A comparative study on the high temperature corrosion of TP347H stainless steel, C22 alloy and laser-cladding C22 coating in molten chloride salts

Corrosion Science 83 (2014) 396–408

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Corrosion Science journal homepage: www.elsevier.com/locate/corsci

A comparative study on the high temperature corrosion of TP347H stainless steel, C22 alloy and laser-cladding C22 coating in molten chloride salts Shunv Liu, Zongde Liu ⇑, Yongtian Wang, Jin Tang Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of Education, North China Electric Power University, Beijing 102206, China

a r t i c l e

i n f o

Article history: Received 1 November 2013 Accepted 4 March 2014 Available online 12 March 2014 Keywords: A. Stainless steel A. Alloy A. Metal coatings C. High temperature corrosion C. Oxidation

a b s t r a c t Isothermal corrosion of TP347H (A1), C22 alloy (A2) and laser-cladding C22 coating (A3) was evaluated by mass loss measurements in molten alkali chloride salts at 450–750 °C. Corrosion mechanisms were characterised by scanning electron microscopy, optical microscopy and X-ray diffraction. A3 exhibited superior corrosion resistance, followed by A2, which results from alloying elements, refined microstructure and Cr–O (CrOx), Co(Fe, Cr)2O4 in the corrosion scale. Severe intergranular corrosion caused failure of A1, slight intergranular corrosion happened in A2 but none in A3. Fe-rich oxides were main products of A1 while NiO of A2 and A3 with Cr2O3 and Mo-containing compositions. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction The alkali metals and chlorine in many biomasses are potentially harmful elements with regard to high temperature corrosion. They originally combine into corrosive compositions in the flue gas and deposit on heat transfer surfaces through mechanisms such as condensation of vapours, sticking to particles and chemical reactions like sintering and slagging [1–3]. Chloride-rich deposits formed on superheater tubes are considered causing the severest corrosion problems in biomass-fired plants [4–7]. A typical superheater tube temperature in some biomass-fired plants is 400– 650 °C. Considering localised overheating, the tube heat exchange surface temperature may be above 700 °C. Higher temperatures increase the amount of molten phases in deposits and vapour phases in the flue gas. This aggravates the corrosion of superheater tubes. However, the superheated steam temperature may be elevated further due to the desire to improve electrical efficiency of biomass-fired plants. Nielsen et al. [4] provided the maximum service temperature of some common superheater tube materials in biomass power plants. The highest service temperature of all materials was below 600 °C. Hence, it is necessary to develop super corrosion resistant materials for superheater tubes in the future biomass-fired plants.

⇑ Corresponding author. Tel./fax: +86 10 61772812. E-mail address: [email protected] (Z. Liu). http://dx.doi.org/10.1016/j.corsci.2014.03.012 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.

Nowadays, austenitic stainless steel TP347H is widely used in superheater tubes due to its high cost performance. With a high content of Ni and Cr in composition, it exhibits a better corrosion resistance. Ni and Cr can both form dense protective oxide scales [8]. And the combination of them suitably matches the n/8 law [9], hence improving the passive performance of the steel dramatically. Ni also succeeds in preventing stress corrosion caused by Cl. Cr is good at preventing intergranular corrosion [10]. The alloying Nb also inhibits intergranular corrosion and stress corrosion. In addition, Ni promotes formation of austenite and then greatly improves mechanical properties such as plasticity and toughness [9]. Li [11] conducted a corrosion test at 650 °C, revealing that TP347H exhibited the best corrosion resistance among conventional superheater tube materials. But Faced with the high temperature molten salt attack induced by biomass combustion, even TP347H failed. In recent works, Hastelloy C22 alloy (one of Ni–Cr–Mo alloys) are widely studied for its excellent corrosion resistance under extreme conditions, especially for stress corrosion [12] and localised corrosion, such as pitting corrosion [13], intergranular corrosion [14] and crevice corrosion [15,16]. Cr is one of the main alloying elements ensuring better corrosion resistance of Hastelloy series. Moreover, high contents of Cr and Mo form a Cr-dominated passive oxide scale and Mo promotes passivation by maintaining a low passive current [17–21]. In addition, Mo also shows resistance to localised corrosion that it can promote fast repassivation by forming Mo insoluble compounds [13,22]. The Co also contributes

S. Liu et al. / Corrosion Science 83 (2014) 396–408

to high temperature oxidation resistance. Literatures have proved that C22 alloy is one of the few materials that are resistant to low temperature corrosion of seawater [23,24]. But few works were aimed at corrosion performance in high temperature environments, especially molten salt corrosion of biomass combustion. Therefore, this paper carries out this work. Considering the practical engineering application, in fact, C22 is not a suitable candidate for biomass plants due to the high cost of 50–60 thousand dollars per ton. To improve this situation, lasercladding technology shows dramatic potential to keep sound balance between the high cost and desired corrosion resistance. Generally, a thin coating of 300–500 lm could obtain an excellent corrosion resistance due to its inherent advantages, such as metallurgical bonding, pore-free, low dilution rate and refined microstructure [25–28]. Thus, to prepare a C22 coating on TP347H by laser-cladding can remarkably lower the cost and improve the corrosion resistance. Furthermore, it could also maintain the superior mechanical properties of substrate TP347H. Laser-cladding technology has been tested to be a success to prepare hot corrosion resistant coatings for gas turbines burning biomass and waste derived fuel gases [29], but few studies are conducted on corrosion resistant alloy coatings for superheater tubes. So In this work, a comparative study is carried out between the laser-cladding C22 coating, C22 alloy and TP347H stainless steel in molten alkali chloride salts, simulating corrosive environment of biomass-fired boilers. In order to systematically investigate the corrosion mechanisms of the three materials from low to high temperatures, a temperature arrange of 450–750 °C is tested. In addition, this work also compares the corrosivity of forestry and agricultural biomasses. The experimental results are discussed thoroughly. 2. Material and experimental methods The laser cladding system used in the current work was a selfassembly computer numerical control (CNC) installation equipped with a pulsed optical-pumped Nd: YAG laser (RH-700, China) as the energy source, a CNC multi-axis cladding workbench and an argon shielding system. The photograph of the cladding system is presented elsewhere [30]. The working parameters: electric current, 500 A; scan speed, 5 mm/s; single clad thickness, 300 lm. It was used for depositing C22 powders offered by Beijing General Research Institute of Mining and Metallurgy (BGRIMM). The sample substrate was prepared from a TP347H superheater tube, which was polished and then degreased with acetone before cladding. The sample A1 (TP347H, 20 mm  10 mm  5 mm) was cut from the tube. The C22 powders (300 meshes) used for laser cladding were deposited on the surface of TP347H superheater tube. The C22 coating is about 3.5 mm thick, prepared by multitrack cladding process, and then cut into the sample A3 (laser-cladding C22 coating, 20 mm  10 mm  2 mm). The sample A2 (C22 alloy, 20 mm  10 mm  3 mm) was cut from a forge piece made of pure C22 alloy. Totally, fourteen specimens were employed for

Table 1 The initial masses of specimens involved in two corrosion reagents at different test temperatures, g. Test sample

Temperature (°C) 450

500

550

600

650

700

750

A1 A1 A2 A2 A3 A3

7.8192 7.9346 5.1173 5.1178 3.4199 3.5576

8.7172 8.7679 5.1033 5.0754 2.7731 3.0008

8.6414 8.5698 5.0764 5.0632 3.4116 3.8324

7.3473 7.3754 5.0999 5.0329 3.5138 3.0481

8.8574 8.0499 4.8118 5.0804 3.4597 3.4363

8.4458 8.8870 5.0214 5.0580 3.4918 3.3167

8.8593 8.8599 5.0688 5.0479 3.2434 3.3344

(R1) (R2) (R1) (R2) (R1) (R2)

397

each sample (one for each corrosion reagent at every test temperature). Table 1 gives the initial masses of all the specimens for three samples. The compositions of TP347H steel, C22 alloy and C22 powder are shown in Table 2. Before corrosion test, all specimens were degreased with acetone in an ultrasonic bath, then cleaned with deionised water and dried in a desiccator. A mixture of 98.6 wt.% KCl and 1.4 wt.% NaCl (solid, marked as R1) was used to simulate the molten salt corrosion induced by forestry biomass combustion, while a mixture of 95.5 wt.% KCl and 4.5 wt.% NaCl (R2) for simulating agricultural biomass combustion [11,31], according to the criterion ASTME1755-01 [32] and ASTME0870-82R98E01 [33]. Before corrosion, the corrosion reagent R1 and R2 were put in two corundum crucible boats (80 mm  20 mm  15 mm) respectively. One specimen A1, A2 and A3 were then embedded in each reagent with a distance of about 5 mm. Owing to considering the experimental condition where the molten salts stayed in a flowing and volatile state, and isothermal area of the resistance furnace was confined to a restricted region, the three samples were placed in the same place in the crucible boat during every corrosion cycle at every test temperature, and the boat was also fixed in the same place at the isothermal area, in order to keep experimental reproducibility and data consistency. The isothermal corrosion was carried out in air using a tubular resistance furnace from 12 h a cycle up to 108 h at 450, 500, 550, 600, 650, 700 and 750 °C. The photograph of the resistance furnace is presented elsewhere [30]. For this furnace, electrothermal alloy 0Cr27Al7Mo2 is used as heating source. The test temperature is set through a temperature control instrument, which is connected to a thermocouple installed at the isothermal area of the resistance furnace. The thermocouple measures the temperature at the isothermal area, and transmits the temperature signal to the temperature control instrument. It then responds immediately to adjust the voltage on or off, maintaining the temperature at the isothermal area to fluctuate within the set temperature ±1 °C. After the set temperature became completely stable, the corundum crucible boat was put at the isothermal area with a self-welded iron carrier. When a corrosion cycle was finished, the crucible boat was taken out immediately and left cooling at room temperature. The melting point of two corrosion reagents was different and the melting situation of them at every test temperature was listed in Table 3. Since the reagents melted at higher temperatures, the reagents should be renewed before each corrosion cycle. The corrosion products were easy to spall and some residual corrosion reagents stuck to corrosion products. Under these circumstances, mass gain cannot be measured accurately. Therefore, mass loss (g/m2) was measured replacing the mass gain to characterise the corrosion resistance of samples according to the equation [34]:

ccorr ¼

Dm A

ð1Þ

where Dm (g) is the cumulative mass loss determined simultaneously with increasing corrosion time, and A is the surface area of samples (0.0007 m2, 0.00058 m2 and 0.00052 m2 for A1, A2 and A3, respectively) measured prior to corrosion in order to provide comparability of the results. The mass was evaluated using an electronic micro-balance (accuracy: ±0.01 mg). As minor corrosion and no spalling were observed in A1 at 450 °C and in A2 and A3 at 450– 500 °C, mass loss measurements were not carried out. After cooling off, the specimens were taken out from the corrosion reagents, blown with a rubber suction bulb and then slightly knocked on the surfaces with tweezers to remove the corrosion products which were collected in beakers. And then the specimens were pickled in the hydrochloric acid solution bath (25 wt.%, 80 °C) for further removing the corrosion products [35]. The hydrochloric acid was selected as pickling solution in order to avoid introducing

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S. Liu et al. / Corrosion Science 83 (2014) 396–408 Table 2 Chemical compositions of test alloys. Alloy

Element (wt.%)

C22 powder C22 forge TP347H

Ni

C

Cr

Mo

Mn

Fe

Si

Co

V

W

Nb

S

P

Bal. Bal. 12.75

0.08 0.001 0.1

21.3 22.0 19.2

13.2 13.8 –

– 0.45 2.0

2.93 5.0 Bal.

– 0.078 0.09

2.0 1.83 –

– 0.3 –

3.0 3.2 –

– – 1.4

0.008 0.03

0.020 0.032

Table 3 Melting situations of corrosion reagent R1 (98.6 wt.% KCl and 1.4 wt.% NaCl) and R2 (95.5 wt.% KCl and 4.5 wt.% NaCl) in air at different temperatures. The melt transforms solid reagents into liquids and compositions do not change. Reagent

R1 R2

Temperature (°C) 450

500

550

600

650

700

750

No melt No melt

No melt No melt

No melt No melt

No melt Slight melt

Slight melt Partial melt

Partial melt Complete melt

Complete melt Complete melt

impurities since the corrosion reagents were chlorides. After corrosion, FexOy was formed as main constituent in the corrosion products of A1, while NiO of A2 and A3 (discussed later). FexOy was easier to be dissolved than NiO by hydrochloric acid. So the pickling time of A1 should be relatively shorter than that of A2 and A3. When samples were corroded at higher test temperatures, the generated corrosion products usually needed longer times to be pickled off. The practical pickling test proved that it was proper to control the pickling time within 4–8 min for A1, 10–25 min for A2 and A3, where the corrosion products can be removed and over pickling can also be avoided. The pickled specimens were then ultrasonically cleaned in deionised water, dried and weighed. After finishing the last corrosion cycle, the corrosion products were not collected and no specimens were pickled in order to observe the morphologies of corrosion products by scanning electron microscopy (SEM, S-4800 HITACHI) in secondary electron (SE) mode. Chemical compositions were determined by the energy dispersive X-ray spectroscopy (EDS, Bruker). The e-beam resolution of EDS used is 123 eV. The acceleration voltage is set at 20 kV and the spot size of e-beam is about 1 lm. It has a detection limit of 0.1%. That is, the detected element generally has a content above 0.1 wt.%. Generally, elements with an atomic number above 10 could be relatively accurately detected while other slight elements cannot be. As for the EDS spot analysis, it detects the elemental composition of a circular area which centres on this point and has a diameter of about 1 lm. Since the EDS spot scan can give a semi-quantitative elementary analysis at selected points, it still has certain research values in this work. The cross-sections of specimens were wet ground through successive SiC papers from 400# to 1200#, followed by polishing, drying, and then observed by optical microscopy (4XBC). Since many residual corrosion reagents were mixed into the collected corrosion products, they were washed with deionised water, dried and then pulverized into powders for phase identification which was conducted by X-ray diffraction (XRD) meter with Cu Ka incident radiation operating at a scan speed of 8°/min and 2h from 10° to 90° at 40 kV and 100 mA. 3. Results 3.1. Corrosion of TP347H stainless steel Fig. 1 presents the mass loss of TP347H stainless steel (A1) corroded in two corrosion reagents in air for 96 h at 500–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 4–8 min. The mass loss curves are almost straight lines at 500–650 °C in two corrosion reagents and an increasing corrosion time leads to a continuous mass loss. As the slope of curves

increases with the increase of temperatures, the mass loss is higher (hence higher corrosion rate). In order to avoid over pickling, a small amount of corrosion products on the surfaces of specimens are not removed, hence decreasing the mass loss. Slagging forms at 650–750 °C (displayed later), which also affects the measurement of mass loss. Hence, the mass loss at 700 °C is lower than that at 650 °C after corrosion in R2 for 72 h, and the mass loss at 750 °C is lower at first than that at 650–700 °C in two reagents. Moreover, Intergranular corrosion happens in TP347H (displayed later). The corrosion products generated at grain boundaries cannot be removed by pickling. And, the surface area decreases much (at 700–750 °C, almost half of the original surface area) after corrosion for 96 h. Therefore, the actual mass lose may be twice higher (more at higher temperatures) than that illustrated in Fig. 1. As can be seen in two reagents at each temperature, a little different mass loss is observed. So the corrosivity of forestry and agricultural biomasses for TP347H is similar. Fig. 2 shows the SEM surface micrographs and EDS spot scan results of the corrosion products of TP347H stainless steel (A1) corroded in R1 in air for 12 h at different temperatures. The corrosion of A1 in R2 exhibits very similar corrosion behaviour and corrosion product compositions compared to that in R1, therefore only the latter is displayed in Fig. 2. At lower test temperatures (Fig. 2a), the specimen surface is covered by an even corrosion scale and some cracks are observed. Then some particle-shaped phases form at 550 °C (Fig. 2b). With the increase of test temperatures, the particle-shaped phase grows larger, and a multi-layered corrosion scale appears (Fig. 2c). According to EDS results, (Fe, Ni)-rich compositions are identified as major phases in the corrosion products. The Cr content in the outer corrosion product layer (spot A) is lower than that in the inner layer (spot B). This indicates that the surface protective Cr2O3 scale may be destroyed and the corrosion media permeate inwards to corrode substrate alloy. Some Cl is also identified in the inner layer, but none in the outer layer. This finding may suggest that some chlorides form, and they may volatilize or serve as intermediate products to be transformed into oxides at last (discussed later). Since no Mo in the composition of TP347H (Table 2), the identified Mo may come from C22 alloy and C22 coating because these three samples are embedded in one reagent near to each other. The element Nb in the corrosion products of C22 alloy while may come from TP347H (presented later, Fig. 8b). When the test temperature reaches 700 °C (for R2 is 650 °C), slagging which is caused by molten or semi-molten ash particles depositing on heating surfaces, forms on A1 (Fig. 3). In this study, molten KCl is observed (EDS results). The similar slagging morphology and composition are also detected in R2. When the slagging happens, the intergranular corrosion (displayed later) aggravates,

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7000

4000

2

5000

6000

mass loss (g/m )

2

mass loss (g/m )

6000

7000

(a)

o

500 C o 550 C o 600 C o 650 C o 700 C o 750 C

3000 2000

5000 4000 3000

(b)

2000

1000

1000

0

0

-12 0

o

500 C o 550 C o 600 C o 650 C o 700 C o 750 C

12 24 36 48 60 72 84 96 108

-12 0

t (h)

12 24 36 48 60 72 84 96 108

t (h)

Fig. 1. Mass loss (Dm/A) of TP347H stainless steel corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air at 500–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 4–8 min.

(a)

(b)

KCl

(c) A

54.11% Fe

53.30% Fe

27.67% Ni

B

40.76% Ni

4.80% Cr

0.86% Cr

1.26% Mn

1.12% Mn

0.28% Mo

0.05% Mo

1.49% Nb

3.39% C

3.90% Si

0.52% O

5.87% O 0.62% Cl

Fig. 2. Surface morphologies of corrosion products of TP347H stainless steel corroded in 98.6 wt.% KCl and 1.4 wt.% NaCl in air for 12 h at (a) 500 °C, (b) 550 °C and (c) 650 °C. EDS spot scan shows elemental compositions (wt.%) at spot A and B.

10.94% Fe

4.82% Fe

B

1.10% Cr 0.66% Mo 0.37% Nb 45.21% K 40.98% Cl 6.63% C 0.13% Si

0.01% Cr 0.90% Ni

0.09% Mn

A

0.76% Mn 0.86% Mo 0.03% Nb 28.91% K 29.52% Cl 0.01% Si 28.06% O

Fig. 3. Slagging morphology of surface corrosion products of TP347H stainless steel corroded in 98.6 wt.% KCl and 1.4 wt.% NaCl in air for 12 h at 700 °C. EDS spot scan shows elemental compositions (wt.%) at spot A and B.

and some corrosion products formed at grain boundaries of the surface substrate alloy are mixed with slag. If the slag is pickled away completely, the surface substrate alloy is destroyed as well. In order

to avoid over pickling, some corrosion products are left. The melting point of R1 and R2 is about 700 °C and 650 °C, respectively (Table 3). The generation of slagging may be consistent with the liquid transformation of two reagents. Fig. 4 displays the optical cross-sectional micrographs of TP347H stainless steel (A1) corroded in two corrosion reagents in air for 108 h at 700–750 °C. The grain boundary attack reveals an obvious intergranular corrosion. It spreads from the alloy surface to inside. This means the failure of TP347H in molten chloride salts is mainly caused by the grain boundary attack. When Fig. 4a and c, b and d are compared respectively, the intergranular corrosion presents a clear increasing trend with the elevated temperature. The grain boundary attack goes deep into TP347H since it spreads much at 750 °C (Fig. 4c and d) in two reagents. So the mass loss of TP347H is higher than that illustrated in Fig. 1, especially at higher test temperatures. In Fig. 5, the EDS results of cross-sections of TP347H stainless steel (A1) corroded in R1 in air for 108 h at 700–750 °C are presented. The corrosion of A1 in R2 exhibits similar EDS results

400

S. Liu et al. / Corrosion Science 83 (2014) 396–408

700 ºC

(a)

700 ºC

(b)

750 ºC

(c)

750 ºC

(d)

Fig. 4. Optical cross-sectional microstructures of TP347H stainless steel corroded in (a and c) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b and d) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 108 h.

compared to that in R1, therefore only the latter is displayed in Fig. 5. The corrosion products at grain boundaries mainly consist of Fe-rich oxides and some (Cr, Ni)-containing oxides is identified by EDS spot scan (spot 0). The generation of metal chlorides cannot be figured out, since no Cl is detected. Similar findings are also observed in EDS line scan. A higher content of O is observed at grain boundaries than in the internal of grains (Fig. 5b), so the oxidation of Fe, Cr and Ni happens. The precipitation of Cr2O3 and NiO consumes Cr and Ni, therefore, leaving (Cr, Ni, Fe)-depleted grain boundaries. The Cr and Ni then diffuse from grains to grain boundaries, and hence a higher content of Cr and Ni is detected at grain boundaries (Fig. 5b). The lower Fe at grain boundaries may be attributed to the severe oxidation since 52. 55 wt.% Fe is detected (spot 0). Apart from oxides, some carbides may also be produced. Fig. 6 displays the X-ray diffraction patterns of corrosion products formed on the surfaces of TP347H stainless steel corroded in two corrosion reagents in air for 96 h at 500–750 °C. An oxide scale, mainly composed of Fe-rich oxides, forms on the surface of A1. At temperatures below 700 °C, the Fe-rich peaks correspond to Fe2O3, Fe3O4, NiFe2O4 or Ni1.43Fe1.7O4. Then Fe3O4 becomes the only iron oxide at 700–750 °C, indicating the possible dissolution of Fe2O3 by molten chlorides. Minor peaks of Cr2O3 are also identified in two reagents at every test temperature below 700 °C. But Cr2O3 disappears at 700 °C, which degrades the corrosion resistance of A1. The Mn-containing composition is identified at higher temperatures between 700 °C and 750 °C. The existence of KCl peaks is attributed to the residual corrosion reagents mixed with the powder corrosion products. Since the corrosion products are washed before XRD analysis and metal chlorides may be dissolved as well, they are not identified. 3.2. Corrosion of C22 alloy Fig. 7 shows the mass loss of C22 alloy (A2) corroded in two corrosion reagents in air for 96 h at 550–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 10– 25 min. The highest mass loss in R1 occurs at 700 °C, while in R2 at 650 °C. The mass losses in R1 at 750 °C and in R2 at 700–750 °C, however, are so small. The increase of test temperature results in

no increasing mass loss, which means the corrosion does not aggravates at higher temperatures. This interesting finding may be attributed to the melting situation of R1 and R2, where they melt completely at 750 °C and 700 °C, respectively (discussed later). After corrosion for 96 h, the surface area of A2 decreases slightly, indicating a better corrosion resistance. The highest mass loss of A1 after corrosion for 96 h is 6391 g/m2 in R1 and 6397 g/m2 in R2 at 750 °C. As to A2, it is 3855 g/m2 in R1 at 700 °C and 2958 g/m2 in R2 at 650 °C, respectively. Considering the corrosion products generated at grain boundaries in A1 and the severe reduction of surface area of A1, A2 exhibits far better corrosion resistance than A1. If the data at 750 °C are compared, the difference in corrosion resistance of A1 and A2 is more obvious. Fig. 8 shows the SEM surface micrographs and EDS spot scan results of the corrosion products of C22 alloy (A2) corroded in R2 in air for 12 h at different temperatures. The corrosion of A2 in R1 exhibits very similar corrosion behaviour and corrosion product compositions compared to that in R2, therefore only the latter is displayed in Fig. 8. At 450 °C, layered (Ni, Cr, Mo)-rich corrosion products are observed apparently on A2 (Fig. 8a). The even surface morphology of corrosion products indicates a slight corrosion of A2. Then the corrosion gets aggravated as the temperature increases. At 600 °C, a needle-shaped Ni-rich corrosion product scale covers the surface of A2 (Fig. 8b). The contents of Cr and Mo decrease very much (spot B) compared to those at spot A, indicating a selective corrosion that some (Cr, Mo)-rich oxides in the surface corrosion layer have been dissolved. The volatilization of dissolved (Cr, Mo)-rich oxides may be then reinforced at higher temperatures, since a unique fibre-shaped aging morphology is presented (Fig. 8c) when the test temperature reaches 650 °C. As aging continues, the fresh substrate alloy beneath the corrosion product scale is unceasingly exposed to the corrosion reagents, and corroded. Multi-layered corrosion products are then formed. Compared with spot B (in the outer layer), the spot C (in the inner layer) reveals high contents of Cr and Mo (especially Cr). This finding suggests that the dissolution and volatilization of some (Cr, Mo)-rich oxides in the outer product layer make it fail to prevent the corrosion media from inwards. Accordingly, the inner fresh

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S. Liu et al. / Corrosion Science 83 (2014) 396–408

(a)

(b) Corrosion products

(c)

52.55% Fe 7.94% Cr 8.54% Ni 0.31% Mn 2.22% Si 24.04% O 4.40% C

Fig. 5. SEM cross-sectional microstructures of TP347H stainless steel corroded in 98.6 wt.% KCl and 1.4 wt.% NaCl in air for 108 h at (a) 700 °C and (c) 750 °C. (b) EDS line scan. EDS spot scan shows elemental compositions (wt.%) at spot 0 (at grain boundary).

Fe3O4

Fe2O3

Ni1.43Fe1.7O4

Cr2O3

KCl

NiFe2O4

NiMn0.5Cr1.5O4

o

500 C

(a)

o

550 C o

600 C o

650 C o

700 C o

750 C o

500 C

(b)

o

550 C o

600 C o

650 C o

700 C o

750 C

0

10

20

30

40

50

60

70

80

90

2θ (deg.) Fig. 6. XRD patterns of corrosion products formed on the surfaces of TP347H stainless steel corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 96 h at 500–750 °C.

S. Liu et al. / Corrosion Science 83 (2014) 396–408

o

550 C,10 × o 600 C o 650 C o 700 C o 750 C

3500

2

mass loss (g/m )

4000 3000 2500

4500

(a)

4000 3500

2

4500

mass loss (g/m )

402

2000 1500 1000

3000 2500

o

550 C,10× o 600 C o 650 C o 700 C o 750 C

(b)

2000 1500 1000 500

500

0

0 -12 0

12 24 36 48 60 72 84 96 108

-12 0

12 24 36 48 60 72 84 96 108

t (h)

t (h)

Fig. 7. Mass loss (Dm/A) of C22 alloy corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air at 550–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 10–25 min. The 550 °C measurements are enlarged as 10 value.

(a)

450

27.80% Ni 22.25% Cr 6.12% Fe

A

10.42% Mo

83.74% Ni

3.34% Pr

1.48% Cr

0.49% K

1.53% Fe

1.34% Cl 20.33% O

KCl

0.06% Mn 4.56% Mo 8.62% C

43.05% Ni

B

42.38% Cr 1.60% Fe 7.55% Pr

A

1.99% Mo 1.42% O 2.01% K

2.28% Sr 4.63% C 0.99% S

(b)

600 66.38% Ni 0.18% Cr

B

2.35% Fe 0.03% Mn 0.08%Nb 25.79% O 5.19% C

(c)

650 34.27% Ni 12.22% Cr 4.08% Fe

C

1.13% Mn 2.72% Mo 25.91% K 15.34% Cl 4.33% O

Fig. 8. Surface morphologies of corrosion products of C22 alloy corroded in 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 12 h at (a) 450 °C, (b) 600 °C and (c) 650 °C. EDS spot scan shows elemental compositions (wt.%) at spot A, B and C.

substrate alloy is corroded. Ni-containing compounds display better stability than Cr and Mo, contributing to the better corrosion resistance of A2 and A3 in molten chloride salts at high temperatures.

Fig. 9. Slagging morphology of surface corrosion products of C22 alloy corroded in 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 12 h at 700 °C. EDS spot scan shows elemental compositions (wt.%) at spot A and B.

When the temperature comes up to 700 °C, slagging is observed on A2 in R2. The slagging temperature of two corrosion reagents is both upon 700 °C. Since the similar morphology and composition are detected by SEM and EDS analysis in R1 compared to those in R2, only the latter is displayed in Fig. 9. When the slagging happens, the fibre-shaped morphology (Fig. 8c) disappears, but the multi-layered corrosion products are still obvious. The EDS results also reveal a Ni-rich phase in the corrosion products of A2 and more content of Cr in the inner corrosion product layer (spot B). Fig. 10 illustrates the optical cross-sectional microstructures of C22 alloy (A2) corroded in two corrosion reagents in air for 108 h at 750 °C. Representative austenitic equiaxed grains are observed and highlighted with 1red dotted lines. The exposed grains indicate that the grain boundaries are attacked, hence the occurrence of intergranular corrosion. This finding means the failure of C22 alloy in molten chloride salts is also mainly caused by the grain boundary attack. Since the cross-sections are not etched with any solution before optical observation, the intergranular corrosion must be attributed to the corrosion of two molten chloride salts. But as can be seen in Fig. 10, the grain boundary attack is not very serious, indicating slight intergranular corrosion of A2 at 750 °C. This finding suggests that the corrosion resistance of C22 alloy is relatively higher than that of TP347H stainless steel. And the corrosion products form at grain boundaries contribute little to the mass loss. So the mass losses illustrated in Fig. 7 are almost the actual data. Fig. 11 displays the X-ray diffraction patterns of corrosion products generated on the surfaces of C22 alloy corroded in two corrosion reagents in air for 96 h at 550–750 °C. NiO peaks are identified and the high quantity of NiO is confirmed by the EDS 1 For interpretation of color in Fig. 10, the reader is referred to the web version of this article.

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Fig. 10. Optical cross-sectional microstructures of C22 alloy corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 108 h at 750 °C.

analysis. The formation of NiO may be an explanation of the better corrosion resistance of A2 and A3. The protective Cr2O3 is also identified at every temperature in two corrosion reagents. At 700–750 °C, Cr1.12Ni2.88 is also identified as major phases, especially in R2. According to EDS results in Fig. 8, there exist some Mo-containing oxides, although the content is significantly lower than that of NiO. At lower temperatures, it corresponds to MoO2 and Cr takes the place of partial Mo (CrMoO3 is identified) as the temperature increases. This finding indicates the decomposition of MoO2 at higher test temperatures. When the phase identification in Fig. 11a and b is compared, an interesting finding appears that the phases identified in R2 are similar to those in R1 which is 50 °C higher than R2. Hence, the similar melting situations of two corrosion reagents (Table 3) result in similar phases generated in the corrosion products. 3.3. Corrosion of laser-cladding C22 coating Fig. 12 shows the mass loss of laser-cladding C22 coating (A3) corroded in two corrosion reagents in air for 96 h at 550–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 10–25 min. The highest mass loss of A3 both occurs at 650 °C in two corrosion reagents after corrosion for 96 h

NiO

Cr2O3

NiFe2O4

KCl

MoO2

Ni1.43Fe1.7O4

(1778 g/m2 in R1 and 1777 g/m2 in R2). Compared to A1 and A2, A3 exhibits the least mass loss (hence best corrosion resistance) after corrosion for 96 h. The mass loss of A3 in R1 at 700 °C is higher than that at 650 °C for corrosion less than 72 h, but it is opposite for longer corrosion times (Fig. 12a). This finding may be caused by the experimental error that some surface corrosion products formed at 700 °C are not removed, reducing the mass loss. The mass losses in R1 at 750 °C and in R2 at 700–750 °C are also small. Apart from the positive role of complete liquid conversion of corrosion reagents, the refined microstructure can also resist to the corrosion (discussed later). After corrosion for 96 h, the surface area of A3 also decreases slightly, indicating a better corrosion resistance of A3 compared to A1. Fig. 13 shows the SEM surface micrographs and EDS spot scan results of the corrosion products of laser-cladding C22 coating (A3) corroded in R2 in air for 12 h at different temperatures. The corrosion of A3 in R1 exhibits very similar corrosion behaviour and corrosion product composition compared to that in R2, therefore only the latter is displayed in Fig. 13. At 450 °C, the corrosion products keep a better integrity (Fig. 13a). Although some cracks are observed, the even surface morphology of corrosion products still indicates a slight corrosion of A3. The corrosion products are mainly composed of (Ni, Cr)-rich oxides. Then the increase of test

CrMoO3

Ni1.12Cr2.88

o

550 C

(a)

o

600 C o

650 C o

700 C o

750 C

(b)

o

550 C o

600 C

o

650 C o

700 C o

750 C

0

10

20

30

40

50

60

70

80

90

2θ (deg.) Fig. 11. XRD patterns of corrosion products of C22 alloy corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 96 h at 550–750 °C.

S. Liu et al. / Corrosion Science 83 (2014) 396–408

o

550 C,10× o 600 C o 650 C o 700 C o 750 C

-12 0

(a) 2

2000 1800 1600 1400 1200 1000 800 600 400 200 0

mass loss (g/m )

2

mass loss (g/m )

404

12 24 36 48 60 72 84 96 108

2000 1800 1600 1400 1200 1000 800 600 400 200 0

o

550 C,10× o 600 C o 650 C o 700 C o 750 C

-12 0

(b)

12 24 36 48 60 72 84 96 108

t (h)

t (h)

Fig. 12. Mass loss (Dm/A) of laser-cladding C22 coating corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air at 550–750 °C after pickling in the hydrochloric acid solution bath (25 wt.%, 80 °C) within 10–25 min. The 550 °C measurements are enlarged as 10 VALUE.

(a)

450

52.56% Ni 5.78% Cr 2.06% Co

A

KCl

B

63.26% Ni

4.83% Fe

10.06% Cr

0.12% Mn

1.33% Cl

A

83.79% Ni

1.86% K

3.40% Cr

12.16% O

0.59% Fe

0.57% Cl

3.16% C

8.99% Mo

0.83% K

4.83% Si

3.24% O

3.70% Mo

4.47% C

0.42% Ge 2.92% S

24.63% O 0.47% Si

(b)

600 63.91% Ni 2.94% Cr

B

1.90% Co 1.32% Fe 4.58% C 25.35% O

(c)

650 87.51% Ni

C

0.82% Cr 5.72% Co 1.58% Fe 0.99% Cl 0.63% K 2.74% O

Fig. 13. Surface morphologies of corrosion products of laser-cladding C22 coating corroded in 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 12 h at (a) 450 °C, (b) 600 °C and (c) 650 °C. EDS spot scan shows elemental compositions (wt.%) at spot A, B and C.

temperatures aggravates the corrosion and some corrosion products break from the specimen surface above 600 °C. At 600 °C, the integrity of surface product scale degrades and small particle-shaped

Fig. 14. Slagging morphology of surface corrosion products of laser-cladding C22 coating corroded in 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 12 h at 700 °C. EDS spot scan shows elemental compositions (wt.%) at spot A and B.

Ni-rich corrosion products (mainly NiO) cover the surface of A3 (Fig. 13b). Compared to spot A, the content of Cr decreases and no Mo is detected at spot B, indicating a selective corrosion, similar with that of A2. When the test temperature reaches 650 °C, an obvious particle-shaped morphology is presented (Fig. 13c). NiO is still the major phase in the corrosion products. According to EDS results, some Co-containing compounds are also identified and they may lead to the better corrosion resistance of A3 than A2, since none are found in the corrosion products of A2 (Fig. 8). In Fig. 14, the slagging morphology and EDS spot scan results of A3 after corrosion in R2 in air for 12 h at 700 °C are present. For A3, the slagging temperature of R1 and R2 is upon 750 °C and 700 °C, respectively. The slagging temperature of A3 is higher than that of A2 and A1, which may be attributed to the better corrosion resistance of A3. Since similar slagging morphology and composition are detected by SEM and EDS analysis in R1 compared tothose in R2, only the latter is displayed in Fig. 14. The layered corrosion morphology is also observed on A3, similar to that on A2 (Fig. 9). The EDS results also reveal a Ni-rich phase in the corrosion products of A3 and more content of Cr in the inner corrosion product layer (spot A). Fig. 15 illustrates the optical cross-sectional micrographs of laser-cladding C22 coating (A3) corroded in two corrosion reagents in air for 108 h at 750 °C. Compared to TP347H stainless steel and C22 alloy, laser-cladding C22 coating exhibits better corrosion resistance since no intergranular corrosion is observed. Without grain boundary attack, only the surface substrate alloy is corroded. Besides, the surface area of A3 decreases little after molten salt corrosion test. Hence, the mass loss illustrated in Fig. 12 is almost the actual situation. Fig. 16 presents the X-ray diffraction patterns of corrosion products generated on the surfaces of laser-cladding C22 coating

S. Liu et al. / Corrosion Science 83 (2014) 396–408

405

Fig. 15. Optical cross-sectional microstructures of laser-cladding C22 coating corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 108 h at 750 °C.

NiO

Cr2O3

CrMoO3

KCl

Co(Fe,Cr)2O4

Cr-O

MoO2

Ni1.12Cr2.88

(a)

o

550 C o

600 C o

650 C o

700 C o

750 C o

(b)

550 C o

600 C o

650 C

o

700 C o

750 C

0

10

20

30

40

50

60

70

80

90

2θ (deg.) Fig. 16. XRD patterns of corrosion products of laser-cladding C22 coating corroded in (a) 98.6 wt.% KCl and 1.4 wt.% NaCl and (b) 95.5 wt.% KCl and 4.5 wt.% NaCl in air for 96 h at 550–750 °C.

corroded in two corrosion reagents in air for 96 h at 550–750 °C. The major diffraction peaks still correspond to NiO. Cr2O3 is identified in two corrosion reagents at every test temperature. No FexOy, NiFe2O4 or Ni1.43Fe1.7O4 is observed. A Co-containing composition (Co(Cr, Fe)2O4) is identified at every test temperature in two reagents, well agreeing with the EDS results in Fig. 13. Furthermore, a Cr–O (CrOx) signal is also observed in two reagents at every temperature except 550 °C while none is identified in the corrosion products of A2 (Fig. 11). In spite of the small amounts, the formation of Cr–O and Co-containing composition may contribute greatly to the better corrosion resistance of A3 than A2.

4. Discussion The molten alkali chloride salt corrosion is a complex process, involving chemical corrosion, electrochemical corrosion as well as the interface reaction and the dissolution of oxides. The corrosion resistance at high temperatures is dependent on the formation of protective oxide scales. Cl/Cl-containing environments are well known to cause accelerated corrosion resulting in increased oxidation, metal wastage, internal attack, and loose non-adherent scales, and hence destroy the protective oxide scales [4,36,37].

4.1. Corrosion at lower test temperatures At 450–600 °C, the two corrosion reagents are generally not melted (Table 3). Chemical corrosion plays a dominated role. Since the corrosion is conducted in air, oxygen also participates in the corrosion. It can freely pass through the solid reagent salts to directly corrode the alloys. Consider the equation below:

xMðsÞ þ ðy=2ÞO2 ðgÞ ! Mx Oy ðsÞ

ð2Þ

where M refers to Fe, Cr and Ni. The reagent salts KCl and NaCl have low melting points and easy to volatize. In this study, large amounts of flocculent substances are observed at the tube mouth and the amount is larger at higher temperatures. They are condensed volatile gaseous KCl and NaCl on the cold tube. They then participate in such reactions:

RClðs; l; gÞ þ H2 OðgÞ ! ROHðs; l; gÞ þ HClðgÞ

ð3Þ

4HClðgÞ þ O2 ðgÞ ! 2Cl2 ðgÞ þ 2H2 OðgÞ

ð4Þ

where RCl corresponds to KCl and NaCl. The partial pressure of HCl and Cl2 in a biomass-derived flue gas is not high enough to cause severe gas phase corrosion, but may increase chlorination of

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superheater tubes. Cl2 has the ability to penetrate the protective oxide scale presumably through pores and cracks to react with the metal alloys. Abels and Strehblow [38] found that for a short period corrosion under HCl–O2 environment, Cl2 is the main corrosion medium, not HCl, according to Eq. (4). The volatile Cl/Cl-containing gases corrode alloys as follows [39–41]:

MðsÞ þ 2HClðgÞ ! MCl2 ðsÞ þ H2 ðgÞ

ð5Þ

MðsÞ þ Cl2 ðgÞ ! MCl2 ðsÞ

ð6Þ

MCl2 ðsÞ ! MCl2 ðgÞ

ð7Þ

where M refers to Fe, Ni and Cr. Occurrence of these reactions depends on the free energy of formation (DG) at different temperatures. Generally, DG for generation of CrCl2 is the smallest, followed by FeCl2, and then NiCl2 [42]. The stability of metal chlorides in terms of the prevailing chlorine partial pressures at a specific temperature can also affect the formation of chlorides. In fact, the vapour pressures of these chlorides are high even at lower temperatures due to their relatively low melting points. A higher temperature gives rise to a stronger volatilization, and hence a severer corrosion takes place [43]. In the diffusing process of chlorides outwards, the chlorides are oxidised for the oxygen concentration increases with an increasing distance from the alloy surface [44]. The liberating Cl2 can diffuse back to the alloy surface and a corrosion cycle forms:

2MClðgÞ þ O2 ðgÞ ! 2MOðsÞ þ Cl2 ðgÞ

ð8Þ

The equilibrium partial pressure varies from one metal chloride/ oxide to another. The oxygen partial pressure needed to transform chromic chloride into chromic oxide is the smallest, while it is an opposite case for nickel chloride. Since DG for generation of chromic chloride is also the smallest, as a result, a selective oxidation of chromium happens and the chromic chloride is oxidised into chromic oxide near the interface of metallic substrate/oxide scale [45]. This corrosion mechanism involves accelerated oxidation of alloys by gaseous Cl/Cl-species, which is often referred to as active oxidation [38,40,46,47]. The oxygen concentration is expected to be the key to active oxidation. Though it causes accelerated oxidation, the formed dense oxide scale composed of Cr2O3, NiO and Fe3O4 can provide protection for the substrate alloy from further attack by covering the surface of substrate. Hence, the three samples all well resist the corrosion at 450–600 °C, especially C22 alloy and C22 coating which form a high amount of more protective NiO and Cr2O3. 4.2. Corrosion at higher test temperatures At 650–750 °C, the two corrosion reagents are generally partial or completely melted (Table 3). Electrochemical corrosion dominates higher temperature corrosion since the liquid phase formed provides an electrolyte for ionic transport. This could be confirmed by the severe and slight intergranular corrosion in TP347H and C22 alloy, respectively. Generally, austenitic stainless steels (>0.03 wt.% C) tend to precipitate (Cr, Fe)23C6 at grain boundaries when improper heating at 425–815 °C. This process is known as sensitization [48]. (Cr, Fe)-depleted grain boundaries hence form [49]. And then more C and Cr, Fe diffuse to grain boundaries from the internal of grains. But it is still the Cr and Fe near grain boundaries are consumed since C has a higher diffusion rate. If the Cr content decreases below a critical value (generally 12 wt.%) needed for passivation [50], a micro-battery is produced where grain boundaries serve as the anode while grains are cathode. And then, grain boundaries are dissolved rapidly while grains remain passive. In addition, when molten salts are partial in liquid, the corrosion may be further enhanced if the protective surface oxide scale

is dissolved by molten salts. Under this circumstance, accelerated corrosion happens for oxygen as well as other aggressive media can still access to the fresh alloy and diffuse in the alloy substrate. Uusitalo et al. [47] conducted a high temperature corrosion experiment in deposited chloride salts of simulated superheater conditions. The destruction of oxides at splat boundaries by chloride and the diffusion of chlorine along grain boundaries were main factors causing failure of austenitic stainless steels. Therefore the corrosion is usually especially serious at 650 °C in R2 and 700 °C in R1. Consider the reactions, i.e. [45,51,52]:

2RClðs; lÞ þ Fe2 O3 ðsÞ þ ð1=2ÞO2 ðgÞ ! R2 Fe2 O4 ðs; lÞ þ Cl2 ðgÞ

ð9Þ

2RClðs; lÞ þ ð1=2ÞCr2 O3 ðsÞ þ ð5=4ÞO2 ðgÞ ! R2 CrO4 ðs; lÞ þ Cl2 ðgÞ

ð10Þ

where RCl corresponds to KCl and NaCl in this study. In addition to the deposited sodium chloride, the presence of NaCl (g) also displays the ability to break down Cr2O3 [53]. Moreover, low melting eutectic compounds may form in molten salts that flux the oxide scale [4]. The melting point of KCl and NaCl is 770 °C and 801 °C, respectively. However, alkali chlorides can react with other chlorides to form low temperature melting eutectic compounds. For example, the eutectic points are: KCl–CrCl2 = 462– 475 °C, KCl–FeCl2 = 340–393 °C, KCl–FeCl3 = 202–220 °C [54]. These low melting eutectic compounds volatilize easily and transfer the alloy from substrate outwards, which accelerates the corrosion. Cr plays an important role in high temperature corrosion resistance. But faced with molten chloride salts, it has lost its advantage according to Eq. (10). Compared to Cr2O3, NiO is not liable to be dissolved. This is mainly associated with the preferential dissolution of Cr2O3 that it consumes oxygen and inhibits the dissolution of NiO [55]. This preferential dissolution also occurs to MoO2. With a NiO skeleton left in the corrosion products, C22 alloy and C22 coating exhibit a higher corrosion resistance than TP347H. Furthermore, it should be figured out that the C22 coating exhibits a superior corrosion resistance than C22 alloy, although they have similar compositions and corrosion products. This is mainly attributed to the refined structure (nanocrystallines) formed by laser cladding. qffiffiffiffi pffiffiffi kp According to N BðminÞ > pVVBOAB , where NB(min) is the required min2D imum concentration of ingredient B to form oxide BO, D is the mutual diffusion coefficient of the alloy, kp is the growth rate constant of BO, VAB and VBO are the molar volumes of alloy and oxide BO, respectively. Refined grains accelerates the short-circuit diffusion of oxidised elements along grain boundaries (a bigger D) [50]. The critical concentration of alloying Cr for selective oxidation decreases. So the selective oxidation of Cr is more likely to happen in C22 coating than C22 alloy, forming a protective chromic oxide scale at alloy surface and reducing the oxygen activity at oxide scale/alloy interface. Ni can then more easily form continuous NiO at a lower concentration. In addition to the chromic oxide, Cr–O (that is CrOx) is also identified in the corrosion products of C22 coating, as illustrated in Fig. 16. When the corrosion reagents are completely melted (Table 3), the Eqs. (9) and (10) are inhibited instead. These could explain the abnormally small mass losses of A2 and A3 at 700–750 °C (Figs. 7 and 12). The oxygen solubility at the melting point is small [50], and it decreases further when the temperature rises more. So we have the opinion that, without sufficient oxygen, the active oxidation and molten salt corrosion are all inhibited. 5. Conclusions Isothermal molten chloride corrosion of TP347H (A1), C22 alloy (A2) and laser-cladding C22 coating (A3) was conducted in 98.6 wt.% KCl and 1.4 wt.% NaCl (R1, simulating forestry biomass)

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and 95.5 wt.% KCl and 4.5 wt.% NaCl (R2, simulating agricultural biomass) in air at 450–750 °C, with interval of 50 °C. The two biomasses showed comparable corrosivity since similar corrosion behaviour and mechanisms happened for one sample. Higher test temperatures aggravated corrosion with larger mass losses. Severe intergranular corrosion caused failure of A1 and slight intergranular corrosion happened in A2, but none in A3. Considering the experimental error caused by corrosion products generated at grain boundaries and the decreased surface area of A1, the highest mass losses indicate that the corrosion resistance of A3 was about twice as much as that of A2, and ten times of A1. Slagging happened due to the melt of reagents. Since oxygen solubility decreased in the liquid reagents, Ni1.12Cr2.88 formed and small mass losses happened to A2 and A3 at 700–750 °C. The selective oxidation of Cr and Mo, and the preferential dissolution of Cr2O3 and MoO2 by molten chlorides left a NiO skeleton which contributes to the better corrosion resistance of A2 and A3 than A1, since FexOy was main phase in the corrosion products of A1. The only identification of Cr–O (CrOx) and Co(Fe, Cr)2O4, as well as the refined microstructure may improve the better corrosion resistance of A3 than A2. Acknowledgements The authors would like to thank the financial support from the National Key Technology Research and Development Program of the Ministry of Science and Technology of China (2011BAE12B03), and the National Natural Science Foundation of China (11372110, 51101056) and the Fundamental Research Funds for the Central Universities (12MS07). References [1] H.P. Michelsen, F. Frandsen, K. Dam-Johansen, O.H. Larsen, Deposition and high temperature corrosion in a 10 MW straw fired boiler, Fuel Process. Technol. 54 (1998) 95–108. [2] A.A. Khan, W. de Jong, P.J. Jansens, H. Spliethoff, Biomass combustion in fluidized bed boilers: potential problems and remedies, Fuel Process. Technol. 90 (2009) 21–50. [3] M. Montgomery, O.H. Larsen, Field test corrosion experiments in Denmark with biomass fuels. Part 2: co-firing of straw and coal, Mater. Corros. 53 (2002) 185–194. [4] H.P. Nielsen, F.J. Frandsen, K. Dam-Johansen, L.L. Baxter, The implication of chlorine-associated corrosion on the operation of biomass-fired boilers, Prog. Energy Combust. Sci. 26 (2000) 283–298. [5] C. Yin, L.A. Rosendahl, S.K. Kær, Grate-firing of biomass for heat and power production, Prog. Energy Combust. Sci. 34 (2008) 725–754. [6] R. Saidur, E.A. Abdelaziz, A. Demirbas, M.S. Hossain, S. Mekhilef, A review on biomass as a fuel for boilers, Renew. Sustain. Energy Rev. 15 (2011) 2262–2289. [7] A. Demirbas, Potential applications of renewable energy sources, biomass combustion problems in boiler power systems and combustion related environmental issues, Prog. Energy Combust. Sci. 31 (2005) 171–192. [8] G. Guo, The Corrosion and Mechanisms of Fe, Cr, Ni and Their Oxides in Molten Salts of NaCl and KCl, M.S. Thesis, Dalian University of Technology, Dalian, CN, 2005. [9] P. Marshall, Austenitic Stainless Steels: Microstructure and Mechanical Properties, Elsevier Applied Science Publishers, 1984. [10] J.K. Kim, Y.H. Kim, J.S. Lee, K.Y. Kim, Effect of chromium content on intergranular corrosion and precipitation of Ti-stabilized ferritic stainless steels, Corros. Sci. 52 (2010) 1847–1852. [11] Z. Li, Study on High Temperature Corrosion of Superheater in Biomass Boiler, M.S. Thesis, North China Electric Power University, Beijing, CN, 2009. [12] K.T. Chiang, D.S. Dunn, G.A. Cragnolino, Effect of simulated groundwater chemistry on stress corrosion cracking of alloy 22, Corrosion 63 (2007) 940– 950. [13] A. Pardo, M.C. Merino, A.E. Coy, F. Viejo, R. Arrabal, E. Matykina, Pitting corrosion behaviour of austenitic stainless steels-combining effects of Mn and Mo additions, Corros. Sci. 50 (2008) 1796–1806. [14] P. Jakupi, J.J. Noël, D.W. Shoesmith, Intergranular corrosion resistance of R3 grain boundaries in alloy 22, Electrochem. Solid-State Lett. 13 (2010) C1–C3. [15] X. He, D.S. Dunn, Crevice corrosion penetration rates of alloy 22 in chloridecontaining waters, Corrosion 63 (2007) 145–158. [16] P. Jakupi, F. Wang, J.J. Noël, D.W. Shoesmith, Corrosion product analysis on crevice corroded alloy-22 specimens, Corros. Sci. 53 (2011) 1670–1679.

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