphous form. Nevertheless, further information about the existing characteristic of WO3 and La2O3 in the tungstenbased coatings should be provided because of the complexity of the SAPS process, which was not the focus of the present work.
3.3. Nano-indentation mechanical properties of the coatings At least twenty indentations were performed on different locations along the coating thickness of each specimen in order to show the properties of the coatings as far as possible by statistical data. Nevertheless, only one typical loading/unloading curve for each coating was illustrated because the testing results were very close and the data error was between 1% and 2%, as shown in Fig. 6a. The areas enclosed by the loading/unloading curves (area I) and under unloading curve (area II) represent plastic deformation energy and elastic deformation energy, respectively, as sketched in Fig. 6b, where dres and dmax refer to the residual deformation and maximum deformation, respectively. Therefore, from measured values of H and E⁎, it is easy to calculate the H/E⁎ ratio. For a stiffer and brittle material, the elastic failure would be formed before the plastic deformation, the H/E⁎ ratio (i.e. a description in terms of ‘elastic strain to failure’) was thought having some merits to show the elastic failure [37].
Fig. 6. Nano-indentation mechanical properties of the SAPS coatings (a) and sketched map of a typical nano-indentation curve (b) where areas I and II represent plastic deformation and elastic recovery, respectively.
Although it has been proposed that a stiffer and brittle material with a higher elastic recovery We (i.e., a value to absorb elastic deformation energy relative to total deformation energy, %) [38,39] might have a higher H/E⁎ ratio, it might be not always the case. This is because the result might also be affected by indented depth and the applied load [38]. For the present stiffer and brittle tungsten-based coatings, it can be found from the indentation results shown in Table 4 that the SAPSW/LaH2 coating with a high H/E⁎ ratio has a higher We value than that of the SAPS-W coating with a low H/E⁎ ratio. However, taking account of the experimental errors and the inevitable microstructure inhomogeneity, it can be proposed that the We values of the two types of coating are almost equal. Therefore, due to the complexity of the evaluation of the ability to resist elastic failure for a stiffer and brittle material [38], further information about the difference in the ability of the two types of coating to resist elastic failure should be provided. 3.4. Failure characteristics of the coatings In order to compare the ability of the SAPS-W and SAPS-W/LaH2 coatings to resist to elastic failure, Vickers indentation tests loaded at as high as 4.96 N were performed on the two types of coating, by which cracks were caused on the tested surfaces. For statistically significant results, twenty indentations were applied on each coating. The
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Table 4 Mechanical properties of SAPS coatings evaluated from nano-indentation tests. Coatings
H, GPa
E⁎, GPa
H/E⁎
dres, nm
dmax, nm
We, %
W coating W/LaH2 coating
9.41 ± 0.88 6.93 ± 0.79
260.8 ± 3.23 143.5 ± 3.85
0.036 ± 0.003 0.048 ± 0.009
385.2 ± 0.4 385.6 ± 0.5
514.5 ± 0.3 528.9 ± 0.8
22.3 ± 0.7 22.7 ± 0.6
representative SEM morphologies of the Vickers indentations on the polished SAPS-W and SAPS-W/LaH2 coatings were shown in Fig. 7a and b, respectively. It can be seen from Fig. 7a that cracks were mainly initiated at all corners of the indented SAPS-W coating. Some corner cracks extended and connected with each other, leading to fracture of the coating surface. In comparison, for the SAPS-W/LaH2 coating, cracks were mainly initiated along the indentation edges and extended through the indentation; no significant fracture occurred on the coating surface, as shown in Fig. 7b. The result indicated that the SAPS-W/LaH2 coating had a higher ability to resist to crack propagation than the SAPSW coating. As shown in Fig. 4b, the bonding strength between the lamellae of the SAPS-W/LaH2 coating would be greater than that of the SAPS-W coating because of the joining effects of the formed La2O3 on the adjacent lamellae. Therefore, the ability of the SAPS-W/LaH2 coating to resist to elastic failure is greater than the SAPS-W coating. It has been pointed out that cracks in brittle solids induced by pyramidal indenters were ideal for toughness evaluation [40–42]. For a
median/radial crack system, the fracture toughness of a given brittle solids was evaluated by the following equation [40]. K C ¼ α ðE=HÞ1=2 P=c3=2
ð7Þ
where KC is the fracture toughness, α is an empirical constant which depends on the geometry of the indenter, E is the Young's modulus, H is the hardness, P is the peak indentation load, and c is the median/radial cracks. In the present work, the values of the peak indentation load P and the averaged lengths of the median/radial cracks c for the SAPS-W coating were 4.96 N and 19.25 ± 0.12 μm, respectively; and those for the SAPS-W/LaH2 coatings were 4.96 N and 17.08 ± 0.25 μm, respectively. Therefore, when the obtained Young's modulus (refers to E⁎ in the present work) and hardness H of the SAPS-W and SAPS-W/LaH2 coatings (Table 4) were substituted into Eq. (7), the fracture toughness of the two types of coating can be evaluated. The calculated KC values for the SAPS-W and SAPS-W/LaH2 coatings are about 309α ± 0.59α MPa·m1/2 and 322α ± 0.63α MPa·m1/2, respectively. It can be seen that the difference of the calculated KC values for the two types of coating is relatively small. Considering the complex and random shapes of the cracks for the two types of coating, it seems to be inappropriate to propose that the fracture toughness of the SAPS-W/LaH2 coating was higher than that of the SAPS-W coating. However, this is not the case from a statistical point of view, which can also be verified by the degree of cracking of the SAPS-W/LaH2 coating was lower than that of the SAPS-W coating, as illustrated in Fig. 7. The joint effect of the formed La2O3 on the adjacent lamellae is one of the main reasons. 4. Conclusions As a kind of rare earth hydride, 1.5 wt.% LaH2 powder were introduced into a tungsten powder and then sprayed by a supersonic atmospheric plasma spraying (SAPS) machine to form a coating. The aim was to form La2O3 in the LaH2-doped tungsten-based coating. The influence of the introduced LaH2 on the microstructure and properties of the tungsten coating was studied. It was obtained that the coatings were mainly composed of lamellar structure and columnar crystalline grains. WO3 formed in the two types of coatings. La2O3 phase, which was distributed at lamellar gaps of the LaH2-doped tungsten-based coating, was successfully prepared. The formed La2O3 had a filling effect on the adjacent lamellae, having a joint effect on the adjacent lamellae. The porosity, density, oxygen content, and thermal conductivity of the LaH2 -doped tungsten-based coating were all lower than those of the tungsten coating without LaH2 introduction. The LaH2-doped tungsten-based coating exhibited a higher ability to resist to elastic failure and a higher fracture toughness as compared with the tungsten coating without LaH2 introduction. Acknowledgements
Fig. 7. SEM micrograph of the indented surface of the SAPS coatings under 4.96 N: (a) W coating; (b) W/LaH2 coating.
This work was supported by Natural Science Foundation of Anhui Province of China (no. 1508085ME95), National Magnetic Confinement Fusion Program of China (no. 2014GB121001), China Postdoctoral
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