Hydrogen Embrittlement of a Boiler Water Wall Tube in a District Heating System

Abstract

A district heating system is an eco-friendly power generation facility with high energy efficiency. The boiler water wall tube used in the district heating system is exposed to extremely harsh conditions, and unexpected fractures often occur during operation. In this study, a corrosion failure analysis of the boiler water wall tube was performed to elucidate the failure mechanisms. The study revealed that overheating by flames was the cause of the failure of the boiler water wall tube. With an increase in temperature in a localized region the microstructure not only changed from ferrite/pearlite to martensite/bainite, which made it more susceptible to brittleness, but it also developed tensile residual stresses in the water-facing side by generating cavities or microcracks along the grain boundaries inside the tube. High-temperature hydrogen embrittlement combined with stress corrosion cracking initiated many microcracks inside the tube and created an intergranular fracture.

1. Introduction

A district heating system is a city-based facility with high energy efficiency and performance that generates large-scale heat from infrastructures to apartments, commercial buildings, and government offices [1,2,3]. Recently, the fourth generation district heating system, which uses recyclable energy has been developed [4]. Steel pipelines are mainly used in power generation and transport facilities to transfer heat sources (boiled water and high-temperature steam) over long distances. However, the materials are frequently exposed to harsh conditions, such as humid environments, high temperature, and chemical solutions. Therefore, cracks and failure occur in tubes (or pipes) because the harsh environment causes wall thinning, residual stress, and hydrogen embrittlement, which result in hydrogen-induced stress corrosion cracking (HISCC) [5].

Stress corrosion cracking (SCC) and hydrogen-induced cracking (HIC) are major issues in steel pipelines and are known as the most important failure mechanisms in such systems [6,7,8,9]. SCC occurs when the combination of tensile stress and corrosive media develops cracks in metals [5], and hydrogen embrittlement (HE) is explained by the result of the degradation of the mechanical properties of metals, which result in a decrease in ductility, strength, and fracture resistance due to the presence of hydrogen within the material [10]. Hydrogen atoms are generated in metals due to the corrosion effect, and hydrogen diffuses into the metals from the surface to the inside of the tube through interstitial sites [11]. Diffused hydrogen accumulates in several structural defects, such as grain boundaries, precipitates, inclusions, and dislocations, resulting in crack initiation and propagation [12,13,14].

Measures to prevent failures caused by the aging of facilities and operating conditions have rarely been investigated. Some studies have reported on the failure analysis of tubes used in power generation facilities. Among various factors for failure, the main causes are the reduction in wall thickness, hydrogen embrittlement, residual stress, and stress corrosion cracking. Luder et al. [15] argued that contaminants such as copper and zinc help to increase galvanic corrosion, which results in a loss of material on the inner surface, leading to a thinning of the tube. Ahmad et al. [16] investigated the SA210-A1 rear water wall tube by visual inspection. They argued that fly-ash erosion and an increase in temperatures induced the localized wall thinning of the rear water wall tube. Dorri et al. [17] reported that copper ions caused a higher corrosion rate, and caustic corrosion gave rise to the perforation of salt-water wall tubes. Munda et al. [18] reported that a failure by overheating in a boiler water wall tube consisting of SA210 Grade C steel exhibited a fish mouth rupture along the longitudinal direction because residual stresses developed in the hoop direction of the tube. Ahmad et al. [19] demonstrated that thermal fatigue, corrosion fatigue, and creep damage accelerated a failure in a heat recovery area. Hong et al. [20] reported that the combination of thermal creep and hydrogen embrittlement led to the crack formation inside the tube, and the tensile residual stresses that developed along the hoop direction accelerated the crack propagation along the grain boundaries in the axial direction. Djukic et al. [6] reported that electrochemical galvanic corrosion caused by copper and the enrichment of hydrogen generated acidic corrosion during boiler operation, and a significant amount of hydrogen provoked high-temperature hydrogen attacks by forming a window-type fracture. Mohtadi-Bonab et al. [7] argued that hydrogen-induced cracks were not generated by oxide inclusions; rather, they mainly propagated in a transgranular mode. Jeong et al. [8] proposed that stress corrosion cracking of X80 steel in seawater environments was highly affected by pits, hydrogen concentration, and hydrogen-induced cracking. Duarte et al. [21] reported that improper water treatment generated corrosion pitting, which resulted in stress concentration sites for crack initiation, and the combination of corrosion pitting and residual stress generated stress corrosion cracking.

Tube failure not only incurs economic losses due to forced shutdown of facilities, but it also threatens human safety; thus, it is very important to elucidate the failure mechanisms and prevent accidents. In order to prevent accidents, customized failure analysis is necessary to elucidate the cause of failure in terms of the facility operating conditions, such as the temperature, operating time, humidity, hydrogen, water quality, fuel, and characteristics of materials. A boiler water wall tube, which had been used for 20 years in a district heating system, was fractured recently during operation. In this study, we conducted a corrosion failure analysis of the water wall tube, from which we revealed the failure mechanisms.

2. Materials and Methods

2.1. Materials and Operating Conditions

The water wall tube studied in this paper was made of ASTM SA210 Grade A1 (Fe-0.27C-0.93Mn-0.1Si-0.035P-0.35S wt.%). The water wall tubes were installed on the wall for heat exchange. One side of the tubes was directly subjected to a flame, and a failure occurred during operation (Figure 1). The temperature of the feed water was 200 °C, the boiler water flowing in the tube was heated to 535 °C, and the steam pressure was 106 kg/cm² during operation.

2.2. Characterization Methods

The preparation of specimens for the microstructural characterization of a normal tube and a failed tube was carried out using electro-discharge machining (EDM) wire cutting. For the investigation of the microstructure, the surface of the specimens was mechanically polished to a 1 $\mathsf{\mu }\mathrm{m}$ diamond suspension to make a mirror plane. Then, the mirror-finished surface was chemically etched by a 3% nital solution for 5~10 s to observe the microstructure of the carbon steel. We performed a visual inspection and microstructural analysis using an optical microscope (OM, OLYMPUS BX51M) and scanning electron microscope (SEM, CARL ZEISS MERLIN COMPACT). A Vickers hardness machine (MITUTOYO HM-200) was used to examine the mechanical properties based on the microhardness.

2.3. Hydrogen Analysis Method

Hydrogen charging experiments were performed electrochemically on the plate-type specimen (10 mm length, 10 mm width, 1 mm thickness) using a 3% NaCl aqueous solution containing 1 g/L CH4N2S as a hydrogen recombination inhibitor with a current density of 5 mA/cm² for 24 h. After the hydrogen charging, thermal desorption spectroscopy (TDS) (BRUKER G4 PHEONIX) analysis was carried out immediately to minimize the loss of hydrogen. TDS was conducted to investigate the H-desorption behavior of the normal tube (ferrite + pearlite) and fractured tube (martensite + bainite) at a constant heating rate of 873 K/h and 1473 K/h. The hydrogen-charged specimen was heated from 300K to 1073 K using a tube furnace, where a high purity N₂ gas was used as a carrier gas. Integrated desorbed hydrogen within the peaks of the hydrogen desorption rate curves was used to measure the amount of hydrogen. The activation energy ($E_a$) of the trapping sites corresponding to each desorption peak was calculated using the Kissinger equation (Equation (1)):

$$\frac{\partial \left(\ln \left(\frac{\mathsf{\Phi }}{{\mathrm{T}}_{\mathrm{p}}^2}\right)\right)}{\partial \left(\frac{1}{{\mathrm{T}}_{\mathrm{p}}}\right)}=-\frac{{\mathrm{E}}_{\mathrm{a}}}{\mathrm{R}}$$

(1)

where $\mathsf{\Phi }$. is the heating rate, ${\mathrm{T}}_{\mathrm{p}}$ is the desorption peak temperature, and R is a gas constant.

3. Results and Discussion

3.1. Visual Inspection

The shape of the damaged (fractured) tube was oval, and the size had a length of ~100 mm along the axial direction, a width of 45 mm along the hoop direction, and a thickness of 6 mm along the radial direction (Figure 2). We could not find any reduction in thickness, but we found macro- and microcracks that originated only on the water-facing side. The cracks were generally propagated along the axial and radial direction, and a few cracks were observed along the hoop direction near the fracture surface (Figure 2). Most cracks were found in the overheated zone, especially near the boiler water-facing side (Figure 2). It is known that the overheated region by the flame induces a nonuniform thermal gradient through the tube, which can generate compressive residual stresses in the air side and tensile residual stresses in the boiler water side [18,20]. In addition, it is recognized that a boiler environment at high temperature can cause hydrogen embrittlement accompanied by decarburization [22].

3.2. Microstructure Analysis

Figure 3 shows the microstructure analysis results of the normal (undamaged) tube, which did not experience overheating from the flame. The microstructure of the normal tube showed a mixture of ferrite (bright color) and pearlite (dark color), which is a typical microstructure of medium carbon steel. Microstructures at different locations, e.g., the air side (Figure 3a–c), the center (Figure 3d,e), and the water-facing side (Figure 3f–h) in the normal tube were similar, which indicated that the microstructures were not dependent on the examined locations. No wall thinning or cracks were found in the normal tube.

Figure 4 and Figure 5 show the microstructure analysis results of the damaged tube in the air side (a), the boiler water side (b), the air side in an overheated zone (c) and the boiler water side in an overheated zone (d). The microstructures in the air side of the damaged specimen were composed of ferrite and pearlite (Figure 4a and Figure 5a). It had a microstructure, similar to the undamaged tube (Figure 3). Pearlite consists of a lamellar structure with ferrite and cementite (Figure 5a), which was formed due to the diffusion of carbon in the steel during slow cooling. However, the boiler water-facing side in the damaged tube (Figure 4b and Figure 5b) had only ferrite, exhibiting that the pearlite had disappeared (Figure 4b and Figure 5b). This was due to the decarburization by hydrogen atoms. The hydrogen atoms diffused easily into the microstructural defects at high temperatures and chemically reacted with the carbon atoms in cementite, forming methane gases.

On the other hand, all of the ferrite and pearlite phases disappeared in the overheated zone, and new phases of martensite and bainite were developed in the air side (Figure 4c and Figure 5c). Martensite was obtained from diffusionless transformation by rapid quenching from the austenite phase, and bainite was formed from a relatively higher cooling rate. Based on the iron–carbon phase diagram, the SA210 Grade A1 carbon steel should undergo heating above A1 temperature in order to change the microstructure from the ferrite/pearlite phase to the martensite phase. This indicated that the materials in the air side (located in Figure 4c and Figure 5c) underwent overheating by the flame.

In Figure 5d, the microstructure in the boiler water side had a mixture of ferrite and a small amount of bainite. Interestingly, in the overheated zone, a large number of micro-cracks were found near the boiler water side (Figure 5). The cracks were mainly propagated along the grain boundaries toward the radial direction (Figure 4d and Figure 5d). It is thought that hydrogen atoms interacted with carbon atoms in the steel forming methane gases, which agglomerated in the grain boundaries inducing an internal pressure high enough to provoke hydrogen embrittlement cracking. Interestingly, the microcracks were even found in locations far away from the water-facing side, as shown in Figure 4d and Figure 5d, indicating that hydrogen easily penetrate into the inside of the tube in the overheated zone.

Both the overheated zone (Figure 4d and Figure 5d) and the non-overheated zone (Figure 4b and Figure 5b) revealed decarburization in the water-facing locations. However, we must ask why the cracks only occurred in the overheated zone. There are two possible reasons; one was because tensile residual stresses in the water-facing side developed along the hoop direction in the overheated zone, and the other is the different microstructures between the overheated zone and non-overheated zone. Hong et al. [20] simulated the residual stress distribution at high temperature for SA210 Grade-A1 steel. They demonstrated that tensile residual hoop stresses developed on the inner surface (water side) of the tube exposed to the flame, which accelerated hydrogen-induced cracking due to the ease of hydrogen invasion into the matrix. Therefore, we deduced that the increase in temperature in the localized region due to the overheating generated a tensile residual stress field in the vicinity of the water side, from which hydrogen easily diffused into the matrix and formed many microcracks in the water side. Park et al. [23] reported that hydrogen trapping efficiency is increased in the order of degenerated pearlite, bainite, and acicular ferrite. Although the trapping efficiency of bainite is lower than that of acicular ferrite, bainite has a more sensitive microstructure to hydrogen-induced cracking than acicular ferrite. From these results, we concluded that the microstructure in the overheated zone was more vulnerable to hydrogen-induced cracking than that in the non-overheated zone.

3.3. Mechanical Properties

Figure 6 shows the microhardness contour mapping results for the normal (a) and damaged (b) tubes. The average hardness value of the normal tube was about 138 HV, and the hardness value was mostly uniform (Figure 6a). However, the damaged tube exhibited a clear difference in hardness, indicating a hardness (328 HV) in the overheated zone about twice as high as the hardness in the non-overheated zone (172 HV) (Figure 6b). As revealed in Figure 5c,d, martensite and bainite were observed only in the overheated zone. Since martensite and bainite have a greater strength than ferrite and pearlite [24,25], the hardness in the overheated zone was higher than in the non-overheated zone. However, phases of martensite and bainite are more brittle compared to phases of ferrite and pearlite [24,26], which are more vulnerable to crack formation.

3.4. Hydrogen Desorption and Trapping Sites

The microstructure is one of the important factors in determining resistance to hydrogen embrittlement. Hydrogen trapping sites in metals are classified as reversible and irreversible depending on the microstructures, and the hydrogen trapping sites can be determined by the activation energy of hydrogen desorption [27,28,29]. Reversible hydrogen trapping sites are lattices, dislocations, and grain boundaries where the activation energy is lower than 20 kJ/mol [27,28,29,30,31]. On the other hand, irreversible hydrogen trapping sites, such as interfaces of different phases, inclusions, and precipitates, have an activation energy higher than 60 kJ/mol [27,28,30,31].

The thermal desorption analysis curves of the normal and fractured tubes are shown in Figure 7. All the pre-charged specimens showed only one hydrogen desorption peak in the range of 300~600 K, and all the peaks moved to higher temperatures as the heating rate increased. The total hydrogen desorption volume of the normal and fractured tubes at a heating rate of 1473 K/h were 0.9 and 21.8 wt.ppm, respectively. Figure 7d shows the calculation of the activation energy of trapping sites based on the Kissinger equation (Equation (1)). The activation energies of the normal tube and the fractured tube were 9.8 kJ/mol and 15.2 kJ/mol, respectively. These results indicate that the main hydrogen trapping sites of both tubes were grain boundaries [29,32], which would account for the formations of microcracks at the grain boundaries in the water-facing side and the resultant intergranular type of fracture (Figure 4d and Figure 5d). Furthermore, it is known that martensite has many sites of structural defects for hydrogen trapping, such as dislocations, interfaces, and grain boundaries, resulting in a higher hydrogen trapping density than ferrite [31,33]. This is why the fractured tube had a much higher hydrogen desorption volume than the normal tube.

3.5. Failure Mechanism

Figure 8 presents the crack initiation and failure mechanism for the boiler water wall tube. When the boiler tube was subjected to overheating by a flame (Figure 8a), the microstructure changed from ferrite/pearlite to martensite/bainite in the overheated zone (Figure 8c), as confirmed by the OM/SEM observations (Figure 4 and Figure 5), and an increased hardness value (Figure 6). Furthermore, TDS analysis demonstrated that diffusible hydrogen was trapped at the grain boundaries of the fractured tube because the hydrogen activation energy of the fractured tube was 15.2 kJ/mol. At the same time, tensile residual stresses developed along the hoop direction in the water side of the overheated zone (Figure 8b). The tensile residual stresses weakened the protective passivation layers, by providing a pathway for caustic corrosion, from which hydrogen atoms more easily diffused into the grain boundaries. (Figure 7 and Figure 8c). The diffused hydrogen reacted with the carbon atoms exhibiting a decarburized microstructure in the water side and forming methane gases at defect sites, which agglomerated at the grain boundaries (Figure 8d). In the overheated zone, hydrogen atoms invaded more deeply into the matrix martensite (which had a much higher trapping density than ferrite) under tensile stresses and generated cavities or microcracks along the grain boundaries. Therefore, the microcracks merged into macrocracks, exhibiting intergranular cracking (Figure 8e). Consequently, high-temperature hydrogen embrittlement combined with stress corrosion cracking initiated many microcracks inside the tube. The micro- and macrocracks propagated along the grain boundaries under the tensile residual stress field toward the radial direction, finally leading to the intergranular type of fracture.

4. Conclusions

We performed a corrosion failure analysis of a boiler water wall tube used in district heating facilities. Our main conclusions are summarized as follows.

(1)

The microstructure of a normal tube initially consists of ferrite and pearlite, but overheating in a localized region changed the initial microstructure into martensite and a small amount of bainite, which have more susceptibility for hydrogen embrittlement.

(2)

The damaged tube exhibited a clear difference in hardness, indicating that the hardness (328 HV) in the overheated zone was about twice as high as the hardness in the non-overheated zone (172 HV). The higher hardness in the overheated zone was due to the microstructural evolution of martensite and bainite with high strength.

(3)

Decarburization was found in the entire water-facing side, whether it was overheated or not. However, micro- and macrocracks located at the grain boundaries were only observed in the overheated zone, which was due to the development of tensile residual stresses in the water side and the microstructural change into martensite, which has a higher hydrogen trapping density than ferrite.

(4)

The hydrogen trapping sites of the fractured tube were at the grain boundaries, which facilitated the formation of microcracks and an intergranular fracture.

(5)

The combined effects of the brittle microstructures of martensite and bainite, high temperature hydrogen embrittlement, and tensile residual stresses accelerated hydrogen-induced stress corrosion cracking, leading to a final fracture.