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Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2SO4-NaCl at 900 °C
Journal Pre-proof Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2 SO4 -NaCl at 900 null Y.F. Yang, Z.L. Liu, P. Ren, Q.W. Wang, Z.B. Bao, S.L. Zhu, W. Li
PII:
S0010-938X(19)32775-1
DOI:
https://doi.org/10.1016/j.corsci.2020.108527
Reference:
CS 108527
To appear in:
Corrosion Science
Received Date:
27 December 2019
Revised Date:
6 February 2020
Accepted Date:
7 February 2020
Please cite this article as: Yang YF, Liu ZL, Ren P, Wang QW, Bao ZB, Zhu SL, Li W, Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2 SO4 -NaCl at 900 x2103;, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108527
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Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2SO4-NaCl at 900 ℃
Y.F. Yang a, Z. L. Liu b, P. Ren a*,Q.W. Wang a*, Z.B. Bao b, S.L. Zhu b, W. Li a a
Institute of Advanced Wear & Corrosion Resistant and Functional Material, Jinan University, West Huangpu Road 601#, Guangzhou 510632, China b
Laboratory of Corrosion and Protection, Institute of Metal Research, Chinese
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Academy of Sciences, Wencui Road 62#, Shenyang 110016, China
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E-mail address: [email protected]
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Corresponding author. Tel: +86-20-85220890; Fax: +86-20-85220890
Co-Corresponding author. Tel: +86-20-85220890; Fax: +86-20-85220890
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E-mail address: [email protected]
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Graphical Abstract:
The schematic diagram shows hot corrosion behaviour of the Pt+Hf co-modified
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NiAl coatings in salt mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C. Firstly, the
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porous metastable oxide scale forms, which inlets tremendous amount of molten salts within the oxide as well as in the coating. Pores nucleates and grows rapidly at the oxide
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scale/coating interface and the prototype of embossments appears simultaneously (stage I). Secondly, metastable-to-stable oxide transformation initiates and HfO2 in Hf rich zone dissolves gradually. Other atoms, like Hf, Ni, Pt, Ti, diffuse outwardly and participate in the oxide scale. Meanwhile, chlorides and sulfides generate, which promote the development of embossments (stage II). As the hot corrosion proceeds
(stage III), the Hf rich zone disappears completely. An exclusive but porous stable αAl2O3 scale covers the coating surface. A continuous strip of slack sulphide forms at the oxide scale/ coating interface, which further aggravates the convex of the embossments.
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Highlights:
The addition of Hf in Pt modified NiAl coating changed its hot corrosion behavior.
The longer staying of metastable oxide in Pt+Hf co-modified NiAl coating led to
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an easier penetration of molten salt and promotes the nucleation and growth of
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pores, which accelerates the hot corrosion process of Pt+Hf co-modified NiAl coating.
The embossed oxide surface formed and developed during the hot corrosion
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procedure for Pt+Hf co-modified NiAl coating due to the formation slack Al2S3
Titanium diffused into the oxide scale due to invalidation of HfO2 diffusion barrier
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and gaseous chloride bubbles.
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by dissolving during the hot corrosion.
Abstract Hot corrosion behavior of Pt+Hf co-modified NiAl coating was evaluated in a salt
mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C in comparing to Pt modified NiAl coating. Worm-like embossments appeared and developed in Pt+Hf co-modified NiAl coating due to the formation of slack Al2S3 and gaseous chloride bubbles. The fast dissolution of HfO2 during hot corrosion led to the disappearing of diffusion barrier effect of Hf-rich zone and elements like, Ni, Ti and Pt diffused into the oxide scale. The hot corrosion mechanism of Pt+Hf co-modified NiAl coating and the effects of Hf on
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this process were discussed in this work.
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Keywords: Platinum; Metal coatings; TEM; EPMA; Hot corrosion
1. Introduction Aluminide coatings are widely used for the sake of protecting the turbine blades
against external oxidative and corrosive environment. Simple aluminide coating, mainly composed of β-NiAl, has drawn much attention due to its attractive properties such as low density, high melting point, and satisfactory mechanical properties at high temperatures. However, one of the greatest weaknesses of β-NiAl is that it is very sensitive to the sulphur-containing environment. Sulphur might exert two influences on the adhesion of oxide film: to enhance pore formation by lowering its nucleation, and
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to weaken the interface by accelerating crack propagation between pores [1-2].
In order to improve the high temperature performance of gas turbine, platinum
modified aluminide coatings have been intensively studied. Yang et al [3] suggested
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that Pt changes the growth mechanism of the oxide scale from anions and cations
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bilateral diffusion to anions unidirectional diffusion, which results in a decrease in the oxidation rate. Haynes and Leyens [4, 5] showed that the addition of Pt improves the
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adherence of oxide scale, however, it cannot change the scale growth rate compared to
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un-doped NiAl. Zhang [6] found that addition of Pt mitigates the detrimental influence of sulfur in the coating and thus improves the adhesion of the oxide scale. Svensson [7]
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proved that the beneficial effect of Pt on the adhesion energy is attributed to the increase of contact areas by theoretical density-functional theory (DFT) calculation.
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Despite the fact that the mechanisms of Pt on aluminide coating’s oxidation
behavior are still in debate, it is widely recognized that Pt could significantly improve the high temperature performance of aluminide coating. However, it also shows some weaknesses, such as the severe surface undulation, which is also known as “rumpling” [8-10]. To alleviate the “rumpling” behavior of NiAl coatings, reactive elements (REs)
such as Hf, Dy and Zr were also added in the coating [11-15]. Ye et al [16] fabricated Hf modified platinum aluminide coatings by introducing HfCl4 during aluminisation, and the oxidation results showed the rumpling extent of the oxide scale decreased. Same conclusion was also drawn in our previous work [13].These results demonstrated that the surface “rumpling” of Pt-Al coating during the exposure in the oxidative atmosphere could be suppressed effectively by Hf addition.
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One might know that land- or marine-based turbines often expose to the high sulphur-containing service environment, and this may lead to a lot of severe hot
corrosion issues [17-19]. According to our previous work [13], the Pt+Hf co-modified
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NiAl coating manifested itself with a satisfactory oxidation resistance compared to the
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(Ni,Pt)Al coating with lower oxidation rate and better oxide scale adherence. However, the hot corrosion resistance of the coating remains unknown. Deodeshmukh et al [20]
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has published some work on Pt and Hf co-modified γ'-Ni3Al+γ-Ni, but there are few
literatures.
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reports on the hot corrosion behavior of Pt and Hf co-modified β-NiAl coating in the
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In this work, efforts have been made to characterize the hot corrosion performance of Pt+Hf co-modified NiAl coating under chloride containing environment, which is
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very similar to the gas turbines’ service environment. The degradation process of Pt+Hf co-modified NiAl coating was analyzed by tracing the microstructure evolution and elements distribution by TEM and EPMA. The effect of Hf on the hot corrosion of (Ni,Pt)Al coating was discussed.
2. Experimental Ni-base single crystal superalloy with composition shown in Table 1 was used as the substrate material. Cylindrical specimens of Φ 15×2 mm were cut from [001]orientated crystal rods using a spark discharging machine. The samples were then grounded with a final 400# SiC sandpaper and humidly blasted using alumina grit with
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300 mesh in size. Before the electroplating of Pt layer and Pt+Hf composite layer, the samples were degreased firstly in boiling NaOH aqueous solution of 50 g L-1 for 10 min and then ultrasonically cleaned in acetone and ethanol for 15 min, respectively.
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Electroplating of Pt was conducted in an acid Pt-plating solution containing
mainly of K2[Pt(NO2)2SO4]. Pt+Hf composite plating layer was successfully obtained
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by adding micro-sized hafnium (0.1-10 μm in diameter) particles into the Pt-plating
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solution. The detailed process of electroplating of Pt and Pt+Hf can be referred to our previous publication [13, 21].
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An annealing treatment in vacuum (< 6×10-3 Pa) was conducted prior to the “above-pack” aluminisation at 600 ℃ for 2 h and 1050 ℃ for 1 h to eliminate the residual
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stress and hydrogen gas induced by electroplating as well as to dilute the concentration
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of Pt near the surface. Samples with Pt layer or Pt+Hf composite layer were aluminized at 1060 °C for 6 h. The preparation of the aluminisation medium can be referred elsewhere [22]. Cyclic hot corrosion resistance of the coatings was tested in static air at 900 °C using a muffle furnace. One hot corrosion cycle was started by depositing saturated aqueous salt-solution of 25 wt.% NaCl + 75 wt. % Na2SO4 onto all surface (including
the side cylindrical surface) of specimens on a heated tab with the total salt film ranged at 1-1.5 mg/cm2. The deposited samples were exposed at 900 °C for a given duration (e.g., 2 h, 5 h, 10 h, 20 h and every 20 h in the following exposure) and then taken out and cooled to room temperature. After each interval, the coating samples were washed in boiling deionized water for 30 min, then dried with flowing hot air. In order to avoid damage the oxide scale as possible, the samples are hanging by hooks during washing
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and drying process. After weighing, fresh salt film was recoated prior to next cycle. A detailed schematic figure showing the process of this hot corrosion test can be refer to
the work of Wang [23]. To ensure the accuracy of experiment, three parallel samples of
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each coating were recorded to obtain an average value of mass change. The sensitivity
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of the electronic balance (Sartorius BP211D) is 0.01 mg.
Microstructure evolution was traced by surface and cross-sectional morphologies
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changing using a field-emission scanning electron microscope (SEM, Inspect F50, FEI
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Co., Hillsboro, OR) equipped with energy dispersive X-ray spectrometer (EDS, Oxford Instruments Co., Oxford, U.K.). Without special instruction, a secondary electron mode
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was used to examine surface morphology, while the back-scattered electron mode was adopted to observe cross-sectional images of coating samples. For preserving corrosion
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products, the samples for cross-sectional observation were deposited with a thin layer of electro-less Ni-plating and mounted in resin. Phase constitution after hot corrosion exposure was examined by an X-ray diffractometer (XRD, X’ Pert PRO, PANalytical, Almelo, Holland). Elemental qualitative distribution of the coating samples was carried out using an electron probe micro-analyzer (EPMA-1610, Shimadzu, J.P.). A
combination of scanning transmission electron microscopy (STEM) and EDS microanalysis were performed by the transmission electron microscope (TEM, JEOL 2100F, J.P.) to identify the microstructure and chemical composition of the oxide scale after hot corrosion. The TEM sample was prepared using a focused ion beam (FIB, FEI Helios 600i Nanolab Dual Beam).
3.1 Microstructure of the as-prepared coatings
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3. Results
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Fig. 1 shows the surface and cross-sectional morphologies of as-prepared Pt and Pt+Hf co-modified β-NiAl. From Fig.1a and Fig.1b, well-defined grain boundaries can
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be observed in the surfaces of both coatings. Compared to the as-prepared Pt modified
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NiAl coating, the average grain size of Pt+Hf co-modified β-NiAl is smaller, which implies that Hf’s grain refinement effects. This is possible because the big hafnium
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atoms could segregate at the grain boundaries and suppress the grain growth of NiAl. Similar effect was found in Dy modified NiAl coating by Guo et al [24].
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The cross-sectional structures of Pt and Pt+Hf co-modified NiAl coating are
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similar as well (Fig. 1c and d).Precipitates are not seen in the outer (Ni,Pt)Al zone for Pt+Hf co-modified NiAl coating. The Hf content in this zone is 0.05 at. % according to the quantitative analysis results by EPMA in table 2. Practically, the content of Hf needs to be regulated at an optimal range (less than 0.5 at. %) to optimize the reactive element effect (REE) [25]. An additional Hf-rich zone could be observed with some dispersed brighter particles in Pt+Hf co-modified NiAl coating, which was proved to be HfO2 and
can act as the diffusion barrier according to one of our unpublished work [26]. The average Hf content in this zone is as high as 3.01 at. % (table 2).
3.2 Hot corrosion kinetics Fig. 2 shows the mass change curves of the two coatings during hot corrosion test in the mixed salt of Na2SO4/NaCl at 900 °C. For Pt modified NiAl coating, a slow and
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steady mass gain sustains until 280 h. Mass change during hot corrosion includes two opposite contributions, i.e., a mass gain owing to the formation of the oxides，and a
mass loss due to the evaporation of the salt and other volatile reaction products or due
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to the spallation of oxide scales. The stable mass gain for Pt modified NiAl coating
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indicates that the oxide spallation or the volatile products were not significant until 280 h. The mass gain for Pt+Hf co-modified NiAl coating is also gradually increasing in
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the first 160 h but mass change rate is much larger than that of Pt modified NiAl coating.
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Moreover, the weight change accelerated at 160 h and then increased linearly.
3.3 Phase constitution of corrosion products
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Fig. 3 shows the phase constitution of the two coatings after hot corrosion at
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900 °C in the mixed salt of Na2SO4/NaCl for 200 h. Exclusive α-Al2O3 formed on both modified NiAl coatings, which indicates a desirable Type I hot corrosion resistance. Peaks of NiAl in both coatings are obvious, and intensified peaks of γ/γ' phase appear as well due to the transformation from β-NiAl to γ'-Ni3Al. However, the strongest peaks remain to be NiAl for Pt modified NiAl coating. In contrast, the peaks of NiAl in Pt+Hf
co-modified coating are relatively weaker, suggesting a quicker Al loss rate for Pt+Hf co-modified NiAl coating.
3.4 Microstructure evolution during hot corrosion 3.4.1 Surface evolution Fig. 4 shows the surface evolution of both Pt modified and Pt+Hf co-modified
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NiAl coating during the hot corrosion. From Fig 4a and Fig. 4b with low magnification, an integral oxide scale formed on both coatings after hot corrosion for only 20 h. The
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oxide on Pt modified coating is grainy as shown in the inserted figure. In contrast, the needle-like oxide is recognized on the ridges of grainy oxide for Pt+Hf co-modified
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NiAl coating, which means the metastable oxide still exists. Another notable difference
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is the roughness of the surface. The surface of Pt modified NiAl is plain while wormlike embossments formed on the surface of Pt+Hf co-modified NiAl coating. When
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looking at the details in Fig 4b, cracks can be found at the top of the embossments, which may be resulted from the large stress concentrated there.
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After hot corrosion exposure for 200 h (Fig. 4c), the surface of Pt modified NiAl
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coating is still smooth but with some oxides spalling. For the Pt+Hf co-modified NiAl coating, the density and size of the surface embossments are further increased. Large areas of oxide spalled at the peak of the embossments owing to further accumulation of stress. It is intriguing to see that the brightness of the spallation zone is not uniform under the BSE mode of SEM (inserted figure in Fig. 4d). The brighter zone (tagged as zone 1#) is the second or multiple oxide spallation zone and the darker zone (tagged as
zone 3#) corresponds to the first oxide spallation zone. The EDS results in table 3 revealed the presence of Hf in these exfoliated areas, which indicates Hf atoms have migrated from the coating into the oxide scale. According to table 3, a higher Ni content was observed in zone 1#, which may account for its higher brightness and indicate that Ni has already participated in the oxide formation when multiple spallation occurs. Meanwhile, chlorine was also found in zone 1#, which comes from the residuals of
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gaseous chloride.
When the hot corrosion test reaches 300 h (Fig. 4e and 4f), more spallation zones can be observed. The newly formed oxide in the spallation zone of Pt modified NiAl is
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cauliflower-like, dense and very little gaps exists between cauliflowers. It is reasonable
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to predicate that fluxing molten salt can be effectively prevented from penetrating into the inner coating when the newly formed oxide grows continuously and successively
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in the following exposure. The surface of Pt+Hf co-modified NiAl coating became
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relatively smoother after exposure for 300 h when compared to samples exposed for 200 h. This superficial re-flatten of the surface might contribute to the further evolution
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of worm-like embossments, and thus lead to a uniform hunch-up of the entire surface. From the detailed spallation zone in Fig. 4f, it can be noticed that the newly formed
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oxide in the spallation zone is porous, which is consistent with the hunch-up of the surface.
3.4.2 Cross-sectional morphologies Fig. 5 shows the cross-sectional morphologies of two modified NiAl coating after
hot corrosion for 300 h. Obviously, the oxide scale formed on Pt modified NiAl coating is thinner than that of Pt+Hf co-modified NiAl, which agrees with the mass change curve in Fig. 2. The oxide scales are indented on both sides for the two coatings. This indicates that the oxide scale dissolved at the molten salt/oxide interface and grew internally at the oxide/coating interface. Low Al phase γ’-Ni3Al formed along the grain boundaries of the coating. According to the BSE morphology, γ’ is darker than β
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because of lower Pt solubility. The density of γ’ in Pt+Hf co-modified NiAl coating is
higher than that in Pt modified NiAl. This means that more Al was consumed for Pt+Hf
co-modified NiAl coating during the hot corrosion, which is consistent with the phase
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constitution results in Fig. 3.
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The internal sulfidation (pointed out in Fig. 5) seems more severe with a deeper and denser sulfidation zone for the Pt+Hf co-modified NiAl coating. The EDS analysis
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in the Hf rich zone (spot 3# in table 4) shows the Hf content is much lower than the as-
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prepared coating (table 2), indicating the degradation of the HfO2 diffusion barrier. The magnified cross-sectional morphologies in Fig. 5c and Fig. 5d displays the details of
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the oxide scale. For Pt modified NiAl coating, the oxide scale is compact with serrated oxide stretched into the coating. In comparison, a higher brightness and slack strip
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formed in the bottom of oxide scale formed on Pt+Hf co-modified coating (pointed out by red arrows in Fig. 5d). The EDS result reveals this strip is rich in S.
3.4.3 Elements distribution after hot corrosion Fig. 6 displayed the elemental mappings for the Pt and Pt+Hf co-modified NiAl
coating after hot corrosion at 900 °C in the mixed salt for 300 h by EPMA. The oxide formed on Pt modified NiAl is pure with little other elements distribution within the oxide scale. Instead, the oxide scale for Pt+Hf co-modified NiAl is mingled with some Cr, S, Hf and Ti in Fig 6b. Cl atoms from the molten salts penetrated into the coating and gathered in β phase. The distribution of S in Pt modified NiAl is uniform with only a few S rich sites (sulfidation zone) near the original coating/substrate interface. For
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Pt+Hf co-modified NiAl coating, an S-rich strip lies between the oxide/coating interface，which resembles the distribution of Al. This suggests that the strip with higher
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brightness in Fig. 5d is aluminum sulfide.
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4. Discussion
The above results manifest that both Pt and Pt+Hf co-modified NiAl coating
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possess a satisfactory Type I hot corrosion resistance in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) when compared with the plain NiAl and simple NiCrAlY
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coating[23,27-28]. The addition of Hf changes the hot corrosion behaviour of Pt
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modified NiAl coating. Therefore, the following discussion will be focused on exploring the hot corrosion mechanism of Pt+Hf co-modified NiAl coating and the
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effect of Hf on hot corrosion resistance.
4.1 Initial stage of Pt+Hf co-modified NiAl coating during hot corrosion There are two processes occurring simultaneously during the hot corrosion
exposure: oxide growth and oxide dissolution [18, 28]. When the coating is exposed to corrosive molten salts, oxides such as Al2O3, NiO, and HfO2 nucleate simultaneously due to the relatively high oxygen solubility in the molten salts. As the consumption of oxygen carries on, only Al2O3 could form selectively. On the other hand, the oxides are dissolved by the molten salt at the salts/oxide scale interface and parts of the molten salts penetrate the coating via the oxide scale [14, 28]. Therefore, the morphology and
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growth rate of the initial oxide scale influence the interactions between the molten salts and oxide scale.
It is generally recognized that Hf addition in aluminide coatings delays the
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metastable-to-stable Al2O3 phase transformation [29]. One of our previous work [14]
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has also proved that oxide transition process from θ-Al2O3 to α-Al2O3 of (Ni,Pt)Al coating is postponed significantly by Hf doping. In general, the metastable Al2O3 scale
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is less protective in comparison with stable α-Al2O3 scale due to the non-dense
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morphology. As a result, the presence of metastable Al2O3 is not conductive to prevent molten salts from entering the interior of coatings. According to the results in Fig. 4b,
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the needle-like θ-Al2O3 oxide remained at the surface of Pt+Hf co-modified NiAl coating after hot corrosion for 20 h. Therefore, the molten salt could penetrate more
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easily into the Pt+Hf coating through its porous metastable oxide scale than into the Pt modified coating. This leads to a severer oxide dissolution at the initial stage of hot corrosion. Besides, the nucleation of pores at the scale/coating interface is the most rapid during a cation-transported-alumina growth dominating process, i.e. θ-Al2O3 growing,
according to P.Y. Hou [30]. The existence of pores provide a direct aisle for the molten salts. Moreover, when S from the molten salts gathers at the scale/coating interface, it will in turn promote the formation of pores [2]. The long-time presence of θ-Al2O3 at Pt+Hf co-modified NiAl coating offers a habitat for pores nucleation and growing. Pores growth involves widening and deepening in the continued hot corrosion and more molten salts invaded into the coating during this growth process. Fig. 7 is the
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STEM morphology and Fig. 8 is the element distribution of the oxide scale formed on
Pt+Hf co-modified NiAl coating after hot corrosion for 200 h. The large-scaled pores
distribute in the oxide scale/coating interface and S permeates not only into the oxide
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scale but also into the coating. Theoretically, this mutual promotion process of molten
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dense and continuous α-Al2O3 scale.
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salts fluxing and pores formation can only be slowed down until the formation of a
4.2 Microstructure evolution of Pt+Hf co-modified NiAl coating
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during hot corrosion
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4.2.1 Formation of chlorides and sulfides The Pt+Hf co-modified NiAl coating also shows a different microstructure
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evolution during the hot corrosion process. The worm-like embossments formed and developed as shown above. The formation of embossments can be ascribed to the following reasons: Firstly, for hot corrosion in NaCl-contained mixed salt, the following reactions will occur concurrently with the dissolving of the oxide:
4NaCl (l)+O2 (g)↔2Na2 O+2Cl2 (g)
(1)
2 2 Cr (s)+Cl2 (g)↔ CrCl3 (g) 3 3 1 1 Hf (s)+Cl2 (g)↔ HfCl4 (g) 2 2 2 2 Al (s)+Cl2 (g)↔ AlCl (g) 3 3 3
(2) (3) (4)
The standard Gibbs free energy for the formation of chloride are calculated (standardized to consuming 1 mol Cl2) by using the software of HSC 6.0 (Fig .9).
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According to Fig. 9, the ΔGθ for all chlorides are negative, indicating that the formation of these products at 900 °C is thermodynamically feasible. The values of ΔGθ reflect
the potential of the reactions and the stability of the products. It is worth noting that the
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value of ΔGθ for the Hf -Cl2(g) reactions is more negative than that of Cr- Cl2 (g) or Al-
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Cl2 (g) reactions. This means that HfCl4 (g) is thermodynamically favourable to form than CrCl3 (g) or AlCl4 (g).
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When the gaseous chlorides form in the inner coating or beneath the oxide, they
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would diffuse outwardly through the cracks or boundaries to the surface. After that, they may be re-oxidized due to the surface’s higher oxygen potential. Actually, this
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process is detrimental to the hot corrosion resistance because gas bubbles will form when these gaseous chlorides diffused outwardly [31, 32]. Moreover, the bubbles
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increase the stress in the oxide scale, which may contribute to the formation of embossments. On the other hand, although most bubbles will release at the oxide scale/molten salt interface, a small part of bubbles are entrapped in the oxide and new pores form within the oxide scale. The distribution of Cl, Cr and Hf in the spalled zone in Fig. 4d verified the remains of chlorides. Moreover, the porous morphologies in Fig.
5c and Fig. 5d strongly support the bubble-assisted formation of pores. Besides, sulfides will also form via reactions between metals in the coating and S flowing into the coating or in the molten salts. The standard Gibbs free energy for the sulfides are also calculated (standardized to consuming 1 mol S) by HSC 6.0 and the results are showed in Fig .9. It can be seen that Al2S3 is the most stable state of sulphides. Combing the element distribution in Fig. 6 and the standard Gibbs free energy, it can
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speculate that the strip rich in S and Al in Fig. 6 is mainly Al2S3. However, Al2S3 is often thought to be non-protective and poorly-adherent [33]. Once it forms at the oxide scale/coating interface, the stress in the oxide scale will increase very quickly, which
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promotes the formation and development of embossments.
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4.2.2 The track of Ti during the hot corrosion process
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It is remarkable to see that the distribution of Ti for Pt and Pt+Hf co-modified NiAl coating are totally different. For Pt modified NiAl coating, the Ti atoms are mainly
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around the original surface of the substrate. However,for Pt+Hf co-modified NiAl coating，the Ti atoms diffuse into the coating as well as into the oxide scale (Fig. 6b
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and Fig. 8). According to Tawancy and Shirvani [34, 35], Pt could act as a diffusion-
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inhibitor for refractory transition elements by restricting their transference in β-NiAl. Hence, the movement of Ti in Pt modified NiAl coating was inactive. However, a strong affinity exists between Hf and Ti and an infinite solid solution
can be formed [36]. Theoretically, the mobility of Ti in Pt+Hf co-modified NiAl coating can be further constrained due to the existence of HfO2 diffusion barrier in the Hf rich zone. This may be true for coatings exposed in the oxidative atmosphere, where Ti was
limited to the coating/substrate interface and Hf rich zone can be maintained even after oxidation at 1100 ℃ for 300 h [13]. However, for Pt+Hf co-modified coating exposed to the hot corrosion environment, the HfO2 in Hf rich zone can be easily dissolved according to the acid/basic fluxing. As a result, the free hafnium ions and Ti atoms diffuse outwardly. As titanium has a strong affinity with O and S [37], trace amounts of
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Ti diffuse quickly to react with O and S and settle down in the oxide scale (Table 5).
4.2.3 Schematic diagram of Pt+Hf co-modified NiAl coating during hot corrosion
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Based on the above discussion, the hot corrosion process of Pt+Hf co-modified NiAl coating can be described as the schematic diagram shown in Fig. 10. Firstly, the
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porous metastable oxide scale forms, which cannot inhibit molten salt from diffusing
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into the oxide and the coating. Pores nucleate and grow rapidly at the oxide scale/coating interface and the prototype of embossments appear (stage I). Secondly,
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metastable-to-stable oxide transformation begins. More molten salts flux into the inner coating via grain boundaries and defects of the coating, leading to a gradual dissolution
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of HfO2 in the Hf rich zone. Other elements, like Hf, Ni, Pt from the coating and even
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Ti from the substrate, diffuse outwardly and participate in the oxide scale formation. Meanwhile, chlorides and sulphides form within the inner coating, which promote the development of embossments (stage II). As the hot corrosion proceeds (stage III), the Hf rich zone will disappear completely. An exclusive stable α-Al2O3 scale covers the coating surface. However, due to the motion of chlorides bubbles, the oxide scale become porous. A continuous strip of slack sulphide forms at the oxide scale/ coating
interface, which further promotes the convex of the embossments. When the stress within the embossments accumulates high enough, spallation occurs.
5. Conclusion Hot corrosion behavior of the Pt+Hf co-modified NiAl coating was evaluated in a salt mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C in comparison with the
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(Ni,Pt)Al coating. Based on the experimental results, the following conclusions could be drawn:
(1) Hf could change the hot corrosion behavior of Pt modified NiAl coating.
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(2) The longer stay of metastable oxide in Pt+Hf co-modified NiAl coating leads
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to an easier penetration of molten salt and promotes the nucleation and growth of pores. This accelerates the hot corrosion process of Pt+Hf co-modified NiAl coating.
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(3) An embossed oxide surface form and develop during the hot corrosion procedure for Pt+Hf co-modified NiAl coating, which is due to the formation of slack
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Al2S3 and gaseous chloride bubbles.
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(4) The HfO2 diffusion barrier degrades quickly during the hot corrosion and titanium from the substrate appears in the oxide scale after hot corrosion exposure for
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200 h.
Author statement
Y.F. Yang: Main contributor and original draft preparation; Z.L. Liu: Data curation; Experiment participator;
P. Ren: Conceptualization; Writing- Reviewing and Editing Q.W. Wang: Conceptualization; Funding acquisition; Z.B. Bao: Writing- Reviewing and Editing; S.L. Zhu: Supervision W. Li: Supervision
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Declaration of interests
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The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgments
This work was supported by “National Key Research and Development Program”
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(2018YFB2002000), the “Fundamental Research Funds for the Central Universities” (21619334), the Scientific Research Funds for the Talents and Innovation Foundation
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of Jinan University.
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Data Availability
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All the data included in this article are available upon request by contact the corresponding author.
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turbine engine applications, Surf. Coat. Technol., 139 (2001) 272-283. [15] C.Y .Jiang, Y. Y.Yang, Z.Y. Zhang, Z.B. Bao, M.H. Chen, S.L. Zhu, F.H. Wang, A Zr-doped single-phase Pt-modified aluminide coating and the enhanced hot corrosion resistance, Corrs. Sci., 133 (2018) 406-416. [16] L.Y. Ye, H.F. Chen, G.Yang, B. Liu, Y.F. Gao, Oxidation behavior of Hf-modified platinum
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turbine engine applications, Surf. Coat. Technol., 139 (2001) 272-283. [28] Y.F.Yang, C.Y. Jiang, Z.Y. Zhang, Z.B. Bao, M.H. Chen, S.L. Zhu, F.H. Wang, Hot corrosion behavior of single-phase platinum-modified aluminide coatings: Effect of Pt content and preoxidation, Corros. Sci., 127 (2017) 82-90. [29] S. Hayashi, B. Gleeson, Early-Stage Oxidation Behavior of Pt-Modified γ'-Ni3Al-Based Alloys with and without Hf Addition, Oxid. Met., 71 (2009) 5-19. [30] P. Y. Hou, Y. Niu, C. Van Lienden, Analysis of Pore Formation at Oxide–Alloy Interfaces—I: Experimental Results on FeAl, Oxid. of Met., 59 (2003) 41-61.
900 and 1000 °C in air, Corros. Sci., 45 (2003) 1217-1226.
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[31] G.M. Liu, M.S. Li, Y.C. Zhou, Hot corrosion of Ti3SiC2-based ceramics coated with Na2SO4 at
[32] G.M. Liu, M.S. Li, Y.C. Zhou, Hot corrosion behavior of Ti3SiC2 in the mixture of Na2SO4– NaCl melts, J Eur. Ceram. Soc., 25 (2005) 1033-1039.
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[33]D.B. Lee, High temperature sulfidation and oxidation of sputter-deposited Nb−Al−Si coatings, Met. and Mater. Inter., 7 (2001) 461-466.
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[34] H. M. Tawancy, A. Ui-Hamid, N. M. Abbas, M. O. Aboelfotoh, Effect of platinum on the oxideto-metal adhesion in thermal barrier coating systems, J Mater.Sci., 43 (2008) 2978-2989.
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[35] K. Shirvani, S. Firouzi, A. Rashidghamat, Microstructures and cyclic oxidation behaviour of Pt-free and low-Pt NiAl coatings on the Ni-base superalloy Rene-80, Corros. Sci., 55 (2012) 378-
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[36] Z. Cai, M. Koike, H. Sato, M. Brezner, Q. Guo, M. Komatsu, O. Okuno, T. Okabe, Electrochemical characterization of cast Ti–Hf binary alloys, Acta Biomater., 1 (2005) 353-356.
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[37] H. Lai, P. Knutsson, K. Stiller, The influence of platinum on the oxidation and sodium sulfate
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induced hot corrosion of NiAl diffusion coatings, Mater. High Temp., 28 (2011) 302-308
Figure Caption Fig. 1 Surface and cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf co-modified NiAl (b, d) coating samples after aluminisation at 1060 °C for 6 h Fig. 2 Mass change of the coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C.
Fig. 3 XRD patterns of different coating specimens corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 200 h. Fig. 4 Surface morphologies of Pt modified NiAl (a-20 h, c-200 h, e-300 h) and Pt+Hf co-modified NiAl (b-20 h, d-200 h, f-300 h) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for different time, showing the evolution of the surface. Fig. 5 Cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf comodified NiAl (b, d) coating samples corroded in the mixed salt of
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Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 300 h. Fig. 6 Elemental mappings for the Pt modified NiAl (a) and Pt+Hf co-modified NiAl
(b) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 300 h.
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Fig. 7 The cross-sectional STEM morphology of the oxide scale formed on Pt+Hf co-
modified NiAl coating after hot corrosion in the mixed salt of Na Na2SO4/NaCl
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(75:25, wt./ wt.) at 900 °C for 200 h.
Fig. 8 The elemental distribution of the oxide scale formed on Pt+Hf co-modified
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NiAl coating after hot corrosion in the mixed salt of Na2SO44/NaCl (75:25, wt./ wt.) at 900 °C for 200 h.
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Fig. 9 Standard Gibbs free energy changes (ΔGθ) of forming sulphides and chlorides as a function of temperature (standardised to consuming 1 mol S or 1mol Cl2). Fig. 10 Schematic diagram showing the hot corrosion process of Pt+Hf co-modified
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NiAl coating in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C
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Fig. 1 Surface and cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf modified NiAl (b, d) coating samples after aluminisation at 1060 °C for 6 h.
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Fig. 2 Mass change of the coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C.
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Fig. 3 XRD patterns of different coating specimens corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 200 h.
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Fig. 4 Surface morphologies of Pt modified NiAl (a-20 h, c-200 h, e-300 h) and Pt+Hf modified NiAl (b-20 h, d-200 h, f-300 h) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for different time, showing the evolution of the surface.
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Fig. 5 Cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf modified NiAl (b, d) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 300 h.
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Fig. 7 The cross-sectional STEM morphology of the oxide scale formed on Pt+Hf co-modified NiAl coating after hot corrosion in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 200 h
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Fig. 8 The elemental distribution of the oxide scale formed on Pt+Hf co-modified NiAl coating after hot corrosion in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 200 h
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Fig. 9 Standard Gibbs free energy changes (ΔGθ) of forming sulphides and chlorides as a function of temperature (standardised to consuming 1 mol S or 1mol Cl2).
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Fig. 10 Schematic diagram showing the hot corrosion process of Pt+Hf co-modified NiAl coating in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C
Tables: Table 1, nominal composition of the single-crystal superalloy (wt.%) elements
Cr
Co
W
Mo
Ta
Ti
Al
Ni
nominal composition
12.0
9.0
3.7
1.9
5.0
4.0
3.6
Bal.
Table 2, EPMA quantitative analysis results for the tagged spots shown in Fig. 1d (at.%) Elements
Co
Al
Ti
Cr
Ni
Mo
Ta
W
Pt
Hf
1#
11.23
17.14
2.56
30.74
27.70
3.10
0.73
1.94
4.82
0.05
2#
5.36
32.49
4.92
4.16
40.93
0.00
0.38
0.00
8.75
3.01
Table 3, EDS results for the tagged spots shown in Fig. 4 (at.%) Al
S
Cl
Ti
Cr
Co
Ni
1#
60.36
31.84
0.32
0.08
0.41
0.79
0.44
2#
66.37
32.57
0.08
--
--
0.66
0.08
3#
62.85
34.30
0.10
--
0.20
1.60
0.03
4#
65.97
30.21
--
0.02
0.32
2.37
0.15
5#
64.89
29.95
--
0.17
0.50
3.26
0.13
Hf
Ta
Pt
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O
5.12
0.06
0.05
0.53
0.22
--
--
0.02
0.87
0.05
--
--
0.93
--
--
0.03
0.97
0.06
--
0.07
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spots
Table 4, EDS results for the tagged spots shown in Fig. 5 (at.%) O
Na
Al
S
Cl
Ti
Cr
Co
Ni
Hf
Ta
W
Pt
1#
7.53
0.53
30.19
0.21
0.23
3.93
4.70
4.30
37.91
--
0.34
0.05
10.09
2#
59.67
0.05
32.08
0.27
--
0.79
0.39
0.63
5.60
0.06
0.14
--
0.32
3#
16.28
2.76
7.08
8.90
0.20
14.59
6.60
6.49
27.92
0.24
2.59
3.65
2.71
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spots
Table 5, EDS results for the precipitates in Fig. 7 (at.%) Ni
precipitates 1
-50.40
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precipitates 2
Pt
Al
Ti
Cr
Co
O
S
--
3.09
55.28
1.91
--
8.44
35.80
18.73
7.26
5.26
4.52
1.28
--
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elements
12.55
PII:
S0010-938X(19)32775-1
DOI:
https://doi.org/10.1016/j.corsci.2020.108527
Reference:
CS 108527
To appear in:
Corrosion Science
Received Date:
27 December 2019
Revised Date:
6 February 2020
Accepted Date:
7 February 2020
Please cite this article as: Yang YF, Liu ZL, Ren P, Wang QW, Bao ZB, Zhu SL, Li W, Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2 SO4 -NaCl at 900 x2103;, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108527
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Hot corrosion behavior of Pt+Hf co-modified NiAl coating in the mixed salt of Na2SO4-NaCl at 900 ℃
Y.F. Yang a, Z. L. Liu b, P. Ren a*,Q.W. Wang a*, Z.B. Bao b, S.L. Zhu b, W. Li a a
Institute of Advanced Wear & Corrosion Resistant and Functional Material, Jinan University, West Huangpu Road 601#, Guangzhou 510632, China b
Laboratory of Corrosion and Protection, Institute of Metal Research, Chinese
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Academy of Sciences, Wencui Road 62#, Shenyang 110016, China
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E-mail address: [email protected]
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Corresponding author. Tel: +86-20-85220890; Fax: +86-20-85220890
Co-Corresponding author. Tel: +86-20-85220890; Fax: +86-20-85220890
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E-mail address: [email protected]
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Graphical Abstract:
The schematic diagram shows hot corrosion behaviour of the Pt+Hf co-modified
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NiAl coatings in salt mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C. Firstly, the
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porous metastable oxide scale forms, which inlets tremendous amount of molten salts within the oxide as well as in the coating. Pores nucleates and grows rapidly at the oxide
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scale/coating interface and the prototype of embossments appears simultaneously (stage I). Secondly, metastable-to-stable oxide transformation initiates and HfO2 in Hf rich zone dissolves gradually. Other atoms, like Hf, Ni, Pt, Ti, diffuse outwardly and participate in the oxide scale. Meanwhile, chlorides and sulfides generate, which promote the development of embossments (stage II). As the hot corrosion proceeds
(stage III), the Hf rich zone disappears completely. An exclusive but porous stable αAl2O3 scale covers the coating surface. A continuous strip of slack sulphide forms at the oxide scale/ coating interface, which further aggravates the convex of the embossments.
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Highlights:
The addition of Hf in Pt modified NiAl coating changed its hot corrosion behavior.
The longer staying of metastable oxide in Pt+Hf co-modified NiAl coating led to
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an easier penetration of molten salt and promotes the nucleation and growth of
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pores, which accelerates the hot corrosion process of Pt+Hf co-modified NiAl coating.
The embossed oxide surface formed and developed during the hot corrosion
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procedure for Pt+Hf co-modified NiAl coating due to the formation slack Al2S3
Titanium diffused into the oxide scale due to invalidation of HfO2 diffusion barrier
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and gaseous chloride bubbles.
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by dissolving during the hot corrosion.
Abstract Hot corrosion behavior of Pt+Hf co-modified NiAl coating was evaluated in a salt
mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C in comparing to Pt modified NiAl coating. Worm-like embossments appeared and developed in Pt+Hf co-modified NiAl coating due to the formation of slack Al2S3 and gaseous chloride bubbles. The fast dissolution of HfO2 during hot corrosion led to the disappearing of diffusion barrier effect of Hf-rich zone and elements like, Ni, Ti and Pt diffused into the oxide scale. The hot corrosion mechanism of Pt+Hf co-modified NiAl coating and the effects of Hf on
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this process were discussed in this work.
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Keywords: Platinum; Metal coatings; TEM; EPMA; Hot corrosion
1. Introduction Aluminide coatings are widely used for the sake of protecting the turbine blades
against external oxidative and corrosive environment. Simple aluminide coating, mainly composed of β-NiAl, has drawn much attention due to its attractive properties such as low density, high melting point, and satisfactory mechanical properties at high temperatures. However, one of the greatest weaknesses of β-NiAl is that it is very sensitive to the sulphur-containing environment. Sulphur might exert two influences on the adhesion of oxide film: to enhance pore formation by lowering its nucleation, and
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to weaken the interface by accelerating crack propagation between pores [1-2].
In order to improve the high temperature performance of gas turbine, platinum
modified aluminide coatings have been intensively studied. Yang et al [3] suggested
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that Pt changes the growth mechanism of the oxide scale from anions and cations
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bilateral diffusion to anions unidirectional diffusion, which results in a decrease in the oxidation rate. Haynes and Leyens [4, 5] showed that the addition of Pt improves the
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adherence of oxide scale, however, it cannot change the scale growth rate compared to
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un-doped NiAl. Zhang [6] found that addition of Pt mitigates the detrimental influence of sulfur in the coating and thus improves the adhesion of the oxide scale. Svensson [7]
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proved that the beneficial effect of Pt on the adhesion energy is attributed to the increase of contact areas by theoretical density-functional theory (DFT) calculation.
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Despite the fact that the mechanisms of Pt on aluminide coating’s oxidation
behavior are still in debate, it is widely recognized that Pt could significantly improve the high temperature performance of aluminide coating. However, it also shows some weaknesses, such as the severe surface undulation, which is also known as “rumpling” [8-10]. To alleviate the “rumpling” behavior of NiAl coatings, reactive elements (REs)
such as Hf, Dy and Zr were also added in the coating [11-15]. Ye et al [16] fabricated Hf modified platinum aluminide coatings by introducing HfCl4 during aluminisation, and the oxidation results showed the rumpling extent of the oxide scale decreased. Same conclusion was also drawn in our previous work [13].These results demonstrated that the surface “rumpling” of Pt-Al coating during the exposure in the oxidative atmosphere could be suppressed effectively by Hf addition.
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One might know that land- or marine-based turbines often expose to the high sulphur-containing service environment, and this may lead to a lot of severe hot
corrosion issues [17-19]. According to our previous work [13], the Pt+Hf co-modified
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NiAl coating manifested itself with a satisfactory oxidation resistance compared to the
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(Ni,Pt)Al coating with lower oxidation rate and better oxide scale adherence. However, the hot corrosion resistance of the coating remains unknown. Deodeshmukh et al [20]
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has published some work on Pt and Hf co-modified γ'-Ni3Al+γ-Ni, but there are few
literatures.
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reports on the hot corrosion behavior of Pt and Hf co-modified β-NiAl coating in the
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In this work, efforts have been made to characterize the hot corrosion performance of Pt+Hf co-modified NiAl coating under chloride containing environment, which is
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very similar to the gas turbines’ service environment. The degradation process of Pt+Hf co-modified NiAl coating was analyzed by tracing the microstructure evolution and elements distribution by TEM and EPMA. The effect of Hf on the hot corrosion of (Ni,Pt)Al coating was discussed.
2. Experimental Ni-base single crystal superalloy with composition shown in Table 1 was used as the substrate material. Cylindrical specimens of Φ 15×2 mm were cut from [001]orientated crystal rods using a spark discharging machine. The samples were then grounded with a final 400# SiC sandpaper and humidly blasted using alumina grit with
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300 mesh in size. Before the electroplating of Pt layer and Pt+Hf composite layer, the samples were degreased firstly in boiling NaOH aqueous solution of 50 g L-1 for 10 min and then ultrasonically cleaned in acetone and ethanol for 15 min, respectively.
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Electroplating of Pt was conducted in an acid Pt-plating solution containing
mainly of K2[Pt(NO2)2SO4]. Pt+Hf composite plating layer was successfully obtained
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by adding micro-sized hafnium (0.1-10 μm in diameter) particles into the Pt-plating
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solution. The detailed process of electroplating of Pt and Pt+Hf can be referred to our previous publication [13, 21].
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An annealing treatment in vacuum (< 6×10-3 Pa) was conducted prior to the “above-pack” aluminisation at 600 ℃ for 2 h and 1050 ℃ for 1 h to eliminate the residual
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stress and hydrogen gas induced by electroplating as well as to dilute the concentration
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of Pt near the surface. Samples with Pt layer or Pt+Hf composite layer were aluminized at 1060 °C for 6 h. The preparation of the aluminisation medium can be referred elsewhere [22]. Cyclic hot corrosion resistance of the coatings was tested in static air at 900 °C using a muffle furnace. One hot corrosion cycle was started by depositing saturated aqueous salt-solution of 25 wt.% NaCl + 75 wt. % Na2SO4 onto all surface (including
the side cylindrical surface) of specimens on a heated tab with the total salt film ranged at 1-1.5 mg/cm2. The deposited samples were exposed at 900 °C for a given duration (e.g., 2 h, 5 h, 10 h, 20 h and every 20 h in the following exposure) and then taken out and cooled to room temperature. After each interval, the coating samples were washed in boiling deionized water for 30 min, then dried with flowing hot air. In order to avoid damage the oxide scale as possible, the samples are hanging by hooks during washing
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and drying process. After weighing, fresh salt film was recoated prior to next cycle. A detailed schematic figure showing the process of this hot corrosion test can be refer to
the work of Wang [23]. To ensure the accuracy of experiment, three parallel samples of
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each coating were recorded to obtain an average value of mass change. The sensitivity
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of the electronic balance (Sartorius BP211D) is 0.01 mg.
Microstructure evolution was traced by surface and cross-sectional morphologies
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changing using a field-emission scanning electron microscope (SEM, Inspect F50, FEI
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Co., Hillsboro, OR) equipped with energy dispersive X-ray spectrometer (EDS, Oxford Instruments Co., Oxford, U.K.). Without special instruction, a secondary electron mode
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was used to examine surface morphology, while the back-scattered electron mode was adopted to observe cross-sectional images of coating samples. For preserving corrosion
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products, the samples for cross-sectional observation were deposited with a thin layer of electro-less Ni-plating and mounted in resin. Phase constitution after hot corrosion exposure was examined by an X-ray diffractometer (XRD, X’ Pert PRO, PANalytical, Almelo, Holland). Elemental qualitative distribution of the coating samples was carried out using an electron probe micro-analyzer (EPMA-1610, Shimadzu, J.P.). A
combination of scanning transmission electron microscopy (STEM) and EDS microanalysis were performed by the transmission electron microscope (TEM, JEOL 2100F, J.P.) to identify the microstructure and chemical composition of the oxide scale after hot corrosion. The TEM sample was prepared using a focused ion beam (FIB, FEI Helios 600i Nanolab Dual Beam).
3.1 Microstructure of the as-prepared coatings
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3. Results
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Fig. 1 shows the surface and cross-sectional morphologies of as-prepared Pt and Pt+Hf co-modified β-NiAl. From Fig.1a and Fig.1b, well-defined grain boundaries can
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be observed in the surfaces of both coatings. Compared to the as-prepared Pt modified
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NiAl coating, the average grain size of Pt+Hf co-modified β-NiAl is smaller, which implies that Hf’s grain refinement effects. This is possible because the big hafnium
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atoms could segregate at the grain boundaries and suppress the grain growth of NiAl. Similar effect was found in Dy modified NiAl coating by Guo et al [24].
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The cross-sectional structures of Pt and Pt+Hf co-modified NiAl coating are
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similar as well (Fig. 1c and d).Precipitates are not seen in the outer (Ni,Pt)Al zone for Pt+Hf co-modified NiAl coating. The Hf content in this zone is 0.05 at. % according to the quantitative analysis results by EPMA in table 2. Practically, the content of Hf needs to be regulated at an optimal range (less than 0.5 at. %) to optimize the reactive element effect (REE) [25]. An additional Hf-rich zone could be observed with some dispersed brighter particles in Pt+Hf co-modified NiAl coating, which was proved to be HfO2 and
can act as the diffusion barrier according to one of our unpublished work [26]. The average Hf content in this zone is as high as 3.01 at. % (table 2).
3.2 Hot corrosion kinetics Fig. 2 shows the mass change curves of the two coatings during hot corrosion test in the mixed salt of Na2SO4/NaCl at 900 °C. For Pt modified NiAl coating, a slow and
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steady mass gain sustains until 280 h. Mass change during hot corrosion includes two opposite contributions, i.e., a mass gain owing to the formation of the oxides，and a
mass loss due to the evaporation of the salt and other volatile reaction products or due
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to the spallation of oxide scales. The stable mass gain for Pt modified NiAl coating
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indicates that the oxide spallation or the volatile products were not significant until 280 h. The mass gain for Pt+Hf co-modified NiAl coating is also gradually increasing in
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the first 160 h but mass change rate is much larger than that of Pt modified NiAl coating.
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Moreover, the weight change accelerated at 160 h and then increased linearly.
3.3 Phase constitution of corrosion products
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Fig. 3 shows the phase constitution of the two coatings after hot corrosion at
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900 °C in the mixed salt of Na2SO4/NaCl for 200 h. Exclusive α-Al2O3 formed on both modified NiAl coatings, which indicates a desirable Type I hot corrosion resistance. Peaks of NiAl in both coatings are obvious, and intensified peaks of γ/γ' phase appear as well due to the transformation from β-NiAl to γ'-Ni3Al. However, the strongest peaks remain to be NiAl for Pt modified NiAl coating. In contrast, the peaks of NiAl in Pt+Hf
co-modified coating are relatively weaker, suggesting a quicker Al loss rate for Pt+Hf co-modified NiAl coating.
3.4 Microstructure evolution during hot corrosion 3.4.1 Surface evolution Fig. 4 shows the surface evolution of both Pt modified and Pt+Hf co-modified
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NiAl coating during the hot corrosion. From Fig 4a and Fig. 4b with low magnification, an integral oxide scale formed on both coatings after hot corrosion for only 20 h. The
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oxide on Pt modified coating is grainy as shown in the inserted figure. In contrast, the needle-like oxide is recognized on the ridges of grainy oxide for Pt+Hf co-modified
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NiAl coating, which means the metastable oxide still exists. Another notable difference
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is the roughness of the surface. The surface of Pt modified NiAl is plain while wormlike embossments formed on the surface of Pt+Hf co-modified NiAl coating. When
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looking at the details in Fig 4b, cracks can be found at the top of the embossments, which may be resulted from the large stress concentrated there.
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After hot corrosion exposure for 200 h (Fig. 4c), the surface of Pt modified NiAl
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coating is still smooth but with some oxides spalling. For the Pt+Hf co-modified NiAl coating, the density and size of the surface embossments are further increased. Large areas of oxide spalled at the peak of the embossments owing to further accumulation of stress. It is intriguing to see that the brightness of the spallation zone is not uniform under the BSE mode of SEM (inserted figure in Fig. 4d). The brighter zone (tagged as zone 1#) is the second or multiple oxide spallation zone and the darker zone (tagged as
zone 3#) corresponds to the first oxide spallation zone. The EDS results in table 3 revealed the presence of Hf in these exfoliated areas, which indicates Hf atoms have migrated from the coating into the oxide scale. According to table 3, a higher Ni content was observed in zone 1#, which may account for its higher brightness and indicate that Ni has already participated in the oxide formation when multiple spallation occurs. Meanwhile, chlorine was also found in zone 1#, which comes from the residuals of
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gaseous chloride.
When the hot corrosion test reaches 300 h (Fig. 4e and 4f), more spallation zones can be observed. The newly formed oxide in the spallation zone of Pt modified NiAl is
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cauliflower-like, dense and very little gaps exists between cauliflowers. It is reasonable
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to predicate that fluxing molten salt can be effectively prevented from penetrating into the inner coating when the newly formed oxide grows continuously and successively
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in the following exposure. The surface of Pt+Hf co-modified NiAl coating became
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relatively smoother after exposure for 300 h when compared to samples exposed for 200 h. This superficial re-flatten of the surface might contribute to the further evolution
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of worm-like embossments, and thus lead to a uniform hunch-up of the entire surface. From the detailed spallation zone in Fig. 4f, it can be noticed that the newly formed
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oxide in the spallation zone is porous, which is consistent with the hunch-up of the surface.
3.4.2 Cross-sectional morphologies Fig. 5 shows the cross-sectional morphologies of two modified NiAl coating after
hot corrosion for 300 h. Obviously, the oxide scale formed on Pt modified NiAl coating is thinner than that of Pt+Hf co-modified NiAl, which agrees with the mass change curve in Fig. 2. The oxide scales are indented on both sides for the two coatings. This indicates that the oxide scale dissolved at the molten salt/oxide interface and grew internally at the oxide/coating interface. Low Al phase γ’-Ni3Al formed along the grain boundaries of the coating. According to the BSE morphology, γ’ is darker than β
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because of lower Pt solubility. The density of γ’ in Pt+Hf co-modified NiAl coating is
higher than that in Pt modified NiAl. This means that more Al was consumed for Pt+Hf
co-modified NiAl coating during the hot corrosion, which is consistent with the phase
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constitution results in Fig. 3.
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The internal sulfidation (pointed out in Fig. 5) seems more severe with a deeper and denser sulfidation zone for the Pt+Hf co-modified NiAl coating. The EDS analysis
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in the Hf rich zone (spot 3# in table 4) shows the Hf content is much lower than the as-
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prepared coating (table 2), indicating the degradation of the HfO2 diffusion barrier. The magnified cross-sectional morphologies in Fig. 5c and Fig. 5d displays the details of
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the oxide scale. For Pt modified NiAl coating, the oxide scale is compact with serrated oxide stretched into the coating. In comparison, a higher brightness and slack strip
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formed in the bottom of oxide scale formed on Pt+Hf co-modified coating (pointed out by red arrows in Fig. 5d). The EDS result reveals this strip is rich in S.
3.4.3 Elements distribution after hot corrosion Fig. 6 displayed the elemental mappings for the Pt and Pt+Hf co-modified NiAl
coating after hot corrosion at 900 °C in the mixed salt for 300 h by EPMA. The oxide formed on Pt modified NiAl is pure with little other elements distribution within the oxide scale. Instead, the oxide scale for Pt+Hf co-modified NiAl is mingled with some Cr, S, Hf and Ti in Fig 6b. Cl atoms from the molten salts penetrated into the coating and gathered in β phase. The distribution of S in Pt modified NiAl is uniform with only a few S rich sites (sulfidation zone) near the original coating/substrate interface. For
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Pt+Hf co-modified NiAl coating, an S-rich strip lies between the oxide/coating interface，which resembles the distribution of Al. This suggests that the strip with higher
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brightness in Fig. 5d is aluminum sulfide.
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4. Discussion
The above results manifest that both Pt and Pt+Hf co-modified NiAl coating
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possess a satisfactory Type I hot corrosion resistance in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) when compared with the plain NiAl and simple NiCrAlY
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coating[23,27-28]. The addition of Hf changes the hot corrosion behaviour of Pt
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modified NiAl coating. Therefore, the following discussion will be focused on exploring the hot corrosion mechanism of Pt+Hf co-modified NiAl coating and the
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effect of Hf on hot corrosion resistance.
4.1 Initial stage of Pt+Hf co-modified NiAl coating during hot corrosion There are two processes occurring simultaneously during the hot corrosion
exposure: oxide growth and oxide dissolution [18, 28]. When the coating is exposed to corrosive molten salts, oxides such as Al2O3, NiO, and HfO2 nucleate simultaneously due to the relatively high oxygen solubility in the molten salts. As the consumption of oxygen carries on, only Al2O3 could form selectively. On the other hand, the oxides are dissolved by the molten salt at the salts/oxide scale interface and parts of the molten salts penetrate the coating via the oxide scale [14, 28]. Therefore, the morphology and
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growth rate of the initial oxide scale influence the interactions between the molten salts and oxide scale.
It is generally recognized that Hf addition in aluminide coatings delays the
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metastable-to-stable Al2O3 phase transformation [29]. One of our previous work [14]
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has also proved that oxide transition process from θ-Al2O3 to α-Al2O3 of (Ni,Pt)Al coating is postponed significantly by Hf doping. In general, the metastable Al2O3 scale
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is less protective in comparison with stable α-Al2O3 scale due to the non-dense
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morphology. As a result, the presence of metastable Al2O3 is not conductive to prevent molten salts from entering the interior of coatings. According to the results in Fig. 4b,
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the needle-like θ-Al2O3 oxide remained at the surface of Pt+Hf co-modified NiAl coating after hot corrosion for 20 h. Therefore, the molten salt could penetrate more
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easily into the Pt+Hf coating through its porous metastable oxide scale than into the Pt modified coating. This leads to a severer oxide dissolution at the initial stage of hot corrosion. Besides, the nucleation of pores at the scale/coating interface is the most rapid during a cation-transported-alumina growth dominating process, i.e. θ-Al2O3 growing,
according to P.Y. Hou [30]. The existence of pores provide a direct aisle for the molten salts. Moreover, when S from the molten salts gathers at the scale/coating interface, it will in turn promote the formation of pores [2]. The long-time presence of θ-Al2O3 at Pt+Hf co-modified NiAl coating offers a habitat for pores nucleation and growing. Pores growth involves widening and deepening in the continued hot corrosion and more molten salts invaded into the coating during this growth process. Fig. 7 is the
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STEM morphology and Fig. 8 is the element distribution of the oxide scale formed on
Pt+Hf co-modified NiAl coating after hot corrosion for 200 h. The large-scaled pores
distribute in the oxide scale/coating interface and S permeates not only into the oxide
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scale but also into the coating. Theoretically, this mutual promotion process of molten
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dense and continuous α-Al2O3 scale.
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salts fluxing and pores formation can only be slowed down until the formation of a
4.2 Microstructure evolution of Pt+Hf co-modified NiAl coating
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during hot corrosion
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4.2.1 Formation of chlorides and sulfides The Pt+Hf co-modified NiAl coating also shows a different microstructure
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evolution during the hot corrosion process. The worm-like embossments formed and developed as shown above. The formation of embossments can be ascribed to the following reasons: Firstly, for hot corrosion in NaCl-contained mixed salt, the following reactions will occur concurrently with the dissolving of the oxide:
4NaCl (l)+O2 (g)↔2Na2 O+2Cl2 (g)
(1)
2 2 Cr (s)+Cl2 (g)↔ CrCl3 (g) 3 3 1 1 Hf (s)+Cl2 (g)↔ HfCl4 (g) 2 2 2 2 Al (s)+Cl2 (g)↔ AlCl (g) 3 3 3
(2) (3) (4)
The standard Gibbs free energy for the formation of chloride are calculated (standardized to consuming 1 mol Cl2) by using the software of HSC 6.0 (Fig .9).
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According to Fig. 9, the ΔGθ for all chlorides are negative, indicating that the formation of these products at 900 °C is thermodynamically feasible. The values of ΔGθ reflect
the potential of the reactions and the stability of the products. It is worth noting that the
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value of ΔGθ for the Hf -Cl2(g) reactions is more negative than that of Cr- Cl2 (g) or Al-
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Cl2 (g) reactions. This means that HfCl4 (g) is thermodynamically favourable to form than CrCl3 (g) or AlCl4 (g).
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When the gaseous chlorides form in the inner coating or beneath the oxide, they
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would diffuse outwardly through the cracks or boundaries to the surface. After that, they may be re-oxidized due to the surface’s higher oxygen potential. Actually, this
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process is detrimental to the hot corrosion resistance because gas bubbles will form when these gaseous chlorides diffused outwardly [31, 32]. Moreover, the bubbles
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increase the stress in the oxide scale, which may contribute to the formation of embossments. On the other hand, although most bubbles will release at the oxide scale/molten salt interface, a small part of bubbles are entrapped in the oxide and new pores form within the oxide scale. The distribution of Cl, Cr and Hf in the spalled zone in Fig. 4d verified the remains of chlorides. Moreover, the porous morphologies in Fig.
5c and Fig. 5d strongly support the bubble-assisted formation of pores. Besides, sulfides will also form via reactions between metals in the coating and S flowing into the coating or in the molten salts. The standard Gibbs free energy for the sulfides are also calculated (standardized to consuming 1 mol S) by HSC 6.0 and the results are showed in Fig .9. It can be seen that Al2S3 is the most stable state of sulphides. Combing the element distribution in Fig. 6 and the standard Gibbs free energy, it can
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speculate that the strip rich in S and Al in Fig. 6 is mainly Al2S3. However, Al2S3 is often thought to be non-protective and poorly-adherent [33]. Once it forms at the oxide scale/coating interface, the stress in the oxide scale will increase very quickly, which
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promotes the formation and development of embossments.
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4.2.2 The track of Ti during the hot corrosion process
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It is remarkable to see that the distribution of Ti for Pt and Pt+Hf co-modified NiAl coating are totally different. For Pt modified NiAl coating, the Ti atoms are mainly
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around the original surface of the substrate. However,for Pt+Hf co-modified NiAl coating，the Ti atoms diffuse into the coating as well as into the oxide scale (Fig. 6b
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and Fig. 8). According to Tawancy and Shirvani [34, 35], Pt could act as a diffusion-
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inhibitor for refractory transition elements by restricting their transference in β-NiAl. Hence, the movement of Ti in Pt modified NiAl coating was inactive. However, a strong affinity exists between Hf and Ti and an infinite solid solution
can be formed [36]. Theoretically, the mobility of Ti in Pt+Hf co-modified NiAl coating can be further constrained due to the existence of HfO2 diffusion barrier in the Hf rich zone. This may be true for coatings exposed in the oxidative atmosphere, where Ti was
limited to the coating/substrate interface and Hf rich zone can be maintained even after oxidation at 1100 ℃ for 300 h [13]. However, for Pt+Hf co-modified coating exposed to the hot corrosion environment, the HfO2 in Hf rich zone can be easily dissolved according to the acid/basic fluxing. As a result, the free hafnium ions and Ti atoms diffuse outwardly. As titanium has a strong affinity with O and S [37], trace amounts of
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Ti diffuse quickly to react with O and S and settle down in the oxide scale (Table 5).
4.2.3 Schematic diagram of Pt+Hf co-modified NiAl coating during hot corrosion
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Based on the above discussion, the hot corrosion process of Pt+Hf co-modified NiAl coating can be described as the schematic diagram shown in Fig. 10. Firstly, the
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porous metastable oxide scale forms, which cannot inhibit molten salt from diffusing
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into the oxide and the coating. Pores nucleate and grow rapidly at the oxide scale/coating interface and the prototype of embossments appear (stage I). Secondly,
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metastable-to-stable oxide transformation begins. More molten salts flux into the inner coating via grain boundaries and defects of the coating, leading to a gradual dissolution
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of HfO2 in the Hf rich zone. Other elements, like Hf, Ni, Pt from the coating and even
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Ti from the substrate, diffuse outwardly and participate in the oxide scale formation. Meanwhile, chlorides and sulphides form within the inner coating, which promote the development of embossments (stage II). As the hot corrosion proceeds (stage III), the Hf rich zone will disappear completely. An exclusive stable α-Al2O3 scale covers the coating surface. However, due to the motion of chlorides bubbles, the oxide scale become porous. A continuous strip of slack sulphide forms at the oxide scale/ coating
interface, which further promotes the convex of the embossments. When the stress within the embossments accumulates high enough, spallation occurs.
5. Conclusion Hot corrosion behavior of the Pt+Hf co-modified NiAl coating was evaluated in a salt mixture of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C in comparison with the
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(Ni,Pt)Al coating. Based on the experimental results, the following conclusions could be drawn:
(1) Hf could change the hot corrosion behavior of Pt modified NiAl coating.
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(2) The longer stay of metastable oxide in Pt+Hf co-modified NiAl coating leads
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to an easier penetration of molten salt and promotes the nucleation and growth of pores. This accelerates the hot corrosion process of Pt+Hf co-modified NiAl coating.
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(3) An embossed oxide surface form and develop during the hot corrosion procedure for Pt+Hf co-modified NiAl coating, which is due to the formation of slack
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Al2S3 and gaseous chloride bubbles.
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(4) The HfO2 diffusion barrier degrades quickly during the hot corrosion and titanium from the substrate appears in the oxide scale after hot corrosion exposure for
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200 h.
Author statement
Y.F. Yang: Main contributor and original draft preparation; Z.L. Liu: Data curation; Experiment participator;
P. Ren: Conceptualization; Writing- Reviewing and Editing Q.W. Wang: Conceptualization; Funding acquisition; Z.B. Bao: Writing- Reviewing and Editing; S.L. Zhu: Supervision W. Li: Supervision
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Declaration of interests
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The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgments
This work was supported by “National Key Research and Development Program”
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(2018YFB2002000), the “Fundamental Research Funds for the Central Universities” (21619334), the Scientific Research Funds for the Talents and Innovation Foundation
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of Jinan University.
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Data Availability
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All the data included in this article are available upon request by contact the corresponding author.
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Figure Caption Fig. 1 Surface and cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf co-modified NiAl (b, d) coating samples after aluminisation at 1060 °C for 6 h Fig. 2 Mass change of the coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C.
Fig. 3 XRD patterns of different coating specimens corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 200 h. Fig. 4 Surface morphologies of Pt modified NiAl (a-20 h, c-200 h, e-300 h) and Pt+Hf co-modified NiAl (b-20 h, d-200 h, f-300 h) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for different time, showing the evolution of the surface. Fig. 5 Cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf comodified NiAl (b, d) coating samples corroded in the mixed salt of
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Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 300 h. Fig. 6 Elemental mappings for the Pt modified NiAl (a) and Pt+Hf co-modified NiAl
(b) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 300 h.
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Fig. 7 The cross-sectional STEM morphology of the oxide scale formed on Pt+Hf co-
modified NiAl coating after hot corrosion in the mixed salt of Na Na2SO4/NaCl
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(75:25, wt./ wt.) at 900 °C for 200 h.
Fig. 8 The elemental distribution of the oxide scale formed on Pt+Hf co-modified
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NiAl coating after hot corrosion in the mixed salt of Na2SO44/NaCl (75:25, wt./ wt.) at 900 °C for 200 h.
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Fig. 9 Standard Gibbs free energy changes (ΔGθ) of forming sulphides and chlorides as a function of temperature (standardised to consuming 1 mol S or 1mol Cl2). Fig. 10 Schematic diagram showing the hot corrosion process of Pt+Hf co-modified
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NiAl coating in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C
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Fig. 1 Surface and cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf modified NiAl (b, d) coating samples after aluminisation at 1060 °C for 6 h.
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Fig. 2 Mass change of the coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C.
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Fig. 3 XRD patterns of different coating specimens corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 200 h.
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Fig. 4 Surface morphologies of Pt modified NiAl (a-20 h, c-200 h, e-300 h) and Pt+Hf modified NiAl (b-20 h, d-200 h, f-300 h) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for different time, showing the evolution of the surface.
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Fig. 5 Cross-sectional morphologies of Pt modified NiAl (a, c) and Pt+Hf modified NiAl (b, d) coating samples corroded in the mixed salt of Na2SO4/NaCl (75:25, wt./wt.) at 900 °C for 300 h.
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Fig. 7 The cross-sectional STEM morphology of the oxide scale formed on Pt+Hf co-modified NiAl coating after hot corrosion in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 200 h
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Fig. 8 The elemental distribution of the oxide scale formed on Pt+Hf co-modified NiAl coating after hot corrosion in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C for 200 h
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Fig. 9 Standard Gibbs free energy changes (ΔGθ) of forming sulphides and chlorides as a function of temperature (standardised to consuming 1 mol S or 1mol Cl2).
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Fig. 10 Schematic diagram showing the hot corrosion process of Pt+Hf co-modified NiAl coating in the mixed salt of Na2SO4/NaCl (75:25, wt./ wt.) at 900 °C
Tables: Table 1, nominal composition of the single-crystal superalloy (wt.%) elements
Cr
Co
W
Mo
Ta
Ti
Al
Ni
nominal composition
12.0
9.0
3.7
1.9
5.0
4.0
3.6
Bal.
Table 2, EPMA quantitative analysis results for the tagged spots shown in Fig. 1d (at.%) Elements
Co
Al
Ti
Cr
Ni
Mo
Ta
W
Pt
Hf
1#
11.23
17.14
2.56
30.74
27.70
3.10
0.73
1.94
4.82
0.05
2#
5.36
32.49
4.92
4.16
40.93
0.00
0.38
0.00
8.75
3.01
Table 3, EDS results for the tagged spots shown in Fig. 4 (at.%) Al
S
Cl
Ti
Cr
Co
Ni
1#
60.36
31.84
0.32
0.08
0.41
0.79
0.44
2#
66.37
32.57
0.08
--
--
0.66
0.08
3#
62.85
34.30
0.10
--
0.20
1.60
0.03
4#
65.97
30.21
--
0.02
0.32
2.37
0.15
5#
64.89
29.95
--
0.17
0.50
3.26
0.13
Hf
Ta
Pt
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O
5.12
0.06
0.05
0.53
0.22
--
--
0.02
0.87
0.05
--
--
0.93
--
--
0.03
0.97
0.06
--
0.07
-p
spots
Table 4, EDS results for the tagged spots shown in Fig. 5 (at.%) O
Na
Al
S
Cl
Ti
Cr
Co
Ni
Hf
Ta
W
Pt
1#
7.53
0.53
30.19
0.21
0.23
3.93
4.70
4.30
37.91
--
0.34
0.05
10.09
2#
59.67
0.05
32.08
0.27
--
0.79
0.39
0.63
5.60
0.06
0.14
--
0.32
3#
16.28
2.76
7.08
8.90
0.20
14.59
6.60
6.49
27.92
0.24
2.59
3.65
2.71
lP
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spots
Table 5, EDS results for the precipitates in Fig. 7 (at.%) Ni
precipitates 1
-50.40
Jo
ur
precipitates 2
Pt
Al
Ti
Cr
Co
O
S
--
3.09
55.28
1.91
--
8.44
35.80
18.73
7.26
5.26
4.52
1.28
--
na
elements
12.55