High Temperature Corrosion and Materials Applications

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26/10/2007 1:51PM Plate # 0

High-Temperature Corrosion And Materials Applications George Y. Lai, editor, p1 DOI: 10.1361/hcma2007p001

Copyright © 2007 ASM International® All rights reserved. www.asminternational.org

CHAPTER 1

Introduction METALS AND ALLOYS will react during high-temperature service with the surrounding environment, resulting in high-temperature corrosion. In gaseous environments, high-temperature corrosion is defined as the corrosion that takes place above the maximum temperature at which acids condense and dew-point corrosion takes place. Although a majority of hightemperature corrosion reactions take place at temperatures above 500 °C (930 °F), severe high-temperature corrosion has been encountered in many cases at temperatures below 500 °C (930 °F). In waste-to-energy boilers, for example, carbon and low-alloy steels have experienced severe fireside corrosion problems in the waterwalls of the boilers at the tube metal temperatures of approximately 260 to 315 °C (500 to 600 °F). This book is intended primarily for engineers and metallurgists who are concerned with hightemperature materials problems in the following industries: aerospace/gas turbine, chemical processing, refining and petrochemical, fossil-fired power generation, coal gasification, waste-toenergy industry, pulp and paper, heat treating, mineral and metallurgical processing, and others. The technical data presented in this book are pertinent to “real” materials problems related to the aforementioned industries. The book will also be useful for both undergraduate and graduate students who are interested in studying or pursuing research on the subject of hightemperature corrosion. The book covers eight basic modes of hightemperature corrosion. A brief description of thermodynamics is included for most chapters to help readers to understand the corrosion reactions. The external stresses (or strains) can cause alloys to suffer preferential corrosion penetration attack in a certain corrosive environment, such as sulfidizing environments. In addition, external stresses or residual stresses can cause the alloy to suffer brittle, intergranular cracking when exposed to the lower end of the intermediate temperatures for certain alloys. This type of cracking is frequently referred to as “reheat

cracking,” “stress-relaxation cracking,” or “strain-age cracking” (for nickel-base alloys). Both of these subjects are covered in Chapter 14, “Stress-Assisted Corrosion and Cracking.” The subject of erosion and erosion/corrosion is also reviewed with an attempt to offer readers general guidance on materials selection and application. Discussion also includes hydrogen attack of carbon steels in boilers and refinery equipment. Finally, extensive discussion on the materials problems in coal-fired boilers, oil-fired boilers, waste-to-energy boilers, and black liquor recovery boilers is included. In summary, the subjects covered extensively in this book include:

Oxidation Nitridation Carburization and metal dusting Corrosion by halogen and hydrogen halides Sulfidation Hot corrosion Molten salt corrosion Liquid metal corrosion and embrittlement Erosion and erosion/corrosion Stress-assisted corrosion and cracking Hydrogen attack Coal-fired boilers Oil-fired boilers and furnaces Waste-to-energy boilers and waste incinerators Black-liquor recovery boilers

The focus of this book is on commercial alloys, including both generic and proprietary alloys. Most data are presented to reveal alloy ranking and thus serve as a general guide to materials selection and application. Engineers can thus use the data and information to compare alloys that are commercially available. The effects of alloying elements, temperature, and environmental conditions on the corrosion behavior of alloys are also discussed, providing information about the capability of an alloy in terms of useful temperature limitation. Trademarks for alloys and alloy manufacturers are listed in Appendix 1. The compositions of alloys are tabulated in Appendix 2.

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High-Temperature Corrosion And Materials Applications George Y. Lai, editor, p3-4 DOI: 10.1361/hcma2007p003


CHAPTER 2

Challenges in Materials Applications for High-Temperature Service IN MANY INDUSTRIAL SYSTEMS, plant operating conditions can be quite complex; it is rather difficult to use laboratory tests to simulate plant conditions. However, laboratory tests can provide good general guidance for making preliminary alloy selection. In situ field testing or field trials of candidate alloys in the operating plant provides the best way of obtaining the corrosion information that can be reliably used for final materials selection. During the preliminary alloy selection process, it is important to evaluate not only the high-temperature corrosion resistance of the alloy, but also its mechanical properties such as tensile and creep-rupture strengths. The microstructural changes at the application temperatures such as thermal stability of the alloy, should also be considered. For example, duplex stainless steels are known to suffer 475 °C (885 °F) embrittlement caused by the formation of alpha prime (α0 ) coherent precipitates. Accordingly, these stainless steels should be avoided for use as a structural component at temperatures approximately, above 340 °C (650 °F). ASME Codes may have lower maximum service temperature limits for these alloys. Consideration should also be given to fabrication issues, such as weldability and welding procedures, annealing heat treatments, postweld heat treatment (PWHT) and stress relieving, and codes and standards requirements. The availability of the alloy can also be an issue. It is not uncommon to find that some alloys are no longer commercially available in stock due to a number of reasons, which may include poor market demands in the past, difficulty in manufacturing, and so forth. In some cases, the alloy may only be available on order for a whole production heat, which can be tens of thousands of pounds of material. Another important factor is the alloy price. A cost

analysis needs to be conducted to balance the material cost with the expected life for the component to ensure the alloy is cost effective. Often the life-cycle cost is a better criterion than the initial material cost in making an alloy selection. Selection of an appropriate filler metal for welding is important for component fabrication involving welding. Normally, it is a simple process when the candidate alloy has a filler metal with matching chemical composition. However, many high-temperature alloys do not have filler metals with matching chemistries. The widely used Fe-Ni-Cr alloy 800H is a good example. Many heat-resistant cast alloys also do not have matching filler metals. Thus, when no matching chemistry filler metal is available for welding, it is critical to select a filler metal that not only possesses excellent weldability but, also exhibits comparable or better high-temperature corrosion resistance along with comparable strengths, thermal stability, and other relevant properties. Some fabricators sometimes use weldability to select a filler metal without considering the resistance of the weld metal to the specific hightemperature corrosive environment in the end application. This can lead to premature failures. For example, because of their good weldability, high nickel filler metals, such as filler metal alloy 82 (ERNiCr-3), are sometimes used for welding the alloys that are to be in service in sulfidizing environments. This can cause preferential sulfidation attack at the weld joint because of the relatively poor sulfidation resistance of high nickel alloys. Welding can still be an issue for some hightemperature alloys even with matching filler metals. This is because some high-temperature alloys contain many alloying elements for various metallurgical reasons, such as improving the resistance to a certain mode of high-temperature

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4 / High-Temperature Corrosion and Materials Applications

corrosion, increasing tensile and creep-rupture strengths, or increasing wear resistance. Increasing the levels of some of these alloying elements can increase the difficulty in the weldability of the alloy. For example, an alloy containing high silicon, high aluminum, high carbon, or very high chromium can be difficult to weld even though a matching filler metal is available. For construction of a component, engineers have the option to consider whether a wrought alloy or a cast alloy will be more suitable metallurgically and/or economically for the intended high-temperature application. Engineers may also consider a totally different approach to address the high-temperature corrosion issue for some existing plant equipment that has suffered corrosion. In refineries, many reactor vessels, such as crude towers, hydrocrackers, and hydrodesulfurizers, are made of clad plates with a corrosion-resistant cladding in original installations. Cladding can be corroded after years of operation. One common approach is to refurbish the corroded vessels by applying a corrosion-resistant weld overlay instead of replacing it with a new construction. This approach has been adopted in the boiler industry in recent years to address the severe corrosion problems with the waterwalls of boilers in waste-to-energy boilers, coal-fired boilers, basic oxygen furnace hoods in steel mills, and so forth. With automatic controls for gas metal arc welding machines, a large scale of weld overlay can be applied in vessels or boilers with engineering quality. Laser cladding can also be applied in the shop on large equipment such as waterwall panels. Coextruded composite tubes with a corrosion-resistant alloy cladding on the outer

diameter have long been available for construction of waterwalls as well as superheaters in boilers. Composite tubes manufactured by a spiral weld overlaying process have been made available in recent years. These composite tubes use the outer diameter cladding for providing corrosion protection and the substrate base tube for the load-bearing structural part. Most of these composite tubes are used in superheaters and reheaters in boilers with metal temperatures being likely less than about 650 °C (1200 °F). Many furnace tubes used in petrochemical processing, such as ethylene cracking furnace tubes, are exposed to temperatures higher than 980 °C (1800 °F) and carburizing gas streams on the internal diameter (ID) of the tube, application of composite tubes with a carburization- and coking-resistant alloy cladding on the tube ID can potentially increase the operating temperature and/or prolong the tube life. Aluminide coatings reportedly have been used in ethylene cracking furnace tubes. At the writing of this book, it appears no commercial companies in the United States provide aluminizing coating services for ethylene furnace tubes or pipes. Another diffusion coating, chromized coating, has also reportedly been used in boilers. Both of these diffusion coatings are very thin. Coatings have been highly successful in providing protection against oxidation and hot corrosion for the high-temperature components, such as airfoils, in gas turbines. The coatings used involve aluminide coatings, overlay MCrAlY coatings by vapor deposition processes (e.g., electron beam physical vapor deposition), and ceramic thermal barrier coatings (e.g., stabilized ZrO2). Coatings are considered sacrificial and are to be replaced periodically.

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CHAPTER 3

Oxidation 3.1 Introduction Oxidation is the most important hightemperature corrosion reaction. Metals or alloys are oxidized when heated to elevated temperatures in air or in highly oxidizing environments, such as combustion atmospheres with excess air or oxygen. Many metallic components, such as retorts in heat treat furnaces, furnace heater tubes and coils in chemical and petrochemical plants, waterwalls and superheater/reheater tubes in boilers, and combustors and transition ducts in gas turbines, are subject to oxidation. For many industrial processes, combustion involves relatively “clean” fuels such as natural gas or No. 1 or No. 2 fuel oil. These fuels generally have low concentrations of contaminants, such as sulfur, chlorine, alkali metals, and vanadium. In many cases, excess air is used to ensure complete combustion of the fuel. The combustion products thus consist primarily of O2, N2, CO2, and H2O. Although alloys in these environments are oxidized by oxygen, other combustion products, such as H2O, may play an important role in affecting the oxidation behavior of the alloy. The presence of N2 in the combustion gas stream can cause significant internal nitridation attack under certain conditions, which is discussed in Chapter 4 “Nitridation.” Oxidation can also take place in a “reducing” environment (i.e., the environment with a low oxygen potential created by the combustion under a substoichiometric condition). When combustion takes place under stoichiometric or substoichiometric conditions, the resultant environment becomes “reducing.” This type of environment is generally characterized by low oxygen potentials. Under this condition, the oxygen potential of the environment is typically controlled by pH2 =pH2 O or pCO =pCO2 ratio, and the oxidation kinetic is generally slow. The development of a protective oxide scale can be

sluggish for most alloys. As a result, the effects of corrosive contaminants can become more pronounced, resulting in other modes of hightemperature corrosion. For example, if the sulfur level in the environment is high, sulfidation then becomes the predominant mode of corrosion, even though oxidation also takes part in the corrosion reaction. Thus, a majority of hightemperature corrosion problems in reducing environments are caused by modes of corrosion attack other than oxidation. Most industrial environments have sufficient oxygen activities (or potentials) to allow oxidation to participate in the high-temperature corrosion reaction regardless of the predominant mode of corrosion. In fact, the alloy often relies on the oxidation reaction to develop a protective oxide scale to resist corrosion attack, such as sulfidation, carburization, hot corrosion, and so forth. The oxidation behavior under these conditions is discussed in other chapters dealing with different modes of corrosion attack. There is a large spectrum of engineering alloys available for applications in different temperature ranges. This chapter presents a large oxidation database for a wide spectrum of engineering alloys, ranging from carbon and Cr-Mo steels, which serve the low end of the temperature spectrum, to superalloys serving the highest temperature regime. The data are organized by alloy groups to help readers to compare alloys within the same alloy group and also to compare alloys between different alloy groups. The focus is to present comparison data, thus allowing readers to consider candidate alloys for applications in the temperature regime of interest. Also included are some important metallurgical and environmental factors that can affect the oxidation behavior of the alloy. It should be noted that a majority of the database has been generated in laboratory tests that were conducted at temperatures higher than those at which the tested alloys would normally be

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pg 6


used. The intent for this approach was to determine the oxidation behavior of alloys within relatively short test durations. For example, many tests have been conducted at 980 to 1200 °C (1800 to 2200 °F) for stainless steels, Fe-Ni-Cr alloys, and Ni-Cr alloys. This may result in the alloy performance ranking based more on scaling (metal loss) than on internal oxidation attack which the alloys would most likely have encountered at lower application temperatures. An example is given in Fig. 3.1, which illustrates an actual field experience with a furnace heater coil made of a Ni-Cr alloy that had been in service for about 4 to 5 years at temperatures less than 900 °C (1650 °F), suffering extensive internal oxidation attack with very little scaling (metal loss) (Ref 1). Few long-term tests have been conducted at temperatures of 650 to 980 °C (1200 to 1800 °F) where most stainless steels, Fe-Ni-Cr alloys, and Ni-Cr alloys are used in high-temperature applications. Furthermore, many test results were presented as weight changes instead of actual measurements of the damage to the metal, such as the total depth of oxidation attack including both metal loss (thickness reduction) and internal penetration.

3.2 Thermodynamic Considerations 3.2.1 Formation of Oxides Thermodynamically, an oxide is likely to form on a metal surface when the oxygen potential (pO2 ) in the environment is greater than the oxygen partial pressure in equilibrium with the oxide. The oxygen partial pressure in equilibrium with the oxide can be determined from the standard free energy of formation of the oxide. Consider the reaction: M+O2 ÐMO2

ð3:1Þ

aMO2 DG =7RT ln aM pO2

ð3:2Þ

Assuming the activities of the metal and the oxide are unity, Eq 3.2 becomes: DG =RT ln pO2

ð3:3Þ

Then

pO2 =eDG

=RT

ð3:4Þ

Equation 3.4 permits the determination of the oxygen partial pressure in equilibrium with the

0.5 mm

Fig. 3.1

A Ni-Cr alloy furnace heater coil suffering extensive internal oxidation attack with little surface scaling after service for 4 to 5 years at temperatures below 900 °C (1650 °F). Source: Ref 1


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pg 7

Chapter 3:

partial pressure of oxygen in equilibrium with Cr2O3 at 1000 °C (1830 °F) is about 10−21 atm from Fig. 3.2. This implies that the formation of Cr2O3 is favored thermodynamically at 1000 °C in environments with oxygen potentials higher than 10−21 atm. When the environment is “reducing” (e.g., the environment generated by stoichiometric or substoichiometric combustion), the oxygen potential is controlled by either pH2 =pH2 O or pCO =pCO2 ratio. The oxygen potential can be determined by the reaction:

oxide from the standard free energy of formation. The standard free energies of formation of selected oxides as a function of temperature are shown in Fig. 3.2. The figure also allows quick determination of the oxygen partial pressure (pO2 ) in equilibrium with the oxide. This oxygen partial pressure can be read by drawing a straight line from the point marked “O” on the left vertical axis of Fig. 3.2 through the free-energy line of the oxide at the intersecting point with the temperature of interest. This line continues to extend until it intersects with the pO2 scale located at the right-hand side and bottom of the Fig. 3.2. The intersecting point shows the oxygen partial pressure in equilibrium with the oxide of interest. If the oxygen partial pressure in the environment is greater than the oxygen partial pressure in equilibrium with the oxide, the oxide is likely to form on the metal surface. Conversely, the oxide is not likely to form. For example, the

2H2 +O2 Ð2H2 O

ð3:5Þ

The standard free energy of formation is related to the partial pressures of hydrogen, oxygen, and water by: p2H2 O 2 pH2 pO2

DG =7RT ln

H2/H2O ratio

10–8

CO/CO2 ratio

10–8

10–6

ð3:6Þ

pO

10–4

10–4

10–6

!

2

10–2 10–2

0

O

1

–100

O4 Fe 3

–200

+O

4

2

=

O3 e2 6F

oO

NiO

+ Ni

O2

M

O2 o+

1 10–2

= 2C

1

2C

=2

10–4

M

2

–300

102 102

∆G °=RT In pO2 (kJ/mole O2)

–400 –500

H C

–600

4- Cr + 3

–700

Si +

M

2 Cr2O 3 =O2 3

O2

M

I 2O 3

O I+

2

B

104

4- A 3

B O O Ca Mg =2 2 B 108 = 2 O + 2 + O 2Ca g 2M Change of state Element Oxide M M Melting point M M B Boiling point B 1010

–900

10–14 106 10–16

M

–1000 –1100 –1200

200 400 600 800 1000 1200 1400 1600 1800 2000 2200 2400 10 Temperature, °C CO/CO2 ratio 10–14 10–12 10

10–10

10–12

106

2- A = 3

10–6 10–8

104

iO 2 =S

–800

10–18 108 10–20

0

OK

10–22

H2/H2O ratio

10–20010–100 10–70 10–60 10–50 10–42 10–38 10–34 10–30 10–28 10–26 10–24

pO

2

Fig. 3.2

Oxidation / 7

Standard free energies of formation of selected oxides as a function of temperature. Source: Ref 2


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pg 8


Rearranging the equation results in:

pO2 =eDG

=RT

1 (pH2 =pH2 O )2

ð3:7Þ

Thus, the oxygen partial pressures at various temperatures can be determined as a function of pH2 =pH2 O values. The pH2 =pH2 O value in equilibrium with the oxide can be read from Fig. 3.2, using the method discussed previously, except that the starting point for the straight line is “H” and the pH2 =pH2 O value is determined from the H2/H2O scale. For example, the oxygen potential, in terms of pH2 =pH2 O , in equilibrium with Cr2O3 at 1000 °C, is about 5 × 103 from Fig. 3.2. Thus, Cr2O3 is likely to form at 1000 °C when the pH2 =pH2 O ratio in the environment is less than 5×103. The equilibrium reaction for an environment whose oxygen potential is controlled by pCO =pCO2 is: 2CO+O2 Ð 2CO2

ð3:8Þ

The corresponding oxygen potential is:

pO2 =eDG

=RT

1 (pCO =pCO2 )2

ð3:9Þ

The pCO =pCO2 value can be read from Fig. 3.2 using the method discussed previously, with the exception that a straight line is drawn from point “C” to the CO/CO2 scale. Thus, it is possible to obtain the oxygen potential of the environment in terms of pO2 , pH2 =pH2 O , pCO =pCO2 , and the oxygen partial pressure in equilibrium with the oxide of interest from Fig. 3.2, to determine whether or not the oxide is likely to form thermodynamically. Figure 3.2 also illustrates the relative stability of various oxides. The most stable oxides have the largest negative values of ΔG°, or the lowest value of pO2 , or the highest values of pH2 =pH2 O and pCO =pCO2 . It is clear from Fig. 3.2 that oxides of iron, nickel, and cobalt, which are the alloy bases for the majority of engineering alloys, are significantly less stable than the oxides of some solutes (e.g., chromium, aluminum, silicon, etc.) in engineering alloys. When one of these solute elements is added to iron, nickel, or cobalt, internal oxidation of the solute is expected to occur if the concentration of the solute is relatively low. As the solute concentration increases to a sufficiently high level, oxidation of the solute will be changed from internal oxidation to external oxidation, resulting in an oxide scale

that protects the alloy from rapid oxidation. This process is known as “selective oxidation.” The majority of iron-, nickel-, and cobalt-base alloys rely on selective oxidation of chromium to form a Cr2O3 scale for oxidation resistance. Some hightemperature alloys use aluminum to form an Al2O3 scale for oxidation resistance. Most oxides exhibit high melting points and remain in a solid state for the temperature range in which the alloys are used. If the oxide is present as a liquid state, catastrophic oxidation can occur. Since many engineering alloys contain many alloying elements for various metallurgical reasons, formation of oxides that become liquid at the service temperature should be prevented. Table 3.1 shows the melting points of selected oxides of alloying elements commonly found in high-temperature alloys. Most oxides remain solid until they reach extremely high temperatures. Oxides of molybdenum (MoO3) and vanadium (V2O5), however, exhibit very low melting points. Vanadium (V), which is a strong carbide former, is often used in alloy steels for increasing the strength of the material. However, the amount used typically is quite small and is not likely to form V2O5. Molybdenum (Mo) is also a strong carbide former and is used in a small amount to strengthen low-alloy steels (e.g., CrMo steels). It is unlikely these steels will be affected by MoO3-related oxidation problems. However, molybdenum is an effective alloying element for improving the resistance of the alloy to aqueous corrosion. Some stainless steel grades contain molybdenum, with superaustenitic stainless steels containing much higher levels of molybdenum. Some nickel-base alloys contain very high levels of molybdenum for either aqueous corrosion resistance or solid-solution

Table 3.1 Melting points of selected oxides for alloying elements commonly found in high-temperature alloys Qxide

αAl2O3 CoO Cr2O3 FeO Mn3O4 MoO3 Nb2O5 NiO SiO2 TiO2 V2O5 WO3 Source: Ref 3

Melting point, °C (°F)

2015 (3659) 1935 (3515) 2435 (4415) 1420 (2588) 1705 (3101) 795 (1463) 1460 (2660) 1990 (3614) 1713 (3115) 1830 (3326) 690 (1274) 1473 (2683)


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pg 9

Chapter 3:

Oxidation / 9

volatile CrO3 (Ref 5–8). Table 3.2 shows the weight-loss data for Cr2O3 pellets after heating to 1000 to 1200 °C (1830 to 2190 °F) in dry O2 due to formation of gaseous CrO3 by oxidation of Cr2O3 (Ref 5). Caplan and Cohen (Ref 5) also observed that moisture promoted volatilization of Cr2O3. Asteman et al. (Ref 9) indicated that high vapor pressure of CrO2(OH)2 can form by reacting Cr2O3 with H2O in O2-containing environments. The theoretical calculated partial pressure of CrO2(OH)2 as a function of temperature for the O2-containing environment with pO2 =0:9 atm and pH2 O =0:1 atm is shown in Fig. 3.4 (Ref 9).

strengthening. The formation of MoO3 and its effect on oxidation are discussed in section 3.2.2 “Volatility of Oxides” and section 3.4.17 “Catastrophic Oxidation.” 3.2.2 Volatility of Oxides Some oxides exhibit high vapor pressures at very high temperatures (e.g., above 1000 °C, or 1830 °F). Oxide scales become less protective when their vapor pressures are high. Figure 3.3 shows vapor pressures of several refractory metal oxides exhibiting high vapor pressures at temperatures above 1000 °C (1830 °F) (Ref 4). Vanadium is typically used in small quantities as a carbide former in alloy steels. Thus, the volatility of VO2 is generally of no concern in oxidation of alloys. Molybdenum (Mo) and tungsten (W) are often used as alloying elements in significant amounts in Ni- or Co-base alloys as solution-strengthening elements. Formation of WO3 or MoO3 may occur under certain conditions in some alloy systems, particularly in alloy systems containing insufficient chromium for forming a protective Cr2O3 scale. A majority of engineering alloys rely on the Cr2O3 scale to provide resistance to oxidation. When heated to very high temperatures (i.e., above 1000 °C), Cr2O3 can react with O2 to form

Table 3.2 Weight loss of Cr2O3 on heating in dry O2 and Ar environments Run

Temperature, °C (°F)

Time, h

Gas

Gas flow, mL/min

Weight loss, mg

1 2 3 4 5 6 7 8 9

1100 (2010) 1200 (2190) 1200 (2190) 1200 (2190) 1200 (2190) 1200 (2190) 1200 (2190) 1200 (2190) 1200 (2190)

20 20 20 20 20 20 42 66 115

Dry O2 Dry O2 Dry O2 Dry O2 Dry O2 Dry O2 Dry O2 Dry Ar Dry Ar

200 10 10 20 200 200 200 200 192

0.6 2.1 1.3 1.8 2.3 2.6 8.0 0 0

Source: Ref 5

Temperature, °C 1

1500

1000

600

10–2

WO3 CrO3

10–4

Vapor pressure, atm

MoO3 10–6 10

VO2

–8

10–10 10–12 10–14 10–16 10–18 0.5

0.6

0.7

0.8

0.9

1

1.1

1.2

Inverse temperature, 103/T (K–1)

Fig. 3.3

Vapor pressures of several refractory metal oxides exhibiting high vapor pressures at temperatures above 1000 °C (1830 °F). Source: Ref 4


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pg 10


–2

–4 CrO2(OH)2, g

Log p, X, atm

–6

–8

–10

–12 CrO3, g

–14

–16 600

700

800

900

1000

1100

1200

1300

Temperature, K Theoretical partial pressures of CrO2(OH)2 and CrO3 as a function of temperature for the environment with pO2 =0:9 atm and pH2 O =0:1 atm: Source: Asteman et al. (Ref 9)

For service temperatures above 1200 °C (2190 °F), the increasing volatility of oxides can progressively cause the oxide to lose its protective capability. SiO2 and Al2O3 are the only two oxides that are capable of forming a very protective barrier against oxidation at temperatures above 1200 °C (Ref 9). However, the SiO2 scale may lose some protective capability by forming gaseous SiO at low oxygen partial pressures (Ref 10).

Inverse log Parabolic Oxide mass, m

Fig. 3.4

Log

Linear

Time, t

3.3 Kinetic Considerations The kinetics of oxidation of metals and alloys generally follow several reaction rates. Most reactions follow a parabolic rate. Some reactions follow a linear rate. Some other reaction kinetics may include logarithmic and inverse logarithmic rates. These reaction kinetics are illustrated schematically in Fig. 3.5 (Ref 11), and a brief summary, based on the article by Danielewski (Ref 11), is presented in sections 3.3.1 to 3.3.3.

Fig. 3.5

Different oxidation kinetics. Source: Ref 11

decreases with increasing time due to the increasing diffusion distance for ions. The oxidation rate is, thus, inversely proportional to the thickness of the oxide scale: X 2 =k′t

ð3:10Þ

where X is the oxide scale thickness, t is the exposure time, and k′ is the parabolic constant; when t = 0, X = 0.

3.3.1 Parabolic Kinetics When the oxide scale forms on the metal surface, the oxidation reaction is controlled by the diffusion of ions through the oxide scale, which is in turn controlled by the chemical potential gradient as a driving force. As the thickness of the oxide scale increases, the rate of oxidation

3.3.2 Linear Kinetics When the oxide scale forming on the metal surface provides no protection barrier due to oxide cracking and spalling, volatile oxides, and molten oxidation products, the oxidation rate generally remains constant with increasing


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pg 11

Chapter 3: Oxidation / 11

time. The linear oxidation kinetic rate can be expressed by:

The inverse logarithmic rate can be expressed by the following equation:

X =kl t

1=X =b7ki log t

ð3:11Þ

where X is the mass (or thickness) of the oxide, t is the exposure time, and kl is the linear rate constant; when t = 0, X = 0. 3.3.3 Logarithmic and Inverse Logarithmic Kinetics

where b and ki are constants.

3.4 Oxidation in Air, O2, and “Clean” Combustion Atmospheres 3.4.1 Carbon and Cr-Mo Steels

At very low temperatures when the oxide film forms on the metal surface, the oxidation rate usually follows either a logarithmic or inverse logarithmic rate. The driving force for the oxidation is the electric field across the oxide film. The logarithmic rate can be expressed by: X =ke log (at+1)

ð3:13Þ

ð3:12Þ

where ke and a are constants.

Carbon and Cr-Mo steels are the most widely used engineering materials and are used extensively for high-temperature applications in power generation, chemical and petrochemical processing, petroleum refining, pulp and paper industry, industrial heating, and metallurgical processing. At temperatures below 570 °C (1060 °F), iron (Fe) oxidizes to form Fe3O4 and Fe2O3. Above

120

110

100 1400 °F

90

Weight-loss, mg/cm2

80

70 1200 °F 60

50

40

30

20

1000 °F

10

800 °F

0

0

100

200

300

400

500

600

700

Time, h

Fig. 3.6

Oxidation behavior of plain low-carbon steel in air at 430, 540, 650, and 760 °C (800, 1000, 1200, and 1400 °F). Source: Ref 12


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570 °C (1060 °F), it oxidizes to form FeO, Fe3O4, and Fe2O3. The oxidation behavior of carbon steel in air at 430, 540, 650, and 760 °C (800, 1000, 1200, and 1400 °F) is summarized in Fig. 3.6 (Ref 12). At 430 and 540 °C (800 and

Average penetration/side, mils

10.5

Calculated continuing penetration rate

53 mpy*

9.0 Carbon steel - 1200 °F

7.5 6.0 4.5

5.3 mpy* A242 type 1 HSLA steel - 1200 °F

3.0

Carbon steel - 1000 °F

1.5

2.3 mpy*

A242 type 1 HSLA steel - 1000 °F 1.0 mpy*

0

200

400

600

800 1000

Exposure time, h

Fig. 3.7

Oxidation of carbon steel and high-strength low-alloy (HSLA) steel in air. Source: Ref 13, reproduced

from Ref 14

1000 °F), carbon steel showed very little weight gain after exposure for 500 h. As the temperature was increased to 650 °C (1200 °F), the oxidation rate was significantly increased. Carbon steel suffered rapid oxidation at 760 °C (1400 °F), exhibiting essentially a linear rate of oxidation attack. Vrable et al. (Ref 13) reported oxidation data for carbon steel and high-strength low-alloy (HSLA) steel, as shown in Fig. 3.7 (Ref 14). At 650 °C (1200 °F), carbon steel suffered an oxidation rate of about 1.3 mm/yr, or 53 mpy (mils per year). The oxidation rate is expected to be much higher when exposed to temperatures higher than 650 °C (1200 °F). Recent test results by John (Ref 15) showed that carbon steel exhibited about 0.25 mm/yr (10 mpy) of oxidation at 604 °C (1120 °F). Figure 3.7 also illustrates that HSLA steel is more oxidation resistant than carbon steel, presumably due to minor alloying elements such as manganese, silicon, chromium, and nickel. Cr-Mo steels are used at higher temperatures than carbon steel because of higher tensile and creep-rupture strengths as well as better microstructural stability. Molybdenum and chromium provide not only solid-solution strengthening but also carbide strengthening. Low-alloy steels with chromium and silicon additions exhibit better oxidation resistance than carbon steel. The beneficial effects of chromium and silicon additions to carbon steel are summarized in Fig. 3.8 (Ref 16). Silicon is very effective in improving the oxidation resistance of Cr-Mo steels. Addition of 1.5% Si to 5Cr-0.5Mo steel significantly improved its oxidation resistance. The most important alloying element for improving oxidation resistance is chromium. As shown in Fig. 3.8, for 0.5% Mo steels, increasing chromium from 1 to 9% significantly increases oxidation resistance. The 7Cr-0.5Mo and 9Cr-1Mo steels showed negligible oxidation rates at temperatures up to 680 °C (1250 °F) and 700 °C (1300 °F), respectively. Further increases in chromium improve oxidation resistance even more. Alloys become martensitic or ferritic grades of stainless steels (400 series) when chromium content is increased to 12% or higher. 3.4.2 Martensitic, Ferritic, and Austenitic Stainless Steels

Fig. 3.8

Effects of chromium and/or silicon on the oxidation resistance of steels in air. Source: Ref 16

The superior oxidation resistance of martensitic and ferritic stainless steels to that of carbon and Cr-Mo steels is illustrated in Fig. 3.9 (Ref 17). As chromium content in the straight

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Fig. 3.9

Oxidation resistance of carbon, low-alloy, and stainless steels in air after 1000 h at temperatures from 590 to 930 °C (1100 to 1700 °F). Source: Ref 17

chromium steels increases from 9 to 25%, resistance to oxidation improves significantly. The 25Cr steel (Type 446) is the most oxidation resistant among the 400 series stainless steels, due to the development of a continuous Cr2O3 scale on the metal surface. In Fe-Cr alloys, it appears that a minimum of approximately 18wt % Cr is needed to develop a continuous Cr2O3 scale against further oxidation attack (Fig. 3.10) (Ref 18). Cyclic oxidation studies conducted by Grodner (Ref 19) also revealed that Type 446 was the best performer in the 400 series stainless steels, followed by Types 430 (14–18Cr), 416 (12–14Cr), and 410 (11.5–13.5Cr) (Fig. 3.11). Figure 3.11 shows that Fe-12Cr steels, such as Types 410 and 416, showed increased rates of cyclic oxidation above 760 °C (1400 °F). At 650 °C (1200 °F) in air, cycling from 650 to 300 °C, 12Cr-1Mo steel (X20 CrMoV 12 1) steel

exhibited a thin, adherent (Fe,Cr)2O3 scale, as observed by Walter et al. (Ref 20). The growth of the (Fe,Cr)2O3 scale as a function of the accumulated isothermal hold time up to 1000 h is shown in Fig. 3.12 (Ref 20). John (Ref 15) reported that Fe-12Cr steel (Type 410) exhibited an air oxidation rate of 0.25 mm/yr (10 mpy) at 832 °C (1530 °F). Another ferritic stainless steel, 18SR (about 18% Cr), was found to be as good as, and sometimes better than, Type 446 (25% Cr), as illustrated in Tables 3.3 and 3.4 (Ref 21). This was attributed to the addition of 2% Al and 1% Si to the alloy. Furthermore, both of these ferritic stainless steels, Type 446 and 18SR, showed better cyclic oxidation resistance than some austenitic stainless steels, such as Type 309 and 310, when cycled to 980 to 1040 °C (1800 to 1900 °F), as shown in Table 3.4.

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pg 14


Fe2O3 Fe3O4 FeO Fe 10–6 Fe2O3 Fe3O4 Parabolic rate constant, g2 · cm4 · s1

FeO 10–7

Fe – 2Cr

Fe/Cr oxide

Fe2O3 (Fe,Cr)2O3

10–8

Fe – 9Cr

Fe2O3

Fe3O4

10–9

Cr2O3 Fe – 16Cr

FeFe(2 ... x)CrxO4

10–10 Cr2O3 Fe – 28Cr 10–11 0

10

20

30

40

50

60

70

80

90

100

Alloy chromium content, wt%

Fig. 3.10

Effect of chromium content on oxidation of Fe-Cr alloys at 1000 °C (1830 °F) in 0.13 atm O2. Source: Ref 18

When the service temperature is above 650 °C (1200 °F), ferritic stainless steels, which have a body-centered cubic (bcc) crystal structure, drastically lose their elevated-temperature strength (both tensile and creep-rupture strength). As a result, the application of ferritic stainless steels becomes limited at higher temperatures. At these temperatures, alloys with a face-centered cubic (fcc) crystal structure are preferred because of their higher creep-rupture strength. Nickel is added to Fe-Cr steels to stabilize the austenitic structure. The austenitic structure is inherently stronger and more creep resistant than the ferritic structure (Ref 22). The 300 series austenitic stainless steels have been widely used for high-temperature components in various industries because of their strength and high-temperature corrosion resistance, including oxidation resistance. These alloys exhibit higher elevated-temperature

strength than do ferritic stainless steels. Furthermore, they do not suffer 475 °C (885 °F) embrittlement or ductility-loss problems in thick sections and in heat-affected zones as do ferritic stainless steels. Nevertheless, some austenitic stainless steels can suffer some ductility loss upon long-term exposure to intermediate temperatures (e.g., 540 to 800 °C, or 1000 to 1470 °F) due to sigma-phase formation. The oxidation resistance of two austenitic stainless steels, Types 309 and 310, is compared with that of a number of ferritic stainless steels in Fig. 3.11. Several austenitic stainless steels are compared in Fig. 3.13 (Ref 23). Nickel improves the resistance of alloys to cyclic oxidation. Moccari and Ali (Ref 24) also observed the similar beneficial effects of nickel in improving the oxidation resistance of alloys. Brasunas et al. (Ref 25) studied the oxidation behavior of about 80 experimental Fe-Cr-Ni alloys exposed


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pg 15


Oxidation resistance of several stainless steels as a function of temperature. Source: Ref 19

Thickness of the oxide layer (d ), µm

Fig. 3.11

40 35 30 25 20 15 10 5 0 0

100

200

300

400

500

600

700

800

900

1000

Testing time (t ), h

Fig. 3.12

Oxidation behavior of 12Cr-1Mo steel (X20 CrMoV 12 1) at 650 °C (1200 °F) in air with every 8 h of exposure specimens being cycled from 650 to 300 °C. Source: Ref 20

to air-H2O mixture at 870 to 1200 °C (1600 to 2190 °F) for 100 and 1000 h. They observed that increases in nickel in excess of 10% in alloys containing 11 to 36% Cr improved the oxidation resistance of the alloys. John (Ref 15) reported that Type 304 and 310 exhibited 0.25 mm/yr (10 mpy) of oxidation attack (both metal loss and

internal oxidation penetration) at 893 and 982 °C (1640 and 1800 °F), respectively, in air. In Fig. 3.13, several high-nickel alloys were found to be more resistant to oxidation than austenitic stainless steels. The oxidation behavior of highnickel Fe-Ni-Cr alloys and Ni-base alloys is discussed in later sections of this chapter.


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Table 3.3 Cyclic oxidation resistance of several stainless steels in air cycling to 870 to 930 °C (1600 to 1700 °F) temperature range Specimen weight changes after indicated cycles, mg/cm2 Alloy

288 cycles

409 + Al Destroyed 430 9.9 22-13-5 0.5 442 0.7 446 0.3 309 0.3 18SR 0.3

480 cycles

750 cycles

958 cycles

… Destroyed −3.0 1.2 0.4 −4.6 0.4

… … −18.8 1.5 0.2 −23.7 0.5

… … −35.7 1.5 0.1 −32.6 0.6

Note: 15 min in furnace and 15 min out of furnace. Source: Ref 21

Table 3.4 Cyclic oxidation resistance of several stainless steels in air cycling to 980 to 1040 °C (1800 to 1900 °F) temperature range Specimen weight changes after indicated cycles, mg/cm2 Alloy

130 cycles

368 cycles

561 cycles

753 cycles

1029 cycles

446 18SR 309 310

0.4 0.7 −24.2 1.5

0.5 1.1 −77.5 −11.3

−0.2 1.5 −178.3 −29.3

7.0 2.2 −242 −62.8

−19.4 3.0 −358 −107

Note: 15 min in furnace and 15 min out of furnace. Source: Ref 21

In evaluating materials for automobile emission-control devices, such as thermal reactors and catalytic converters, Kado et al. (Ref 26) and Michels (Ref 27) have carried out cyclic oxidation tests on various stainless steels. In cyclic oxidation tests performed by Kado et al. (Ref 26) in still air at 1000 °C (1830 °F) for 400 cycles (30 min in the furnace and 30 min out of the furnace), Types 409 (12Cr), 420 (13Cr), and 304 (18Cr-8Ni) suffered severe attack. Type 420 (13Cr) was completely oxidized after only 100 cycles, although the sample did not show any weight changes. Alloys that performed well under these conditions were Types 405 (14Cr), 430 (17Cr), 446 (25Cr), 310 (25Cr-20Ni), and DIN 4828 (19Cr-12Ni-2Si), as illustrated in Fig. 3.14. When cycled to 1200 °C (2190 °F) for 400 cycles (30 min in the furnace and 30 min out of the furnace), all the alloys tested except F-1 alloy (Fe-15Cr-4Al) suffered severe oxidation attack (Fig. 3.15). This illustrates the superior oxidation resistance of alumina formers (i.e., alloys that form Al2O3 scales when oxidized at elevated temperatures). The data also illustrate that for temperatures as high as 1200 °C (2190 °F), Cr2O3 oxide scales can no longer provide adequate oxidation resistance. Oxidation of Fe-Cr-Al alloys is discussed in Section 3.4.7.

Kado et al. (Ref 26) also investigated oxidation behavior in a combustion environment that simulated the gasoline engine. Their test involved air-to-fuel ratios of 9 to 1 and 14.5 to 1 and regular gasoline that contained 0.01wt% S. Exhaust gas taken from the exhaust manifold was mixed with air before being piped into a furnace retort where tests were performed. Test specimens were exposed to the mixture of exhaust gas and air. The gas mixture consisted of 72.4% N2, 9.7% H2O, 9.93% O2, 8% CO2, and 507 ppm NOx, when an air-to-fuel ratio of 14.5 to 1 was used for combustion, while that coming from the combustion using an air-to-fuel ratio of 9 to 1 consisted of 70.6% N2, 13.7% H2O, 3.21% O2, 12.5% CO2, and 34 ppm NOx. The total accumulated test duration was 200 h (400 cycles with 30 min in the hot zone and 30 min out of the hot zone). Test results along with air oxidation data are summarized in Fig. 3.16. There were no significant differences between air and exhaust gas test environments when tested at 800 °C (1470 °F). All the alloys tested showed negligible attack except Type 304, which exhibited much more severe attack in the exhaust environment. When the test temperature was increased to 1000 °C (1830 °F), all the 400 series stainless steels with less than 17% Cr (i.e., Types 409, 405, 410, and 430) and Type 304 exhibited significantly more oxidation attack in the exhaust environment. Type 310, Type 446, DIN 4828 (Fe-19Cr-12Ni-2Si), F-1 alloy (Fe15Cr-4Al), A-1 alloy (Fe-16Cr-13Ni-3.5Si), and A-2 alloy (Fe-20Cr-13Ni-3.5Si) performed well. At 1200 °C (2190 °F), all the alloys tested suffered severe oxidation attack with the exhaust environment being more aggressive than air. These authors attributed the enhanced attack to the presence of sulfur in the exhaust gas environment, although low-sulfur (0.01%) gasoline was used for testing. Sulfur segregation to the scale/metal interface was detected. In a study by Michels (Ref 27), the engine combustion atmosphere was also found to be significantly more corrosive than the air-10% H2O environment. The engine combustion exhaust gas contained about 10% H2O along with 2% CO, 0.33 to 0.55% O2, 0.05 to 0.24% hydrocarbon, and 0.085% NOx. The balance was presumably N2 (not reported in the paper). The engine exhaust gas was piped into a furnace retort where the tests were performed. The results, which were generated in the air-10%H2O and the engine exhaust environment, are shown in Fig. 3.17. After exposure to the air-10%H2O

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pg 17


10

625 600 80Ni-20Cr 60Ni-15Cr

0

800 22Cr-32Ni

–10 Type 310 25Cr-20Ni

– 20

Change in weight, %

19Cr-14Ni 20Cr-25Ni Type 309 23Cr-13Ni – 30

– 40 Type 347 18Cr-8Ni(Cb)

– 50 Type 304 18Cr-8Ni

– 60

– 70 0

200

400

600

800

1000

Hours of cyclic exposure (15 min heating – 5 min cooling)

Fig. 3.13

Cyclic oxidation resistance of several stainless steels and nickel-base alloys in air at 980 °C (1800 °F). Source: Ref 23

environment at 980 °C (1800 °F) for 102 h, Type 309, Type 310, 18SR, alloy OR-1 (Fe-13Cr-3Al), alloy 800, and alloy 601 were all relatively unaffected. On the other hand, only 18SR and alloy 601 were relatively unaffected by the engine exhaust gas environment, with alloy OR-1, Type 309, Type 310, and alloy 800 suffering severe oxidation attack. The sulfur content

in the gasoline used in this test was not reported. The relatively high gas velocity, about 6.1 to 9.2 m/s (20 to 30 ft/s) was considered by the author to be one of the possible factors responsible for much higher oxidation attack in the engine exhaust gas test. Oxidation data generated in combustion atmospheres is relatively limited. No systematic


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pg 18


Fig. 3.14

Cyclic oxidation resistance of several ferritic and austenitic stainless steels in still air at 1000 °C (1830 °F) for up to 400 cycles (30 min in furnace and 30 min out of furnace). Source: Ref 26

409 0 F–1 (Fe-15Cr-4AI)

Weight change, mg/cm2

430 –100

DIN 4828 (19Cr-12Ni-2Si)

420 –200

310

304

446

–300

0

100

200

300

400

Number of cycles

Fig. 3.15

Cyclic oxidation resistance of several ferritic and austenitic stainless steels in still air at 1200 °C (2190 °F) for up to 400 cycles (30 min in furnace and 30 min out of furnace). Source: Ref 26


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pg 19


1200 °C (2190 °F)/400 cycles 100

80

60

40

Thickness loss, %

20

0

1000 °C (1830 °F)/400 cycles

100

80

60

40

20

0 800 °C (1470 °F)/400 cycles 20

) .5

Si

) 3N

(F e-

20

C

r-1

3N (F e-

i-3

.5 i-3

31 A– 2

16

C

r-1

D

A– 1

In air In exhaust gas (R = 9)

Si

0S

0

IN

48

31

28

4 30

6

0

44

r-4

F–

1

(F e-

15

C

43

) AI

0 41

5 40

40

9

0

Fig. 3.16

Comparison of cyclic oxidation resistance between air and gasoline engine exhaust gas environments at 800, 1000, and 1200 °C (1470, 1830, 2190 °F) for 400 cycles (30 min in hot zone and 30 min out of hot zone). Alloy F-1 suffered localized attack at 1200 °C in engine exhaust gas. Source: Ref 26

studies have been reported that varied combustion conditions, such as air-to-fuel ratios. In combustion atmospheres, the oxidation of metals or alloys is not controlled by oxygen only. Other combustion products, such as H2O, CO, CO2, N2, hydrocarbon, and others, are expected to influence oxidation behavior. When air is used for combustion, nitridation in conjunction with oxidation can occur in combustion atmospheres under certain conditions. This nitridation/ oxidation behavior of alloys is discussed in

Chapter 4 “Nitridation.” The presence of water vapor can be an important factor in affecting oxidation behavior of alloys. The effect of water vapor on the oxidation resistance of alloys is covered in Section 3.4.15. 3.4.3 Surface Chemistry versus Bulk Chemistry It is important to point out that the surface chemistry may not be the same as the reported


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Fig. 3.17

Cyclic oxidation resistance of several ferritic and austenitic stainless steels in (a) air-10H2O at 980 °C (1800 °F) cycled every 2 h, and (b) gasoline engine exhaust gas at 980 °C (1800 °F) cycled every 6 h. Source: Ref 27

“bulk” chemistry. Some of the manufacturing processes involved in the production of an alloy product, such as plate, sheet, or tubular products, may result in lower chromium contents at or near the surface of the product. This is particularly important for austenitic stainless steels, since the chromium specification range for austenitic stainless steels can be at or near the borderline for forming a continuous Cr2O3 scale. For example, ASTM A 213/A 213M (or ASME SA213/SA-213M) specification for the chromium range is 18.0 to 20.0% for Type 304, 16.0 to 18.0% for Type 316, 17.0 to 20.0% for Type 321, and 17.0 to 20.0% for Type 347. The lower end of the chromium content in these alloys is essentially at the minimum level that is considered to be required for forming a continuous Cr2O3 scale when exposed to elevated temperatures. A slight surface depletion in chromium due to some manufacturing processes may result in the condition that a continuous, protective Cr2O3 scale cannot be formed, thus resulting in accelerated oxidation due to formation of nonprotective iron oxides. Some manufacturing processes are prone to producing the final finished product with surface depletion of chromium. This surface depletion

becomes more critical for sheet products or thinwall tubular products because of much higher percentage of the surface-depletion zone in the overall thickness of the component. In alloy manufacturing, annealing is required after each cold-rolling step for sheet product manufacturing (or cold pilgering for reduction in thickness and sizes for tubular product manufacturing) to soften the metal for further cold-reduction steps until a final product is produced. Stainless steels are typically annealed in the temperature range of 1010 to 1120 °C (1850 to 2050 °F) (Ref 28). When annealing is performed in air, heavy chromium oxide scales form on the metal surface. As a result, the matrix immediately underneath the oxide scales can be depleted in chromium. Figure 3.18 shows the concentration profile of chromium near the surface of the plate of alloy AL-6XN (Fe-21Cr-24Ni-6.5Mo-0.2N) after annealing in air at 1120 and 1175 °C (2050 and 2150 °F), respectively (Ref 29). Oxidation during annealing at either temperature resulted in a significant chromium depletion near the surface of the plate. It is quite common to perform annealing in air during manufacturing of stainless steels. This is commonly referred to as “black annealing,” as opposed to “bright

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pg 21


annealing,” which is performed in a protective atmosphere, such as hydrogen atmosphere. The oxide scales on the alloy plate, sheet, or tubing are generally removed by acid pickling. This manufacturing process can often produce flat products (plate or sheet) as well as tubular products with surface depletion of chromium. When bright annealing is performed, the alloy surface is protected from oxide scale formation. As a result, chromium depletion at and near the product surface is minimized. Surface depletion in chromium becomes increasingly critical as the thickness of the component, such as sheet or tubular product,

decreases. Figure 3.19 shows an example of surface depletion of chromium in a thin-gage commercial heat-exchanger tube made from Type 321 in the as-fabricated condition (Ref 30). The manufacturing process involved was not known. The Type 321 tube is shown to exhibit depletion in chromium near the tube surface when analyzed by energy-dispersive x-ray spectroscopy (EDX). The analysis was terminated at approximately about 0.5 μm from the tube surface. It is expected that the chromium content would decrease further if the analysis was performed at locations closer to the surface. After service for 6 months as a recuperator tube at

22 As hot rolled

Chromium content, wt%

20

18

Annealed 1175 °C Annealed 1120 °C

16 AL-6XN 14 2

0

8

6

4

10

12

14

16

18

20

Distance, µm

Fig. 3.18

Surface deletion of chromium near the surface of the plate of alloy AL-6XN (Fe-21Cr-24Ni-6.5Mo-0.2N) after annealing in air at 1120 and 1175 °C (2050 and 2150 °F), respectively. Source: Ref 29

18.4

Cr concentration, wt%

18.2 18 17.8 17.6 17.4 17.2 17 16.8 0

1

2

3

4

5

6

7

8

9

10

Distance from tube OD surface, µm

Fig. 3.19

Surface depletion of chromium observed in a thin-gage commercial heat-exchanger tube in the as-fabricated condition made from Type 321. Source: Ref 30


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metal temperatures of approximately 620 to 670 °C (1150 to 1240 °F) for preheating air, significant oxide spalling and scaling was observed on the air side of the heat-exchanger tube. Figure 3.20 shows heavy oxide scales formed on the side of the tube exposed to the incoming air after 6 months of service (Ref 30). The oxide scales were analyzed by scanning electron microscopy/energy-dispersive x-ray spectroscopy (SEM/EDX) analysis; the results are shown in Fig. 3.21. The analysis showed that the outermost oxide layer was essentially iron oxides with very little chromium. Thus, the stainless steel recuperator no longer exhibited adequate resistance to oxidation because of its failure to form and maintain a protective Cr2O3 scale. Once a protective Cr2O3 scale is no longer present on the stainless steel surface, iron oxides then take over. This results in scaling and accelerating oxidation. This is often referred to as “breakaway” oxidation (or corrosion). “Breakaway” oxidation is discussed in Section 3.4.13. Stainless steels manufactured by different producers can exhibit different chemical compositions and different surface characteristics. This is illustrated in Fig. 3.22 (Ref 30). Two Type 321 heat-exchanger tubes, manufactured by two different suppliers (supplier A and B), were tested in the field in the same recuperator as described previously for preheating air at approximate metal temperature of 620 to 670 °C (1150 to 1240 °F) for about 1008 h. Figure 3.22 (a) shows the formation of mushroom-type oxide nodules on the surface of the Type 321 tube, produced by Supplier A. This is considered the initiation of breakaway oxidation. As shown in Fig. 3.22(a), the mushroom-type oxide nodule consisted of two layers of oxides, light grayish outer oxide scale and darkish inner oxide layer. The rest of the metal surface was still protected

by a thin, adherent oxide scale, which was not clearly revealed in the figure at about 400× magnification. Once the formation of the light, grayish outer oxide scale formed on the metal surface, the inner oxidation took place with accelerated growth. The outer oxide scale was found to be iron-rich oxides with little chromium, as shown in Fig. 3.23 (Ref 30). With the formation of the iron-rich oxide scale at the outer layer, which failed to provide protection, the inner Fe-Cr oxides were found to penetrate into the metal, causing a relatively massive internal oxidation penetration (Fig. 3.23). In the same test, the other Type 321 heat-exchanger tube (produced by supplier B), which was subjected to the same test condition and duration, was found to exhibit a thin, adherent oxide scale with no evidence of mushroom-type oxide nodules, as shown in Fig. 3.22(b). The oxide scale was analyzed by SEM/EDX, showing an Fe-Cr oxide scale formed on the metal surface (Fig. 3.24). Bulk chemical compositions of two tubes were analyzed; results are shown in Table 3.5. The chemical analysis results showed that the chromium in the Type 321 tube from supplier A was essentially at the lower end of the specification range. Any depletion in chromium near the surface could result in a chromium level that is below the specification limit. On the other hand, the tube from supplier B contained much higher chromium and was found to be much more resistant to oxidation under the same test conditions. It appears that some stainless steel manufacturers produce their products with leaner chemistry in terms of major alloying elements, such as chromium and nickel. For hightemperature oxidation and other modes of corrosion, stainless steel with the chromium in the lower end of the specification range can be potentially more prone to breakaway oxidation

25 µm

Fig. 3.20

Heavy oxide scales formed on the side of Type 321 recuperator tube that was exposed to the incoming air after 6 months of service with the metal temperatures approximately 620 to 670 °C (1150 to 1240 °F). This tube was from the same batch of tubes that shows surface chromium depletion (Fig. 3.19). Source: Ref 30

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pg 23


A

B C D

E 25 µm (a)

30 µm

Fig. 3.21

Scanning electron micrograph (backscattered electron image) showing the oxide scales formed on the outside diameter of the heat-exchanger tube (from the same batch of tubes that showed surface chromium depletion) exposed to air for 6 months. Energy-dispersive x-ray spectroscopy (EDX) analysis was performed to determine the chemical compositions at different locations, marked “A” to “E.” The results (wt%) of the EDX analysis are summarized below (minor elements not included). Source: Ref 30 A: 1.9% Cr, 97.6% Fe. B: 44.2% Cr, 44.4% Fe, 6.8% Ni. C: 48.8% Cr, 40.4% Fe, 4.4% Ni. D: 28.3% Cr, 46.4% Fe, 21.0% Ni. E: 38.7% Cr, 45.3% Fe, 10.2% Ni.

or corrosion. A brief discussion of this important issue is presented in Section 3.4.5. 3.4.4 Surface Conditions As discussed in Section 3.4.3 “Surface Chemistry versus Bulk Chemistry,” the concentration of chromium at and near the surface of the alloy product plays a significant role in the oxidation of stainless steels such as Types 304, 316, 321, 347, and so forth. This is because the chromium concentration of these stainless steels is at the lower end of the chromium range that is generally required to form a continuous Cr2O3 scale when heated to elevated temperatures.

25 µm (b)

Fig. 3.22

Type 321 heat-exchanger tubes, which were manufactured by two different alloy suppliers, were tested in the same facility as described previously for preheating air at approximate metal temperature of 620 to 670 °C (1150 to 1240 °F) for about 1008 h. (a) Supplier A. (b) Supplier B. Note the tube from supplier A showed the initiation of accelerated oxidation attack (a) as opposed to the tube from supplier B showing no sign of accelerated oxidation attack (b).

Manufacturing processes can greatly influence the surface chemistry of an alloy product. Stainless steels can be finished into the final product by bright annealing (i.e., annealing is performed in a protective atmosphere, such as hydrogen environment or dissociated ammonia environment). This process generally produces a product with minimal depletion of chromium at or near the surface. On the other hand, when the alloy product is finished by black annealing (i.e., annealing is performed in air or combustion atmosphere in the furnace) and followed by acid pickling, there is a good chance that the alloy product may exhibit a surface depletion of chromium. This is particularly important for thin-gage sheet products or thin tubular products.


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pg 24


In most laboratory oxidation tests, the test specimens are typically prepared by grinding and polishing with different grits of emery papers prior to testing. The objective of grinding and polishing the test specimens to a certain surface finish condition is to keep the surface condition of all test specimens constant in order to compare the oxidation behavior of different alloys. However, the mechanical forces of grinding and polishing can produce a thin cold-worked layer on the specimen surface. This cold-worked structure at the surface layer can significantly enhance the diffusion of chromium from the interior to the surface of the metal to form

1

chromium oxide scales when heated to elevated temperatures, thus increasing the oxidation resistance of the alloy. For some critical applications involving a thin-gage sheet (or foil) product or a thin tubular product, testing should be carried out on the specimen that retains the surface condition of the product without prior surface grinding or mechanical polishing. Electropolishing, which is not commonly used to improve the surface finish of the alloy product for high-temperature services, may cause surface depletion of chromium for the product. Nevertheless, some investigators may use electropolishing to prepare the surface condition of the test specimens. Table 3.6 shows some comparison oxidation data generated in wet O2 between the wet ground and electropolished surface conditions for several stainless steels (Ref 31). 3.4.5 Today’s Stainless Steels Some stainless steel producers may manufacture stainless steels at the bottom of the

2

1 3

2 4

13 µm 13 µm

Fig. 3.23

Scanning electron micrograph (backscattered electron image) showing the oxide scales formed on the outside diameter of Type 321 tube (from supplier A) exposed to air at approximately 620 to 670 °C (1150 to 1240 °F) for 1008 h. Energy-dispersive x-ray spectroscopy (EDX) analysis was performed to determine the chemical compositions at different locations as marked 1 to 4 in the oxides. The results (wt%) of the EDX analysis are (minor elements not included): 1: 12% Cr, 86% Fe. 2: 52% Cr, 24% Fe, 18% Ni. 3: 32% Cr, 25% Fe, 38% Ni. 4: 37% Cr, 52% Fe, 7% Ni.

Table 3.5

Fig. 3.24

Scanning electron micrograph (backscattered image) showing the oxide scales formed on the outside diameter of Type 321 tube (from supplier B) exposed to air at approximately 620 to 670 °C (1150 to 1240 °F) for 1008 hours. EDX analysis was performed to determine the chemical compositions at different locations, marked as No. 1 and 2, in the oxides. The results (wt%) of the EDX analysis are (minor elements not included): 1: 35% Cr, 50% Fe, 6% Ni. 2: 21% Cr, 59% Fe, 4% Ni.

Chemical compositions of Type 321 tubes from Suppliers A and B Composition, wt%

Supplier

A B

C

Cr

Ni

Ti

Mn

Si

Mo

Cu

P

S

Fe

0.043 0.072

17.01 17.77

8.85 8.81

0.32 0.38

1.04 0.94

0.54 0.76

0.35 0.21

0.26 0.09

0.032 0.035

6.35

>6.35

>6.35

>6.35

RA85H

617

HR120

800HT

Average internal attack Metal loss

6 5 4 3 2 1 0

214

HR160

230

601

Fig. 3.47 Oxidation data in terms of metal loss, resulting from external oxide scales, and internal attack, resulting from internal oxide and/or void formation, for alumina-former alloy 214 and chromia/silica-former alloy HR160 along with several other nickeland iron-base alloys, generated at 1200 °C (2200 °F) in air for 360 days. Source: Ref 49

(a)

20 µm

(b)

(c)

(d)

Fig. 3.48

A high-silicon Ni-Cr-Fe alloy, 45TM (Ni-27Cr-23Fe-2.7Si), after oxidation testing in air for 1056 h at (a) 850 °C (1560 °F), (b) 1000 °C (1830 °F), (c) 1100 °C (2010 °F), and (d) 1200 °C (2190 °F). For testing at 850, 1000, and 1100 °C, specimens were cycled to room temperature from the test temperature every 16 h, with 2 h of heating and 6 h of cooling. For 1200 °C testing, specimens were removed from the hot zone every 16 h. Magnification bar represents 20 μm for all micrographs. Source: Ref. 52. Courtesy of ThyssenKrupp VDM

cobalt-base superalloys, is summarized in Table 3.12 (Ref 68). Tests were conducted in flowing air (30 cm/min) at 980, 1095, 1150, and 1200 °C (1800, 2000, 2100, and 2200 °F) for

1008 h. The specimens were cooled to room temperature for visual examination once a week (168 h). Specimens from sheet products were surface ground to maintain the uniform surface

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Fig. 3.49

Long-term oxidation tests (10,000 h) in air at 815 °C (1500 °F) with 1000 h cycles (a total of 10 cycles to room temperature for the entire test) for iron-, nickel-, and cobalt-base alloys. Also included is the upper limit of the metal loss for isothermal tests (i.e., 10,000 h without cycling to room temperature) for same alloys. Source: Ref 67

condition for all test specimens. Some of the observations of the data are summarized:

Types 304 and 316 specimens were consumed completely at 1095, 1150, and 1200 °C (2000, 2100, and 2200 °F). Type 304 was found to be much more resistant than Type 316 at 980 °C (1800 °F). Type 446 was not as good as Type 310 at all test temperatures. Type 446 specimens were consumed at 1150 and 1205 °C (2100 and 2200 °F). Type 310 appeared to be slightly better than RA330 and 800H. For applications at high temperatures, many superalloys contain numerous alloying elements for increasing the elevated-temperature strength of the alloy. Molybdenum and tungsten are common alloying elements for providing solidsolution strengthening for increasing the creeprupture strength of the alloy. Two iron-base superalloys, Multimet alloy (Fe-20Ni-20Co21Cr-3Mo-2.5W-1.0Nb+Ta) and alloy 556

(Fe-20Ni-18Co-22Cr-3Mo-2.5W-0.6Ta-0.02La0.02Zr), are good examples. However, the oxides of both molybdenum and tungsten (MO3 and WO3) exhibit high vapor pressures at very high temperatures, as shown in Fig. 3.3. Multimet alloy suffered rapid oxidation attack at 1150 and 1200 °C (2100 and 2200 °F), with specimens completely consumed at both temperatures. However, formation of the volatile oxides of MO3 and WO3 can be minimized by modification of some key alloying elements in Multimet alloy. The development of alloy 556 was aimed at improving the oxidation resistance of Multimet alloy without losing the elevated-temperature strength by making some modification of alloying elements in Multimet alloy. The modification involved a slight increase in chromium, a decrease in cobalt, replacement of niobium with tantalum, and addition of a rare-earth element, lanthanum, and a reactive element, zirconium, but the same amounts of molybdenum and tungsten were kept. The result was a much more oxidation-resistant alloy, alloy 556, at 1095

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and 1150 °C (2000 and 2100 °F), although the alloy still suffered rapid oxidation at 1200 °C (2200 °F). As shown in the Table 3.12 nickel- and cobaltbase alloys containing molybdenum and/or tungsten were also found to suffer rapid oxidation at very high temperatures (i.e., specimens were consumed during the tests). Specimens of nickel-base alloys that were consumed at 1200 °C (2200 °F) were alloy S (14% Mo), alloy X (9% Mo, 0.6% W), and alloy 625 (9% Mo, 3.5% Nb). Some of those nickel-base alloys containing molybdenum and/or tungsten that were not consumed at 1200 °C (2200 °F) were alloy 230 (14% W), alloy 617 (9% Mo), and RA333 (3% Mo, 3% W). From these two different sets of oxidation behavior at very high temperatures, one can design nickel-base alloys (relying on chromium oxide scales) containing molybdenum and tungsten for elevatedtemperature strengthening to resist oxidation resistance at very high temperatures by adjusting other alloying elements.

For nickel-base alloys containing high levels of molybdenum and/or tungsten, it is believed that increasing chromium is probably the most important factor in suppressing rapid oxidation involving molybdenum and/or tungsten. Some nickel-base precipitation-strengthened alloys containing high titanium as well as molybdenum that were consumed at 1200 °C (2200 °F) were Waspaloy (4.3% Mo, 3.0 Ti), René 41 (10% Mo, 3.0 Ti), and alloy 263 (6% Mo, 2.2Ti). Titanium was found to be very active in oxide-scale formation. Figure 3.50 illustrates the oxide scale formed on alloy 263 after exposure to air for 1 h at 1200 °C (2200 °F), showing mainly titanium-rich oxides and Cr-Ti oxides. The formation of titanium-rich oxides apparently disrupts the Cr2O3 scale. Nagai et al. (Ref 69) found that titanium was detrimental to the oxidation resistance of Ni-20Cr alloy. In the Fe-CrAl system, however, the addition of 1% Ti to Fe18Cr-6Al was found to improve resistance in cyclic oxidation in air at 950 °C (1740 °F) (Ref 70). It is not clear whether the beneficial

Table 3.12 Results of oxidation tests for various alloys at indicated temperatures in flowing air (30 cm/min) for 1008 h 980 °C (1800 °F)

Alloy

Metal loss, mm (mils)

1095 °C (2000 °F)

Average metal affected, mm (mils)


1150 °C (2100 °F)



1205 °C (2200 °F)




214

0.0025

(0.1) 0.005

(0.2)

0.0025

(0.1)

0.0025

(0.1)

0.005

(0.2)

0.0075

(0.3)

0.005

(0.2)

0.018

(0.7)

601 600

0.013 0.0075

(0.5) 0.033 (0.3) 0.023

(1.3) (0.9)

0.03 0.028

(1.2) (1.1)

0.067 0.041

(2.6) (1.6)

0.061 0.043

(2.4) (1.7)

0.135 0.074

(5.3) (2.9)

0.11 0.13

(4.4) (5.1)

0.19 0.21

(7.5) (8.9)

230 S 617 333 X 671

0.0075 0.005 0.0075 0.0075 0.0075 0.0229

(0.3) (0.2) (0.3) (0.3) (0.3) (0.9)

0.018 0.013 0.033 0.025 0.023 0.043

(0.7) (0.5) (1.3) (1.0) (0.9) (1.7)

0.013 0.01 0.015 0.025 0.038 0.038

(0.5) (0.4) (0.6) (1.0) (1.5) (1.5)

0.033 0.033 0.046 0.058 0.069 0.061

(1.3) (1.3) (1.8) (2.3) (2.7) (2.4)

0.058 0.025 0.028 0.05 0.11 0.066

(2.3) (1.0) (1.1) (2.0) (4.5) (2.6)

0.086 0.043 0.086 0.1 0.147 0.099

(3.4) 0.11 (1.7) >0.8I (3.4) 0.27 (4.0) 0.18 (5.8) >0.9 (3.9) 0.086

(4.5) 0.20 (31.7) >0.8I (10.6) 0.32 (7.1) 0.45 (35.4) >0.9 (3.4) 0.42

(7.9) (31.7) (12.5) (17.7) (35.4) (16.4)

625 Waspaloy R-4I 263

0.0075 0.0152

(0.3) 0.018 (0.6) 0.079

(0.7) (3.1)

0.084 0.036

(3.3) (1.4)

0.12 0.14

(4.8) (5.4)

0.41 0.079

(16.0) (3.1)

0.46 0.33

(18.2) >1.2 (13.0) >0.40

(47.6) >1.2 (15.9) >0.40

(47.6) (15.9)

0.0178 0.0178

(0.7) 0.122 (0.7) 0.145

(4.8) (5.7)

0.086 0.089

(3.4) (3.5)

0.30 0.36

(11.6) (14.2)

0.21 0.18

(8.2) (6.9)

0.44 0.41

(17.4) >0.73 (16.1) >0.91

(28.6) >0.73 (35.7) >0.91

(28.6) (35.7)

188 25 150 6B

0.005 0.01 0.01 0.01

(0.2) (0.4) (0.4) (0.4)

0.015 0.018 0.025 0.025

(0.6) (0.7) (1.0) (1.0)

0.01 0.23 0.058 0.35

(0.4) (9.2) (2.3) (13.7)

0.033 0.26 0.097 0.39

(1.3) 0.18 (10.2) 0.43 (3.8) >0.68 (15.2) >0.94

(7.2) 0.2 (16.8) 0.49 (26.8) >0.68 (36.9) >0.94

(8.0) (19.2) (26.8) (36.9)

(21.7) (37.9) (46.1) (36.8)

(21.7) (37.9) (46.1) (36.8)

556 Multimet

0.01 0.01

(0.4) 0.028 (0.4) 0.033

(1.1) (1.3)

0.025 0.226

(1.0) (8.9)

0.067 0.29

(2.6) 0.24 (11.6) >1.2

(9.3) 0.29 (47.2) >1.2

(11.6) >3.8 (47.2) >3.7

800H RA330 310 316 304 446

0.023 0.01 0.01 0.315 0.14 0.033

(0.9) (0.4) (0.4) (12.4) (5.5) (1.3)

0.046 (1.8) 0.14 0.11 (4.3) 0.02 0.028 (1.1) 0.025 0.36 (14.3) >1.7 0.21 (8.1) >0.69 0.058 (2.3) 0.33

(5.4) 0.19 (0.8) 0.17 (1.0) 0.058 (68.4) >1.7 (27.1) >0.69 (13.1) 0.37

(7.4) 0.19 (7.5) 0.23 (6.7) 0.041 (1.6) 0.22 (2.3) 0.075 (3.0) 0.11 (68.4) >2.7 (105.0) >2.7 (27.1) >0.6 (23.6) >0.6 (14.5) >0.55 (21.7) >0.55

>0.55 >0.96 >1.I7 >0.94

>0.55 >0.96 >1.I7 >0.94

(150.0) >3.8 (146.4) >3.7

(8.9) 0.29 (11.3) 0.35 (8.7) 0.096 (3.8) 0.21 (4.4) 0.2 (8.0) 0.26 (105.0) >3.57 (140.4) >3.57 (23.6) >1.7 (68.0) >1.73 (21.7) >0.59 (23.3) >0.59

(150.0) (146.4) (13.6) (8.3) (10.3) (140.4) (68.0) (23.3)

Note: 3304 cm3/min of flow rate in a 1.75 in. diam furnace tube. The moisture was removed from the air by a filter prior to entering into the furnace tube. Specimens were cathodically descaled for measurement of the metal loss. The average metal affected is the sum of the metal loss and the depth of internal attack. The depth of internal attack was measured by metallography. Source: Ref 68


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effect of titanium for oxidation resistance is only for alumina formers such as in this case (Fe18Cr-6Al), but not for chromia formers in Ni20Cr alloy. Niobium is another alloying element that may be detrimental to alloy oxidation resistance at very high temperatures. The relatively poor oxidation resistance of alloy 625 at 1095 and 1150 °C (2000 and 2100 °F) can be attributed to niobium. Cobalt-base alloys with tungsten, such as alloy 188 (Co-22Cr-22Ni-14W-0.04La), alloy 25 (Co-20Cr-10Ni-15W), and alloy 6B (Co-30Cr4.5W-1.2C), suffered rapid oxidation at 1205 °C (2200 °F). A cobalt-base alloy, alloy 150 (Co27Cr-18Fe), containing no tungsten also suffered rapid oxidation attack at 1205 °C (2200 °F). Again, the oxidation of a cobalt-base alloy can be significantly improved with some modification of alloying elements. Alloy 25 with 15% W exhibits excellent creep-rupture strengths at high temperatures. However, because of the high level of tungsten, the alloy suffers high oxidation rates at very high temperatures, such as 1095 and 1150 °C (2000 and 2100 °F). With slight increase in chromium and nickel along with the addition of lanthanum, the result of the modification was alloy 188. As shown in Table 3.12, alloy 188 exhibits significantly better oxidation resistance than alloy 25 at 1095 and 1150 °C (2000 and 2100 °F). The best alloy among those investigated was an alumina former, alloy 214 (Ni-16Cr-4.5Al-Y). The depth of oxidation attack (average metal affected) was found to be less than 0.025 mm (1.0 mils) after 1008 h at temperatures up to

Fig. 3.50 Scanning electron micrograph showing the early stage of oxidation in air at 1200 °C (2200 °F) for 1 hour for alloy 263, revealing titanium-rich and Cr-Ti oxides on the outermost oxide scale. Area 1: 28.1% Cr, 70.9% Ti, 0.8% Co, 0.2% Ni. Area 2: 57.0% Cr, 36.8% Ti, 2.8% Co, 25% Ni, 1.0% Fe.

1205 °C (2200 °F). The oxidation resistance of alloy 214 is also presented in Fig. 3.42 to 3.45 and 3.47. The oxidation data presented in Table 3.12 were generated with a weekly cycle (168 h). When the cyclic frequency was increased to 25 h cycles, oxidation rates were increased for all the alloys tested. However, some alloys are more sensitive to cyclic oxidation than others. The effect of thermal cycling on the oxidation resistance of various alloys in air at 1095 °C (2000 °F) is illustrated in Table 3.13, which compares once-a-week (168 h) cycle data (Ref 68) with 25 h cycle data (Ref 71). All the data are presented in terms of the average depth of metal affected, which represents the metal loss plus the depth of internal oxidation attack. Both sets of the data were generated under the same test conditions using the same test furnaces and test procedures except the differences in cyclic frequencies. Some chromia formers, such as alloys 230, S, 188, 556, and 310, showed good resistance to thermal cycling. A long-term oxidation test program was undertaken to test alloys up to 2 years at 980, 1095, and 1150 °C (1800, 2000, and 2100 °F)

Table 3.13 Comparative oxidation resistance of various alloys in flowing air between 168 h and 25 h cycles at 1095 °C (2000 °F) Total depth of attack, mm (mils)

Extrapolated oxidation rate, mm/yr (mpy)

Alloy

1008 h/168 h

1050 h/25 h

168 h cycles

25 h cycles

214 601 600 671 230 S G-30 617 RA333 625 Waspaloy 263 188 25 150 6B 556 Multimet 800H RA330 310 446

0.003 (0.1) 0.066 (2.6) 0.041 (1.6) 0.061 (2.4) 0.003 (1.3) 0.003 (1.3) 0.122 (4.8) 0.046 (1.8) 0.058 (2.3) 0.122 (4.8) 0.137 (5.4) 0.361 (14.2) 0.003 (1.3) 0.259 (10.2) 0.097 (3.8) 0.394(15.5) 0.066 (2.6) 0.295 (11.6) 0.188 (7.4) 0.170 (6.7) 0.058 (2.3) 0.368 (14.5)

0.025 (1.0) 0.297 (11.7) 0.185 (7.3) 0.584 (23.0) 0.086 (3.4) 0.061 (2.4) 0.203 (8.0) 0.267 (10.5) 0.130 (5.1) 0.414 (16.3) 0.414 (16.3) 0.478 (18.8) 0.058 (2.3) 0.490 (19.3) 0.353 (13.9) >0.800 (31.5) 0.117 (4.6) 0.381 (15.0) 0.406 (16.0) 0.442 (17.4) 0.112 (4.4) 0.655 (25.8)

0.025 (1) 0.58 (23) 0.36 (14) 0.53 (21) 0.28 (11) 0.28 (11) 1.07 (42) 0.41 (16) 0.51 (20) 1.07 (42) 1.19 (47) 3.12 (123) 0.28 (11) 2.26 (89) 0.84 (33) 3.43 (135) 0.58 (23) 2.57 (101) 1.63 (64) 1.47 (58) 0.51 (20) 3.20 (126)

0.20 (8) 2.49 (98) 1.55 (61) 4.88 (192) 0.71 (28) 0.51 (20) 1.70 (67) 2.24 (88) 1.09 (43) 3.45 (136) 3.45 (136) 3.99 (157) 0.48 (19) 4.09 (161) 2.95 (116) >6.68 (263) 0.97 (38) 3.18 (125) 3.40 (134) 3.68 (145) 0.94 (37) 5.46 (215)

Note: 3304 cm3/min of flow rate in a 1.75 in. diam furnace tube. The moisture was removed from the air by a filter prior to entering into the furnace tube. Specimens were cathodically descaled for measurement of the metal loss. The total depth of attack is the sum of the metal loss and the depth of internal attack. Metal loss was measured by cathodically descaling the oxide scale prior to measurement of the specimen thickness. Source: Ref 71


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(Ref 49). Test specimens, which were cut from plate products, had dimensions of 1.27 cm (thick) by 2.54 cm (width) by 2.54 cm (length). Tests were conducted in a box furnace with still air; specimens were removed from the furnace every 30 days to allow air cooling. Weights were then measured to determine the time to initiation of breakaway oxidation. For block specimens (e.g., 1.25 cm, or 0.5 in., thick specimens), oxide scales were cracking, breaking, and spalling with “bing” noises as soon as the specimens were removed from the hot zone in the furnace. Cracking noises would not stop until almost all the oxide scales were broken and spalled off. Thus, no cathodic descaling was necessary to remove the oxide scales for measurement of the specimen thickness at the end of the test. When similar oxidation testing was performed with thin test coupons (about 3.2 mm, (0.125 in.) or

200

Weight gain, mg/cm2

0 –200 –400 –600 –800 –1000

0 30 60 90 120 150 180 210 240 270 300 330 360 Days

Fig. 3.51

The oxidation behavior of alloy 800H tested in still air at 1095 °C (2000 °F) involving a thick, blocky specimen (1.25 cm, or 0.5 in., thick) cycling to room temperature every 30 days for weight measurement, showing the alloy was under protective scales initially for about 30 days and then suffered breakaway oxidation. Courtesy of Haynes International, Inc.

thinner), no “cracking” noises were heard during specimen cooling from the hot zone, and oxide scales remained on the specimen surface for these thin test coupons. Figure 3.51 shows weight change data for alloy 800H in this longterm testing at 1095 °C (2000 °F) using thick, blocky specimens, showing breakaway oxidation with linear weight loss after 60 days of exposure. Some alloys showed no breakaway oxidation even after 2 years of testing. Figure 3.35 shows no breakaway oxidation for HR120 after 2 years of testing at 980 °C (1800 °F), while 800H and RA85H suffered breakaway oxidation. HR160 and 601 showed no breakaway oxidation after one year of testing at 1095 °C (2000 °F), while alloy 800H suffered breakaway oxidation (Fig. 3.46). At the end of testing for 2 years (720 days) at 980 °C (1800 °F) and 1 year (360 days) at 1093, 1150, and 1200 °C (2000, 2100, and 2200 °F), specimens were cut, mounted, and polished for metallographic determination of the depth of internal attack. Metal loss was determined by subtracting the original specimen thickness from the thickness after testing. The data generated from blocky specimens are summarized in Tables 3.14 to 3.16. The annual oxidation rates in terms of the total depth of oxidation attack are included in Tables 3.14 to 3.16 at 980, 1090, and 1150 °F (1800, 2000, and 2100 °F), respectively. These oxidation rate values would be considered to be quite reasonable, since the test duration was almost 2 years for 980 °C (1800 °F) and about 1 year for 1090 and 1150 °C testing. At 980 °C (1800 °F), alloys 230, 617, HR120, 556, and HR160 exhibited oxidation rates of less than 10 mpy. At 1090 °C (2000 °F), only alloy 230 exhibited about 10 mpy of oxidation rate, while other alloys tested exhibited more than 20 mpy. At 1150 °C (2100 °F), all alloys tested exhibited more than 30 mpy of

Table 3.14 Oxidation of several high temperature alloys in still air at 980 °C (1800 °F) for 720 days with specimens cycling to room temperature every 30 days Alloy

230 617 HR120 556 HR160 601 RA85H 800HT



Total depth of attack, mm (mils)

Oxidation rate, mm (mpy)

−1.4 1.0 −33.7 −19.8 −51.2 −9.9 −122.2 −417.8

0.00254 (0.1) 0 0.04060 (1.6) 0.02286 (0.9) 0.0635 (2.5) 0.0127 (0.5) 0.16002 (6.3) 0.52578 (20.7)

0.14732 (5.8) 0.23876 (9.4) 0.30988 (12.2) 0.38608 (15.2) 0.42418 (16.7) 0.56896 (22.4) 1.36398 (53.7) 2.02692 (79.8)

0.0508 (2) 0.127 (5) 0.1524 (6) 0.2032 (8) 0.2286 (9) 0.2794 (11) 0.6858 (27) 1.0414 (41)

Note: Total depth of attack = metal loss + internal attack. Source: Ref 49


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Table 3.15 Oxidation of several high-temperature alloys in still air at 1095 °C (2000 °F) for 360 days with specimens cycling to room temperature every 30 days Alloy

230 556 RA330 HR160 HR120 601 800HT RA85H




Oxidation rate, mm/yr (mpy)

−42.1 −298.1 −405.0 −73.1 −665.7 −110.5 −893.9 −348.1

0.04826 (1.9) 0.36322 (14.3) 0.50800 (20.0) 0.09144 (3.6) 0.82804 (32.6) 0.13716 (5.4) 1.12522 (44.3) 0.45466 (17.9)

0.27178 (10.7) 0.53848 (21.2) 0.60706 (23.9) 0.73660 (29.0) 0.96520 (38.0) 1.14554 (45.1) 1.29540 (51.0) 2.03962 (80.3)

0.27940 (11) 0.55880 (22) 0.60960 (24) 0.73660 (29) 0.99060 (39) 1.16840 (46) 1.32080 (52) 2.05740 (81)


Table 3.16 Oxidation of several high-temperature alloys in still air at 1150 °C (2100 °F) for 360 days with specimens cycling to room temperature every 30 days Alloy

230 617 HR120 HR160 800H 601 RA85H




Oxidation rate, mm/yr (mpy)

−249.7 −452.7 −894.2 −155.5 −1315.6 −258.7 −389.2

0.28194 (11.1) 0.54102 (21.3) 1.10998 (43.7) 0.19304 (7.6) 1.65608 (65.2) 0.32004 (12.6) 0.50800 (20.0)

0.83680 (34.0) 0.94488 (37.2) 1.34620 (53.0) 1.49098 (58.7) 1.78562 (70.3) 1.84912 (72.8) 2.40792 (94.8)

0.88900 (35) 0.96520 (38) 1.37160 (54) 1.52400 (60) 1.80340 (71) 1.87960 (74) 2.4384 (96)


oxidation rate. At 1200 °C (2200 °F), aluminaformer alloy 214 showed little or no oxidation attack (Fig. 3.47). These data are valuable in providing readers with the oxidation data in terms of the total depth of oxidation attack in air based on the actual measurements of the specimens after 1 to 2 years of testing. Since the data were generated from thick, blocky specimens, caution should be used when the data are being considered for application in thin-gage sheets or foils. This is related to the reservoir effect of a solute alloying element for the formation of a protective oxide scale. The discussion of the reservoir issue and the oxidation in thin foils is presented later. In this test program (Ref 49), in addition to the metal loss caused by formation of external oxide scales and internal attack caused by formation of internal oxides and/or voids, the weight-loss values were also determined. The weight loss of the specimens is related mostly to the metal loss resulting from the removal of the external oxide scales and is not significantly affected by the formation of internal oxides and/or voids. The weight-loss values of the alloys tested are plotted against their corresponding metal-loss values at 980, 1090, and 1150 °C (1800, 2000, and 2100 °F), revealing a nice straight line correlation (Fig. 3.52). Alloys tested were 230, 617, 601, 556, HR160, HR120 RA330, 800HT, and

RA85H. These alloys are primarily chromia formers. This correlation may be useful in making rough estimates of the metal loss for an alloy that showed only weight-loss data. The depth of oxidation attack was also investigated by John (Ref 15) for a wide variety of commercial alloys in isothermal air oxidation testing. Table 3.17 summarizes his data in terms of the temperature at which the oxidation rate reaches 10 mpy. Figure 3.53 illustrates some oxidation data in terms of oxide penetration as a function of test temperature in air after 1 year for some alloys (Ref 15). Table 3.18 shows the depth of oxidation attack of various heat-resistant alloys after cyclic oxidation tests at 1100 °C (2010 °F) in air + 5% H2O for 504 h with specimens cycling out of the furnace every 15 min (Ref 72). Lai et al. (Ref 73) reported the oxidation data generated from a field test inside a radiant tube fired with natural gas with an average temperature of 1010 °C (1850 °F) (Table 3.19). The test rack containing coupons of various alloys was exposed for about 3000 h. Many chromia formers, such as alloys 601, 230, 556, 310, 600, and RA330, were found to perform well, with extrapolated oxidation rates of less than 0.5 mm/ yr (20 mpy). Type 304, however, suffered severe oxidation attack with an extrapolated oxidation rate of more than 4.4 mm/yr (>175 mpy). The


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70 60

1.50

1.00

40 30 20

Metal loss, mm

Metal loss, mils

50

0.50

10 0 0

200

400

600

800

1000

1200

1400

Weight loss, mg/cm2 Correlation between weight loss (mg/cm2) and the depth of metal loss (mils) for commercial alloys that are primarily chromia formers tested in air at 980 °C (1800 °F)/720 days, 1095 °C (2000 °F)/360 days, and 1150 °C (2100 °F)/360 days. Alloys tested were 230, 617, 601, 556, HR160, HR120, RA330, 800HT, and RA85H. 1.0 mil = 0.0254 mm

Fig. 3.52

Table 3.17 The oxidation rate of total depth of attack was reached after 1 year in air Maximum temperature for 0.25 mm/yr (10 mpy), °C (°F)

Carbon steel Copper Nickel 9Cr-1Mo 410 304 617 803 625 800H 601GC DS 230 310 RA330 446 556 HR120 253MA 602CA MA956 214

… C11000 N02270 S50400 S41000 S30400 N06617 … N06625 N08810 … … N06230 S31000 S33000 S44600 R30556 … S30815 … S67956 N07214

604 (1120) 677 (1250) 782 (1440) 799 (1470) 832 (l530) 893 (1640) 938 (1720) 954 (1750) 960 (1760) 966 (1770) 977 (1790) 977 (1790) 982 (1800) 982 (l800) 999 (1830) 1010 (1850) 1010 (1850) 1010 (1850) 1082 (1980) 1121 (2050) >1150 (>2100) >1150 (>2100)

AISI 410

AISI 304 10

0.3 Alloy 617 Nickel AISI 310 Alloy 800 H

Penetration, mils

UNS No.

2.5

9Cr -1Mo Carbon steel

Penetration, mils

Alloy

102

1 0.03

0.1 1000

1200

1400

1600

1800

2000

Temperature, °F

Source: Ref 15

Fig. 3.53

alumina former, alloy 214, showed little or no oxidation attack with an extrapolated oxidation rate of about 0.076 mm/yr (3 mpy). In another field test (Ref 74), a test rack containing coupons of various alloys was placed in a natural-gas-fired furnace used for reheating ingots and slabs of nickel- and cobalt-base alloys. The test was conducted for about 113 days at temperatures varying from 1090 to 1230 °C (2000 to 2250 °F), with frequent cycles to 540 °C

Oxidation penetration (metal loss + internal attack) as a function of test temperature for 1 year in air for a variety of commercial alloys. Source: Ref 15

(1000 °F) during furnace idling. The results are summarized in Table 3.20. All the chromia formers tested suffered severe oxidation attack. The alumina former (alloy 214), however, exhibited little attack. Examination of the oxide scale formed on alloy 214 was found to consist of essentially aluminum-rich oxides.


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Table 3.18 Cyclic oxidation resistance of various heat-resistant alloys at 1100 °C (2010 °F) for 504 h in air-5H2O Specific weight change (descaled), mg/cm2 Alloy

ACI grade HK 310SS 800 601 617 X RA333 IN-814 188 MA-956

Mean

Range


Maximum attack, mm (mils)

−105.8 −149.0 −168.6 −11.0 −13.5 −20.0 −30.5 −3.5 −25.0 −1.0

−98 to −124 −92 to −235 −83 to −223 −6.3 to −17.2 −6.5 to −17.5 −10.0 to −29.5 … −2.5 to −4.9 −13.0 to −40.5 −0.3 to −1.5

0.25 (9.9) 0.31 (12.2) 0.39 (15.4) 1.5 (60)(b)

0.076 (3) 0.23 (9) 0.30 (12) 0.30 (12) 0.30 (12) 0.46 (18) 0.46 (18) 0.89 (35) 1.5 (58) >4.4 (175)

(a) Metal loss + maximum internal penetration. (b) Sample was consumed. Source: Ref 73

Table 3.20 Results of field test in a natural-gas-fired furnace for reheating nickeland cobalt-base alloy ingots and slabs for 113 days at 1090 to 1230 °C (2000 to 2250 °F) with frequent cycles to 540 °C (1000 °F) Alloy

214 RA330 601 600 800H 310SS 304SS 316SS 446SS


Maximum metal affected(a), mm (mils)

0.013 (0.5) 0.39 (15.5) 0.18 (7.2) 0.64 (25.0) >0.79 (31.0)(b) >1.0 (41.0)(b) >1.5 (60.0)(b) >1.6 (63.0)(b) >0.61 (24.0)(b)

0.11 (4.5) 0.65 (25.5) 0.95 (37.2) 1.1 (45.0) >0.79 (31.0)(b) >1.0 (41.0)(b) >1.5 (60.0)(b) >1.6 (63.0)(b) >0.61 (24.0)(b)

(a) Metal loss + internal penetration. (b) Samples were consumed. Source: Ref 74

3.4.10 Oxide-Dispersion-Strengthened (ODS) Alloys Oxide-Dispersion Strengthened alloys use very fine oxide particles that are uniformly distributed throughout the matrix to provide excessive strengthening at very high temperatures. These oxide particles, typically yttrium oxide, do not react with the alloy matrix so no coarsening or dissolution occurs during the exposure to very high temperatures, thus maintaining the strengthening of the alloy. This group of superalloys is produced using specialty powders that are manufactured by the mechanical alloying process. These powders are essentially composite powders with each particle containing a uniform distribution of submicron oxide particles in an alloy matrix. The process of producing these ODS powders involves repeated fracturing and rewelding of a mixture of powder particles in vertical attritors or horizontal ball mills (Ref 75). Alloy powders are then canned, degassed, and hot extruded, followed by hot working and

annealing to produce a textured microstructure (Ref 75). Alloys are available in mill products such as bar, plate, sheet, and so forth, or custom forgings. Some ODS alloys are shown in Table 3.21 (Ref 75). The oxidation behavior of some of these ODS alloys tested in air containing 5% H2O at 1200 °C (2190 °F) is shown in Fig. 3.54 (Ref 75). The oxidation behavior of MA956 compared with those of several ironand nickel-base alloys at 1100 °C (2010 °F) is shown in Fig. 3.55 (Ref 76). Additional oxidation data for some ODS alloys is presented in Section 3.4.12. 3.4.11 Effect of Oxygen Concentration on Oxidation Air atmosphere consists primarily of oxygen and nitrogen with some water vapor and small amounts of inert gases, such as argon, neon, and helium. Dry air consists of essentially 21% O2


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Table 3.21

Nominal chemical compositions of several ODS alloys

Alloy

Ni

Fe

Cr

Al

Ti

W

Mo

Ta

Y2O3

B

Zr

MA754 MA758 MA760 MA6000 MA956

bal bal bal bal …

… … … … bal

20 30 20 15 20

0.3 0.3 6.0 4.5 4.5

0.5 0.5 … 2.5 0.5

… … 3.5 4.0 …

… … 2.0 2.0 …

… … … 2.0 …

0.6 0.6 0.95 1.1 0.5

… … 0.01 0.01 …

… … 0.15 0.15 …

Composition, wt%

Note: All alloys contain 0.05% C. Source: Ref 75

50 Mass change, mg/cm2

MA 760 MA 754

–50 –100

0 –0.5 –1.0 –1.5 –2.0

–150 MA 6000 –200 –250 0

10

20

30

40

50

–2.5 –3.0 –3.5 60

Mass change, Ib/in.2 × 10–3

0.5

MA 956 0

Exposure time, days

Fig. 3.54

Oxidation behavior of several ODS alloys in air containing 5% H2O at 1200 °C (2190 °F).

Source: Ref 75

and 78% N2 with about 1% inert gases. In combustion, the concentration of oxygen may vary. Also, in some processes, oxygen may be the only gaseous component to which the equipment is exposed. Thus, the effect of oxygen concentration on the oxidation behavior of alloys may need to be evaluated for some applications. John (Ref 15, 77) investigated the oxidation behavior of a wide variety of commercial alloys in N2-O2 mixtures with the concentration of oxygen varying from 1 to 100%. The data generated at 871 °C (1600 °F) are shown in Fig. 3.56 (Ref 77), and data generated at 927 °C (1700 °F) are shown in Fig. 3.57 (Ref 15). The data presented in Fig. 3.56 were based on the 1152 h testing, while the data in Fig. 3.57 were based on tests after 1 year. At 871 °C (1600 °F), Type 304 was the only alloy that showed significant increase in oxidation attack from about 0.25 mm/yr (10 mpy) in N2-21%O2 to close to 2.5 mm/yr (100 mpy) in 100% O2. All other alloys in the figure showed about 10 mpy or less of oxidation attack at three different levels of O2 concentrations (1%, 21%, and 100%). Figure 3.57 shows the oxidation behavior of a number of alloys at 927 °C (1700 °F). The alloys that were found to increase oxidation attack with increasing oxygen

Fig. 3.55

Cyclic oxidation resistance of ODS alloy MA956 compared with alloy 601, HK alloy, alloy 800, and Type 310. Source: Ref 76

concentration included 9Cr-1Mo steel, 410, 304, and 617, while carbon steel, nickel, 800H, and 310 were relatively unaffected by oxygen concentrations. In Fig. 3.57, Type 304 was found to exhibit approximately 0.25 mm/yr (10 mils) of attack at 927 °C (1700 °F) after 1 year in 100% O2, while close to 2.5 mm/yr (100 mpy) of attack was extrapolated based on 1152 h exposure at 871 °C (1600 °F) in 100% O2, as shown in Fig. 3.56. Extrapolation from short-term tests here showing a higher oxidation rate at lower temperature could be an issue here. More longterm tests are needed. It is also of practical interest to perform long-term tests to evaluate the oxidation behavior in 100% O2 environments for some alloys that are to be used in chemical processes involving 100% O2. 3.4.12 High-Velocity Combustion Gas Streams Oxidation of alloys can significantly increase under high-velocity gas streams. Combustors and transition ducts in the gas turbine are subject to such conditions. These components are also


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103

AISI 304 AISI 310

102

956MA

Penetration rate, mpy

10

253MA HR120

1

800HT AISI 446

10–1

617 10–2

230 214

10–3

601GC 10–4 10–2

1

10–1 Partial pressure of O2, atm

Fig. 3.56

602CA 556

Effect of oxygen concentration in the N2-O2 mixture on the oxidation penetration (metal loss + internal attack) at 871 °C (1600 °F) for 1152 h. 1.0 mil = 0.025 mm. Source: Ref 77

subject to severe thermal cycling, particularly for gas turbines in airplane engines. Laboratory burner rigs have been developed to evaluate this type of oxidation, often referred to as “dynamic oxidation,” under the condition of very high gas velocities. Some of these dynamic oxidation burner rigs are described elsewhere (Ref 78–83). Lai (Ref 82) investigated a wide range of alloys—from stainless steels to superalloys—in a burner rig that generated a combustion gas stream with 0.3 Mach (100 m/s) velocity. The specimens were held in a carousel-type holder rotating at 30 rpm with respect to the combustion gas stream. Every 30 min the carousel was withdrawn from the hot zone, and quenched to less than 260 °C (500 °F) by a blast of cold air, and then automatically reinserted back into the hot zone. The specimens were subject to severe thermal cycling. Combustion was generated using No. 2 fuel oil with an air-to-fuel ratio of 50 to 1, producing a high-velocity (0.3 Mach, or 100 m/s) test gas. The tests were conducted at 1090 °C (2000 °F) with 30 min cycles, and the results are tabulated in Table 3.22. At 1090 °C (2000 °F) with a high-velocity gas stream plus severe thermal cycling, most alloys suffered significant metal loss, which constituted a large portion of the total depth of oxidation attack. With protection by an aluminum oxide scale, alloy 214 suffered very little

attack. The alloy showed no sign of breakaway oxidation after 500 h. Figure 3.58 shows the oxide scale formed after testing for 500 h (1000 cycles) (Ref 82). The scale consisted of aluminum-rich oxides. After 1000 h (2000 cycles) of testing, the scale remained aluminumrich. The maximum metal affected (metal loss +maximum internal penetration) remained about the same after 1000 h compared to after 500 h (Ref 82). The test results generated at 980 °C (1800 °F) for 1000 h (2000 cycles) are shown in Table 3.23 (Ref 82). Unlike 1090 °C (2000 °F) testing (Table 3.22), testing at 980 °C (1800 °F) resulted in internal oxidation and nitridation in addition to metal loss. Internal nitridation penetrated deeper into metal interior than internal oxidation penetration. Table 3.23 included only internal oxidation penetration data (i.e., maximum metal affected = metal loss + internal oxidation penetration). Limited tests were conducted to determine the effect of thermal cycling by testing at 980 °C (1800 °F) for 1000 h with 30 min cycling and without thermal cycling in dynamic oxidation testing (Ref 82). As expected, thermal cycling primarily contributed metal loss portion of the oxidation attack. The results are summarized in Table 3.24 (Ref 82). Total attack presented in Table 3.24 was based on metal loss and internal oxidation penetration. Since internal

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10

3

Carbon steel

10

9Cr-1Mo

2

Penetration, mils

10

Penetration, mm

AISI 410 1 Nickel Alloy 800 H

10

AISI 310

10

AISI 304

–1

Alloy 617

–2

1 10

–3

10

–2

10

–1

1

10

pO , atm 2

Fig. 3.57

Effect of oxygen concentration in the N2-O2 mixture on the oxidation penetration (metal loss + internal attack) after 1 year at 927 °C (1700 °F) for various commercial alloys. 1.0 mils = 0.025 mm. Source: Ref 15

nitridation attack was found to penetrate deeper into the alloy than internal oxidation attack does for many alloys, the total depth of attack for many alloys under dynamic oxidation test conditions was more than that reported in Tables 3.23 and 3.24. The oxidation/nitridation behavior of various alloys under dynamic oxidation test conditions is discussed in Chapter 4 “Nitridation.” Hicks (Ref 83) performed dynamic oxidation tests with 170 m/s gas velocity at 1100 °C (2010 °F) with 30 min cycles for several wrought chromia-former superalloys and an ODS alumina-former (MA956). Alumina former MA956 was found to be considerably better than chromium formers, such as alloys 230, 86, 617, 188, and 263. His results are shown in Fig. 3.59.

MA956 along with some ODS alloys was tested by Lowell et al. (Ref 78) with 0.3 Mach gas velocity at 1100 °C (2010 °F) with 60 min cycles. ODS alloys tested included MA956 (Fe19Cr-4.4Al-0.6Y2O3), HDA8077 (Ni-16Cr4.2Al-1.6Y2O3), TD-NiCr (Ni-20Cr-2.2ThO2) and STCA264 (Ni-16Cr-4.5Al-1Co-1.5Y2O3). Also included in the test was physical vapor deposition (PVD) coating of Ni-15Cr-17Al-0.2Y on MAR-M-200 alloy (Ni-9Cr-10Co-12W-1Nb5Al-2Ti). Their results are shown in Fig. 3.60. MA956 and HDA8077 as well as PVD Ni-Cr-AlY coating were found to perform well. No explanation was offered in the paper for STCA264, which did not perform as well as HDA8077 although both alloys had similar chemical compositions.

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Table 3.22 Dynamic oxidation resistance of iron-, nickel-, and cobalt-base alloys in high-velocity combustion gas stream (0.3 mach velocity) with 30 min cycles at 1090 °C (2000 °F) for 500 h Alloy

214 230 RA333 188 556 X RA330 S 600 310 601 617 800H 625 Multimet



0.013 (0.5) 0.056 (2.2) 0.10 (4.0) 0.19 (7.5) 0.22 (8.7) 0.23 (9.0) 0.28 (10.9) 0.30 (11.8) 0.44 (17.2) 0.54 (21.2) 0.27 (10.7) 0.32 (12.4) 0.77 (30.5)(b) >0.79 (31.0)(c) 1.25 (49.1)(d)

0.046 (1.8) 0.15 (5.7) 0.22 (8.7) 0.27 (10.7) 0.30 (11.7) 0.34 (13.5) 0.35 (13.6) 0.39 (15.2) 0.53 (20.7) 0.61 (24.1) 0.61 (24.0) 0.61 (24.0) 0.86 (34.0)(b) >0.79 (31.0)(c) 1.42 (55.8)(d)

Note: Gas velocity was 0.3 mach (100 m/s, or 225 mph); samples were cycled to less than 260 °C (500 °F) once every 30 min; 50 to 1 air-to-fuel ratio; two parts No. 1 fuel oil and one part No. 2 fuel oil. Internal nitridation occurred in some alloys, but is not included in the current data. See section 4.3.3 in Chapter 4. (a) Metal loss + maximum internal, penetration. (b) Extrapolated from 400 h; sample was about to be consumed after 400 h. (c) Sample was consumed in 500 h. (d) Extrapolated from 225 h; sample was about to be consumed after 225 h. Source: Ref 82

Table 3.23 Dynamic oxidation resistance of iron-, nickel-, and cobalt-base alloys in high-velocity combustion gas stream at 980 °C (1800 °F) for 1000 h Alloy

214 230 188 556 X S RA333 625 617 RA330 Multimet 800H 310 600 601 304 316



0.010 (0.4) 0.020 (0.8) 0.028 (1.1) 0.043 (1.7) 0.069 (2.7) 0.079 (3.1) 0.064 (2.5) 0.12 (4.9) 0.069 (2.7) 0.20 (7.8) 0.30 (11.8) 0.31 (12.3) 0.35 (13.7) 0.31 (12.3)(b) 0.076 (3.0) >9.0 (354)(c) >9.0 (354)(c)

0.031 (1.2) 0.089 (3.5) 0.107 (4.2) 0.158 (6.2) 0.163 (6.4) 0.17 (6.6) 0.18 (7.0) 0.19 (7.6) 0.27 (10.7) 0.30 (11.8) 0.38 (14.8) 0.39 (15.3) 0.42 (16.5) 0.45 (17.8)(b) 0.51 (20.0) >9.0 (354)(c) >9.0 (354)(c)

Note: Gas velocity was 0.3 mach (100 m/s, or 225 mph); samples were cycled to less than 260 °C (500 °F) once every 30 min; 50 to 1 air-to-fuel ratio; two parts No. 1 fuel oil and one part No. 2 fuel oil. Internal nitridation occurred in some alloys, but is not included in the current data. See section 4.3.3 in Chapter 4. (a) Metal loss + maximum internal penetration. (b) Extrapolated from 917 h; sample was about to be consumed after 917 h. (c) Extrapolated from 65 h; sample was consumed in 65 h. Source: Ref 82

3.4.13 Breakaway Oxidation In Fe-Cr, Fe-Ni-Cr, Ni-Cr, and Co-Cr alloy systems, the formation of an external Cr2O3 oxide scale provides the oxidation resistance for

Fig. 3.58

Oxide scales formed on alloy 214 in a high-velocity gas stream (0.3 Mach velocity) with 30 min cycles at 1090 °C (2000 °F) for 500 h. Area 1: 96.5% Al, 1.5% Cr, 0.1% Fe, 1.9% Ni. Area 2: 75.2% Al, 6.2% Cr, 2.6% Fe, 16.0% Ni. Area 3: 95.8% Al, 1.0% Cr, 0.1% Fe, 3.1% Ni. Area 4: 53.0% Al, 2.8% Cr, 9.2% Fe, 35.0% Ni. Source: Ref 82

Table 3.24 Effect of thermal cycling in dynamic oxidation behavior of several nickel-base alloys at 980 °C (1800 °F) for 1000 h No thermal cycling

Thermal cycling

Alloy

Metal loss, mm

Total attack(a), mm

Metal loss, mm

Total attack(a) mm

230 617 X 263

0.04 0.03 0.03 0.07

0.11 0.16 0.12 0.21

0.07 0.17 0.16 0.32

0.16 0.24 0.23 0.42

(a) Metal loss + internal oxidation. Source: Ref 82

the alloy. The growth of the Cr2O3 oxide scale follows a parabolic rate law as the exposure time increases. As the temperature increases, the oxide scale grows faster. The growth of the Cr2O3 oxide scale requires a continuous supply


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0

Total depth of oxidation, mm/side

MA 956

230

0.1

617 0.2 86 188 (1971 Cast)

0.3

X

263

188 (1980 Cast) 0.4 (C) 0

200

400

600

800

Number of cycles

Fig. 3.59

Dynamic oxidation resistance of several wrought superalloys including MA956 alloy in highvelocity combustion gas stream (170 m/s) at 1100 °C (2010 °F) with 30 min cycles. Source: Ref 83

Specific weight change, mg/cm2

10

Coated MAR-M-200

0 MA 956 –10

HDA8077

–20 –30 –40

STCA-264

–50 0

400 800 1200 1600 2000 2400 2800 3200 Cycles

Fig. 3.60

Dynamic oxidation tests at 1100 °C (2010 °F) in a Mach 0.3 gas stream with each cycle consisting of 1 h at temperature followed by quenching to ambient temperature for 3 min. Source: Ref 78

of chromium from the alloy interior diffusing to the oxide/metal interface. Continued oxidation can eventually deplete chromium in the alloy matrix immediately underneath the oxide scale. When the chromium concentration in the alloy matrix immediately beneath the oxide scale is reduced to below a critical concentration, the alloy matrix no longer has adequate chromium to reform a protective Cr2O3 oxide scale when the scale cracks or spalls due to oxide growth stresses

or thermal cycling. Once this occurs, fastgrowing, nonprotective iron oxides, or nickel oxides or cobalt oxides (i.e., oxides of base metal) form and grow on the alloy surface. Breakaway oxidation, thus, initiates, and the alloy begins to undergo oxidation at a rapid rate. This is illustrated in Fig. 3.51. The alloy thus requires the level of chromium immediately underneath the chromium oxide scale to have a critical level to allow the chromium oxide scale to reheal. Gleeson (Ref 84) presented air cyclic oxidation data for three chromia formers tested at 982 °C (1800 °F) for up to 360 days. Also presented were the corresponding chromium concentration analyzed by EDX on the surface of the metal when the oxide scale was spalled off from the test specimen. The data are presented in Fig. 3.61. In this test program, thick, blocky specimens (13 mm thick, 25 mm wide, and 25 mm long) instead of typically thin coupons were used. The specimens were cycled to room temperature once every 30 days for weight measurement. Oxide scales were found to completely spall off while the specimens were removed from the furnace for cooling to room temperature with “popping” noises being heard during cooling (Ref 85). Figure 3.61 shows that alloy 230 exhibited very little weight loss with no evidence of breakaway oxidation after 360 days at 982 °C (1800 °F). The chromium concentration of the alloy immediately underneath the spalled oxide scale was found to remain at about 16% with no sign of decreasing with increasing exposure time. For HR120 alloy, the weight loss data also revealed no evidence of breakaway oxidation up to 360 days of exposure. The corresponding chromium concentration of the alloy on the surface underneath the spalled oxide scale remained approximately about 18 to 20% up to 240 days, and then dropped to about 13% after 360 days. Alloy 800HT, on the other hand, showed breakaway oxidation after 180 days. The corresponding chromium concentration of the alloy on the surface immediately underneath of the spalled oxide scale after 180 days was found to be about 10%. With continuing oxidation, alloy 800HT suffered linear weight loss and continued the decrease in chromium concentration to 8% when the exposure reached to 360 days. However, when alloy 800HT was oxidized after 90 days showing no sign of weight loss, the chromium concentration of the alloy underneath the oxide scale was about 11%. The data appear to suggest

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that the critical chromium concentration of the alloy underneath the oxide scale is approximately 11% for alloy 800HT. When the chromium concentration underneath the oxide scale is below 11%, the reformation of a protective chromium-rich oxide scale is not possible, thus resulting in breakaway oxidation. In the same oxidation test program at 1150 °C (2100 °F), similar analysis on the chromium concentration profile underneath the oxide scale was performed for alloys 230, HR160 and HR120 with the data presented in Fig. 3.62 and 3.63 (Ref 86). Figure 3.62 shows the weight-loss data for three alloys up to 360 days. The chromium concentration profile from the surface immediately underneath the

230 Weight change, mg/cm2

0.0 HR120

–100.0

–200.0 800HT

–300.0 0

100

(a)

200 Time, days

300

400

spalled oxide scale to the alloy interior after 360 days of testing was determined using EDX analysis with the results shown in Fig. 3.63. The chromium concentration immediately underneath the spalled oxide scale for both alloys 230 and HR120 was well below 10%, while that of alloy HR160 was about 10%. Alloy HR120 showed sign of breakaway oxidation after 90 days of exposure. Continuing oxidation testing resulted in a linear weight-loss rate. Alloy 230 showed signs of breakaway oxidation after 240 days. HR160, however, showed a linear weight loss up to 360 days, although suffering the least weight-loss rate, with no clear sign of breakaway oxidation. Oxidation of alloy HR160 is involved the formation of Cr2O3 and SiO2. This may explain that HR160, although continuing to lose weight, still showed no sign of breakaway oxidation after 360 days at 1150 °C (2100 °F). In a long-term oxidation study of Fe-20Cr-25Ni alloy in CO2 containing 1% CO, 300 ppm H2O, and 300 ppm CH4 at 1023 to 1173 K (750 to 900 °C), Evans et al. (Ref 87) found the critical chromium level for rehealing of chromium oxide scales to be about 16%. The minimum level of chromium needed to maintain a protective chromium oxide scale to prevent breakaway oxidation may vary from environment to environment and from alloy to alloy. To prolong the time for the initiation of breakaway oxidation, it is necessary to have an adequate reservoir for chromium immediately below the oxide scale to provide adequate chromium to maintain a protective chromium oxide

28.0


Surface Cr concentration, wt%

0

HR160

24.0 HR120 20.0 230 16.0

12.0

–250.0 230 –500.0 HR120 –750.0

800HT 8.0 (b)

0

100

200

300

400

–1000.0 0

Time, days

60

120

180

240

300

360

Time, days

Fig. 3.61

Weight changes as a function of exposure time in long-term cyclic oxidation tests in air at 982 °C (1800 °F) for alloys 800HT, HR120, and 230 (a), and the corresponding changes in the surface chromium concentration (measured after the scale was spalled off) as a function of exposure time (b). Source: Ref 84

Fig. 3.62

Weight changes as a function of exposure time for alloys 230, HR160, and HR120 in air oxidation tests at 1150 °C (2100 °F) with thermal cycling to room temperature for weight measurement once every 30 days. Source: Ref 86


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80.0 70.0 Ni Concentration, wt%

60.0 50.0 40.0 30.0 Cr 20.0 W

10.0 0 0

200

400

600

800

1000

1200

Distance from surface, µm

(a) 50.0

Ni

Concentration, wt%

40.0

Co

30.0

20.0

Cr

10.0

0 0

400

800

1200

1600

2000


(b) 60.0

Concentration, wt%

50.0 Ni

40.0

30.0

Fe

20.0

Cr

10.0

0 0 (c)

Fig. 3.63

200

400

600

800

1000


Concentration profiles for (a) alloy 230, (b) HR160, and (c) HR120 after air oxidation tests for 360 days at 1150 °C (2100 °F), as shown in Fig. 3.62. Source: Ref 86

scale or to reheal the oxide scale that suffered local cracking or failure. Brady et al. (Ref 88) proposed that chromium carbides might provide a reservoir of chromium for maintaining the growth and rehealing of chromium oxide scales. These authors observed that an as-cast Fe-15Cr0.5C formed a nonprotective iron-rich oxide scale when exposed in O2 at 850 °C (1562 °F). When the alloy was forged at 1150 °C (2100 °F) to produce a uniformly distributed fine carbide phase, the forged alloy showed a thin protective Cr2O3 scale under the same test condition (Ref 88). During oxidation testing of this forged sample, the fine chromium-rich carbides are dissolved into the underlying alloy substrate to continue supplying chromium to maintain and reheal the chromium oxide scales (Ref 88). For alumina formers, such as Fe-Cr-Al alloys, and Fe-Cr-Al-base and Ni-Cr-Al-base ODS alloys, breakaway oxidation occurs when aluminum concentration underneath the Al2O3 scale has reduced to a critical level such that healing of the Al2O3 is no longer possible, thus resulting in the formation of nonprotective, fast-growing oxides of base metals (e.g., iron oxides or nickel oxides). The breakaway oxidation due to rapid growth of iron oxides or nickel oxides becomes essentially a life-limiting factor. This critical aluminum concentration was found to be about 1.0 to 1.3% for Fe-Cr-Al-base ODS alloys (e.g., MA956, ODM751) at 1100 to 1200 °C (2012 to 2192 °F) (Ref 89, 90). These values were obtained from foil specimens (0.2 to 2 mm thick) tested in still air at 1100 to 1200 °C. For the nonODS Fe-20Cr-5Al alloy, this critical aluminum concentration was found to be higher (about 2.5%) at 1200 °C (Ref 89). Since the breakaway oxidation is related to aluminum reservoir in the alloy, and the aluminum reservoir becomes a critical issue when the component is made of thin sheet or foil. Because of excellent oxidation resistance at very high temperatures, there is increasing interest in looking at alumina formers for products that require thin foils, such as honeycomb seals in gas turbines, metallic substrates for automobile catalyst converters, and recuperators in microturbines. The oxidation behavior of several commercial alumina formers in thin foils is summarized in Section 3.4.14. For alumina formers to improve their resistance to breakaway oxidation, yttrium is frequently used to increase the adhesion of the aluminum oxide scale. Other alloying elements that are known to increase the adhesion of the aluminum oxide scale include zirconium and

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hafnium. Quadakkers (Ref 91) shows that both MA956 (Fe-20Cr-4.5Al-0.5Y2O3) and Aluchrom (Fe-20Cr-5Al-0.01Y) exhibited much more cyclic oxidation resistance than Fe-20Cr5Al when tested at 1100 °C in synthetic air with a hourly cycle to room temperature (Fig. 3.64). Addition of Y2O3 to an alumina former has a similar beneficial effect as yttrium alloying element. Klower and Li (Ref 92) studied the oxidation resistance of Fe-20Cr-5Al alloys in 10 different compositions containing various amounts of yttrium ranging from 0.045 to 0.28%. All 10 compositions contained 0.002% S, and eight compositions contained 0.04 to 0.06% Zr with two compositions containing no zirconium. Cyclic oxidation tests were performed at 1100 and 1200 °C (2012 and 2192 °F), respectively, with each cycle consisting of 96 h at temperature and rapid air cooling to room temperature. These authors concluded that the yttrium addition of about 0.045% was sufficient to prevent the oxide scales from spalling and when the yttrium concentration was increased to more than 0.08%, substantial internal oxidation could occur, resulting in rapid metal wastage, as shown in Fig. 3.65 (Ref 92). Sulfur in the alloy is known to play a very significant role in the adhesion of the aluminum oxide scale to the alloy substrate for alumina formers. The role of yttrium is believed to prevent the preferential segregation of sulfur in the alloy to the scale/metal interface to weaken the adhesion of the oxide scale (Ref 93–95). Reducing the concentration of sulfur in a Ni-Cr-Al alloy can significantly improve the oxidation resistance of the alloy. Smeggil (Ref 96) compared cyclic oxidation resistance between the

normal purity Ni-Cr-Al alloys (approximately 30 to 40 ppm S) with the high-purity Ni-Cr-Al alloys (approximately 1 to 2 ppm S), showing a significant improvement in cyclic oxidation resistance when sulfur in the alloy was significantly reduced. This is illustrated in Fig. 3.66 (Ref 96). Also demonstrated in the figure is the beneficial effect of yttrium addition to the normal purity Ni-20Cr-12Al alloy, showing significant improvement in the cyclic oxidation resistance of the alloy without reducing the sulfur content in the alloy. Sulfur has been found to segregate to the oxide/alloy interface during oxidation in FeCr-Al alloys (Ref 97, 98). The role of yttrium is believed to tie up sulfur at the oxide/metal interface, thus improving the oxide-scale adhesion (Ref 96). 3.4.14 Thin Foils There are some industrial applications that require thin-gage sheet materials or thin foils for construction of some critical components. As the component thickness decreases, oxidation becomes a major limiting factor for its service life. When the component is made of thin foil, prolonging the incubation time before the initiation of breakaway oxidation is the controlling factor for extending the service life of the component. Thus, as applications are being pushed toward higher and higher temperatures, alloys that form aluminum oxide scales can offer tremendous advantages in performance over those alloys that form chromium oxide scales. In gas turbine applications, one important component made of a thin foil is a turbine seal ring assembly that controls the turbine tip

Depth of internal oxidation, µm


4 Fe-20Cr-AI MA956 Aluchrom

3 2 1 0 –1 0

200

400

600

800

1000

1200

800

1100 °C 1200 °C

600 400 200 0 0

Fig. 3.64

Cyclic oxidation resistance of MA956, Aluchrom, and Fe-20Cr-5Al tested in synthetic air at 1100 °C (2012 °F) with an hourly cycle (each cycle consisted of 56 min heating and 4 min cooling). Source: Ref 91

0.1

0.2

0.3

Yttrium, wt%

Time, h

Fig. 3.65

Maximum internal oxidation depth as a function of yttrium content in the alloys after 3000 h of cyclic oxidation tests at 1100 and 1200 °C (2012 and 2192 °F) in air. Source: Ref 92


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pg 54


Mass change/unit area, mg/cm2

0

High-purity NiCrAI –8 Normal-purity NiCrAIY –16

–24

Normal-purity NiCrAI

–32 0

20

40

60

80

100

Number of 1 h cycles

Fig. 3.66

Cyclic oxidation resistance of the normal purity Ni-20Cr-12Al (30 to 40 ppm S), the high-purity Ni-20Cr-12Al (1 to 2 ppm S) and the normal purity Ni-20Cr-12Al-Y at 1180 °C. Source: Ref 96

clearance for improving thermal efficiency. The seal ring assembly is typically constructed out of a honeycomb seal brazed onto a superalloy casing. The traditional alloys used for honeycomb seals are chromia formers, such as nickel-base alloy X. Lai (Ref 99) evaluated honeycomb samples made of wrought alloy 214 (Ni-16Cr-3Fe-4.5Al-Y) and alloy X (Ni-22Cr18.5Fe-9Mo) in dynamic burner rig testing that simulated gas turbine hot gas conditions. Tests were conducted at 954 °C (1750 °F) in a highvelocity combustion gas stream (0.3 Mach or 100 m/s) with cycling every 30 min. Both alloy X and alloy 214 honeycomb samples were made of 0.076 mm (3 mils) foils. The alloy X honeycomb sample was completely oxidized (destroyed) after 154 h of testing, while the alloy 214 honeycomb sample was unaffected after 317 h when the test was terminated. Figure 3.67 shows the condition of both samples after tests. Simms et al. (Ref 100) conducted extensive oxidation studies on several commercial foil materials with different thicknesses (from 0.05 to 0.127 mm) in a simulated combustion environment (nominally N2-14%O2-3%CO2-7.8%H2O) at 950 to 1250 °C (1742 to 2282 °F). Test specimens were cycled to room temperature every 100 h when tested at 950 and 1050 °C (1742 and 1922 °F), every 40 h at 1150 °C (2102 °F), every 20 h at 1250 °C (2280 °F). Each specimen was

contained in an individual alumina crucible so the spalled oxides could be included in the weight measurement. All the specimens were preoxidized in air at 1050 °C (1922 °F) for 1 h prior to oxidation testing. The alloys tested included Kanthal AF (Fe-20Cr-5Al-0.05Y0.08Zr wrought alloy), Aluchrom YHf (Fe-20Cr5.8Al-0.04Y-0.05Zr-0.36 (14)(c) 0.005 (0.2) 0.058 (2.3) 0.11 (4.5)

>0.38 (15)(b) >0.36 (14)(c) 0.0075 (0.3) 0.086 (3.4) 0.147 (5.8)

Weight gain, mg/cm2



8.8 61.5 1.3 81.2 169(d)

0.02 (0.8) >0.36 (14)(c) 0.005 (0.2) 0.11 (4.5) >0.9 (35.4)(e)

>0.38 (15)(b) >0.36 (14)(c) 0.018 (0.7) 0.20 (7.9) >0.9 (35.4)(e)

Note: Flowing air 30 cm/min (472 cm3 min in a 1.75 in. diam tube); cycled to room temperature once a week (168 h cycles). (a) Metal loss + average internal penetration. (b) Internal penetration through thickness. (c) Sample was completely oxidized. (d) Specimen weight gain after 504 h. (e) Extrapolated from 504 h; specimen was completely oxidized (consumed) in 504 h. Source: Ref 55

the only oxide formed on the sample; no Al2O3 was detected. Burner rig dynamic oxidation tests under a high-velocity combustion gas stream (0.3 Mach or 100 m/s) were also performed on IC-50 compared with commercial alloys (Ref 55). The nickel aluminide IC-50 suffered significantly more severe oxidation attack than nickel-base alloys 214 (alumina former) and 230 (chromia former) after testing at 1090 °C (2000 °F) for 500 h with 30 min cycles. IC-50 suffered more than 0.38 mm (15 mils) of oxidation attack, compared with 0.05 mm (1.8 mils) for alloy 214 and 0.14 (5.7 mils) for alloy 230. The data for other commercial alloys tested under the same conditions are shown in Table 3.22. The nickel aluminide was very susceptible to internal oxidation. Significant internal oxidation was observed after only 50 h of testing. SEM/EDX analysis of the scale formed on the

50 h tested sample revealed mainly nickel-rich oxides. 3.4.17 Catastrophic Oxidation As temperature increases, metals and alloys generally suffer increasingly higher rates of oxidation. When the temperature is excessively high, metals and alloys can suffer rapid oxidation. There is, however, another mode of rapid oxidation that takes place at relatively lower temperatures. This mode of rapid oxidation, which is often referred to as “catastrophic oxidation,” is associated with the formation of a liquid oxide. The liquid oxide disrupts and dissolves the protective oxide scale, causing the alloy to suffer rapid oxidation at relatively low temperatures. Leslie and Fontana (Ref 119) observed an unusually rapid oxidation for Fe25Ni-16Cr alloy containing 6% Mo when heated


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Original Sample Thickness

33.3 µm

Fig. 3.76

Nickel-rich oxides formed on nickel aluminide IC218 after 1008 h at 1150 °C (2100 °F) in air with 168 h cycles. Area 1: 11.7% Al, 80.0% Ni, 8.3% Cr. Area 2: 18.8% Al, 49.0% Ni, 32.2% Cr. Area 3: 22.7% Al, 57.1% Ni, 20.2% Cr. Source: Ref 55

200 µm

Fig. 3.75

Specimen cross section of IC-50 after oxidation testing in air at 1150 °C (2100 °F) for 1008 h with 168 h cycles. Samples were descaled prior to metallographic mounting. Source: Ref 55

in static air to 900 °C (1650 °F). The same alloy exhibited good oxidation resistance when heated to the same temperature in flowing air. They postulated that the rapid oxidation was due to the accumulation of gaseous MoO3 on the metal surface. The oxidation is accelerated by the thermal dissociation of MoO3 into MoO2 and O. Meijering and Rathenau (Ref 120), Brasunas and Grant (Ref 121), and Brennor (Ref 122), however, attributed this to the presence of a liquid oxide phase. The MoO3 oxide melts at about 795 °C (1463 °F). The 19Cr-9Ni steel suffered catastrophic oxidation attack in the presence of MoO3 at 770 °C (1420 °F). This temperature was very close to the eutectic temperature of MoO2-MoO3-Cr2O3, which was reported to be 772 °C (1420 °F) (Ref 123).

Other mixed oxides involving MoO3 that exhibit low melting points are MoO2-MoO3 (778 °C, or 1435 °F), MoO2-MoO3-NiO (764 °C, or 1400 °F), Fe2O3-MoO3 (730 °C, or 1345 °F), Fe2O3-Fe-MoO3 (725 °C, or 1335 °F), V2O5MoO3 (610 to 718 °C, or 1130 to 1330 °F), Na2O-MoO3 (499 °C, or 930 °F), and Cu2OMoO2-MoO3 (470 °C, or 880 °F) (Ref 123). Other oxides, such as PbO and V2O5, can also cause metals or alloys to suffer catastrophic oxidation in air at intermediate temperatures of 640 to 930 °C (1200 to 1700 °F) (Ref 124). PbO and V2O5 melt at 888 and 690 °C (1630 and 1270 °F), respectively. The deleterious effect of lead oxide was believed to be related to exhaust-valve failures in gasoline engines. Gasoline additives were a primary source for lead compounds. Vanadium is an important contaminant in residual or heavy fuel oils. Therefore, V2O5 plays a significant role in oil ash corrosion in oil-fired boilers, which is discussed in Chapter 11. Sawyer (Ref 124) indicated that accelerated oxidation of Type 446 stainless steel in the presence of lead oxide can proceed at temperatures where the liquid phase does not exist. Experiments carried out by Brasunas and Grant (Ref 125) showed that 16Cr-25Ni-6Mo alloy specimens placed adjacent to, but not in contact with, 0.5 g samples of WO3 oxides suffered accelerated oxidation attack when tested in air at 868 °C (1585 °F), which is well below the melting point of WO3 (i.e., 1473 °C, or 2683 °F).


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Because molybdenum and tungsten are very important solid-solution-strengthening alloying elements, many superalloys containing either or both have been developed since Leslie and Fontana first observed catastrophic oxidation in 1948. Some of these alloys, including alloys X, R-41, 625, 617, S, 230, 25, and 188, have been used successfully in service for the hot section of gas turbine engines. Many of them have also been used successfully in heat treating, chemical processing, and related industries. The most effective way to alleviate the potential catastrophic oxidation problem is to avoid a stagnant condition for the gaseous atmosphere.

3.5 Oxidation/Nitridation in Air and Combustion Atmospheres Nitridation can take place in conjunction with oxidation in oxidizing atmospheres including air and combustion environments. This can result in significant internal nitridation penetration and affect the mechanical properties of the alloy. This topic is covered in detail in Chapter 4 “Nitridation.”

3.6 Summary Oxidation data for carbon and low-alloy steels, ferritic stainless steels, austenitic stainless steels, Fe-Cr-Ni alloys with 20-25Cr/30-40Ni, Fe-Cr-Al alloys, and superalloys including ODS alloys are presented. The data are presented in such a way for readers to make comparisons between alloys within the same alloy group or between alloys from different alloy groups. The chapter focuses on long-term oxidation behavior. Oxidation resistance of some nickel aluminides is also compared with that of several commercial nickel-base superalloys. A large portion of commercial hightemperature alloys relies on chromium for forming a protective chromium oxide scale to resist oxidation attack. The temperature range in which the chromium oxide scale is effective in providing adequate oxidation resistance varies from 540 to 1090 °C (1000 to 2000 °F). This group of alloys is frequently referred to as chromia formers. An adequate supply of chromium from the alloy interior to the metal surface to form and maintain a continuous, protective chromium oxide scale is necessary for an alloy to maintain its oxidation-resistant capability under

the service condition. Once the chromium supply drops below the critical level to maintain and reheal the chromium oxide scale, breakaway oxidation is initiated followed by rapid growth of oxides of base metals, such as iron, nickel, or cobalt oxides. Detailed discussion on the effect of chromium on the oxidation behavior and breakaway oxidation of various alloys is presented. Effects of other minor elements on the oxidation behavior of chromia formers are also discussed. A relatively small number of commercial high-temperature alloys rely on aluminum for forming a protective aluminum oxide scale to resist oxidation attack at very high temperatures. This group of alloys that form aluminum oxide scales are typically referred to as alumina formers. The oxidation behavior of alumina formers that include Fe-Cr-Al, Ni-Cr-Al, and ODS alloys is presented. The effect of aluminum as well as yttrium and sulfur on the oxidation resistance of alumina formers is discussed. Discussion also includes oxidation under high-velocity gas streams, oxidation of thin foils, effect of surface depletion of chromium in austenitic stainless steels, effect of water vapor, and catastrophic oxidation (oxidation under molten oxides). Most oxidation data were generated at 980 to 1200 °C (1800 to 2200 °F). However, many industrial applications are in the temperature range of 650 to 980 °C (1200 to 1800 °F), which are below the test temperatures at which most data were generated. More long-term oxidation data need to be generated at 650 to 980 °C (1200 to 1800 °F) for stainless steels, Fe-Ni-Cr and some simple Ni-Cr alloys to provide a more reliable database at intended application temperatures.

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97. P.Y. Hou and J. Stringer, Oxide Scale Adhesion and Impurity Segregation at the Scale/Metal Interface, Oxid. Met., Vol 38 (No. 5/6), 1992, p 323 98. P.Y. Hou, Compositions at Al2O3/ FeCrAl Interfaces after High Temperature Oxidation, Mater. Corros., Vol 51, 2000, p 329 99. G.Y. Lai, “Several Modern Wrought Superalloys for Gas Turbine Applications,” Paper 96-TA-030, presented at ASME Turbo Asia ’96 (Jakarta, Indonesia), Nov 5–7, 1996 100. N.J. Simms, R. Newton, J.F. Norton, A. Encinas-Oropesa, J.E. Oakey, J.R. Nicholls, and J. Wilber, Mater. High Temp., Vol 20 (No. 3), 2003, p 439 101. J. Klower, Factors Affecting the Oxidation Behaviour of Thin Fe-Cr-Al Foils, Mater. Corros., Vol 49, 1998, p 758 102. P.J. Maziasz, B.A. Pint, R.W. Swindeman, K.L. More, and E. Lara-Curzio, “Advanced Stainless Steels and Alloys for High Temperature Recuperators,” DOE/CETC/ CANDRA Workshop on Microturbine Applications (Calgary, Alberta, Canada), Jan 21–23, 2003 103. B.A. Pint, “The Effect of Water Vapor on Cr Depletion in Advanced Recuperator Alloys,” GT2005-68495, ASME Turbo Expo 2005 (Reno-Tahoe, Nevada), June 6–9, 2005 104. B.A. Pint, Stainless Steels with Improved Oxidation Resistance for Recuperators, J. Eng. Gas Turbines Power, Vol 128, 2006, p 1 105. B.A. Pint, “The Effect of Water Vapor on Cr Depletion in Advanced Recuperator Alloys, GT2005-68495,” ASME Turbo Expo 2005, (Reno-Tahoe, Nevada), June 6–9, 2005 106. C.W. Tuck, M. Odgers, and K. Sachs, Scaling Rates of Pure Iron and Mild Steel in Oxygen, Steam, Carbon Dioxide in the Range 850°–1000 °C, Anti-Corrosion, June 1966, p 14 107. S.C. Stultz and J.B. Kitto, Ed., Steam and Its Generation and Use, 40th ed., Babcock & Wilcox, 1992, p 13-2 108. S.C. Stultz and J.B. Kitto, Ed., Steam and Its Generation and Use, 40th ed., Babcock & Wilcox, 1992, p T-17 109. K. Segerdahl, J.E. Svensson, and L.G. Johansson, The High Temperature Oxidation of 11% Chromium Steel: Part

110.

111.

112. 113.

114.

115.

116.

117. 118. 119.

120. 121. 122. 123.

124. 125.

I—Influence of pH2 O , Mater. Corros., Vol 53, 2002 p 247 K. Segerdahl, J.E. Svensson, and L.G. Johansson, The High Temperature Oxidation of 11% Chromium Steel: Part II— Influence of Flow Rate, Mater. Corros., Vol 53, 2002, p 479 B.A. Pint, R. Peraldi, and P.F. Tortorelli, The Effect of Alloy Composition on the Performance of Stainless Steels in Exhaust Gas Environments, Paper No. 03499, Corrosion/2003, NACE International, 2003 B.A. Pint, private communication, 2007 R.L. McCarron and J.W. Schultz, in Proc. Symp. High Temperature Gas-Metal Reactions in Mixed Environments, AIME, 1973, p 360 C.C. Clark and W.R. Hulsizer, Superalloys Development for Gas Turbines Operating in the Marine Environment, Conf. Proc., Gas Turbine Materials Conference, Naval Ship Engineering Center, 1972, p 35 C. Sarioglu et al., The Adhesion of Alumina Films to Metallic Alloys and Coatings, Mater. Corros., Vol 51, 2000, p 358 K. Onal, M.C. Maris-Sida, G.H. Meier, and F.S. Pettit, Water Vapor Effects on the Cyclic Oxidation Resistance of Alumina Forming Alloys, Mater. High Temp., Vol 20 (No. 3), 2003, p 327 G. Welsch and P.D. Desai, Ed., Oxidation and Corrosion of Intermetallic Alloys, Purdue University, 1996 K. Natesan, Oxid. Met., Vol 30 (No. 1/2), 1988, p 53 W.C. Leslie and M.C. Fontana, Paper No. 26, 30th Annual Convention of ASM (Philadelphia, PA), Oct 25–29, 1948 J.K. Meijering and G.W. Rathenau, Nature, Vol 165, Feb 11, 1950, p 240 A.D. Brasunas and N.J. Grant, Iron Age, Aug 17, 1950, p 85 S.S. Brennor, J. Electrochem. Soc., Vol 102 (No. 1), Jan 1955, p 16 J.H. DeVan, “Catastrophic Oxidation of High Temperature Alloys,” ORNL-TM-51, Oak Ridge National Laboratory, Oak Ridge, TNn, Nov 10, 1961 J.W. Sawyer, Trans. TMS-AIME, Vol 221, 1961, p 63 A. de S. Brasunas and N.J. Grant, Trans. ASM, Vol 44, 1950, p 1133

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CHAPTER 4

Nitridation 4.1 Introduction In air or combustion atmospheres containing nitrogen, nitridation can take place under certain exposure conditions. In most cases, oxidation dominates the high-temperature corrosion reaction. However, nitridation can take place for some alloys when oxide scales no longer provide protection. The alloys that are particularly susceptible to oxidation/nitridation attack are those containing strong nitride formers, such as titanium and aluminum. Many high-temperature nickel-base alloys containing both aluminum and titanium are strengthened by γ′ phase, Ni3(Al,Ti). For these alloys, nitridation by forming internal nitrides of aluminum and titanium can deplete the surface layer with aluminum and titanium, thus weakening the alloy. Under a high-velocity combustion gas stream with severe thermal cycling, similar to the conditions in “flying” gas turbines (aircraft engines), nitridation can be particularly severe in oxidation/nitridation attack. In nitrogen-base atmospheres, such as N2 or N2-H2, metals and alloys can also suffer nitridation attack. This type of atmosphere is often used as a protective atmosphere in heat treating and sintering operations. Molecular nitrogen can be severely nitriding for many metals and alloys, particularly when temperatures are sufficiently high. Ammonia (NH3) is a commonly used nitriding gas for case hardening at temperatures from 500 to 590 °C (925 to 1100 °F) (Ref 1). Furnace equipment and components repeatedly subjected to these service conditions frequently suffer brittle failures as a result of nitridation attack. Carbonitriding is another important method of case hardening that produces a surface layer of both nitrides and carbides. The process is typically carried out at 700 to 900 °C (1300 to 1650 ° F) in ammonia, with additions of carbonaceous gases, such as CH4 (Ref 2). Thus, the heat treat retort, fixtures, and other furnace equipment are subject to both nitridation and carburization.

Cracked ammonia (i.e., ammonia that is completely dissociated into H2 and N2) provides an economical protective atmosphere for processing metals and alloys. Many bright annealing operations for stainless steels use a protective atmosphere consisting of N2 and H2, generated by dissociation of ammonia. With three parts H2 and one part N2 produced in cracked ammonia, nitridation is less critical for the heat treating equipment. In the chemical processing industry, nitriding environments are generated by processes employed for production of ammonia, nitric acid, melamine, and nylon 6-6 (Ref 3, 4). Ammonia is produced by reacting nitrogen with hydrogen over a catalyst at temperatures of typically 500 to 550 °C (930 to 1020 °F) and pressures of 200 and 400 atm. Commercial processes for ammonia synthesis are discussed in detail in Ref 5. The converter, where the ammonia synthesis reaction takes place, may suffer nitridation attack. Brittle failure of the welds for the waste heat boiler of an ammonia plant has been reported by Van der Horst (Ref 6). These tube-to-tube sheet welds were high-nickel alloy 182, which suffered nitridation attack. Production of nitric acid involves the oxidation of ammonia over a platinum gauze catalyst at temperatures of about 900 °C (1650 °F) (Ref 5). The catalyst grid support structure and other processing components in contact with ammonia may also be susceptible to nitridation attack. Figure 4.1 shows the nitrided structure of a nickel-base alloy catalyst grid support after two years of service in a nitric acid plant.

4.2 Thermodynamic Considerations When metal is exposed to nitrogen gas at elevated temperatures, nitridation proceeds according to: 1=2N2 (gas)=[N] (dissolved in metal)

ð4:1Þ


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[%N] = k(pN2 )1=2

ð4:2Þ

where k is the equilibrium constant and pN2 is the partial pressure of N2 in the atmosphere. In nitrogen-base atmospheres, the nitriding potential is proportional to (pN2 )1=2 . Increasing the nitrogen partial pressure (or nitrogen concentration) increases the thermodynamic potential for nitridation. Molecular nitrogen is less aggressive than ammonia in terms of nitridation of metals. However, when metal is heated to excessively high temperatures (e.g., 1000 °C (1830 °F) or higher), nitridation by molecular nitrogen can become a serious material issue. Under oxidation/nitridation conditions, nitrogen molecules permeate through cracks and pores and reach the metal underneath the oxide scales when the oxide scale is no longer protective. Nitridation then proceeds by dissociation of

100 µm

nitrogen molecules and absorption of nitrogen atoms by the metal following Reaction 4.1. Nitridation often takes place in the metal at the vicinity of cracks developed under creep conditions in air or N2-containing combustion atmospheres. In this case, oxides are often associated with the crack under creep deformation in air. Oxidation consumes the oxygen molecules from air, which penetrates into the crack, thus depleting oxygen and increasing nitrogen potential (or concentration) around the crack. As a result, nitridation takes place in the vicinity of the creep crack. When the environment is ammonia (NH3) or contains NH3, metals or alloys may undergo rapid nitridation reactions. It is precisely for this reason that NH3 is frequently used for case hardening. Ammonia is metastable and dissociates into molecular N2 and molecular H2 when heated to elevated temperatures. Once NH3 is completely dissociated into N2 and H2, the nitriding potential is defined by Reaction 4.3. To increase nitrogen absorption by steel, molecular NH3 should be allowed to dissociate on the steel surface, thus allowing dissociated atomic nitrogen to be dissolved at the metal surface (Ref 7, 8). Thus, to increase nitridation reactions, it is necessary to bring as much fresh, uncracked NH3 as possible in contact with the surface of the metal to minimize the production of molecular nitrogen. At temperatures below 600 °C (1110 °F) and at high gas flow rates, the production of nitrogen is minimized and the nitrogen solubility at the surface of iron is determined by (Ref 8):

(a)

100 µm (b)

Fig. 4.1

Alloy X (Ni-22Cr-18.5Fe-9Mo-0.6W) catalyst grid support structure bar after 2 years of service in a nitric acid plant. (a) Internal nitride precipitates (about 20 mils in depth) containing mainly chromium-rich nitrides along with some carbides formed during thermal aging. (b) Microstructure in the unaffected interior containing mainly carbides due to thermal aging at the service temperature

NH3 $ 3=2 H2 +[N] (dissolved in Fe)

ð4:3Þ

[%N]=k(pNH3 =(pH2 ))3=2

ð4:4Þ

where k is the equilibrium constant, and pNH3 , and pH2 are partial pressures of NH3 and H2, respectively. The nitriding potential is proportional to pNH3 =( pH2 )3=2 . Increasing ammonia partial pressure (or concentration) in the atmosphere increases the thermodynamic potential for nitridation. When nitrogen in the metal exceeds its solubility limit, nitrides will then precipitate out. The nitrides of important alloying elements for engineering alloys are tabulated in Table 4.1 (Ref 9). For iron, nickel, and cobalt, three important alloy bases for high-temperature alloys, only iron forms stable nitrides. No nitrides of nickel and cobalt have been reported. The relative stabilities among various nitrides can


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Chapter 4: Nitridation / 69

alloys with nickel ranging from 0 to 35% (Ref 19). In Fe-18Cr-Ni-N system, increasing nickel reduces the solubility of nitrogen, as shown in Fig. 4.3 (Ref 11). It is also shown in Table 4.2 that the solubility of nitrogen in stainless steels increased with increasing temperature. Alloys with higher nitrogen solubilities generally exhibit less resistance to nitridation attack. e 4N

+10

2F

N H

3

0

2

N o2

–10

2M

2 2C

2 1/

–40

Iron Chromium Molybdenum Tungsten Aluminum Titanium Niobium Tantalum Zirconium Hafnium Silicon Vanadium Boron Manganese Magnesium

Fe4N CrN MoN WN AlN TiN NbN TaN ZrN HfN Si3N4 VN BN Mn4N Mg3N

bN

–60 2

a

3N 2

–70

VN N2

C

AI

N 2

2

Ta

–100 N

2

–120

Ti

N

2

–130

Zr

2M + N2 = 2MN

–140 Nb4N3 Ta3N5

0

500

1000

Hf4N3

Fig. 4.2 Mn3N

102

N2 partial pressure, bar

Nickel, %

20

γ + Cr2N

γ 10 α+γ 0.1

0.2

0.3

1000 °C (1830 °F)

γ + CrN

1 γ + Cr2N

0.1 10 –2

γ

γ+π

10 –3 10 –4 0

α + γ + Cr2N

2000

Standard free energy of formation for selected nitrides. Source: Ref 10

10

0

1500

Temperature, °C

Source: Ref 9

0 α

BN

g3

N

–90

V 2N Mn2N

2

M

–110

Hf3N2

N

2

Fe2N Cr2N Mo2N W 2N Ti2N Nb2N Ta2N

N4 Si 3

–50

Table 4.1 Nitrides of important alloying elements for engineering alloys Nitrides

Cr

r 2N

–30

–80

Element

N

–20

∆G°, Kcal

be compared in terms of their free energies of formation, as illustrated in Fig. 4.2 (Ref 10). Physical-chemical properties of some of these nitrides can be found in Ref 9. The types of nitrides that are likely to form in the alloy can be predicted by examining the phase-stability diagram. Figure 4.3 shows a phase-stability diagram of Fe-18Cr-Ni-N system at 900 °C (1650 °F), indicating phase regions of Cr2N in γ (or α+γ) phase as a function of nickel and nitrogen contents (Ref 11). Nitride phases formed in alloys are also dependent on nitrogen partial pressure ( pN2 ), as shown in Fig. 4.4 for Ni-Cr-N system at 1000 °C (1830 °F) (Ref 12). Nitrogen solubility in the alloy is important in affecting the nitridation resistance of the alloy. Table 4.2 summarizes some nitrogen solubility data for iron, stainless steels, and nickel alloy (Ref 13–18). Iron and stainless steels exhibit significantly higher nitrogen solubility than a nickel-base alloy (Ni-20Fe). Nickel has been found to decrease the nitrogen solubility in Fe-Ni

0.4

γ+α 10

20

30

40

50

60

Cr concentration, wt %

Nitrogen, %

Fig. 4.3

Phase stability diagram for Fe-18Cr-Ni-N system at 900 °C (1650 °F). Source: Ref 11

Fig. 4.4

Phase stability diagram for the Ni-Cr-N system as a function of N2 partial pressure at 1000 °C (1830 °F). Source: Ref 12


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Table 4.2

Nitrogen solubility in metals and alloys

Nitrogen, wt%

0.06 0.26 max 0.02 (at pN2 : 1 atm) 0.125 0.177 0.190 0.258 0.281 0.18 0.18 0.21 0.26 0.26 0.0001 (at pN2 : 1 atm)

Metal or alloy

α-Fe γFe Fe-10%Ni Type 304 Type 304 Type 304 Type 304 Type 304 Fe-18Cr-12Ni-2Ti Fe-18Cr-12Ni-2Ti Fe-18Cr-12Ni-2Ti Fe-18Cr-12Ni-2Ti Fe-18Cr-12Ni-2Ti Ni-20%Fe


502 (936) γ region of Fe-C 1000 (1832) 538 (1000) 593 (1100) 927 (1700) 954 (1749) 981 (1800) 985 (1805) 1040 (1905) 1093 (2000) 1150 (2100) 1210 (2210) 1000 (1832)

Ref

13 14 15 16 17 17 17 17 18 18 18 18 18 15

4.3 Internal Nitridation in Oxidizing Environments In air or oxidizing combustion environments, oxidation usually dominates high-temperature corrosion reactions. However, under certain conditions, alloys can suffer internal nitridation attack along with oxidation. Internal nitridation attack, when it occurs, can penetrate farther into the metal interior than oxidation, thus significantly affecting the creep-rupture behavior of the alloy by accelerating the creep crack growth. Discussion of internal nitridation under no external stresses and under creep conditions is presented in sections 4.3.1 and 4.3.2. 4.3.1 Internal Nitridation in Air under No External Stress Some high-temperature alloys that contain strong nitride formers such as aluminum and titanium can suffer internal nitridation even in air environments. In an air oxidation study for two nickel-base alloys, IN939 (Ni-22Cr-20Co-3.8Ti1.4Al-2W-1Nb-1.3Ta) and IN738LC (Ni-16Cr9Co-3.5Ti-3.3Al-1.8Mo-1Nb-1.8Ta) at 700, 900, and 1100 °C (1290, 1650, and 2010 °F). Litz et al. (Ref 20) observed internal titanium nitrides (needle shape) formed in front of internal aluminum oxides that formed underneath the external oxide scales. In an oxidation study of alloy 800HT (Fe-21Cr-32Ni-0.5Al-0.5Ti) in air at 980 °C (1800 °F) for about 2 years (720 days), Harper et al. (Ref 21) observed Widmanstätten acicular chromium-rich nitrides along with aluminum nitrides that formed below chromiumrich oxides. Lai (Ref 22) observed internal aluminum nitrides (needle shape) that formed

underneath the external oxide scale and internal oxides in alloy 601 exposed to a furnace oxidizing atmosphere for about 4 to 5 years at temperatures probably between 760 and 870 °C (1400 and 1600 °F), as shown in Fig. 4.5. Severe thermal cycling that causes cracking and spalling of oxide scales can also result in severe internal nitridation. Han and Young (Ref 23) conducted cyclic oxidation tests by heating the specimens to 1100 °C (2010 °F) in still air for 1 h followed by cooling to room temperature for 15 min then repeating the cycle again for 260 cycles. The alloys investigated were Ni-24 to 38%Cr-14 to 25%Al. The specimens suffered severe oxide scale spallation. The internal nitridation attack was found to be extensive, and the nitridation zone consisted of AlN beneath Cr2O3 and Al2O3, then AlN+Cr2N, and then AlN in the deepest region (Ref 23). Douglas (Ref 24) indicated that the diffusivity of nitrogen appears to be two orders of magnitude greater than that of oxygen in nickel or nickel alloys. Table 4.3 summarizes the diffusion coefficients of nitrogen in nickel and iron alloys compared with those of oxygen and carbon in nickel, based on the diffusivity data from Rubly and Douglas (Ref 25, 26), Grabke and Peterson (Ref 27), Park and Alstetter (Ref 28), and Gruzin et al. (Ref 29). The diffusivity of nitrogen is also on the same order of magnitude as that of carbon as shown in Table 4.3. It is thus not surprising to find internal nitrides were advancing in front of internal oxides. 4.3.2 Internal Nitridation at Creep Cracks in Air Environment During creep testing in air, extensive internal nitridation can develop in the vicinity of cracks. Brickner et al. (Ref 30) found that types 302, 304, and 310 stainless steels showed significant nitridation after creep-rupture testing in air at 870 °C (1600 °F) in less than 1000 h. Acicular nitrides (believed to be chromium nitrides) in a Widmanstätten pattern were found to form extensively in the vicinity of microcracks, as shown in Fig. 4.6 (Ref 30). Extensive nitridation was confirmed by the chemical analysis of the tested specimens for nitrogen, which showed the nitrogen content was increased from about 0.058% before testing to 0.30 to 0.53% after creep-rupture testing (Ref 30). Extensive internal nitrides were also observed in the vicinity of creep cracks in alloy 253MA (Fe-21Cr-11Ni)

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(a) 0.1 mm

1 2

(c) 10 µm

(b) 0.0010 in Formation of internal aluminum nitrides beneath external oxide scales and internal oxides in alloy 601 after exposing to a furnace oxidizing atmosphere for approximately 4 to 5 years in a temperature range of 760 to 870 °C (1400 to 1600 °F). (a) Optical micrograph showing the external oxide scales and the internal oxides, and then the chromium denuded zone immediately below, followed by internal nitrides underneath the denuded zone. (b) Optical micrograph at higher magnification showing internal nitrides. (c) SEM (backscattered electron image) showing internal aluminum nitrides and the EDX analysis of nitrides. Results of the semiquantitative EDX analysis (at.%) on internal aluminum nitrides are summarized as:

Fig. 4.5

Phase 1 Phase 2

Table 4.3

41.5% Al, 24.7% Ni, 10.6% Cr, 6.8% Fe, 5.5% Ti, and 10.0% N 58.0% Al, 13.1% Ni, 6.1% Cr, 3.6% Fe, 0.8% Ti, and 17.3% N

Diffusion coefficients of nitrogen, oxygen, and carbon. Diffusion coefficient, cm2/s


700 (1290) 800 (1470) 900 (1650)

N in Ni-Alloys(a)

N in Fe-20Ni(b)

O in Ni(c)

C in Ni(d)

9.5 × 10−9 to 2.3 × 10−8 3.2 × 10−8 to 8.5 × 10−8 1.4 × 10−7 to 4.0 × 10−7

1.17 × 10−8 3.86 × 10−8 1.47 × 10−7

7.4 × 10−11 5.05 × 10−10 2.3 × 10−9

3.19 × 10−9 1.47 × 10−8 0.55 × 10−7

(a) Ref 25, 26. (b) Ref 27. (c) Ref 28. (d) Ref 29

after creep-rupture testing at 900 °C (1652 °F) for 11,800 h in air (Ref 31), in 800H after creep-rupture testing at 900 to 1000 °C (1650 to 1830 °F) in air (Ref 32), and in alloy 800H during the creep crack growth testing at 1000 °C (1830 °F) (Ref 33). The nitrides identified were Cr2N in 253MA (Ref 31), Cr2N (major) and CrN (minor) in 800H creep-ruptured specimens

(Ref 32), and Cr2N and AlN in 800H creep-crack growth specimens (Ref 33). Hoffman and Lai (Ref 34) investigated an alloy 800HT pigtail that suffered cracking after about 7.5 years of service in a hydrogen reformer. The pigtail section, which was exposed to air at approximately 850 °C (1565 °F) at the outside of the reformer furnace, was found to show


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extensive blocky precipitates along grain boundaries and acicular precipitates in the matrix in the vicinity of cracks at the tube outer-diameter side (exposed to air). Samples were cut from this pigtail section and solution-annealed in a furnace for 1 h at 1093, 1149, and 1204 °C (2000, 2100, and 2200 °F), respectively. Microstructural examination of these samples indicated that both blocky, grain-boundary phases and acicular phases in the matrix remained in the microstructure and were not put back into solution, suggesting those phases were nitrides instead of carbides. Also, chemical analysis of the samples from the pigtail indicated that carbon content remained about the same as that of the material before service (about 0.07%), while nitrogen content was about 0.27 wt% with a nominal nitrogen content of about 0.02% prior to service. The process gas in the tube contained essentially no nitrogen (typically about 0.01%). Thus, the nitrogen ingress into the tube was primarily from air from the outside diameter side of the pigtail. Using scanning electron microscopy with energy-dispersive x-ray spectroscopy (SEM/EDX) analysis, acicular phases were found to be enriched in aluminum, while blocky phases were enriched in chromium; the former was believed to be aluminum nitride and the latter chromium nitride. Figure 4.7 shows the acicular aluminum nitrides and blocky chromium nitrides that remained in the microstructure after solution annealing at 1150 °C (2100 °F) for 1 h for the sample from the straight section of the pigtail (Ref 34). Figure 4.8 shows extensive nitride formation in the vicinity of creep cracks in

the bend section of the pigtail from another hydrogen reformer (Ref 34). When creep cracks initially develop at the metal surface during creep testing in air, oxidation occurs at the crack surface including the crack tip. The oxide scales formed on the crack surface become nonprotective due to creep deformation, thus causing the oxygen potential to decrease significantly with concurrent increase in nitrogen potentials at the oxide/metal interface. As a result, nitrogen is absorbed by the metal and is diffused into the metal in the vicinity of cracks to form internal nitrides.

100 µm

Fig. 4.7

Acicular aluminum nitrides and blocky chromium nitrides, which formed in the vicinity of the creep cracks in alloy 800HT pigtail in a hydrogen reformer, were not dissolved into solution after the sample was resolution annealed at 1150 °C (2100 °F) for 1 h. Source: Ref 34

100 µm

Fig. 4.6 Acicular nitrides (believed to be chromium nitrides) in a Widmanstätten pattern formed in the vicinity of creep cracks in Type 302SS after creep-rupture testing at 870 °C (1600 °F) in less than 1000 h. Original magnification, 500×. Source: Ref 30

Fig. 4.8

Extensive aluminum and chromium nitrides formed in the vicinity of creep cracks in the bend section of an alloy 800H pigtail in another hydrogen reformer. Source: Ref 34


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4.3.3 Oxidation and Nitridation in Combustion Atmospheres High-temperature alloys that are exposed to a high-velocity, oxidizing combustion gas stream at high temperatures are susceptible to internal nitridation attack. In investigating transition duct component failures in a land-based gas turbine, Swaminathan and Lukezich (Ref 35) observed that alloy 617 (Ni-22Cr-12.5Co-9Mo-1.2Al) had suffered severe oxidation and nitridation attack from both the air side (outside diameter side of the transition duct) and the combustion side (inside diameter side of the duct) after service for slightly less than 2 years (14,000 h). Extensive internal nitridation from both air and combustion gas sides of alloy 617 transition duct is shown in Fig. 4.9 (Ref 35). Alloy 230 (Ni-22Cr-14W2Mo-0.3Al-La) was also tested for 16,000 h as a transition duct, suffering similar oxidation/nitridation attack (Ref 35). However, no aluminum nitrides were observed in alloy 230. Significant nitrogen pickup was observed from both transition ducts. Results of the chemical analyses of nitrogen from samples at the exit end of the transition duct and at the location far away from the exit for both alloy 617 and 230 transition ducts are shown in Table 4.4 (Ref 35). Lai (Ref 34) used a high-velocity dynamic burner rig test to simulate a gas turbine combustion environment. The simulated combustion gas stream was generated by burning fuel oil (a mixture of two parts No. 1 fuel and one part No. 2 fuel) with an air-to-fuel ratio of approximately 50 to 1 in a laboratory burner rig. Most of the air for combustion was from a compressor. When combusted with fuel oil, a

OD

ID

high-velocity combustion gas stream with about 0.3 Mach (100 m/s) was generated. Specimens were loaded in a carousel specimen holder that rotated at 30 rpm during testing to ensure all the specimens were subjected to the same test conditions. Furthermore, the specimens were subjected to severe thermal cycling once every 30 min by lowering the carousel from the test chamber followed by rapid fan-air cooling to below 260 °C (500 °F) for 2 min before returning the carousel back to the test chamber. A schematic of this dynamic burner rig is shown in Fig. 4.10. The combustion gas was determined to consist of 76% N2, 13% O2, 6% CO2, and 5% H2O. The test on alloy 617 produced severe internal nitridation, with the microstructure very similar to that observed by Swaminathan and Lukezich (Ref 35) from the transition duct in a land-based gas turbine power plant. Figure 4.11 shows the microstructure of an alloy 617 specimen after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. Extensive needle-shape aluminum nitrides were observed. Some blocky chromium nitrides were observed to form right below the external oxide scales. Aluminum nitrides were found to penetrate farther into the metal interior than chromium nitrides. The oxide scales were found to be porous and nonprotective. Alloy 230 was included in the test and found to show less nitridation attack under the same test condition. Nitridation in alloy 230 involved only the formation of internal chromium nitrides below internal chromium oxides and chromium denuded zone with no aluminum nitrides. Two other common combustor alloys, alloys X and 263, were also included in the test. Figures 4.12 and 4.13 show the microstructures of alloys X and 263, respectively, after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. Alloy X showed mainly internal chromium nitrides, while alloy 263 showed mainly tiny needle-shaped nitrides, presumably titanium

Table 4.4 Nitrogen contents at different locations of alloy 617 and alloy 230 transition ducts

Fig. 4.9

Cross section (2.5 mm, or 0.1 in.) of an alloy 617 (Ni22Cr-12.5Co-9Mo-1.2Al) transition duct after service for less than 2 years (about 14,000 h) in a land-based gas turbine, showing extensive formation of both aluminum and chromium nitrides from both air side (outside diameter of the transition duct) and the combustion gas side (inside diameter of the duct). Source: Ref 35

Transition duct/service

Location

Alloy 617/14,000 h

Exit Far away from exit Exit Far away from exit

Alloy 230/16,000 h Source: Ref 35

Nitrogen, wt%

0.24 0.004 0.22 0.05


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Thermocouple for steady-state control

Thermocouple for recording specimen temperature history

50 mm square insulated flame tunnel Specimen temperature measured by pyrometer

Compressed inlet air 425 °C (800 °F)

Fuel

Combustor Rotating shaft

Blower (thermal shock)

Thermal cycle

Fig. 4.10

The dynamic burner rig used by Lai (Ref 36) for simulating a gas turbine combustion environment in evaluating the oxidation/nitridation behavior of gas turbine combustor alloys. Courtesy of Haynes International, Inc.

50 µm

50 µm

Fig. 4.11

Some blocky chromium nitrides and extensive acicular aluminum nitrides formed in alloy 617 after testing in the dynamic burner rig at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. The combustion gas stream with about mach 0.3 (100 meter/s) consisted of 76% N2, 13% O2, 6% CO2, and 5% H2O. Source: Ref 36

nitrides. Similar to both alloy 617 and 230, alloys X and 263 also exhibited nonprotective, porous oxide scales.

Fig. 4.12

Extensive internal chromium nitrides formed in alloy X after testing in the dynamic burner rig at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. The combustion gas stream with about mach 0.3 (100 m/s) consisted of 76% N2, 13% O2, 6% CO2, and 5% H2O. Source: Ref 36

X-ray diffraction analysis of the oxide scales formed on alloys 230, 617, and X was performed with the results summarized in Table 4.5. The results showed that NiO and NiCr2O4 along with


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Cr2O3 oxides made up the external oxide scales for alloy 230, and NiO and Cr2O3 made up the oxide scales for alloys 617 and X. With the formation of NiO oxides, the alloys were no longer protected by Cr2O3 oxide scales. An electrolytic extraction technique was used to extract precipitate phases in the tested specimens for analysis, and the results are summarized in Table 4.5, revealing essentially nitride phases in all three alloys. Alloy 230 also showed M6C carbides, which were the carbides in the alloy in the as-solution-annealed condition. Bulk nitrogen

contents for the original samples (before testing) and those after testing are summarized in Table 4.6, showing significant nitrogen adsorption for alloys 263, 617, and X. Alloy 230 showed little nitrogen adsorption. Bulk carbon contents for the samples before and after testing were also determined, and the results clearly showed that carburization was not involved (Table 4.7). The overall test results in terms of weight loss (due to oxidation), metal loss (due to oxidation), internal oxidation, internal nitridation, and total depth of attack are summarized in Table 4.8. In continuing his testing program for the same simulated gas turbine environment involving the same four combustor alloys (i.e., 230, 617, 263, and X) at the same test temperature and duration Table 4.6 Results of bulk nitrogen analysis before and after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling Alloy

Original nitrogen, wt%

Nitrogen after testing, wt%

230 263 617 X

0.05 0.004 0.03 0.04

0.06 0.42 0.52 0.57

Source: Ref 36

50 µm

Fig. 4.13

Extensive internal nitrides (believed to be titanium nitrides) formed in alloy 263 after testing in the dynamic burner rig at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. The combustion gas stream with about mach 0.3 (100 m/s) consisted of 76% N2, 13% O2, 6% CO2, and 5% H2O. Source: Ref 36

Table 4.5 Results of x-ray diffraction analysis on oxide scales and extracted precipitate phases for alloys 230, 617, and X after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling Alloy

Surface oxide scales

230

NiCr2O4 (strong) Cr2O3 (medium) NiO (medium) NiO (strong) Cr2O3 (medium) NiO (strong) Cr2O3 (weak)

617 X Source: Ref 36

Table 4.7 Results of bulk carbon analysis before and after testing in the dynamic burner rig at 980 °C (1800 °F) for 1000 h with 30 minute thermal cycling Alloy

Original carbon, wt%

Carbon after testing, wt%

230 263 617 X

0.09 0.06 0.05 0.08

0.09 0.03 0.04 0.01

Source: Ref 36

Table 4.8 Test results in terms of weight loss, depth of oxidation penetration, depth of nitridation, and total depth of attack after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling Alloy


Metal loss, mm

Internal oxidation, mm

230 617 263 X

−6.8 −80.1 −219.4 −107.1

0.07 0.17 0.32 0.16

0.09 0.07 0.10 0.07

Extraction residues

M6C (strong) Cr2N (medium) AlN (strong) TiN (medium weak) CrN (medium strong) Cr2N (medium strong)

Internal nitridation, mm

0.17 >0.41(b) 0.29 >0.40(b)

Total attack(a), mm

0.24 >0.58(b) 0.61 >0.56(b)

Note: 1.0 mm=39.4 mils. (a) Metal loss + internal oxidation or internal nitridation (whichever is greater). (b) Internal nitridation through thickness. Source: Ref 36


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(i.e., 980 °C for 1000 h), Lai (Ref 37) examined the effect of the thermal cycling on the internal nitridation. In this test, no thermal cycling was involved. The results of this test (Ref 37) were then compared with those in the earlier test (Ref 36). The results of the high-velocity dynamic burner rig test at 980 °C (1800 °F) for 1000 h without thermal cycling are summarized in Table 4.9. The results of chemical analysis showing nitrogen content before and after testing are summarized in Table 4.10. In comparing the test results with thermal cycling and those without thermal cycling, thermal cycling significantly accelerated oxidation attack by causing oxide spallation, as shown in Fig. 4.14. Thermal cycling was also found to accelerate nitridation attack, as shown in Fig. 4.15. Among the four combustor alloys, alloy 230, however, was least affected by thermal cycling. The test program was extended to include some iron-base alloys under the same test conditions using the same dynamic burner rig (Ref 38). The results showed that iron-base alloys suffered significantly more nitridation attack than nickel-base alloys. Figure 4.16(a) shows the microstructure of alloy 556 (Fe-22Cr20Ni-18Co-3Mo-2.5W-0.6Ta-0.2N-La), revealing significant internal nitridation attack with formation blocky chromium nitrides after testing

at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling. The nitrogen content was found to increase to 1.27% after testing from the original 0.13% before testing. Figure 4.16(b) shows the microstructure of Type 310 stainless steel (SS) after testing at 980 °C (1800 °F) for 1000 h with 30 min thermal cycling, revealing significant internal nitridation attack with formation of blocky chromium nitrides. The

Table 4.9 Test results in terms of weight loss, depth of oxidation penetration, depth of nitridation, and total depth of attack after testing at 980 °C (1800 °F) for 1000 h without thermal cycling

Fig. 4.14

Metal loss, mm

Internal oxidation, mm

Internal nitridation, mm

Total attack(a), mm

230 617 263 X

−3.7 −12.1 −19.3 −5.0

0.039 0.031 0.072 0.030

0.065 0.126 0.135 0.086

0.18 >0.55(b) 0.16 0.10

0.22 >0.58(b) 0.23 0.13

Note: 1.0 mm = 39.4 mils. (a) Metal loss +internal oxidation or internal nitridation (whichever is greater). (b) Internal nitridation through thickness. Source: Ref 37

Table 4.10 Results of bulk nitrogen analysis before and after testing at 980 °C (1800 °F) for 1000 h without thermal cycling Alloy

Original nitrogen, wt%

Nitrogen after testing, wt%

230 263 617 X

0.05 0.004 0.03 0.04

0.065 0.097 0.179 0.075

Source: Ref 37

263


–40

–60

–80

–100

–219 Non cyclic test Cyclic test Comparison weight change data between the thermal cycling test (30 min cycles) and no thermal cycle test during the dynamic burner rig testing at 980 °C (1800 °F) for 1000 h for alloys 230, X, 617, and 263. Source: Ref 36, 37

Nitrogen, wt%


617

X

–20

Cyclic test

0.6 0.4 0.2 0

Nitrogen, wt%

Alloy

230

0

230

X

617

263

Non cyclic test

0.4 0.2 0

230

X

617

263

Original After testing

Fig. 4.15

Comparison nitrogen gain data between the thermal cycling test (30 min cycles) and no thermal cycle test during the dynamic burner rig testing at 980 °C (1800 °F) for 1000 h for alloys 230, X, 617, and 263. Source: Ref 36, 37


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nitrogen content increased from the original 0.03% before testing to 1.69% after testing. Alloy 800H was also shown to suffer severe nitridation attack with chromium nitrides and needle-shaped aluminum nitrides (Fig. 4.16c). The oxide scales formed on these iron-base alloys after testing were porous and nonprotective. At 870 °C (1600 °F), internal nitridation attack was found to be significantly reduced.

Figure 4.17 shows the microstructures of alloys 230, 617, and X after testing at 870 °C (1600 °F) for 2000 h with 30 min cycles. Some aluminum nitrides and chromium nitrides were observed in alloy 617, and only some chromium nitrides were observed in alloy X, while no nitrides were observed in alloy 230, as shown in Fig. 4.17

(a)

(a)

20 µm

50 µm

(b)

(b)

(c) (c)

Fig. 4.16

Extensive internal blocky chromium nitrides formed in alloy 556 (a), Type 310 (b), and alloy 800H (c) after the dynamic burner rig testing at 980 °C (1800 °F) for 1000 h with 30 min cycles. Courtesy of Haynes International, Inc.

Fig. 4.17

Alloys 230 (a), 617 (b), and X (c) after the dynamic burner rig testing at 870 °C (1600 °F) for 2000 h with 30 min cycles. Alloy 230 revealed no nitrides, alloy 617 showed both chromium nitrides (blocky phases) and aluminum nitrides (needle phases), and alloy X showed only blocky chromium nitrides. Internal oxides were observed for all three alloys, and all three alloys showed porous and nonprotective external oxide scales. Courtesy of Haynes International, Inc.


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(Ref 38). Oxide scales for three alloys were found to be porous and nonprotective. Iron-base alloys, such as Type 310SS, on the other hand, were found to suffer severe internal nitridation attack, as shown in Fig. 4.18 (Ref 38). The difference in the internal nitridation attack between nickel-base and iron-base alloys is likely to be caused by the differences in nitrogen solubilities in two different alloy systems with much lower nitrogen solubilities in nickel-base alloys. The above test data were generated from Ni-Cr and Fe-Ni-Cr alloys. These alloys are chromia formers (i.e., alloys forming Cr2O3 oxide scale). These chromia formers are susceptible to oxidation/nitridation attack in varying degrees under gas turbine combustion conditions. For very high temperatures and harsh oxidizing combustion conditions, alumina formers (i.e., alloys forming Al2O3 oxide scale) are better performers. Lai (Ref 38) investigated two alumina formers; one was wrought alloy 214 (Ni-16Cr-3Fe-4.5Al-Y) and the other oxidedispersion-strengthened alloy (produced by powder metallurgy) MA956 (Fe-20Cr-4.5Al0.5Y2O3). Due to much more tenacious aluminum oxide scales, these two alloys were tested at 1150 °C (2100 °F) for 200 h with 30 min cycles. The test results showed no nitridation in either alloy. Figure 4.19 shows the cross section of

alloy 214 tested specimen, and Fig. 4.20 shows the cross section of alloy MA956 tested specimen. Both alloys showed fingerlike preferential oxidation penetration. The preferential oxidation penetration was the result of thermal stresses developed from severe thermal cycling from 1150 to less than 260 °C (2100 to 0.58 (23.0)

Note: 100% NH3 in the inlet gas and less than 5% NH3 (detection limit) in the exhaust gas. Source: Ref 41

Table 4.18 Nitridation resistance of various alloys in ammonia at 1090 °C (2000 °F) for 168 h Alloy

Table 4.16 Nitridation resistance of various alloys in ammonia at 650 °C (1200 °F) for 168 h

Alloy base

Nitrogen absorption mg/cm2

600 214 S 230 25 617 188 HR-160 601 RA330 625 316 304 X 150 556 446 6B MULTIMET 825 RA333 800H 253MA 310

Alloy base

Nitrogen absorption, mg/cm2

Depth of nitride penetration, mm (mils)

Nickel Nickel Nickel Nickel Cobalt Nickel Cobalt Nickel Nickel Iron Nickel Iron Iron Nickel Cobalt Iron Iron Cobalt Iron Nickel Nickel Iron Iron Iron

0.2 0.2 1.0 1.5 1.7 1.9 2.0 2.5 2.6 3.1 3.3 3.3 3.5 3.8 4.1 4.2 4.5 4.7 5.0 5.2 5.2 5.5 6.3 9.5

0 0.02 (0.7) 0.34 (13.4) 0.39 (15.3) >0.65 (25.5) >0.56 (22) >0.53 (21) 0.46 (18) >0.58 (23) >0.56 (22) >0.56 (22) >0.91 (36) >0.58 (23) >0.58 (23) 0.51 (20) >0.51 (20) >0.58 (23) >0.64 (25) >0.64 (25) 0.58 (23) >0.71 (28) >0.76 (30) >1.5 (60) >0.79 (31)

Note: 100% NH3 in the inlet gas and less than 5% (detection limit) in the exhaust gas. Source: Ref 41


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15

8



10

6

4

2

0

0

20

40

60

10

5

0

80

0

20

40

60

80

Ni + Co, wt%

Ni + Co, wt%

Fig. 4.28

Fig. 4.27

Effect of the Ni + Co content in iron-, nickel-, and cobalt-base alloys on nitridation resistance at 650 °C (1200 °F) for 168 h in ammonia (100% NH3 in the inlet gas and 30% NH3 in the exhaust). Source: Ref 41

Effect of the Ni + Co content in iron-, nickel-, and cobalt-base alloys on nitridation resistance at 980 °C (1800 °F) for 168 h in ammonia (100% NH3 in the inlet gas and 3.81 (150) >3.81 (150) >3.81 (150) >3.81 (150) >3.81 (150)

2.16 (85) >3.81 (150) >3.81 (150) … >3.81 (150) … …

(a) Specimens were cycled to room temperature once every 24 h for the first 3 days and then weekly for the remainder of the test. (b) Specimens were cycled to room temperature once every 96 h (4 days). (c) Isothermal exposure. Source: Ref 48

Table 4.21 Major phases formed in the nearsurface region of the test specimens after exposure to 100% N2 at 980 °C (1800 °F) for 1008 h, as determined by x-ray diffraction

Table 4.23 Results of x-ray diffraction analysis of extraction residues obtained from specimens after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h

Alloy

Alloy

Major phases

Type 314SS Type 330SS Alloy 800 Alloy 601 Alloy 600

(Cr,Fe)2N (Cr,Fe)2N (Cr,Fe)2N, AlN (Cr,Fe)2N, AlN CrN

Source: Ref 49

Table 4.22 Nitrogen absorbed (mg/cm2) and the average depth of internal nitridation for iron-, nickel-, and cobalt-base alloys after exposure in 100% N2 at 1090 °C (2000 °F) for 168 h Alloy

214 600 230 HR160 X 617 601 188 150 RA330 RA85H 556 HR120 253MA 800H 800HT Type 310 SS

Nitrogen absorbed, mg/cm2

Depth of internal nitridation, mm

0.2 1.1 2.7 3.9 6.0 5.1 7.2 3.7 9.0 6.6 8.5 9.0 9.6 10.0 10.3 11.4 12.3

0.0 0.41 0.46 1.19 0.63 >0.58 >0.59 >0.51 >0.80 >1.52 >1.44 >1.52 >0.86 >1.50 >1.50 >1.46 >0.79

214 (a) 230 600 601 617 HR160 188 150 RA85H

Phases detected

AlN, Al2O3 Cr2N, (Cr,Mo)12(Fe, Ni)8-xN4-z, M6C Cr2N, TiN Cr2N, AlN Cr2N, AlN CrN, Cr2N Cr2N Cr2N Cr2N, AlN

(a) Surface analysis. Source: Ref 50

Source: Ref 50

Fig. 4.33

with AlN/Al2O3 surface scales, as illustrated in Fig. 4.35. Alloy 150 (Co-27Cr-18Fe) suffered nitridation attack as severe as that experienced by some iron-base alloys. Figure 4.36 shows a through-thickness nitrided alloy 150 compared with Type 310SS. Extensive Cr2N nitrides were

Optical micrographs showing a through-thickness nitridation attack for alloy 617, a nickel-base alloy containing about 1.3%Al, after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h. Note extensive blocky chromium nitrides and needle-shaped aluminum nitrides. Magnification bar represents 200 μm. Courtesy of Haynes International, Inc.

observed throughout the alloy 150 specimen cross section.


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Nickel-base alloys are in general more resistant to nitridation attack than iron-base alloys. This is illustrated in Fig. 4.37 comparing alloy X with 253MA after 168 h in 100% N2 at 1090 °C (2000 °F). Similar findings were observed in nitridation studies in nitrogen atmospheres by Smith and Bucklin (Ref 48) and Tjokro and Young (Ref 51). Tjokro and Young (Ref 51) investigated a number of commercial alloys in N2-5%H2 at 1100 and 1200 °C. Their results

10 µm

showed the nitridation rate constants decreased with increasing nickel concentration, as illustrated in Fig. 4.38.

(a)

200 µm

(b)

200 µm

Fig. 4.34

Scanning electron micrograph (backscattered image) showing internal chromium nitrides (blocky phases) and aluminum nitrides (long needle-shaped phases) formed in alloy 601, a nickel-base alloy containing about 1.3% Al, after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h. Source: Ref 50

4.50 230 alloy

N absorbed, mg/cm2

4.00 3.50 3.00 2.50 2.00 1.50 1.00 0.50 0.00 0.00

214 alloy 200.00

400.00

600.00

Time, h

Fig. 4.35

Nitridation kinetic data for alloy 214 (nickel-base alloy containing 4.5% Al) and alloy 230 (nickelbase alloy containing little aluminum) after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h. Source: Ref 50

Fig. 4.36

Optical micrographs showing through-thickness nitridation attack for (a) Type 310SS (Fe-25Cr20Ni) and (b) alloy 150 (Co-27Cr-18Fe) after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h. Courtesy of Haynes International, Inc.


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4.7 Nitridation Kinetics between NH3 and N2 Atmospheres Nitridation data generated in 100% NH3 (Ref 41) and those generated in 100% N2 (Ref 50) were compared. Since both test programs were generated using the same test apparatus and procedures, and both tests were carried out by same technicians, the laboratoryto-laboratory variation was significantly minimized. Thus, the comparison between these two sets of test results could yield a more meaningful comparison in terms of the difference in environments. The test results generated at 1090 °C

(a)

200 µm

(2000 °F) for 168 h in 100% NH3 (Ref 41) and 100% N2 (Ref 50) are tabulated in Table 4.24. The data are also presented in terms of nitrogen absorption as a function of Ni + Co content in the alloy (Fig. 4.39). The results clearly indicated that the nitrogen atmosphere was a more severe nitriding environment than the ammonia environment at 1090 °C (2000 °F). The amount of nitrogen absorbed in N2 environment was more than double that in NH3 environment for many alloys. Figure 4.40 shows two nitrided alloy 601 specimens, one exposed to NH3 environment and the other to N2 environment. The N2 environment caused significantly more internal nitride formation than for the NH3 environment. More chromium nitrides (blocky shaped) and aluminum nitrides (needle shaped) formed in the N2 environment than in the NH3 environment. Ammonia readily dissociates to one part N2 and three parts H2 at 1090 °C (2000 °F). With the test system used in the study by Barnes and Lai (Ref 41), 100% NH3 was fed into the alumina test tube with no test specimens inside, and the exhaust gas was measured to contain less than 5% NH3, which was the detection limit of the the apparatus used for measuring NH3 (Table 4.11). It is believed most, if not all, of the ammonia had been dissociated into H2 and N2 before the test gas was in contact with the test specimens. The NH3 test environment was essentially a cracked ammonia, which was dissociated into H2 and N2. Thus, the nitridation potential ( pN2 ) in the NH3 test environment (0.25 atm) was much lower than that in the N2 test environment (1.0 atm). As a result, the N2 test environment was found to produce more severe nitridation attack for most of the alloys tested (Table 4.24 and Fig. 4.39).

4.8 Summary

(b)

Fig. 4.37

200 µm

Optical micrographs showing a through-thickness nitrided Fe-20Cr-10Ni-1.7Si-Ce alloy 253MA (a) and a better nitridation resistant Ni-22Cr-9Mo-18Fe-0.6W alloy X after exposure to 100% N2 at 1090 °C (2000 °F) for 168 h. Courtesy of Haynes International, Inc.

Nitridation behavior of metals and alloys in (a) air, (b) gas-turbine combustion gas, (c) NH3-H2O, (d) NH3, and (e) N2 environments is reviewed. Nitridation attack can occur in air and oxidizing, combustion environments. Under certain conditions, alloys can suffer oxidation/ nitridation attack. Internal nitridation attack is much more prevalent in a high-velocity combustion gas stream with thermal cycling. In NH3H2O environments, alloys appear to behave differently under nitridation attack. Extensive review is carried out on the behavior of metals

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6000 Intragranular 1000 °C (1830 °F) 1100 °C (2010 °F)

309S

4000

kp, µm2/h

310S

153MA

2000

RA330 AC66 800

253MA

353MA IN601 0

0

0.200

0.400

0.600

0.800

1.000

XNi' (a) 12 Intergranular 1000 °C (1830 °F) 1100 °C (2010 °F)

153MA

kp, 103 µm2/h

253MA

800

309S 310S

6

RA330 AC66

353MA IN601

0

0

0.200

0.400

0.600

0.800

1.000

XNi' (b)

Fig. 4.38

Nitridation rate constants as a function of the alloy’s nickel concentration when tested in N2-5%H2 at 1000 and 1100 °C (1830 and 2010 °F). Source: Ref 51

and alloys in NH3 and N2 environments. Comparative resistance to nitridation attack for a wide variety of alloys is presented.

REFERENCES

1. Metals Handbook, Vol 2, 8th ed., American Society For Metals, 1964, p 149 2. Metals Handbook, Vol 2, 8th ed., American Society For Metals, 1964, p 119 3. G.L. Swales, Behavior of High Temperature Alloys in Aggressive Environments, Proc.

1979 Petten International Conference, I. Kirman et al., Ed., The Metals Society, London, 1980, p 45 4. K. Rorbo, Environmental Degradation of High Temperature Materials, Series 3, No. 13, Vol 2, The Institution of Metallurgists, London, 1980, p 147 5. R.N. Shreve, The Chemical Process Industries, McGraw-Hill, 1956 6. J.M.A. Van der Horst, Corrosion Problems in Energy Conversion and Generation, C.S. Tedmon, Jr., Ed., The Electrochemical Society, 1974


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Table 4.24 Nitrogen absorption in NH3 and N2 environments at 1090 °C (2000 °F) for 168 h Nitrogen absorption, mg/cm2 Alloy

214 600 230 617 601 X 188 150 556 RA330 800H 253MA

NH3

N2

0.2 0.3 1.5 1.9(a) 2.6(a) 3.8(a) 2.0(a) 4.1 4.2(a) 3.1(a) 5.5 6.3(a)

0.0 1.1 2.7 5.1(a) 7.2(a) 6.0(a) 3.7(a) 9.0(a) 9.0(a) 6.6(a) 10.3(a) 10.0(a)

(a) Nitrided all the way through specimens. Source: Ref 41, 50

13 1090 °C (2000 °F) / 168 h N2 12

(a)

200 µm

(b)

200 µm

NH3


10

8

6

4

2

0

0

20

40

60

80

Ni or Ni + Co, wt%

Fig. 4.39

Nitrogen absorption as a function of Ni + Co content in the alloy for 100% NH3 and 100% N2 environments at 1090 °C (2000 °F) for 168 h. Source: Ref 41, 50

7. M.B. Bever and C.F. Floe, Source Book on Nitriding, American Society For Metals, 1977, p 125 8. B.J. Lightfoot and D.H. Jack, Source Book on Nitriding, American Society For Metals, 1977, p 248 9. K.N. Strafford, Corros. Sci., Vol 19, 1979, p 49 10. T. Rosenquist, Principles of Extractive Metallurgy, McGraw-Hill, 1974

Fig. 4.40

Optical micrographs showing both chromium nitride and aluminum nitride (needle-shaped phase) formed in alloy 601 after exposure to (a) 100% NH3 and (b) 100% N2 at 1090 °C (2000 °F) for 168 h. Courtesy of Haynes International, Inc.

11. T. Masumoto and Y. Imai, J. Jpn. Inst. Met., Vol 33, 1969, p 1364 12. H.J. Christ, S.Y. Chang, and U. Krupp, Thermodynamic Characteristics and Numerical Modeling of Internal Nitridation of Nickel Base Alloys, Mater. Corros., Vol 54 (No. 11), 2003, p 887


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13. N.S. Corney and E.T. Turkdogan, The Effect of Alloying Elements on the Solubility of Nitrogen in Iron, J. Iron Steel Inst., Aug 1955, p 344 14. D.H. Jack and K.H. Jack, Carbides and Nitrides in Steel, Mater. Sci. Eng., Vol 11, 1973, p 1 15. H.A. Wriedt and O.D. Gonzalez, Trans. AIME, Vol 221, 1961, p 532 16. J.F. Eckel, T.P. Floridis, and B.N. Ferry, Nitrides in Type 304 Stainless Steel, Virginia J. Sci., Vol 17, 1966, p 325 17. J.F. Eckel and T.B. Cox, J. Mater., Vol 3, 1968, p 605 18. L.E. Kindlimann and G.S. Ansell, Kinetics of the Internal Nitridation of Austenitic Fe-Cr-Ni-Ti Alloys, Metall. Trans., Vol 1, 1970, p 163 19. A.J. Heckler and J.A. Peterson, The Effect of Nickel on the Activity of Nitrogen in Fe-Ni-N Austenite, Trans. Metall. Soc. AIME, Vol 245, 1969, p 2537 20. J. Litz, A. Rahmel, M. Schorr, and J. Weiss, Scale Formation on the Ni-Base Superalloys IN 939 and IN 738LC, Oxid. Met., Vol 32, 1989, p 167 21. M.A. Harper, J.E. Barnes, and G.Y. Lai, Long-Term Oxidation Behavior of Selected High Temperature Alloys, Paper No. 132, Corrosion/97, NACE International, 1997 22. G.Y. Lai, unpublished results, 2003 23. S. Han and D.J. Young, Simultaneous Internal Oxidation and Nitridation of NiCr-Al Alloys, Oxid. Met., Vol 55, 2001, p 223 24. D.L. Douglass, Anomalous Behavior During Internal Oxidation and Nitridation, JOM, Nov 1991, p 74 25. R.P. Rubly and D.L. Douglass, Oxid. Met., Vol 35, 1991, p 269 26. R.P. Rubly and D.L. Douglass, Internal Nitridation of Ni-Cr-Al Alloys, Proc. Int. Symp. On Solid-State Chemistry of Advanced Materials: High-Temperature Corrosion Workshop, 1992 27. H.J. Grabke and E.M. Peterson, Scr. Met., Vol 12, 1978, p 1111 28. J.-W. Park and C. J. Alstetter, Metall. Trans. A, Vol 18A, 1987, p 43 29. P.L. Gruzin, Y.A. Polikarpov, and G.B. Federov, Fiz. Metal. I Metalloved., Vol 4 (No. 1), 1957, p 94 30. K.G. Brickner, G.A. Ratz, and R.F. Domagala, Creep-Rupture Properties of Stainless Steels at 1600, 1800, and 2000 °F,

31. 32.

33.

34. 35.

36.

37.

38. 39. 40. 41.

Advances in the Technology of Stainless Steels and Related Alloys, STP 369, ASTM, 1965, p 99 M. Yu, R. Sandstrom, B. Lehtinen, and C. Westman, Scand. J. Metall., Vol 16, 1987, p 154 V. Guttmann and R. Burgel, CreepStructural Relationship in Steel Alloy 800H at 900–1000 °C, Met. Sci., Vol 17, 1983, p 549 M. Welker, A. Rahmel, M. Schutze, Oxidation and Nitridation of Alloy 800H at a Growing Creep Crack and for Unstressed Samples, Metall. Trans. A, Vol 20A, 1989, p 1541 J.J. Hoffman and G.Y. Lai, Paper No. 5402, Corrosion 2005, NACE International, 2005 V.P. Swaminathan and S.J. Lukezich, Degradation of Transition Duct Alloys in Gas Turbines, Advanced Materials and Coatings for Combustion Turbines, Proc. ASM 1993 Materials Congress Materials Week (Pittsburgh, PA), Oct 17–21, 1993, V.P. Swaminathan and N.S. Cheruvu, Ed., ASM International, 1994, p 99 G.Y. Lai, Nitridation of Several Combustor Alloys in a Simulated Gas Turbine Combustion Environment, Advanced Materials and Coatings for Combustion Turbines, Proc. ASM 1993 Materials Congress Materials Week (Pittsburgh, PA), Oct 17–21, 1993, V.P. Swaminathan and N.S. Cheruvu, Eds., ASM International, 1994, p 113 G.Y. Lai, Nitridation Attack in a Simulated Gas Turbine Combustion Environment, Materials for Advanced Power Engineering, Part II, D. Coutsouradis et al., Ed., Kluwer Academic Publishers, The Netherlands, 1994, p 1263 G.Y. Lai, unpublished results, Haynes International, Inc., 1995 Y.M. Park and R.E. Sonntag, Int. J. Energy Res., Vol 14, 1990, p 153 H.J. Grabke, S. Strauss, and D. Vogel, Nitridation in NH3-H2O Mixtures, Mater. Corros., Vol 54 (No. 11), 2003, p 895 J.J. Barnes and G.Y. Lai, High Temperature Nitridation of Fe-, Ni-, and Co-base Alloys, Corrosion & Particle Erosion at High Temperatures, Proc. TMS-ASM Symposium, V. Srinivasan and K. Vedula, Ed., The Minerals, Metals & Materials Society, 1989, p 617

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42. K.M. Verma, H. Ghosh, and J.S. Rai, Brit. Corros. J., Vol 13 (No. 4), 1978, p 173 43. V. Cihal, Corrosion Mechanisms in Ammonia Synthesis Equipment, Conf. Proc., First International Congress on Metallic Corrosion (London, U.K.), April 10–15, 1961, L. Kenworthy, Ed., Butterworths, London, 1962, p 591 44. J.J. Moran, J.R. Mihalisin, and E.N. Skinner, Corrosion, Vol 17 (No. 4), 1961, p 191t 45. D.W. McDowell, Jr., Mater. Protect., Vol 1 (No. 7), 1962, p 18 46. K.H. Jack, High Temperature Gas-Metal Reactions in Mixed Environments, S.A. Jansson and Z.A. Foroulis, Ed., The Metallurgical Society of AIME, 1973, p 182 47. H. Schenck, M.G. Frohberg, and F. Reinders, Stahl Eisen, Vol 83, 1963, p 93

48. G.D. Smith and P.J. Bucklin, Some Observation on the Performance of Nickel-Containing Commercial Alloys in Nitrogen-Based Atmospheres, Paper No. 375, Corrosion/86, NACE, 1986 49. P. Ganesan and G.D. Smith, Performance of Selected Commercial Alloys in Nitrogen Based Sintering Atmospheres, Paper No. 278, Corrosion/90, NACE, 1990 50. J.J. Barnes and G.Y. Lai, Factors Affecting the Nitridation Behavior of Fe-Base, NiBase and Co-Base Alloys in Pure Nitrogen, J. Physique IV, Colloque C9, supplemental au Journal de Physique III, Vol 3, 1993, p 167 51. K. Tjokro and D.J. Young, Comparison of Internal Nitridation Reactions in Ammonia and in Nitrogen, Oxid. Met., Vol 44, 1995, p 453

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CHAPTER 5

Carburization and Metal Dusting 5.1 Introduction Metals and alloys are susceptible to carburization when exposed to an environment containing CO, or CH4 or other hydrocarbon gases, such as ethane (C2H6), propane (C3H8), and so forth, at elevated temperatures. Carburization attack generally results in formation of internal carbides, which often cause the alloy to suffer embrittlement as well as other mechanical property degradation. Carburization problems are quite common to heat treating equipment, particularly furnace retorts, baskets, fans, and other components used for case hardening of steels by gas carburizing. A common commercial practice for control of gas carburizing is to use an endothermic gas as a carrier enriched with one of the hydrocarbon gases, such as CH4, C3H8, and so forth (Ref 1). An endothermic gas enriched with about 10% natural gas (CH4) is a commonly used atmosphere (Ref 2). The typical endothermic gas consists of 39.8% N2, 20.7% CO, 38.7% H2, and 0.8% CH4, with a dew point of −20 to −4 °C (−5 to +25 °F) (Ref 1). Gas carburizing occurs typically at 840 to 930 °C (1550 to 1700 °F). Furnace equipment and components repeatedly subjected to these service conditions frequently suffer brittle failures as a result of carburization attack. In the petrochemical industry, carburization is one of the major modes of high-temperature corrosion for processing equipment. The pyrolysis furnace tubes for production of ethylene and olefins are a good example (Ref 3–5). Ethylene is formed by cracking petroleum feedstock, such as ethane and naphtha, at temperatures up to 1150 ° C (2100 °F). This generates a strong carburizing gas stream inside the tubes. As a result, carburization was found to be a major mode of tube failure in a survey of ethylene and olefin pyrolysis furnaces conducted by Moller and Warren (Ref 3).

Production of carbon fibers also generates carburizing atmospheres in a furnace. As a result, the furnace’s retorts, fixtures, and other components require frequent replacement because of carburization attack. Metal dusting, a form of catastrophic carburization, can occur at intermediate temperatures when a process gas stream consists of primarily H2, CO, and CO2 along with some hydrocarbons with high carbon potentials (ac > 1). Metals or alloys can suffer rapid metal wastage in a form of pitting or general thinning of the cross-sectional thickness of a metallic component. Metal dusting typically occurs at temperatures between 430 and 900 °C (800 and 1650 °F). Materials failures associated with metal dusting have been encountered in refining and petrochemical processing, such as production of syngas in hydrogen, ammonia, and methanol plants, heat treating, and other industrial processes (Ref 5–10).

5.2 Carburization 5.2.1 Carburization—Thermodynamic Considerations Whether an alloy is likely to be carburized or decarburized depends on the carbon activity (ac) in the environment and that of the alloy. The thermodynamic condition that dictates either carburization or decarburization can be described simply. The alloy is likely to be carburized when: (ac )environment 4 (ac )metal

The alloy is likely to be decarburized when: (ac )environment 5 (ac )metal

Thus, in order to predict whether an alloy will be carburized, one needs to know the carbon activities of both the environment and the alloy.

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pg 98


Carburization can proceed by one of the following reactions when the environment contains CH4, CO, or H2 and CO: CO+H2 =C+H2 O

ðEq 5:1Þ

2CO=C+CO2

ðEq 5:2Þ

CH4 =C+2H2

ðEq 5:3Þ

Assuming that carburization follows Reaction 5.1, the carbon activity in the environment can be calculated by: DG =7RT ln

ac pH2 O pCO pH2

ðEq 5:4Þ

ac =e7DG

=RT

pCO pH2 pH2 O

DG =7RT ln

ac =e7DG

=RT

ðEq 5:5Þ

p2CO pCO2

ðEq 5:6Þ

ðEq 5:7Þ

pCH4 p2H2

!

ðEq 5:8Þ

Carbon activities as a function of ( pCH4 =p2H2 ) are plotted in Fig. 5.3. Reactions 5.1 and 5.2 have a similar characteristic, showing lower carbon activities with increasing temperature (Fig. 5.1 and 5.2).

H2 = C + H2O

Carbon activity (ac) as a function of gaseous composition in terms of (pCO pH2 =pH2 O ) ratios based on Eq 5.1 for various temperatures. Also plotted are carbon activities for carbon steel (in equilibrium with Fe3C), and for 2.25Cr-1Mo and austenitic stainless steels (both measured ac).

ac pCO2 p2CO

Plots of carbon activities as a function of gas compositions in terms of (p2CO =pCO2 ) for various temperatures are shown in Fig. 5.2. When carburization follows Reaction 5.3, the carbon activity in the environment is: ac =e

From Eq 5.5, one can construct graphs of carbon activity as a function of gaseous composition in terms of ( pCO pH2 =pH2 O ) ratios for various temperatures, as shown in Fig. 5.1.

Fig. 5.1

7DG =RT

Rearranging the equation changes it to:

Similarly, if carburization follows Reaction 5.2, the carbon activity of the environment can also be calculated:

2CO = CO2 + C

Fig. 5.2

Carbon activity (ac) as a function of gaseous com2 position in terms of (pCO =pCO2 ) based on Eq 5.2 for various temperatures. Also plotted are carbon activities for carbon steel (in equilibrium with Fe3C), and for 2.25Cr-1Mo and austenitic stainless steels (both measured ac).


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Chapter 5:

Reaction 5.3, on the other hand, shows increased carbon activities with increasing temperature (Fig. 5.3). If the environment contains CH4, the carbon activity of the environment at higher temperatures is likely to be dominated by Reaction 5.3. When no CH4 is present in the environment, Reaction 5.1 and/or 5.2 will dictate the carbon activity. The carbon activity maps shown in Fig. 5.1 to 5.3 were previously described by Mazandarany and Lai (Ref 11) in assessing the carburizationdecarburization behavior of alloys in hightemperature gas-cooled helium environments containing H2, CO, CO2, CH4, and H2O. These activity maps provide a simple means of estimating an environment’s carbon activity for predicting whether or not the environment is thermodynamically capable of carburizing an alloy. When the gas stream contains many gaseous components, such as H2, CO, CO2, CH4, and

pg 99

Carburization and Metal Dusting / 99

H2O, and under very dynamic conditions with a high gas velocity such that the gaseous components do not have time to react to reach a thermodynamic equilibrium (i.e., nonequilibrium conditions), the gas-metal reaction can be reasonably assumed to follow the dominating reaction from one of those shown in Reaction 5.1, 5.2, or 5.3, and thus, one of the activity maps shown in Fig. 5.1, 5.2, or 5.3. Data for carbon activities of commercial alloys at temperatures below 1200 °C (2190 °F) is very limited. Natesan (Ref 12) reported that ac for 2.25Cr-1Mo steel is in the range of 1×10−1 to 10−2 from 550 to 750 °C (1020 to 1380 °F). Natesan and Kassner (Ref 13) reported the carbon activities of Fe-18Cr-8Ni alloys. These values are superimposed in Fig. 5.1 to 5.3. For carbon steels, carbon activity can be estimated by assuming that it is in equilibrium with cementite (Fe3C): ðEq 5:9Þ

3Fe+C=Fe3 C DG =7RT ln DG =7RT ln

aFe3 C (ac ) (aFe )3

1 ac

ðEq 5:10Þ

ðEq 5:11Þ

where aFe3 C and aFe are assumed to be unity. ac =eDG

CH4 = 2H2 + C

Fig. 5.3

Carbon activity (ac) as a function of gaseous com2 position in terms of (pCH4 =pH ) based on Eq 5.3 for 2 various temperatures. Also plotted are carbon activities for carbon steel (in equilibrium with Fe3C), and for 2.25Cr-1Mo (measured ac). Carbon activities of austenitic stainless steels are below 10−2 at 800 to 1000 °C (1470 to 1830 °F).

=RT

ðEq 5:12Þ

The ac values for carbon steel based on Eq 5.12 are plotted in Fig. 5.1 to 5.3. Using Fig. 5.1, 5.2, or 5.3 one can make a quick determination as to whether an environment has a carbon potential (or activity) high enough to carburize the alloy of interest. Even though in cases where the gas mixture may not reach an equilibrium condition, it will be of great benefit to better understand the gas-metal reaction in terms of the thermodynamic equilibrium condition in multicomponent gases environments. The thermodynamic equilibrium gaseous composition along with its thermodynamic potentials, such as carbon activity (ac), oxygen potential ( pO2 ), and other potentials, can be determined using a commercial software program such as “HSC Chemistry for Windows” (Ref 14) and ChemSage (Ref 15). The environment can also be characterized in terms of ac and pO2 to determine the relative severity of its carburization potential. Both carbon activity and oxygen potential can be calculated by a computer program. The environment


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2

Log PCO /PCO

2

During carburization, the relative stabilities of these carbides can be best described by a stability diagram, such as the one shown in Fig. 5.5. If the carbon and oxygen activities of the environment are in the Cr3C2 region, conditions will favor formation of Cr3C2 on the surface and/or in the underlying metal. As carbon diffuses farther into the alloy’s interior, carbon activities will be lowered, thus favoring Cr7C3. Moving even farther into the interior, carbon activities will be further reduced, favoring the formation of Cr23C6. These chromium carbides can incorporate other alloying elements depending on alloy system. For example, in Fe-Ni-Cr system, iron with very little nickel can be incorporated into these chromium carbides. The combined metal elements in the carbide are then represented by “M,” as M3C2, M7C3, and M23C6. An example of the metallic compositions of M7C3 and M23C6 formed in carburization of Type 304L is illustrated in Fig. 5.6 (Ref 20). Both M7C3 and M23C6 contain essentially Cr and Fe with negligible amount of Ni. For many high-temperature alloys, particularly superalloys, there are other alloying elements, such as Ti, Ta, Nb (or Cb), Mo, and W, that can form carbides. The carbides of these alloying elements are important to the physical metallurgy of high-temperature alloys in that they provide an important strengthening mechanism. A general review of binary metallic Log PCH /PH 2 4 2

can then be presented in a stability diagram of a metal-carbon-oxygen system. The stability diagrams of Fe-C-O and Cr-C-O systems are shown in Fig. 5.4 and 5.5 (Ref 16). From the stability diagram, the possible phases that the alloy may form at the gas/metal interface can be predicted. As the activities of both carbon and oxygen are decreasing from the gas/metal interface to the metal interior, the possible phases that the alloy may form beneath the gas/metal interface can also be predicted. For carbon and alloy steels with low concentrations of chromium, ingress of carbon into the metal or alloy may result in the formation of iron carbides. Several forms of iron carbides have been reported (Ref 17), with compositions ranging from Fe4C to Fe2C. They are ξ phase (Fe4C), θ phase (Fe3C), χ phase (Fe2.2C), and ε phase (Fe2-3C). Fe3C (cementite) is the most stable iron carbide. Other iron carbides are less stable. Browning et al. (Ref 18) found that χ phase, which formed by carburizing αFe with butane at 275 °C (530 °F), was converted to Fe3C when heated to 500 °C (930 °F). The ε phase (Fe2-3C) is a transition phase that forms in martensite during tempering of steel (Ref 19). In ferritic and austenitic stainless steels and nickel- and cobalt-base alloys, ingress of carbon into the alloy results in the formation of mainly chromium carbides. There are three forms of chromium carbides: Cr23C6, Cr7C3, and Cr3C2.

Log PCO /PCO –18 –16 –14 –12 –10 –8 –6 –4 –2 2 Log P /P –18 –16 –14 –12 –10 –8 –6 –4 –2 H2O

H2

0

2

4

6

8

10 12

0

2

4

6

8

10 12

0 2

C(s)

0

–2 0

–2

–4

Fe3C(s)

–2

Log aC

–4

–6 –4

–6 –8 –10

–8 Fe(s)

Fe3C4(s)

–6 –10

Fe0.95O(s)

–8 –12 –10

–12

–14

–14

–16

–16

–12 –14 –50 –45 –40 –35 –30 –25 –20 –15 –10 Log PO , atm 2

Fig. 5.4

Fe2O3(s)

Stability diagram of Fe-C-O system at 870 °C (1600 °F). Source: Ref 16

–5

0

5

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26/10/2007 12:21PM Plate # 0

pg 101

H2

Cr3C2(s)

0 –2

10 12 14

8

6

8

10 12 14

C(s)

Cr7C3(s)

–4 Cr23C6(s)

Log aC

–6 –8

Cr2O3(s)

–10

2

2

H2O

6

2

4

Log PCO /PCO –14 –12 –10 –8 –6 –4 –2 0 2 4 2 –14 –12 –10 –8 –6 –4 –2 0 2 4 Log P /P

Log PCO /PCO

Log PCH /PH2

Chapter 5: Carburization and Metal Dusting / 101

0

0

–2

–2

–4

–4

–6

–6

–8

–8

–10 –10

Cr(s)

–12

–12 –12

–14

–14 –14 –16 –16

–16 –50 –45 –40 –35 –30 –25 –20 –15 –10

–5

0

5

Log PO , atm

(a)

H2O

H2

0 –2

0

2

4

6

8 10 12

0

2

4

6

8 10 12

0 2

Cr3C2(s)

C(s)

–2 0

Cr7C3(s)

–4 –2

–4

–6

Log aC

Cr23C6(s)

–4

–6

–8 –6

–8 –10

Log PCH /PH 2

Log PCO /PCO –18 –16 –14 –12 –10 –8 –6 –4 –2 2 –18 –16 –14 –12 –10 –8 –6 –4 –2 Log P /P

4 2 2 Log PCO /PCO 2

2

–10

Cr2O3(s)

Cr(s)

–8 –12 –10

–12

–14

–14

–16

–12 –14

–16 –50 –45 –40 –35 –30 –25 –20 –15 –10 (b)

Fig. 5.5

–5

0

5

Log PO , atm 2

Stability diagrams of Cr-C-O system at (a) 620 °C (1150 °F), (b) 870 °C (1600 °F), and (c) 1090 °C (2000 °F). Source: Ref 16

carbides can be found elsewhere (Ref 21, 22). The relative stabilities of some binary carbides are shown in Fig. 5.7 (Ref 22). When the environment contains oxygen and carbon activities, temperature is an important factor in determining whether the oxide or carbide will be thermodynamically stable. Considering a chemical reaction, such as 3Cr2O3 + 4C = 2Cr3C2 + 9/2 O2, Cr3C2 will remain stable when the reaction goes from left to right. In order

to keep the reaction going from left to right (i.e., keeping Cr3C2 stable), the pO2 of the environment shall be lower than the equilibrium pO2 associated with the above reaction. On the other hand, if the pO2 of the environment is higher than the equilibrium pO2 associated with the above reaction, Cr2O3 will become stable. Temperature can be a significant factor in determining whether chromium oxide or chromium carbide is stable, and thus significantly affects the carburization


26/10/2007 12:21PM Plate # 0

H2

0

Cr3C2(s)

4

6

8

10

4

6

8

10

4 –2

C(s)

2

Cr7C3(s)

–2

–4 0 –6

Cr23C-6(s)

–4 Log aC

2 2

–2 –8

–6

–4

–8

–10

Cr2O3(s)

Cr(s)

–6 –12

–10

–8 –14

–12

–10 –16

–14

–12 –18

–16 –50 –45 –40 –35 –30 –25 –20 –15 –10

Concentration, wt.%

1123 k 150 h

80

M23C6 Cr

40

Fe

Ni

M7C3

0 0.2

5

(c) 1090 °C (2000 °F). Source: Ref 16

100

0

0

2

Fig. 5.5 (continued)

20

–5

Log PO , atm

(c)

60

2

2

H2O

0 0

2

4

Log PCO /PCO –20 –18 –16 –14 –12 –10 –8 –6 –4 –2 2 Log P /P –20 –18 –16 –14 –12 –10 –8 –6 –4 –2

Log PCO /PCO

Log PCH /PH 2


0.4

0.6

0.8

1

Distance to surface, mm

Fig. 5.6

Compositions of the metallic components of M7C3 and M23C6 formed in Type 304L after carburizing at 1123 K (850 °C) in H2-2.6CH4 (ac = 0.9) for 150 h. Source: Ref 20

behavior of an alloy. This can be nicely illustrated in a plot that contains the oxygen potentials of the environment (Boudouard reaction [2CO = C + CO2] is used for this example) and those in equilibrium with Cr2O3/Cr3C2 as a function of temperature. This is shown in Fig. 5.8 (Ref 23). The figure shows the equilibrium pO2 line of Cr2O3/Cr3C2 intersecting with the environment’s pO2 lines (pCO = 1.0 bar, pCO = 0.5 bar, and pCO = 0.25 bar). The intersections are between 1000 and 1200 °C. On the right side of the intersections (i.e., lower temperatures), pO2 (environment) is greater than pO2 (Cr2O3/Cr3C2).

Thus, Cr2O3 is stable. On the left side of the intersections (i.e., higher temperatures), pO2 (environment) is lower, thus favoring the formation of carbides, but not oxides. Accordingly, higher temperatures favor carburization thermodynamically and lower temperatures favor formation of oxides thus retarding carburization. Nishiyama et al. (Ref 24) examined the effect of the temperature on the stability of chromium oxide and carbides based on the ethylene pyrolysis environment, which is generated by the reaction of naphtha with steam for the production of ethylene (C2H4) and propylene (C3H6) at temperatures of approximately 900 to 1100 °C (1650 to 2012 °F). In their calculation of the oxygen potential for the reaction of naphtha with steam, three naphtha feedstocks were used with the steam/naphtha weight ratios of 0.4 and 0.5. The calculated pO2 for the pyrolysis environment as a function of temperature is plotted in Fig. 5.9. Also plotted in Fig. 5.9 are pO2 values in equilibrium with Cr3C2/Cr2O3 and those in equilibrium with Cr7C3/Cr2O3. The results in Fig. 5.9 are very similar to those shown in Fig. 5.8, where the environment was calculated from CO-CO2 reaction (Boudouard reaction). The calculation by Nishiyama et al. showed that the chromium oxide was stable up to 1030 to 1040 °C (1886 to 1904 °F) and became unstable above those temperatures. At temperatures above 1030 to

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Fig. 5.7

Standard free energies of formation for carbides. Source: Ref 22

1040 °C (1886 to 1904 °F), carbides such as Cr3C2 and Cr7C3 became stable. This oxidecarbide transition temperature can vary depending on the steam/naphtha ratio, as illustrated in Fig. 5.10 (Ref 24). As shown in the figure, when the steam/naphtha ratio is decreased to 0.1 from

the common operating ratios of 0.35 to 0.5, the oxide-carbide transition temperature is decreased to about 970 °C (1778 °F). In Reactions 5.1 to 5.3, carbon deposition (coking) can occur when the carbon activity (ac) in the environment is greater than 1.0 (ac = 1

pg 103


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in equilibrium with graphite). Many laboratory carburization tests have been conducted in H2-CH4 mixtures. Figure 5.11 shows the carbon activities of H2-1%CH4 and H2-2%CH4 as a function of temperature (Ref 25). Grabke

600

800

1000

Cr3C2/Cr2O3

–10

PCO–1 bar

Cr3C2 –15

PCO–0.5 bar PCO–0.25 bar

2

Log PO , bar

1200

1800 1600 1400

T, °C

(Ref 25) recommended that the H2-1%CH4 mixture with ac 1 for testing to correspond to the processes where coking is taking place. The test temperature is thus recommended to be higher than 1000 °C (1832 °F) (see Fig. 5.11). Ethylene pyrolysis environment is known to develop coking on the internal surface of the pyrolysis furnace tubes. The deposition of carbon is the result of decomposition of ethylene (C2H4) in the reaction as described in Eq 5.13 (Ref 26). Significant coking in the internal surface of the ethylene pyrolysis furnace tube and the repeated decoking

–20 C/CO-CO2

Cr2O3

–25 5.10–4

75.10–4

10.10–4

12.5.10–4

1/T, K–4

Fig. 5.8

The oxygen potentials of the environment based on the Boudouard reaction (2CO = C + CO2) and those in equilibrium with Cr2O3/Cr3C2 as a function of temperature. Source: Ref 23

T, °C 0,150 1,100 1,050 1,000 –17

900

850

Equilibrium pO2 of the environment based on ethylene pyrolysis of naphtha I with various steam/ naphtha weight (S/O) ratios as a function of temperature, and pO2 in equilibrium with Cr3C2/Cr2O3 and Cr7C3/Cr2O3 as a function of temperature. Source: Ref 24

Fig. 5.10

Cr-carbides + metastable Cr2O3

–18

Cr2O3+ metastable Cr-carbides –19

ac

2

Log PO .atm

950

–20

) =1 (a c 3 ) r 2O =1 /C C2 (a c 3 Cr 3 r 2O /C C3 Cr 7

II (0

Na

.5)

pht

–21 Cr-carbides –22

5

7

7.5

8

ha

and

III (

)

.4)

8.5

2% CH4

4

0.5

I (0

Carbon activity in H2-CH4 1 bar

3

9

1/T × 104, 1/K Equilibrium pO2 of the environment based on ethylene pyrolysis of naphtha with steam/naphtha ratios of 0.4 and 0.5 (in parentheses) as a function of temperature, and pO2 in equilibrium with Cr3C2/Cr2O3 and Cr7C3/Cr2O3 as a function of temperature. Naphtha I, II, and III represent different naphtha feed stocks. Source: Ref 24

1% CH4

2 1

Fig. 5.9

800

Fig. 5.11 Ref 25

850

900

950 1000 1050 1100 1150 °C

Carbon activities (ac) of H2-1%CH4 and H2-2%CH4 are plotted as a function of temperature. Source:

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operation to remove this coke deposit can significantly degrade the tube life (Ref 27). C2 H4 =2C+2H2

ðEq 5:13Þ

5.2.2 Resistance to Carburization Carburization attack generally results in the formation of internal carbides in the alloy matrix as well as at grain boundaries. The gravimetric method has been widely used for studying carburization kinetics. This method can sometimes produce a misleading result when the environment exhibits an oxygen potential high enough to form oxides of some active alloying elements. The weight gain, in this case, is the result of both carbon ingress and oxide formation. Measurements of carburization depth have also been used by some investigators. Different alloy systems can produce significant differences in the concentration profile for the carburized layer. Thus, one alloy may exhibit a large carburization depth with only a slight concentration gradient, while another alloy may show a narrow carburization depth with a steep concentration profile. Furthermore, measurements of carburization depth by the metallographic method can be difficult when separating the carbides formed by carburization from those formed by thermal aging. Some investigators measured the total amount of carbon in the alloy after the exposure. The measurement of the carbon concentration profile as a function of distance from the metal surface may be an excellent method for characterizing

Carbon content, %

5

A

4 B

3

2 C 1

0 ID

Distance from bore of tube, mm

OD

Wall thickness

Fig. 5.12

Three possible carbon concentration profiles of carburized alloys. Source: Ref 28

the carburized alloy. Each evaluation method has its merit. It certainly will be beneficial to use as many evaluation methods as possible for characterizing the carburized alloy. With respect to the impact of carburization on an alloy’s performance, Krikke et al. (Ref 28) believed that not only the total amount of carbon absorbed but also the maximum carbon level and the maximum carbon concentration gradient are the most important factors. Figure 5.12 illustrates three possible carbon concentration profiles, as suggested by Krikke et al. (Ref 28). They considered profile A with a steep concentration profile to be the most damaging. Heubner (Ref 29) tested various commercial alloys in H2-CH4 gas mixture (ac = 0.8) at 1000 °C (1832 °F) and observed a steep carbon concentration profile in Fe-Ni-Cr alloys and a low flat concentration profile in Ni-Cr alloys, as shown in Fig. 5.13. One Ni-Cr alloy (alloy 45TM), which contained relatively high Fe and high Si, was an exception showing a steep carbon concentration profile similar to Fe-Ni-Cr alloys such as 800H, AC66, and DS. The Fe-Ni-Cr alloys were found to have suffered more room-temperature impact toughness drop in general than Ni-Cr alloys (Ref 29). However, the relative room-temperature impact toughness loss (%) was found to increase with increasing total carbon pickup (Ref 29). For carburization, the real issue is the effect of carburization on the alloy’s mechanical properties, such as creep-rupture properties and toughness or ductility. This type of data, however, is quite limited and is inadequate as a basis for making an informed materials selection. Thus, this chapter reviews mainly the carburization data in terms of mass gain, mass of carbon absorption, carburization depth, and concentration profile of the carburized layer. When the environment is such that no protective oxide scale (e.g., Cr2O3 scale) is formed on the metal surface, carburization is controlled by diffusivity and solubility (Ref 30). The ingress of carbon will be greatly reduced when a chromium oxide scale is developed. Carburization kinetics in this case are then controlled by the diffusion of carbon through the oxide scale. Wolf and Grabke (Ref 31) demonstrated that there was no detectable solubility of carbon in Cr2O3 oxides by equilibrating the oxides with CO2-CO mixtures tagged with radioactive 14C at 1000 °C (1832 °F). Thus, carbon permeation is not possible through the perfectly dense Cr2O3 oxide layer, unless the oxide layer contains pores and fissures (Ref 31).

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alloy compared with the as-cast surface finish (Fig. 5.17). Norton and his colleagues (Ref 35–37) at Petten Laboratories have conducted a series of studies on the effects of silicon, niobium, chromium, and iron in Fe-Ni-Cr alloys, including four commercial alloys (HK-40, HP-40Nb, Type 314 SS, and alloy 800H) and three experimental alloys. Their test environments had fixed carbon activities (ac) of 0.3 and 0.8, with various oxygen potentials (pO2 ) at temperatures from 825 to 1050 °C (1520 to 1920 °F), as illustrated in Fig. 5.18. The oxygen potentials of the test environments were below that in equilibrium with Cr2O3 (Fig. 5.18). That means that no chromium oxide scales should have formed on the metal surface. However, a SiO2 scale was likely to form at 825 °C (1520 °F), but not at 1000 °C (1830 °F). The test results at 825 °C (1520 °F) showed that Type 314 stainless steel (2.04% Si) was significantly more carburization

Resistance to carburization is an important factor in the performance of pyrolysis furnace tubes as well as pigtails for ethylene and olefin plants. Furnace tubes are typically constructed of Fe-Ni-Cr cast alloys, such as HK (Fe-25Cr-20Ni or 25/20), HP (Fe-25Cr-35Ni or 25/35) and their variants. Some of these variants involved additions of niobium (or columbium), tungsten, molybdenum, silicon, and titanium. Some also involved increases in nickel and/or chromium. Some of the modified alloys are referred to as “microalloyed” castings. It has been found that these additions and increases improve carburization resistance as well as creep-rupture strengths. Figures 5.14 to 5.17 illustrate the carburization resistance of some of these modified alloys compared with HK alloy (Ref 32–34). Also shown in Fig. 5.17 is the effect of the surface finish on carburization resistance of the alloy. The machined-finished surface significantly reduced the carbon ingress into the 3.5 AC 66

Carbon concentration, %

3 2.5

alloy 800 H

2

alloy DS

1.5 1 0.5 0 0

1

2

3

4

5

∞

5

∞

Distance to surface, mm

(a) 2.5

45 TM

Carbon concentration, %

2

1.5

1

alloy 600H

0.5 alloy 617 alloy 602 CA

0 0

1

2

3

4

Distance to surface, mm (b)

Fig. 5.13

Carbon concentration profile for alloys after testing at 1000 °C (1832 °F) for 1008 h in a H2-CH4 mixture (ac = 0.8) for (a) Fe-Ni-Cr and (b) Ni-Cr alloys. Source: Ref 29

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resistant than HK-40 (1.35% Si), HP-40Nb (1.29% Si), and alloy 800H (0.4% Si), as shown in Fig. 5.19. Also revealed by the test results was that the model alloys were significantly less resistant to carburization than the commercial alloys with the same chromium, nickel, and iron. The model alloys had much lower silicon levels (0.13–0.28%) as well as manganese, aluminum, titanium, and so forth. The beneficial effect of silicon on carburization resistance was clearly demonstrated when the data were replotted (Fig. 5.20). The presence of a SiO2 scale was confirmed by Van der Biest et al. (Ref 36). When the test temperature was increased to 1000 °C (1832 °F), where the pO2 was below that in equilibrium of SiO2 (Fig. 5.18), Type 314 SS was found to be similar to HK-40, alloy 800H, and HP-40Nb (Fig. 5.21). Under these conditions, SiO2 was no longer thermodynamically stable. Thus, the silicon effect was diminished. The model alloy 50/50 (50Ni-50Cr with very low silicon level) was among the best performers in the alloys tested. The ratio of Ni to Cr + Fe is an important factor (Ref 35) in governing carburization resistance, as shown in Fig. 5.22. Decreases in Cr+ Fe in Fe-Ni-Cr alloys improved carburization resistance. Both chromium and iron are carbide

formers, constituting the major elements in M7C3 and M23C6 carbides resulting from carburization. Harrison et al. (Ref 37) found that the surface carbides removed from HK-40 sample after testing at 1000 °C (1830 °F) contained 55 to 58% Cr and 41 to 43% Fe, with very little nickel (approximately Cr4Fe3C3). Nickel reduces the diffusivity of carbon in Fe-Ni-Cr alloys, as demonstrated by Demel et al. (Ref 38) in Fig. 5.23. Nickel also decreases the solubility of carbon in Fe-Ni alloy system as shown in Fig. 5.24 (Ref 39). Decreases in carbon diffusivity and solubility can result in increases in carburization resistance. Grabke et al. (Ref 40) observed that increasing nickel improved carburization resistance in Fe-Ni-Cr alloys, with the maximum resistance achieved when the ratio of Ni to Fe was 4 to 1. This is in general agreement with the product of carbon solubility and diffusivity (Ref 41). High-nickel alloys are generally more resistant to carburization than Fe-Ni-Cr alloys. This is illustrated by the test results of Klower and Heubner (Ref 29), as shown in Fig. 5.25. The beneficial effect of nickel on carburization resistance can also be clearly revealed in Fig. 5.26 when the data of Fe-Ni-Cr alloys

Fig. 5.14

Fig. 5.15

Carburization resistance of HK (25Cr-20Ni) and several HP alloys (Cr/Ni) as a function of temperature in pack carburization tests. Source: Ref 32

Carbon concentration profiles for HK (25Cr-20Ni) and several HP alloys (Cr/Ni) carburized at 1100 °C (2010 °F) for 520 h in pack carburization tests. Source: Ref 32

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and Ni-Cr alloys generated at 1000 °C for 1008 h were plotted as a function of iron content (i.e., decreasing nickel content) (Ref 29). The oxygen partial pressure (pO2 ) of the test environment—which was 10−27.92, 10−25.25, and 10−23.88 for 850, 1000, and 1100 °C, respectively—was below the pO2 in equilibrium of Cr2O3 (Ref 29). Accordingly, no chromium oxide scale would be present to provide protection for the alloys in these tests. In an extensive study undertaken by Steel and Engel (Ref 42) on the relative influence of nickel and chromium on the carburization resistance of Fe-Ni-Cr alloys, standard heats of ASTM grades varying from HC to HX cast alloys were investigated, along with many experimental cast alloys. They found nickel to be beneficial, as illustrated in Fig. 5.27. The role of chromium, however, appeared to be different for different levels of nickel, as shown in Fig. 5.28. For ironbase alloys with 25% or less nickel, increasing chromium significantly reduced carbon pickup. A slight decrease in carbon pickup with increasing chromium was noted for alloys

Fig. 5.16

containing 26 to 45% Ni. For alloys containing 46 to 70% Ni, increasing chromium resulted in an increase in carbon pickup. In this study, although there is no mention of the oxygen potential of the test environment, it is believed that the chromium oxide scale was not involved in the carburization reaction. Wolfe (Ref 43) also found that a higher Ni-containing Fe-Cr-Ni cast alloy, alloy HU (Fe-19Cr-39Ni-0.5C), was significantly better than Type 304 (19Cr-9Ni) and Type 321 (18Cr12Ni) with similar amounts of chromium but lower nickel. Even a low Cr-containing Fe-Cr-Ni cast alloy, alloy HT (Fe-15Cr-35Ni-0.5C) was significantly better than both Types 304 and 321 (Ref 43). These data are shown in Fig. 5.29. Small additions of some minor elements, such as titanium, niobium, tungsten, and rare earth elements, may also improve an alloy’s resistance to carburization in test environments of H2-8.6%CH4-7%H2O and H2-12%CH4-10% H2O. This is illustrated in Table 5.1 (Ref 44). The superior carburization resistance of TMA 4750 alloy to HK-40 (2% Si) can be attributed to small additions of titanium, niobium, tungsten,

Carbon concentration profiles of HK and HP alloys tested at 1050 °C (1920 °F) for 1200 h in 37%N2-40%H2-20%CO-3% CH4 (ac = 1.0, pO2 = 3:4 · 10720 atm). Source: Ref 33

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and rare earth elements. In both test environments, the oxygen potentials, although not discussed, are believed to be high enough to form oxide scales. Thus, improvements in carburization resistance may be partly due to the improved oxide scale. Silicon has also been found to be very effective in improving carburization resistance. The beneficial effect of silicon on carburization resistance has been reported by Norton and Barnes (Ref 35), Steinkusch (Ref 45), Wolfe (Ref 46), and Van den Bruck and Schillmoller (Ref 47), and the data are illustrated in Fig. 5.19, and Fig. 5.30 to 5.32. Kane (Ref 48) investigated a large number of centrifugally cast tubes of HK (Fe-25Cr-20Ni) and HP (Fe-25Cr-35Ni) alloys from four producers. The tubes contained silicon varying from about 1% to more than 2%. Tests were conducted at 980 and 1090 °C (1800 and 2000 °F). A unit carbon activity (ac = 1.0) was maintained for all the test environments. Oxygen potentials of the test environments were varied by injecting different levels of H2O. With no H2O injection, the environment’s pO2 was below

that in equilibrium with SiO2 (i.e., SiO2 could not form). Thus, silicon played no role in carburization resistance. At high oxygen potentials (i.e., 1% and 10% H2O injections) where SiO2 was stable, silicon improved carburization resistance. The focus thus far has been primarily on alloys used in ethylene cracking and steam hydrocarbon reforming operations. Most of the alloys are cast alloys used for furnace tubes. A variety of wrought alloys of stainless steels, Fe-Ni-Cr

Fig. 5.18

Oxygen potentials of the test environments used by Norton and his colleagues in carburization studies at Petten Laboratories. Source: Ref 37

Fig. 5.17

Carbon concentration profiles of several centrifugally cast alloys in (a) the as-cast surface condition and (b) the machined surface condition after 1 year of field testing in an ethylene cracking furnace. Source: Ref 34

Fig. 5.19

Carburization rate constants of several Fe-Ni-Cr alloys at 825 °C (1520 °F) in the test environment with a carbon activity of 0.8 and an oxygen potential such that SiO2 is stable (but not Cr2O3), as shown in Fig. 5.18. Source: Ref 35

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Fig. 5.20

Carburization rate constants as a function of silicon content in the alloy for several Fe-Ni-Cr alloys tested at 825 °C (1520 °F) in the test environment with a carbon activity of 0.8 and an oxygen potential such that SiO2 is stable (but not Cr2O3), as shown in Fig. 5.18. Source: Ref 35

Fig. 5.22

Carburization rate constants of several Fe-Ni-Cr alloys at 1000 °C (1830 °F) in the test environment with a carbon activity of 0.8 and an oxygen potential such that SiO2 is not stable as shown in Fig. 5.18. Source: Ref 35

Carburization rate constants as a function of Ni to Cr + Fe ratio [Ni/(Cr + Fe)] for several Fe-Ni-Cr alloys tested at 1000 °C (1830 °F) in the test environment with a carbon activity of 0.8 and an oxygen potential such that SiO2 is not stable as shown in Fig. 5.18. Source: Ref 35

alloys, and Ni-Cr alloys have been widely used in various industries, including heat treating and chemical processing. Mason et al. (Ref 49) investigated various stainless steels by performing pack carburization tests. Their results are summarized in Table 5.2. Silicon was again

noted for its beneficial effect, as illustrated by Type 330 (0.47% Si) versus Type 330 (1.0% Si) and Type 304 (0.39% Si) versus Type 302B (2.54% Si). Chromium was found to be beneficial in Fe-Cr alloys, as shown by Type 446 (27% Cr) versus 430 (16% Cr). Small additions

Fig. 5.21

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Fig. 5.23

Effect of nickel content on the diffusion coefficient of carbon in Fe-15Cr-Ni alloys. Source: Ref 38

of titanium or niobium appeared to be beneficial when comparing Type 321 and Type 347 to Type 304. Several nickel-base alloys along with HK-40 were investigated by Kane and Hosier (Ref 50). Tests were conducted in environments with unit carbon activity and various oxygen potentials. Different rankings were obtained at different oxygen potentials. Test results for two environments are summarized in Table 5.3. In the test environment of H2-12%CH4-10%H2O with 1.3 ×10−20 atm of pO2 , where SiO2 was thermodynamically stable, HK-40 (1.19% Si) performed the best. When pO2 was reduced to 1.9 × 10−24 atm in H2-0.1%C7H13OH, where SiO2 was not thermodynamically stable, and the carbon

activity was maintained at unity, alloys containing aluminum, such as alloys 601 and 617, were much more resistant to carburization than HK-40. The authors (Ref 50) attributed this to the formation of an Al2O3 scale, although alloys 601 and 617 contain only about 1.3% Al. Other investigators (Ref 24, 51) also showed beneficial effect of silicon in 25Cr-30/35Ni type alloys. Nishiyama et al. (Ref 24) performed their study of carburization in a simulated ethylene pyrolysis environment (H2-15%CH4-3%CO2) for Fe-Ni-Cr alloys with high Si (1.7%) and low Si (0.3–0.5%). When tested at 1000 °C, where Cr2O3 was stable (see Fig. 5.9), chromium, not silicon, was more effective in improving carburization resistance as shown in Fig. 5.33

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pg 112


of alloy 214 compared with those of alloys X, 601, and 150 after exposure to the test environment at 980 °C (1800 °F) for 55 h (Ref 52). The

Mass change, g/m2

× alloy 602CA

100

×

50

200

0

alloy 617 ∗ alloy 601 ∗ 45TM

∗

∗

∗

0

400

(a)

Mass change, g/m2

× alloy 800H AC66

×

150

600

800

1000

1200

Time, h 350 300 250 200 150 100 50 0

× AC66 alloy 800H × × • • alloy DS • ×• • alloy 600H × × ∗ alloy 601 ∗ × ∗ 45TM ∗ ∗ alloy 602CA alloy 617

×• 0

(b)

Mass change, g/m2

(Ref 24). Alloy B (26Cr-35Ni-0.5Si) was significantly better than alloy D (21Cr-31Ni-0.3Si), while alloy B (26Cr-35Ni-0.5Si) was similar to alloy A (25Cr-37Ni-1.8Si). At 1150 °C, where Cr2O3 oxide was not stable, high-Si alloys (alloy A: 25Cr-37Ni-1.8Si, alloy C: 32Cr-43Ni1.7Si) exhibited significantly better carburization resistant than low-Si alloys (alloy B: 26Cr-35Ni0.5Si, alloy D: 21Cr-31Ni-0.3Si), as shown in Fig. 5.34. Aluminum is the most effective alloying element in improving an alloy’s carburization resistance at high temperatures. Lai (Ref 52) showed that when tested at 870, 930, and 980 °C (1600, 1700, and 1800 °F) alloy 214 (Ni-16Cr3Fe-4.5Al-Y) was the most carburizationresistant alloy among more than 20 commercial wrought alloys, ranging from stainless steels and Fe-Ni-Cr alloys to nickel- and cobalt-base superalloys in the test environments, which were characterized by a unit carbon activity and oxygen potentials such that Cr2O3 was not expected to form on the metal surface. Oxides of silicon, titanium, and aluminum were expected to be stable under the test conditions. The excellent carburization resistance of this alloy was attributed to the Al2O3 oxide scale formed on the metal surface. Figure 5.35 illustrates the microstructure

200

400

600

800 1000 1200 1400

Time, h

320

×

240

× ×

160 80

×

0

∗• 0

× ∗

∗

∗ • 200

•

• 400

(c)

∗ 45TM

∗

∗

• alloy DS

•

•

600

× AC66

×

alloy 601

alloy 617 alloy 602CA alloy 600H

800

1000

1200

Time, h

Fig. 5.25

Carbon pick-up after 1000 h of exposure at 1000 °C g/cm2

Results of carburization tests in H2-CH4 mixtures (ac = 0.8) at (a) 850 °C, (b) 1000 °C, and (c) 1100 °C for Fe-Ni-Cr alloys (800H, AC66, and DS) and nickel-base alloys (alloys 600H, 601, 602CA, 617, and 45TM). Source: Ref 29

400 AC66 300

•

200

alloy 600H

•

•

45TM

•

•

•

alloy 800H

•

•

alloy DS

HPM

alloy 625 alloy 601

• alloy 602CA •alloy 617

100 0

0

10

20

30

40

50

Fe, %

Fig. 5.24

Carbon solubility as a function of nickel content at different carbon activities in Fe-Ni alloy system at 1000 °C (1830 °F). Source: Ref 39

Fig. 5.26

Weight gain as a function of iron content (i.e., a function of nickel) in Fe-Ni-Cr and Ni-Cr alloys tested at 1000 °C for 1008 h in a H2-CH4 mixture (ac = 0.8 and 725:25 pO2 =10 ). Source: Ref 29


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Fig. 5.27

Effect of nickel content on the carburization resistance of Fe-Ni-Cr alloys. Source: Ref 42

Fig. 5.29

Fig. 5.28

Carburization resistance of several wrought and cast Fe-Cr-Ni alloys (Type 304, 321, HT, HU, and HK) after testing in dry ethane (C2H6) for 24 h at temperatures from 880 to 1000 °C. Source: Ref 43

existence of this oxide scale was confirmed by Auger analysis (Ref 52). Similar test results were also observed by Lai et al. (Ref 53) when tested in

H2-2%CH4 at 982 °C (1800 °F) for 96 h. The carbon activity for the test environment at the test temperature was greater than 1.0 (see Fig. 5.11). Their results are summarized in Fig. 5.36. Alloy

Effect of chromium on the carburization resistance of Fe-Ni-Cr alloys after testing in H2-2.5%CH4 at 1050 °C (1920 °F) for 100 h. Source: Ref 42

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pg 114


214 showed no evidence of carburization, while Type 310 and alloys 800H, 25-35NbMA, 25-35Nb, 803, and 602CA showed different degrees of carburization attack. Alloy 602CA with about 2% Al (Ni-25Cr-10Fe-2Al-Y-Zr), although not as good as alloy 214, was significantly better than other alloys. A similar observation was also made by Kane et al. (Ref 54), showing that the alumina-forming MA956 Table 5.1 Weight gain (mg/cm2) for several cast alloys after 100 h at 1090 °C (2000 °F) in H2-CH4-H2O mixtures Weight gain, mg/cm2 Alloy(a)

HK-40 (1% Si) HK-40 (2% Si) TMA-4750 HP-45 TMA-6350

H2-8.6CH4-7H2O

H2-12CH4-10H2O

25.0 16.8 2.0 19.0 3.8

21.8 10.2 1.0 4.3 2.3

Table 5.2 Results of pack carburization tests at 980 °C (1800 °F)(a) for various stainless steels Nominal composition

Si content, %

Increase in C content(b), %

Fe-21Cr-34Ni Fe-15Cr-35Ni Fe-15Cr-35Ni-Si Fe-25Cr-20Ni Fe-25Cr-20Ni-Si Fe-25Cr-12Ni Fe-18Cr-8Ni-Nb Fe-18Cr-8Ni-Ti Fe-18Cr-8Ni Fe-18Cr-8Ni-Si Fe-28Cr Fe-16Cr

0.34 0.47 1.00 0.38 2.25 0.25 0.74 0.49 0.39 2.54 0.34 0.36

0.04 0.23 0.08 0.02 0.03 0.12 0.57 0.59 1.40 0.22 0.07 1.03

Alloy

800 330 330 310 314 309 347 321 304 302B 446 430

(a) 40 cycles of 25 h each cycle at 980 °C (1800 °F). Carburizer was renewed after each cycle. (b) Bulk analysis. Source: Ref 49

(a) HK-40 (1% Si): 0.43C-0.60Mn-0.96Si-25.4Cr-20.7Ni. HK-40 (2% Si): 0.41C0.60Mn-1.98Si-25.0Cr-20.7Ni. TMA-4750: 0.44C-0.69Mn-1.99Si-24.9Cr20.8Ni-0.11Ti-0.29Nb-0.30W-REM. HP-45: 0.51C-0.54Mn-1.65Si-25.5Cr36.1Ni. TMA-6350: 0.50C-0.70Mn-1.84Si-25.1Cr-38.4Ni-0.13Ti-0.28Nb-0.27WREM. REM denotes rare earth metals. Source: Ref 44

Fig. 5.31

Effect of silicon on the carburization resistance of cast Fe-20Ni-Cr alloys tested at 1090 °C (2000 °F) for 24 h in wet ethane (C2H6). Source: Ref 46

Fig. 5.30 Ref 45

Effect of silicon on the carburization resistance of 25Cr-20Ni and 35Cr-25Ni-Nb alloys. Source:

Fig. 5.32

Effect of silicon on the carburization resistance of HK-40 with different silicon levels tested at 1100 °C (2010 °F) for 520 h in carbon granulate (pack carburization test). Source: Ref 47


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alloy (Fe-20Cr-4.5Al-0.5Y2O3) performed significantly better than alloys 601, 800H, Type 310, and HK alloy (Table 5.4). Aluminum, which has been used for alloy addition to form external Al2O3 scale in ironand nickel-base wrought alloys, has been used in centrifugally cast alloys for oxidation or carburization resistance. Recently a commercial, centrifugally cast nickel-base alumina-forming alloy, alloy 60HT, containing approximately 25% Cr, 11% Fe, 0.4% C, and Al, was developed (Ref 55). In the paper published by Kirchheiner et al. (Ref 55), no carburization data were reported. However, the resistance to coking was studied on alloy 60HT containing three levels of aluminum (i.e., 2.35, 3.55, and 4.81%). They found significant reduction in coking rates for the samples containing 3.55 and 4.81% Al, with the sample containing 2.35% Al exhibiting only slight reduction of coking rates compared with the conventional HP-40 alloy (4852). Coking, which is an important phenomenon in ethylene cracking, develops on the internal surface of the pyrolysis tube and reduces the heat transfer. The ethylene cracking operation has to be interrupted Table 5.3 Weight gain (mg/cm2) for several Fe-Ni-Cr and Ni-base alloys after ten 24 h cycles at 1100 °C (2010 °F) in H2-12%CH4-10%H2O and H2-0.1%C7H13OH environments Weight gain, mg/cm2 Alloy

HK-40 (1.19% Si) 800H 600 617 601 690

H2-12CH4-10H2O(a)

H2-0.1C7H13OH(b)

6.97 23.46 12.00 16.84 22.74 25.54

54.38 46.60 4.85 0.50 1.81 54.38

(a) Inlet gas mixture: ac = 1.0, pO2 ¼ 1:3 · 10720 atm at 1100 °C. (b) Inlet gas mixture: ac = 1.0, pO2 =1:9 · 10724 atm at 1100 °C. Source: Ref 50

Fig. 5.33

Carbon profiles of high-Si (HSi) and low-Si (LSi) Fe-Ni-Cr alloys after testing at 1000 °C (1832 °F) for 96 h in H2-15%CH4-3%CO. Source: Ref 24

to allow decoking, typically with steam and air to remove the coked layer. This decoking operation can have detrimental effects on the tube properties, thus reducing the tube’s service life. Other factors such as surface finish have been found to be very important in affecting carburization reactions. Machining the metal surface to improve the surface finish can significantly increase an alloy’s carburization resistance. It is common practice to bore or hone the internal diameter of a centrifugally cast tube to remove surface shrinkage pores. Figures 5.17, 5.37, and 5.38 illustrate the significant improvement in carburization resistance as a result of surface machining (Ref 34, 56). A cast metal surface with shrinkage pores can generate stagnant conditions in crevices, which are very conducive to carburization attack. In addition, a machined surface exhibits a cold-worked layer, which tends to accelerate the diffusion process and results in rapid formation of oxide scale or film, thus slowing subsequent carbon ingress. Compared to the machined or ground surface, the electropolished surface exhibited accelerated carburization (Fig. 5.39). Norton and Barnes (Ref 35) considered the electropolished surface a work-free surface that failed to develop a surface oxide film as readily as the cold-worked ground surface. Not mentioned by Norton and Barnes in their paper (Ref 35), electropolishing can produce a surface layer that may be depleted in chromium, and the surface depletion of chromium may result in the observed accelerated carburization. In their study of oxidation/ carburization of a high-temperature gas-cooled reactor (HTGR) helium environment containing parts per million (ppm) levels of H2, CO, CO2, CH4, and H2O, Mazandarany and Lai (Ref 57) observed that Type 316SS specimen (obtained

Fig. 5.34

Carbon profiles of high-Si (HSi) and low-Si (LSi) Fe-Ni-Cr alloys after testing at 1150 °C (2100 °F) for 48 h in H2-15%CH4-3%CO. Source: Ref 24

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from sheet material) with the as-received surface condition suffered severe carburization after testing at 649 °C (1200 °F) for 5000 h in He-1500 µatm H2-450 µatm CO-50 µatm CH4-50 µatm H2O. The 316 specimen with the as-ground specimen, on the other hand, showed no evidence of carburization tested in the same retort under the same test condition. This observation was supported by both microstructure and microhardness profile results, as shown in Fig. 5.40 (Ref 57). The oxygen potentials of the test environment were reported to be high enough to form chromium oxides and too low to form iron and nickel oxides. The unusual

Fig. 5.35

observation for the 316 specimen with the asreceived surface condition after testing at 649 °C (1200 °F) in He-1500 µatm H2-450 µatm CO-50 µatm CH4-50 µatm H2O was the formation of an Fe-Ni metallic scale on the specimen surface and internal oxides underneath the metallic scale. This is shown in Fig. 5.41 (Ref 57). The 316SS specimen with the as-ground surface condition after testing under the same condition showed surface oxide scales with no outer metallic scale and internal oxides. The authors thus believed that the initial surface of the 316SS sheet was depleted in chromium at a significant degree, such that not enough chromium was available

Optical micrographs showing the microstructures of (a) alumina-former alloy 214 (Ni-16Cr-3Fe-4.5Al-Y), and several chromia-former alloys (b) 601 (Ni-23Cr-14Fe-1.4Al), (c) X (Ni-22Cr-18Fe-9Mo), and (d) 150 (Co-27Cr-18Fe) after testing at 980 °C (1800 °F) for 55 h in Ar-5%H2-5%CO-5%CH4 (ac = 1.0,pO2 =9 · 10722 atm). Source: Ref 52

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Fig. 5.36

Weight gain as a function of cycles (24 h cycle) tested at 982 °C (1800 °F) for 96 h with 24 h cycles in H2-2%CH4 for Type 310SS, 800H (N08810), 25-35Nb (25Cr-35Ni-1.3Nb-0.4C-2.0max Si), 25-35NbMA (microalloyed, 25Cr-35Ni-1.1Nb0.4C-1.4Si, Ti, REM), 803 (S35045: Fe-25Cr-35Ni), 602CA (N06025: Ni-25Cr-10Fe-2.1Al-Y-Zr), and 214 (N07214: Ni-16Cr-3Fe-4.5AlY). Note: All four alloy 800H (N08810) specimens were control samples. Source: Ref 53

Table 5.4 Carburization resistance of Fe-Ni-Cr, Ni-base, and ODS alloys in H2-2%CH4 at 1000 °C (1830 °F) for 100 h Alloy

MA956 601 800 310SS HK (1.07% Si) HK (2.54% Si)

Weight gain, mg/cm2

0.5) to avoid decarburization during testing, while asreceived and pre-aged specimens were tested in

air. Precarburization was carried out at 1000 °C (1832 °F) for 1200 h in a H2-CH4 mixture with ac being 0.8, which resulted in a fully carburized condition containing only M7C3 (about 30 vol %). Partially carburized specimens were carburized at 1000 °C (1832 °F) for 100 h in a H2-CH4 mixture with ac being 0.8, which produced a carburized depth of about 1 mm (equivalent to 56% of cross-sectional area being carburized). Pre-aging treatment consisted of heating the specimen at 1000 °C (1832 °F) for 100 h in air to produce a comparable structure to the uncarburized core of the partially carburized specimens. Their creep-rupture data are summarized in Fig. 5.52. Fully carburized specimens were found to exhibit significantly higher creeprupture strengths than the as-received specimens tested in air. Partially carburized or pre-aged specimens showed strengths comparable to those of the as-received specimens. As for rupture ductility, the fully carburized specimens showed lower rupture ductility for shorter rupture times (at high stresses), but comparable to those of as-received specimens and pre-aged as well as partially carburized specimens for longer rupture times (i.e., at lower stresses). This is illustrated in Fig. 5.53. The authors also investigated the effect of the temperature on creep-rupture ductility of alloy 800H that was fully carburized. The alloy’s rupture ductility (when fully carburized) was found to decrease significantly with decreasing temperature. The carburized specimen failed in a completely brittle manner on loading when tested at 600 °C (1112 °F). This is illustrated in Fig. 5.54.

Fig. 5.50

Fig. 5.51

Comparative creep curves of HK-40 tested at 1000 °C (1832 °F) and 15 MPa in air and H2-1% CH4 (ac = 0.8). Source: Ref 63

1% creep strengths of HK-40 and HK-30 tested at 1000 °C (1832 °F) in air, H2-1%CH4 (ac = 0.8), and for precarburized specimens tested in H2-1%CH4. Source: Ref 63

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5.3 Metal Dusting 5.3.1 Metal Dusting—Thermodynamic Considerations Metal dusting occurs in carburizing atmospheres at intermediate temperatures (approximately between 430 and 900 °C (800 and 1650 °F) with the maximum rate of attack occurring at about 600 to 700 °C (1112 to 1292 °F), depending on alloys and environments. It is now well understood that metal dusting occurs in environments that exhibit highcarbon activities (i.e., ac > 1) (Ref 66–75). As discussed in section 5.2.1, three chemical

Fig. 5.52

Stress rupture data of precarburized (fully carburized and partially carburized) specimens tested at 800 °C (1472 °F) in a carburizing environment (to prevent decarburization) is compared with that of as-received specimens and pre-aged specimens tested in air at the same temperature. Source: Ref 65

Fig. 5.53

Comparison rupture ductility of alloy 800H tested at 800 °C (1472 °F) for fully and partially carburized specimens in comparison with as-received and pre-aged specimens. Source: Ref 65

reactions (Eq 5.1–5.3) that can cause carburization can also cause metal dusting attack. Metal dusting most often occurs in syngas environments, which consist of primarily H2 and CO. In the earlier discussion of Eq 5.1: H2 + CO = C + H2O, where graphical presentation of the chemical reaction is presented as a function of carbon activity (ac), temperature, and the gas composition in terms of (pCO pH2 = pH2 O ), as shown in Fig. 5.1. Figure 5.1 illustrates that the carbon activity of the environment decreases with increasing temperature. Thus, at high temperatures, many gaseous compositions are not thermodynamically favored for causing metal dusting. Conversely, lower temperatures with increasing carbon activity are thermodynamically favored for causing metal dusting. This is why metal dusting attack diminishes as the temperature increases after passing the peak reaction temperature due to decreasing carbon activity. For example, for the gas mixture with (pCO pH2 =pH2 O ) ratio being 1.0, its carbon activity varies from more than 102 to approximately 10−1 as the temperature increases from 400 to 800 °C (752 to 1472 °F). Metal dusting attack diminishes as the temperature decreases after dropping below the peak temperature. This is due to the reaction’s kinetics that decreases with decreasing temperature. The reaction Eq 5.2, as graphically presented in Fig. 5.2, shows a similar fashion. When the environment consists of CO and CO2, metal dusting attack follows that involves H2 and CO. However, when the environment consists of H2, CO, and CO2, the Boudouard reaction (CO-CO2

Fig. 5.54

Effect of temperature on rupture ductility of fully carburized alloy 800H tested at different temperatures with stresses to cause rupture in 200 h. Source: Ref 65

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reaction) is believed to be slower in reaction kinetics than that of Eq 5.1, thus making Eq 5.1 to be the dominant reaction for the metal dusting (Ref 76, 77). The reaction 5.3 (Eq 5.3), on the other hand, increases the carbon activity with increasing temperature. The graphic presentation of this reaction is shown in Fig. 5.3. The carbon activity is generally too low for many environments containing CH4 where metal dusting occurs at 600 or 700 °C (1112 or 1292 °F). The presence of CH4 in the gas mixture containing CO and CO2 may help bring up the carbon activity at higher temperatures, thus increasing the upper temperature limit at which metal dusting can take place.

ac (gas) CO + H2 H2O + C (diss.) ac″ (Fe3C1Fe)

ac=1

(a)

Fe3C

C (diss.) + Fe

CO + H2 H2O + C (in Fe3C) (b)

CO + H2 H2O + C (graphite)

5.3.2 Mechanisms of Metal Dusting Attack Significant advancement in understanding the mechanisms of metal dusting has been achieved in the past decade, with most of the contribution being made by Professor Grabke and his group (Ref 68–75). Two models have been proposed with one model explaining the metal dusting behavior of iron and low-alloy steels and the other model explaining the behavior of Crcontaining high-temperature alloys (i.e., Fe-Cr, Fe-Ni-Cr, and Ni-Cr alloys). These two models, which are summarized by Grabke (Ref 73), are briefly described here. The model for iron and low-alloy steels is summarized in Fig. 5.55. Based on the model, the prerequisite for metal dusting to occur is a high carbon activity environment (ac > 1) that causes the oversaturated metal with the carbon activity being more than 1. This condition promotes the formation of cementite (Fe3C) on the metal surface. As shown in Fig. 5.1, the calculated carbon activities in equilibrium with cementite are higher than 1.0 at 700 °C (1292 °F) and lower temperatures. Figure 5.56 shows a metastable Fe-C-O phase diagram at 600 °C indicating that metastable cementite forms at ac > 1 (Ref 78). Once graphite deposits on the cementite with the carbon activity of the graphite at 1, cementite in contact with graphite becomes unstable and then decomposes into iron and carbon with iron migrating outward into the graphite layer. As a result, metal wastage takes place as a repeated reaction of formation of metastable cementite and decomposition of cementite into iron and carbon. For iron and low-alloy steels with no protective oxide scales, metal wastage attack in general follows uniform

(c)

CO + H2 H2O + C (graphite)

(d)

3 Fe + C

Fe3C

Fig. 5.55

Schematic showing mechanism of metal dusting for iron and low alloy steels with the following steps: (a) The metal is oversaturated with carbon (ac > 1) due to carbon transfer from a high carbon activity environment (ac > 1) to the metal, (b) thus resulting in the formation of cementite (Fe3C) at the surface, (c) and later the formation of graphite on top of cementite and then decreases in ac to one (ac = 1 for graphite) at the cementite in contact with graphite, (d) thus, cementite becomes unstable and decomposes into iron and carbon, with iron migrating into graphite layer to form iron particles (embedded in graphite), which act as catalysts for more carbon deposition and coking (d). Source: Ref 73

Fig. 5.56

Fe-C-O metastable phase stability diagram showing that metastable cementite (Fe3C) forms under high carbon activities (ac > 1) at 700 °C and lower. Source: Ref 78

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metal thinning (Ref 73). A layer of cementite (Fe3C) was found to occur on 1Cr-0.5Mo steel after exposure in metal dusting conditions with a layer of coke on top of the cementite (Ref 71). For chromia formers in Fe-Cr, Fe-Ni-Cr, and Ni-Cr alloys, a different mechanism was proposed, which was summarized in Fig. 5.57 (Ref 73). These high-chromium alloys form a protective chromium oxide scale. Metal dusting attack is initiated when the oxide scale develops local defects, allowing carbon transfer from a high carbon activity environment (ac > 1) to metal causing oversaturation of carbon, and subsequent formation of graphite after formation of metastable M3C carbides in low nickel alloys

Cr2O3 alloy

5.3.3 Alloy Resistance to Metal Dusting Graphite

(a) (d)

C

Internal carbides

3M+C

and C

M3C

'Coke'

(b)

(e) C

(or direct graphite formation in high nickel alloys), thus resulting in lower carbon activity (ac = 1) and decomposition of metastable carbides into metal particles and carbon. As a result, the attack starts locally and often leads to the formation of hemispherical pits (Ref 73). This type of morphology is shown in Fig. 5.58. Pippel et al. (Ref 79) found graphite on cementite that formed on Fe-5Ni alloy, when exposed to metal dusting conditions. For Fe-10Ni and Fe-25Ni to Fe-30Ni alloys, graphite was found to grow into metal with no metastable carbide formed on the alloy (Ref 79). Graphite was also observed to grow into metal directly in high-nickel Fe-Ni alloys, such as Fe-40Ni, Fe-50Ni, and Fe-80Ni (Ref 72).

Process equipment failures due to metal dusting were reported in the refinery industry during the 1950s. Eberle and Wylie (Ref 80) reported metal wastage of uncooled components, such as soot blower elements, made of Types 347SS and 310SS in the waste heat boiler of a synthesis gas reactor. The synthesis gas, predominantly CO and H2 with some water vapor and carbon particles, was produced by combustion of methane with oxygen. The wastage took place at temperatures between 480 and 900 °C (900 and 1650 °F). Both 347SS and 310SS suffered severe metal wastage after only 3 weeks of service. Prange (Ref 8) reported that tubes containing a chromia-alumina catalyst at a temperature of about 590 °C (1100 °F) in a butane dehydrogenation system, where butane was converted to butene, suffered metal loss problems. The oxide

Metastable carbide

(c)

Fig. 5.57

Mechanism of metal dusting for chromia formers with the following steps: (a) Development of local defects in the oxide scale, allowing carbon transfer from the environment to the metal, (b) carbon ingress into the metal resulting in the formation of stable carbides, M23C6 and M7C3, (c) continued carbon ingress increases the carbon activity to more than one (ac > 1), resulting in the formation of metastable M3C carbides in low-nickel alloys, or resulting in direct growth of graphite in high-nickel alloys, (d) graphite deposition occurs decreasing the carbon activity to one, thus, causing the decomposition of metastable M3C into metal particles and carbon with metal particles catalyzing more coking (e). Source: Ref 73

1cm

Fig. 5.58

Metal dusting attack in alloy 800 in synthesis gas environment at 550 °C (1022 °F). Source: Ref 75


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dusts resulting from metal dusting contaminated the catalyst and caused undesirable side effects for the operation. The alloys reported to perform poorly were 12Cr steel, 18Cr-11Ni, and alloy 600. The alloys that performed well included Fe-27Cr, 25Cr-20Ni, and Type 302B (18Cr-8Ni with 2.4% Si). Severe metal loss in the form of pitting was also observed by Hoyt and Caughey (Ref 81) for Type 310SS equipment exposed to a gas mixture rich in H2 and CO at temperatures of 650 to 700 °C (1200 to 1300 °F) in a plant that converted coal to gasoline and other products. Metal dusting problems were also encountered in reforming plants (Ref 5), where a synthesis gas (i.e., H2 + CO) for methanol manufacturing was produced. Type 304SS and 310SS reformer outlet tubes were perforated by severe pitting attack at 650 to 725 °C (1200 to 1340 °F). Hydrodealkylation units (Ref 5), acetic acid cracking furnaces (Ref 5), and coal gasification plants (Ref 6) were reported to suffer metal dusting problems. Perkins et al. (Ref 6) reported pitting attack of alloy 800 tubes in a preheater for gasifier recycle gas rich in H2 and CO with some H2O. The tubes, with a wall thickness of 0.38 cm (0.15 in.), were perforated in a few thousand hours. The attack occurred at 540 to 870 °C (1000 to 1600 °F). Dunmore (Ref 82) reported a failure of the waste heat boiler in an ammonia plant. The alloy 800 exit ferrules suffered severe metal wastage, with large uniform circular pits penetrating through the tubes and black sooty deposits on the pitted surface. The attack occurred at a location where the metal temperature was about 600 °C (1110 °F). The environment was highly enriched in H2 and CO, along with large quantities of steam. This was unusual in that steam was considered an effective additive to mitigate the metal dusting problem (Ref 7). Grabke and Spiegel (Ref 83) discussed several industrial failure cases that were caused by metal dusting. These cases involved an alloy 800 heat exchanger for synthesis gas production, HK-40 gas heaters in a direct iron reduction plant, 5Cr and 9Cr steels in catalyst regeneration units in a refinery, and an alloy 601 heat exchanger in an ammonia plant. Metal dusting has also been encountered in the heat treating industry (Ref 84). Refractory anchors, fan housing assemblies, and other components in carburizing furnaces frequently suffer metal dusting problems. Alloys typically used include stainless steels, such as Type 310, nickel-base alloys, such as alloys X and 333, and iron-base alloys, such as Multimet alloy

(or N-155). Metal dusting typically occurred at temperatures between 540 and 820 °C (1000 and 1500 °F). Severe attack frequently took place after 1 year of service and at locations where the gaseous environment became stagnant. Favorite locations included the interface with refractories, where small gaps or “dead” spaces were created. Figure 5.59 (Ref 84) shows a sample of Multimet alloy obtained from a furnace fan housing in a carburizing furnace. The fan housing suffered metal dusting at the metal

1

2

3

4

Cm

Fig. 5.59

Multimet alloy fan box suffering metal dusting attack in a carburizing furnace. (a) General view of failed the sample. Note the perforated edge of the fan box. (b) Cross section of the sample showing pitting attack and severe metal thinning. (c) Severe carburization attack beneath the pitted and wasted area. Attack was initiated in the “dead” spaces created by the refractory and the fan box. Source: Ref 84


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surface that was in contact with the furnace refractory lining. The 1.9 cm (0.75 in.) thick fan housing was perforated in about a year. The nearby metal surface that was exposed to the circulating carburizing gas, on the other hand, showed no sign of metal dusting (Fig. 5.60). A chromium-rich oxide scale was observed on the metal surface. One common feature associated with metal dusting is carburization beneath the pitted area (Fig. 5.59). Manufacturing of carbon fibers for carbon composite materials can also generate high temperature carburizing environments for causing metal dusting problems to some furnace components. Carbon fibers are manufactured from polyacrylonitrile (PAN) (Ref 85). The last step of manufacturing, “carbonization,” is carried out at about 900 °C (1650 °F) in an environment enriched with CO, CO2, CH4, N2, HCN, H2O, and so forth. Furnace components made of nickel- and iron-base alloys were found to suffer general metal thinning and pitting attack. An example is shown in Fig. 5.61. The general characteristics of metal dusting observed in the industrial environments were reproduced in many laboratory tests involving primarily gas mixtures of H2-CO-H2O. Accordingly, laboratory testing has become an important tool for generating relevant data to allow performance ranking among engineering alloys. Laboratory data generated by various laboratories are summarized below. Maier and Norton (Ref 86) investigated 9Cr-1Mo steel (P91), 12Cr steel (Type 410),

Fig. 5.60

Oxide scale and no evidence of metal dusting on the surface of a Multimet alloy fan box exposed to flowing carburizing gas (same fan box as that shown in Fig. 5.59). The sample was plated with nickel before mounting for metallographic examination.

Model alloy (Type 410 + 2.75Si), and Type 310 (Fe-25Cr-20Ni) in H2-24.4CO-2.4H2O at 560 °C (ac > 1). The specimens were annealed in H2 at 1000 °C (1832 °F) for 1 h to eliminate surface deformation structure and produced grain coarsening. The test specimens were ground. Their results are summarized in Fig. 5.62. In this short-term test, alloys with low chromium contents (9Cr for P91 and 12Cr for 410SS) showed metal dusting attack. Type 310SS (25Cr-20Ni) showed no sign of metal dusting attack up to 200 h. Silicon was found to be a very effective alloying element in improving the resistance to metal dusting. This is illustrated by the data comparing “model” alloy (410SS + 2.75Si) with 410SS. Grabke et al. (Ref 69) investigated commercial alloys, which included Fe-Cr, Fe-Ni-Cr, Ni-base alloys, and silicon-containing alloys, at 650 °C (1200 °F) for 7 days in H2-24.7CO-1.9H2O (ac = 15). All specimens were annealed in dry hydrogen at 1000 °C (1832 °F) for 1 h to achieve large grains, followed by grinding and polishing. All test specimens were in as-ground surface conditions. The data are summarized in Fig. 5.63. The

Fig. 5.61

Type 310SS furnace component suffering metal dusting in a furnace used for manufacturing carbon fibers. (a) General view of the failed component. (b) Cross section of the sample showing pitting and thinning

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hatched data show mass gain of coke deposits on the test specimen, while the solid data represent the either weight loss due to metal dusting or weight gain due to oxidation and/or carbon gain (no metal dusting attack). Sicromal (Fe-18Cr1Al-1Si) contained not very high chromium level, but showed no metal dusting. This was due

to the presence of some aluminum and silicon in the alloy. Both aluminum and silicon are effective alloying additions in improving the metal dusting resistance. X18 CrN28 (Fe-28Cr) showed significantly better performance than Sanicro 28 (Fe-27Cr-30Ni). Both alloys contained about the same level of chromium, but one

250 560 °C, 1.5 bar

Mean metal loss, µm

200

150

100

50

0 0

50

100

150

200

Exposure time, h

Fig. 5.62

Average metal loss as a function of exposure time (hours) for P91, 410, 310, and model alloy (410 + 2.75Si). Source: Ref 86

1000 100 10

Mass gain, mg/cm

2

1 0.1 0.01 0.001 –0.01 –0.1 –1 –10

Removable deposits

Fig. 5.63

Inconel 600

AC 66

Sanicro 28

HK 40

X15 CrNiSi 2012

253 MA

12 CrMoV

Alloy 410

X18 CrN28

Sicromal

–100

Cleaned specimen

Mass gain (hatched data) due to coke deposits and metal loss of specimen (solid data) due to metal dusting as well as mass gain of specimen (solid data) due to possible oxidation/carburization after exposure at 650 °C (1200 °F) in H2-24.7CO-1.9H2O (ac = 15). X18 CrN28 (Fe-28Cr); Sanicro 28 (Fe-27Cr-30Ni); 253MA (Fe-21Cr-11Ni-2Si-0.05Ce); X15 CrNiSi 2012 (Fe-20Cr-12Ni-2Si). Source: Ref 69


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Fe25Cr2.5Ni 0

Fe60Cr

Mass change, mg/mm2

–0.02

Fe25Cr

–0.04 –0.06 Fe25Cr5Ni

–0.08 –0.1 –0.12 –0.14

Fe25Cr25Ni

–0.15 Fe25Cr10Ni –0.18 0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 160 170 180 190 200 210 Exposure cycles, h

Fig. 5.64

Weight loss of specimens after removing carbon deposits as a function of thermal cycles for Fe-Cr and Fe-Cr-Ni alloys in metal dusting tests at 680 °C (1256 °F) in H2-68%CO-6%H2O (ac = 2.9 and pO2 =2 · 10723 atm) cycling specimens to room temperature every 60 min in the test environment. Source: Ref 90

was a ferritic alloy (Fe-Cr) while the other was an austenitic alloy (Fe-Ni-Cr). The diffusivity of chromium in ferritic alloys is generally about two orders of magnitude higher than that in austenitic Fe-Ni-Cr alloys (Ref 87–89). As a result, ferritic stainless steels form chromium oxide scales much more readily than austenitic stainless steels, as was observed in Fig. 5.63. Both 12Cr steels (410SS and 12CrMoV) showed some metal dusting attack. Again, silicon was found to greatly improve metal dusting resistance. Both alloy 253MA (Fe-21Cr-11Ni-2Si-0.05Ce) and X15 CrNiSi 2012 (Fe-20Cr-12Ni-2Si) showed no metal dusting attack. Another siliconcontaining alloy, HK-40 (Fe-25Cr-20N-2Si), showed some metal dusting attack. The alloys that showed the worst attack in this test were Fe-Ni-Cr alloys (Sanicro 28 and AC66), and a low-chromium nickel alloy (alloy 600). The results of Toh et al. (Ref 90) further confirmed that Fe-Cr alloys exhibit better metal dusting resistance than Fe-Ni-Cr alloys. They investigated Fe-25Cr alloy and Fe-25Cr with 5, 10, and 25% Ni. All alloys were annealed in Ar-10%H2 at 1050 °C (1922 °F) for 100 h. Test specimens were then cut, ground, and polished to 3 µm before being electrolytically polished. Testing was conducted at 680 °C (1256 °F) in H2-68%CO-6%H2O (ac = 2.9 and pO2 =2 · 10723 atm). Specimens were cycled by heating them to the test temperature for 60 min for each cycle. Their test results are summarized in Fig. 5.64. Fe-25Cr alloy was found to be significantly more resistant to metal dusting than Fe-25Cr alloys with 5, 10, and 25% Ni. Fe-25Cr

with 2.5% Ni appeared to perform well. It was likely that the alloy with only 2.5% Ni remained a ferritic structure. The test also included Fe-60Cr, which showed good metal dusting resistance. Ferritic Fe-Cr alloys are more resistant to metal dusting than austenitic Fe-Ni-Cr alloys presumably due to much faster chromium diffusion rates, thus resulting in faster formation of chromium oxide scales, and then better metal dusting resistance. Rapid diffusion of chromium to the metal surface to form a protective chromium oxide scale is important in retarding metal dusting attack. Accordingly, for the same alloy, a fine-grained material can form a surface oxide scale more readily than a coarse-grained material, thus better metal dusting resistance. This is illustrated in Fig. 5.65, where the fine-grained Type 304SS (as-received from the supplier) with a grain size of ASTM No. 10 (average size of 10 µm) showed no metal dusting attack while the coarse-grained material with a grain size of about ASTM No. 3 to 7 (about 30–100 µm) suffered severe attack (Ref 91). The coarsegrained material was obtained by annealing the sample to 1000 °C (1832 °F) for 1 h. Metal dusting testing was conducted at 600 °C (1112 °F) in H2-24CO-2H2O. More discussion on the effects of surface conditions on the metal dusting behavior is presented later. Fe-Ni-Cr alloys, which showed less resistance to metal dusting than Fe-Cr alloys due to lower chromium diffusivity, are also much less resistance to metal dusting as compared with Ni-base alloys. This is illustrated in Fig. 5.66, which were generated on the commercial alloys without prior

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annealing treatments. Specimens were in the asreceived condition from the supplier and were ground to a 120-grit finish prior to testing in H2-90%CO for 672 days at 482, 566, 649, and 732 °C (900, 1050, 1200, and 1350 °F) (Ref 92).

10 18Cr-8Ni-Steel

Mass gain, mg/cm2

8

6 Coarse grain 4

2 Fine grain 0

0

50

100

150

200

Time, h

Fig. 5.65

Effect of grain size on the metal dusting behavior of Type 304SS tested at 600 °C (1112 °F) in H224CO-2H2O. Source: Ref 91

Fig. 5.66

Type 304, 310, and 800H—which are Fe-Ni-Cr alloys—were found to suffer metal dusting attack (i.e., the reaction rates in weight loss) while two nickel-base alloys (alloy 601 and RA333) showed no metal dusting attack (i.e., the reaction rates in weight gain). The figure also shows that alloy 85H (Fe-18.5Cr-14.5Ni-3.5Si-1Al) showed no metal dusting attack. This was due to beneficial effect of silicon along with aluminum in the alloy. Beneficial effects of silicon and aluminum are demonstrated by additions of these two elements to alloy 800 (Fe-20Cr-32Ni), as shown in Fig. 5.67 (Ref 93). Also in this study by Strauss and Grabke (Ref 93), other alloying elements, such as niobium, molybdenum, and tungsten, were found to have some beneficial effects as well. Klower et al. (Ref 94) investigated a large number of commercial alloys including several Fe-Ni-Cr alloys, such as alloys 800H, HK-40, HP-40, and DS, and many nickel-base alloys. Test results are summarized in Table 5.5. Fe-Ni-Cr alloys were much less resistant to metal dusting than nickel-base alloys. High-siliconcontaining alloy DS (Fe-18Cr-35Ni-2.2Si) was much better than other Fe-Ni-Cr alloys. This is nicely illustrated in Fig. 5.68. In Ni-Cr-Fe

Metal dusting resistance of several Fe-Ni-Cr alloys and Ni-base alloys tested in H2-90%CO for 672 days at 482, 566, 649, and 732 °C (900, 1050, 1200, and 1350 °F). The alloys tested were Type 304 (S30403), Type 310 (S31000), alloy 85H (S30615), alloy 800H (N08810), alloy 601 (N06601), and RA333 (N06333). Source: Ref 92

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system, increases in chromium along with aluminum additions can significantly improve the alloy’s metal dusting resistance, as illustrated in Fig. 5.69. Alloy 601 was more resistant than alloy 600 because of higher chromium and aluminum addition in alloy 601. Alloy 602CA was better than alloy 601 because of higher aluminum content. There were two different surface finishes involved in this test, with alloys 600 and 601 in as ground surface, and alloy 602CA in “black” surface finish. There was no discussion about the procedure for preparing the “black” surface finish in the paper. If the 602CA specimens were black annealed (heat treated in air), with about 20 +4.5%AI

+2.5%Si + +

Mass loss, mg/cm2

+3%Mo +3%Nb

–20

+3%W

Fe-32%Ni-20%Cr –60

+0.05%Ce –100

0

500

Fig. 5.67

1000 Time, h

1500

Effects of Al, Si, Nb, Mo, and W additions to a model alloy (alloy 800, Fe-20Cr-32Ni) on the metal dusting behavior of modified experimental alloys. Source: Ref 93

2.3% Al in alloy 602CA, alloy 602CA specimens might have preformed an Cr/Al-rich oxide scales prior to metal dusting testing. The effects of surface finishes are discussed further. In Table 5.5, it is surprising to find that alloy 214 with about 4.5% Al did not perform well compared with 602CA (2.3% Al), 617 (1.3% Al), and 690 (30% Cr, no Al). It is believed that with only about 16% Cr, alloy 214 did not contain enough chromium to rapidly develop a protective chromium oxide scale at such a low temperature. And, at this low temperature, development of an exclusive Al2O3 is difficult in short time. Thus, a Table 5.5 Metal wastage rates after exposure to H2-CO-H2O gas mixtures at 650 °C (1200 °F) for various commercial alloys Alloy

Surface condition

Total time, h

Wastage rates, mg/cm2h

800H HK-40 HP-40 DS 600H 601 601 601 C-4 214 HR160 45TM 602CA 617 690

Ground … … Ground Ground Black Polished Ground Ground Ground Ground Black Black Ground Ground

95 190 190 1,988 5,000 6,697 1,988 10,000 10,000 9,665 9,665 10,000 10,000 7,000 10,000

0.21 0.04 0.038 4.3 × 10−3 3.3 × 10−2 7.3 × 10−3 4.9 × 10−3 5.8 × 10−4 1.1 × 10−3 1.2 × 10−3 6.3 × 10−4 1.0 × 10−5 1.1 × 10−5 3.7 × 10−6 2.0 × 10−6

Note: Gas mixtures H2-24CO-2H2O (ac = 14 at temperature) for the first 5000 h of testing and H2-49CO-2H2O for the subsequent 5000 h. Source: Ref 94

102

Metal wastage rate, mg/cm2h

10 1 10–1 10–2 10–3 10–4 10–5

Fig. 5.68

Metal dusting resistance of Fe-Ni-Cr alloys (800H and HP40) in comparison with Ni-base alloy 600H tested at 650 °C (1200 °F) in H2-24CO-2H2O (ac = 14). Source: Ref 94


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pg 133


high-chromium nickel-base alloy, alloy 690 (30% Cr), performed the best in this test. It is also surprising to find that HR-160 alloy with both high Cr (28%) and high Si (2.75%) did not perform as well as some other nickel-base alloys. Alloy 45TM, also with high Cr (27%) and high Si (2.7%) performed better than HR-160 alloy. The

alloy 45TM specimen was in a “black” surface condition. The “black” surface condition might be resulted from a preoxidation treatment that might have produced chromium- and silicon-rich oxide scales prior to metal dusting testing. In conducting metal dusting testing that involved only the specimens with as-ground

10–1

Metal wastage rate, mg/cm2h

Alloy 600 10–2 Alloy 601 10–3

10–4

Alloy 602CA

10–5

10–6 0

2,000

4,000

6,000

8,000

10,000

Exposure time, h

Fig. 5.69

Increasing Cr along with addition of aluminum improves the metal dusting resistance of the alloy in Ni-Cr-Fe alloys. Source: Ref 94

Fig. 5.70

Metal dusting behavior of various Ni-base alloys tested at 593 °C (1100 °F) in H2-18.4CO-5.7CO2-22.5H2O at 14.3 atm of pressure. Source: Ref 95


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surface finish, Natesan and Zeng (Ref 95) performed tests at 593 °C (1100 °F) in H2-18.4CO5.7CO2-22.5H2O at 14.3 atm of pressure. Alloy 45TM suffered rapid metal dusting attack with about 59.1 mg/cm2 of weight loss after only 3300 h while HR-160 alloy exhibited only 7.3 mg/cm2 of weight loss after 9700 h. Their results are summarized in Fig. 5.70. The bestperforming alloy was alloy 693 (30% Cr, 3.3Al), followed by alloy 602CA (25% Cr, 2.3% Al), and HR-160 (28% Cr, 2.75% Si) and alloy 690 (30% Cr). The worst performer was alloy 45TM, followed by alloys 617, 214, and 601. These tests conducted in Argonne National Laboratory were under a high pressure (14.3 atm, or 210 psi), and all other data generated in other laboratories were under gas pressure close to 1 atm (atmosphere pressure). Natesan and Zeng (Ref 95) observed that high pressure could significantly reduce the time to initiate metal dusting attack for some alloys. Tests were performed at 593 °C (1100 °F) for 246 h in 1, 14.3, and 40.8 atm pressures. Specimens were then examined for signs of metal dusting attack. Their results are summarized in Table 5.6, showing alloys 601, 690, 617, and 214 suffering metal dusting attack at high pressures but not at 1 atm pressure after 246 h. On the other hand, alloys 45TM, 602CA, and HR160 showed no metal dusting attack at both low and high pressures after 246 h. They also measured the maximum pit size and average pit depth in addition to weight loss. This is illustrated in Table 5.7. In terms of the pitting depth, HR160 was found to perform the best, while alloy 690 with little weight loss showed significant pit depth. The data suggested that the total weight loss was not correlated well with pitting depth. Alloy 214 was found to show uniform metal wastage. This was apparently due to insufficient chromium content (16%) at such a low temperature to form a continuous Cr2O3

scale. The Raman spectra of the alloy 214 tested specimen showed low intensity for the Cr2O3 band (Ref 95). In manufacturing of sheet products (or thin gage tubular products) in nickel-base alloys, the manufacturing process in the later stage typically consists of cold working (cold pilgering) and bright annealing (i.e., annealing in H2 atmosphere). The hydrogen atmosphere in bright annealing typically contains some moisture, thus it is typically characterized with a certain dew point (i.e., a certain oxygen potential, pO2 ). In general, annealing furnaces exhibit low enough dew points such that chromium oxide scales do not form on the metal surface when processing stainless steels or nickel-base alloys. Sheet products of these alloys typically look shiny. However, the dew points or oxygen potentials in hydrogen-annealing furnaces generally are high enough that Al2O3 or SiO2 will form on the metal surface of a bright-annealed sheet of a nickel alloy containing a sufficient level of aluminum or silicon. Klarstrom and Grabke (Ref 96) tested bright-annealed sheet specimens of alloy 214 (alumina former) and alloy HR160 (silica former) along with alloy 230 and HR120 in the bright-annealed condition. Also included in testing were black annealed and pickled sheet specimens of alloys 800H and 601 (no brightannealed sheet samples were available for these two alloys at the time of testing). No surface grinding for bright-annealed sheet specimens (alloys 214, HR160, 230, and HR120). Surface grinding to a 120-grit surface finish was performed for black-annealed and pickled sheet specimens of alloys 800H and 601. Testing was performed at 650 °C (1200 °F) for up to 10,000 hours in H2-49CO-2H2O (ac =18.9). Test results are summarized in Table 5.8. HR160 alloy was found to perform very well, showing no evidence of metal dusting after 10,000 h

Table 5.6 Surface conditions of alloys after testing at 593 °C (1100 °F) for 246 h in H2-18.4CO-5.7CO2-22.5H2O at 1, 14.3, and 40.8 atm pressures

Table 5.7 Pit size, pit depth, and weight loss for alloys after testing at 593 °C (1100 °F) for 9700 h in H2-18.4CO-5.7CO2-22.5H2O at 14.3 atm pressure

Condition at pressure Alloy

601 690 617 602CA 214 45TM HR160 Source: Ref 95

Alloy

1 atm

14.3 atm

40.8 atm

Clean surface Clean surface Clean surface Clean surface Clean surface Clean surface Clean surface

Pits Pits Pits Clean surface Pits Clean surface Clean surface

Pits Pits Pits Clean surface Pits Clean surface Clean surface

601 690 617 602CA 214 45TM(b) HR160 693

Weight loss, mg/cm2

Pit depth, µm

Pit diameter, µm

19.5 6.5 35.1 2.1 25.6 59.1 7.3 0.1

110 147 201 96 (a) 141 13 37

450 440 887 374 (a) 600 210 99

(a) Specimen uniformly corroded. (b) Exposed for only 3300 h. Source: Ref 95


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of exposure. The alloy, however, showed slight carburization underneath the metal surface. Alloy 230 showed slight metal dusting attack with a few spots 1 to 2 mm in diameter on each side of the specimen after 10,000 h of exposure. Alloy 601 suffered significantly more metal dusting attack than alloy 230. Typical surface conditions for these three alloys after 10,000 h of testing are shown in Fig. 5.71. Alloy 214, however, did not perform well, suffering metal Table 5.8 Final metal wastage rates for various alloys tested at 650 °C (1200 °F) for up to 10,000 h in H2-49CO-2H2O (ac = 18.9) Alloy

HR120 800H 214 601 230 HR160

Total exposure time, h

Wastage rate, mg/cm2h

190 925 5,707 10,000 10,000 10,000

4.1 × 10−2 2.7 × 10−3 1.0 × 10−3 2.5 × 10−3 3.2 × 10−4 0.0(a)

dusting attack with fairly high wastage rates. Both Fe-Ni-Cr alloys, alloys 800H, and HR120, performed the worst among the alloys tested. Figure 5.72 shows typical surface conditions of alloys 214, HR120, and 800H after 5707 h, 190 h, and 925 h, respectively. Baker and Smith (Ref 97, 98) and Baker et al. (Ref 99) investigated about 20 commercial alloys in H2-80CO at 621 °C (1150 °F) for times up to 16,000 h. All materials were cold rolled and annealed sheets (i.e., bright annealed products) except alloys K-500 and 617, which were hotrolled, annealed plates (i.e., black annealed and pickled products), and alloy DS, which was extruded and annealed tubing. Nevertheless, test coupons were all ground to a 120-grit finish. For exposure times less than 10,000 h, weight changes as a function of exposure time for Fe-Ni-Cr alloys as well as 9Cr steel and Monel in comparison with some nickel-base alloys are

(a) Attack too small for analysis. Source: Ref 96

Fig. 5.71

Optical micrographs showing typical surface conditions for (a) HR160 alloy, (b) alloy 230, and (c) alloy 601 after testing at 650 °C (1200 °F) for 10,000 h in H249CO-2H2O (ac = 18.9). Magnification bar represents 20 µm for (a), 200 µm for (b) and (c). Source: Ref 96

Fig. 5.72

Typical surface conditions for (a) alloy 214 after 5707 h, (b) alloy HR120 after 190 h, and (c) alloy 800H after 925 h at 650 °C (1200 °F) in H2-49CO-2H2O (ac = 18.9). Source: Ref 96

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summarized in Fig. 5.73. Nickel-base alloys except alloy 600 performed much better than Fe-Ni-Cr alloys. Nickel-base alloy 600 performed poorly because of its low chromium content (15–16%). MA956 (Fe-20Cr-4.5AlY2O3 ODS alloy) performed very well. The test results are also presented in Fig. 5.74 in terms of

maximum pitting depth as a function of exposure time. Alloys 690, 617, and MA956 showed the minimum pitting attack. Figure 5.75 shows the test results for some additional alloys and some better performing alloys tested up to 16,000 hours (Ref 99). Alloys 690, 263, and MA956 were observed to show pits with

20 263

617 0 825

864

MA 956

690 601

600 –40 9Cr-1Mo

–60

K-500

Mass change, mg/cm2

–20

DS 803

800

–80

–100 330 –120 0

1000

2000

3000

4000

5000

6000

7000

8000

9000

10000

Exposure time, h

Fig. 5.73

Mass change as a function of exposure time for various alloys including 9Cr steel, Ni-Cr alloy (Monel K-500), Fe-Ni-Cr alloys, and Ni-base alloys tested at 621 °C (1150 °F) in H2-80CO. All surfaces were ground to a 120-grit finish prior to testing. Source: Ref 98

762

30 800 803

601

o -1M

20

508

330

15

381 864

10

254

5

82 5

690 617

600

127

Maximum pitting depth, µm

635

DS

9Cr

Maximum pitting depth, mils

25

MA956 0

0 K500 –5 0

1000

2000

3000

4000

5000

6000

7000

8000

–127 9000

Exposure time, h

Fig. 5.74

Maximum pit depths as a function of exposure time for various alloys including 9Cr steel,Fe-Ni-Cr alloys and Ni-base alloys tested at 621 °C (1150 °F) in H2-80CO. All surfaces were ground to a 120-grit finish prior to testing. Source: Ref 98


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25 24 23 22 21 20 19 18 17 16 15 14 13 12 11 10 9 8 7 6 5 4 3 2 1 0 –1

800

than that of alloy 693 (Fig. 5.76). It should be cautioned in interpreting pit depth data, when the specimen surface has shown general wastage without the original specimen surface surrounding the pit, pitting depth data thus measured then becomes questionable. This would particularly be the case after very long exposure when pits have spread throughout the specimen surface. In reviewing the data summarized in Fig. 5.75 and 5.76, one alloy of interest was alloy 671 (Ni-46Cr), which was found to exhibit excellent metal dusting resistance. Unfortunately, the test on this alloy was terminated at close to 10,000 h. This alloy exhibited low pit depths and very low metal loss rates (lower than alloy 693). Grabke et al. (Ref 100) also found that Ni-50Cr alloy was very resistant to metal dusting, showing a thin chromium oxide scale (about 3 µm thick) after testing in H2-49CO-2H2O at 650 °C (1200 °F) (ac = 18.9) for 10,000 h with no evidence of metal dusting or carburization. Only minute amounts of coke were observed on the specimen surface. Under this test condition, alloy 601 had already suffered metal dusting attack after only 1993 h of exposure. Similarly, a very thin chromium oxide scale was observed in

DS

610

803

559 601

508

33

0

956 263

457

690

356 305

4

254

86 5

82

203

TD 758 LCE

625

152

C–276

671

102

754

617

693

600 0

Pit depth, µm

406

602CA

400

Pit depth, mils

significant pit depths after 14,000–16,000 h of exposure. Alloys 617 and 693, on the other hand, still exhibited pits with insignificant pit depths. Alloy 693 was a recently developed commercial alloy targeting for metal dusting environments by adding about 3% Al to alloy 690 with same amount of Cr (about 30%) but slightly lowered Fe. The alloy was found to perform much better than alloy 690. It is understandable that high Cr and high Al in nickel-base alloys should perform well in metal dusting environments. However, it was surprising to find that alloy 617 with only 22% Cr (quite normal level chromium for hightemperature alloys) and fairly low aluminum content (about 1.3%) exhibited a pitting depth similar to that of alloy 693. In metal dusting testing by Natesan and Zeng (Ref 95) involving several nickel-base alloys including alloy 617, the test results, which were presented as weight loss, showed alloy 617 was worse than alloy 601, 602CA, and several other nickel-base alloys (Fig. 5.70). When Baker and Smith (Ref 98) presented their test results in terms of weight loss rate as a function of exposure time, alloy 617 was found to exhibit metal loss rate (in the range of alloys 263 and MA956) significantly higher

51 0

2000

4000

6000

8000

10000

12000

14000

16000

18000

Exposure time, h

Fig. 5.75

Maximum pit depths as a function of exposure time for various alloys including 9Cr steel, Fe-Ni-Cr alloys, and Ni-base alloys tested at 621 °C (1150 °F) in H2-80CO for exposure time up to 16,000 h. All surfaces were ground to a 120-grit finish prior to testing. Source: Ref 99


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pg 138


0 80

6 0 82 00 5

10–1

33

40

0

100

1

80

3

4 75

864 DS

60

Mass loss rate, mg/cm2h

10–2

625

10–3

690

C 276

58

263

602CA

7

MA

TD

617

956

10–4

10–5 693 671 0

2000

4000

6000

8000

10000

12000

14000

16000

18000

Exposure time, h

Fig. 5.76

Metal loss rate as a function of the exposure time for various alloys tested at 621 °C (1150 °F) in H2-80CO for exposure time up to 16,000 h. All surfaces were ground to a 120-grit finish prior to testing. Source: Ref 99

Cr-5Fe-1Y2O3 after 10,000 h of exposure in the same test, showing no evidence of metal dusting attack. 5.3.4 Effects of Surface Conditions and Finish From the discussion so far on metal dusting, formation of a good, protective chromium oxide scale is a very effective way to provide protection against metal dusting attack. Since the temperature range for metal dusting attack is quite low and chromium diffusivities are relatively low at these temperatures, rapid formation of a good, protective chromium oxide scale is critical. Accordingly, the concentration of chromium at the surface of the metal becomes important. Higher chromium concentration at the metal surface provides faster formation of a protective chromium oxide scale. Thus, alloys with higher bulk chromium concentration (higher surface chromium concentration) are more resistant to metal dusting as was discussed in the previous section. Also, for the same alloy, the fine-grained structure exhibits better metal dusting resistance than the coarse-grained structure, as illustrated in

Fig. 5.65. This is because grain boundaries provide fast diffusion paths for chromium to reach to the metal surface. A fine-grain-sized material will have more grain boundaries and, thus, more chromium reaching to the metal surface to form a better chromium oxide scale faster than a coarsegrained material, thus, resisting metal dusting attack much better. The surface of the metal can also be prepared by grinding or machining to produce a thin coldworked layer with high density of dislocations, which also provide fast diffusion paths for chromium to reach to the metal surface to form a protective chromium oxide scale. As a result, metal dusting is greatly improved when the surface is ground or machined. The beneficial effect of the ground surface condition is illustrated Fig. 5.77 and 5.78 in a thermogravimetric study by Grabke et al. (Ref 101). For both alloy 800 and Type 310SS, the as-ground specimen was most resistant to metal dusting compared with the as-received surface (cold rolled) and electropolished surface. Electropolished surface condition was the worst because of possible surface depletion in chromium during electropolishing. Both alloy 800 and Type 310SS were


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tested under the same condition; Type 310 was much more resistant to metal dusting than alloy 800. This is primarily because Type 310SS contains more chromium. Specimens with ground surface finish, alloy 601, showed better metal dusting resistance than those with electropolished surface finish, as shown in Fig. 5.79 (Ref 94). In this study, the “black” sample (believed to be a black-annealed sample) performed as poorly as the electropolished

10 Alloy 800 8 Mass gain, mg/cm2

Electropolished Cold rolled 6

4

Ground

2

0

0

5

10

15

20

25

Time, h

Fig. 5.77

Thermogravimetric testing in H2-24CO-2H2O at 600 °C (1112 °F) for alloy 800 (Fe-21Cr-32Ni) in three different surface conditions: as-ground surface (to 600-grit), cold-rolled (as-received surface), and electropolished surface. Source: Ref 101

sample. The black-annealed sample, as in electropolished specimen, exhibited surface chromium depletion due to annealing in air (i.e., black annealing). The formation of chromium oxide scales during annealing in air results in some chromium depletion in the surface layer. Similar results were obtained by Baker and Smith (Ref 98) in their study on alloy 601 with different surface conditions, as shown in Fig. 5.80 and 5.81. In this study, samples from as-received sheet manufactured from blackannealing and pickling (i.e., annealing in air followed by acid pickling to remove oxide scales), which were identified as as-produced (annealed+ pickled), were slightly better than those black-annealed samples and electropolished samples. Black annealing and acid pickling can often result in surface depletion in chromium. These processes are often involved in producing sheet products by hot rolling. When sheet products are produced via cold rolling, the sheet is typically bright annealed. Under this condition, no chromium oxide scales formed on the sheet metal; thus no chromium depletion occurred at the metal surface. For highaluminum or high-silicon alloys, bright-annealing atmospheres typically exhibit high enough oxygen potentials (or dew points) that a thin aluminum or silicon oxide film can form, which can be exceedingly beneficial to the subsequent service in metal dusting environments. This was discussed earlier for HR160. Alloy 214, unfortunately with too low a chromium level, showed no beneficial effect from bright annealing in its resistance to metal dusting attack.

6

Mass gain, mg/cm2

25Cr-20Ni - Steel

5.3.5 Metal Dusting Behavior of Weldments Selection of an appropriate filler metal for metal dusting environments is also critical, since

Electropolished

4

2 Cold rolled

Ground 0 0

Fig. 5.78

5

10 15 Time, h

20

25

Thermogravimetric testing in H2-24CO-2H2O at 600 °C (1112 °F) for Type 310SS (Fe-25Cr-20Ni) in three different surface conditions: as-ground surface (to 600grit), cold-rolled (as-received surface), and electropolished surface. Source: Ref 101

Fig. 5.79

Metal dusting behavior at 650 °C (1200 °F) in H2-49CO-2H2O for alloy 601 in three different surface conditions: as-ground surface, elecropolished surface, and black-annealed surface. Source: Ref 94

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not every wrought alloy has a matching filler metal. Thus, it is necessary to select a filler metal that is at least as good as, but preferably better than, the wrought alloy selected for the application. The suitable filler metal requires not only good resistance to metal dusting but also good

weldability. Also, in many cases, machining or grinding the weld joint may not be possible, particularly in tubular butt welds where inside diameter (ID) grinding or machining is not feasible after joining. Furthermore, alloys that resist metal dusting contain high aluminum or

Ground

Fig. 5.80

Metal dusting behavior in terms of weight change tested at 621 °C (1150 °F) in H2-80CO for alloy 601 in various surface conditions. Source: Ref 98

Fig. 5.81

Metal dusting behavior in terms of weight loss rate tested at 621 °C (1150 °F) in H2-80CO for alloy 601 in various surface conditions. Source: Ref 98


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silicon along with high chromium, as discussed in previous sections. If matching filler metals are available for these alloys, the weldability of these high-aluminum or high-silicon filler metals can be an issue. Grabke et al. (Ref 101) investigated the metal dusting behavior of different weldments involving base metals of alloys 800H, 600H, 601, and 602CA using different filler metals, such as alloy 82 (Nicrofer S 7020 or UTP 068 HH) and alloy 602CA (Nicrofer 6025 or UTP 6225 Al). Weldment specimens included alloy 800H to alloy 800H with alloy 82 filler metal, alloy 600H to alloy 600H with alloy 82 filler metal, alloy 601H to alloy 601H with alloy 602CA, and alloy 800H to alloy 602CA with alloy 602CA as a cap layer and alloy 82 filler metal for root pass and filler. Welding processes involved were listed as TIG welding (i.e., gas-tungsten arc welding) and MMAW (probably shielded metal arc welding). The surface finish for most weldment specimens 1.0 no H2S

0.1 ppm H2S 0.3 ppm H2S 0.5 ppm H2S 0.75 ppm H2S

∆m/A, mg/cm2

0.8 0.6 0.4

1 ppm H2S

0.2 0.0

0

50

100

150

200

250

were “brushed,” with one weldment being ground, one being sandblasted, and one being pickled. Tests were conducted at 600 and 650 °C (1112 and 1200 °F) in H2-24CO-2H2O. In almost all cases, metal dusting was initiated at the interface between the weld metal and the base metal (i.e., the heat-affected zone surface). Alloy 82 weld metal was found to be more resistant to metal dusting than alloy 800H and alloy 600H. The authors (Ref 101) also concluded that TIG welding led to a better resistance to metal dusting than “hand-welding,” and grinding led to a modest delay in initiation of metal dusting attack. 5.3.6 Effects of Sulfur on Metal Dusting Sulfur, which has a strong tendency to segregate to the surface and grain boundaries, can act as an inhibitor to metal dusting. Hochman (Ref 67) first reported beneficial effect of sulfur against metal dusting attack. Extensive investigations on the effect of sulfur on metal dusting behavior of iron were carried out by Grabke and Muller-Lorenz (Ref 102), Schneider et al. (Ref 103), Schneider et al. (Ref 104), and Schneider and Grabke (Ref 105). The mechanism for sulfur to inhibit metal dusting attack, as proposed by Grabke and Muller-Lorenz (Ref 102), is the absorption of sulfur on the surface of cementite suppresses the nucleation of graphite, thus inhibiting metal dusting attack. Figure 5.82 shows the beneficial effect of H2S on the kinetics of metal dusting of iron at 500 °C

300

Time, h

Fig. 5.82

Effect of H2S on metal dusting behavior of iron at 500 °C (932 °F) in H2-CO-H2O-H2S gas mixture (ac = 100). Source: Ref 103

1.50

no H2S 0.7 ppm H2S 0.1 ppm H2S 1 ppm H2S 0.5 ppm H2S

∆m/A, mg/cm2

1.25 1.00 0.75

5 ppm H2S

0.50

15 ppm H2S

0.25 0.00

T=700 °C ac = 100

0

Fig. 5.83

50

100

150 200 Time, h

250

300

350

Effect of H2S on metal dusting behavior of iron at 700 °C (1292 °F) in H2-CH4-H2S gas mixture (ac = 100). Source: Ref 104

Fig. 5.84

Effect of H2S on metal dusting behavior of iron in terms of pH2 S =pH2 versus 1/T. Open data points represent that the onset of metal dusting was retarded for more than 48 h, while the solid data points represent metal dusting without retardation. The hatched region represents a transition to an iron surface saturated with sulfur. Source: Ref 74

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(932 °F) in H2-CO-H2O-H2S gas mixture (ac = 100) (Ref 103), and Fig. 5.83 shows the similar effect for iron at 700 °C (1292 °F) in H2-CH4H2S gas mixture (ac = 100) (Ref 103). Grabke (Ref 74) provides metal dusting behavior of iron in terms of pH2 S =pH2 ratio as a function of temperature (1/T) in Fig. 5.84. The figure shows the region when metal dusting was avoided by sulfur injection under certain pH2 S =pH2 ratios at different temperatures from about 500 to 1000 °C (932 to 1832 °F). Injecting a right amount of sulfur into the environment to retard metal dusting without causing accelerated sulfidation attack is a balancing act.

5.4 Summary Metals and alloys are generally susceptible to carburization when exposed to environments containing CO, CH4, or other hydrocarbon gases at elevated temperatures. Carburization typically results in the formation of internal carbides in the matrix as well as boundaries, causing the alloy to lose its room-temperature ductility and/or creep-rupture strengths. Fe-Ni-Cr alloys are widely used for processing equipment to resist carburization in the petrochemical industry. The cast 25Cr-20Ni, HK40, was once the workhorse of pyrolysis furnace tubes in ethylene cracking operations. Many modifications based on HK40 have been developed and used now with improved carburization resistance as well as increased creep-rupture strengths. These alloy modifications involve the use of alloying elements such as, titanium, niobium, tungsten, molybdenum, and silicon, as well as increases in nickel and/or chromium. Increasing nickel in Fe-Ni-Cr alloys improves carburization resistance. Nickel reduces the diffusivity of carbon in Fe-Ni-Cr alloys. Nickel also reduces the solubility of carbon in Fe-Ni alloys. Among these alloying elements, silicon is the most effective in improving carburization resistance. This is attributed to the formation of SiO2 scale, which is more impervious to carbon ingress than Cr2O3 scale. However, when the silicon content in the alloy is too high, the weldability of the alloy can become a serious issue. Aluminum is another alloying element that can significantly improve carburization resistance. However, effectiveness generally requires about 4% or higher aluminum, the amount needed to form a continuous Al2O3 scale.

Alumina formers (i.e., alloys forming Al2O3 scale, such as alloy 214 and MA956) are significantly better than chromia formers (i.e., alloys forming Cr2O3 scale). Surface finish plays an important role in carburization resistance. For cast products, machining the metal surface can significantly reduce carburization attack. Injecting sulfur compounds (e.g., 50 to 100 ppm) into the processing gas stream is also effective in reducing carburization attack. Metal dusting is another form of carburization attack; it typically causes an alloy to suffer pitting attack and/or thinning. The metal beneath the pitted area generally shows carburization. The corrosion products typically consist of carbon soots, metal particles, carbides, and oxides. The environment in which metal dusting occurs generally contains H2, CO, CO2, and H2O with high carbon activities (i.e., aC > 1). Stagnant gas conditions can be conducive in initiating metal dusting attack. The metal temperatures at which metal dusting occurs are between 430 and 900 °C (800 and 1650 °F). Metal dusting data for various commercial alloys are presented. Nickel-base alloys containing high chromium and high aluminum (e.g., alloys 602CA and 693) or containing high chromium and high silicon (e.g., alloy HR160) showed excellent resistance to metal dusting attack. Surface conditions also play an important role in metal dusting resistance. Sulfur may also retard metal dusting attack.

REFERENCES

1. Metals Handbook, 8th ed., Vol 2, American Society for Metals, 1964, p 93 2. L.J. Haga, Heat Treat., Dec 1986, p 6 3. G.E. Moller and C.W. Warren, Paper No. 237, Corrosion/81, NACE, 1981 4. A.J. McNab, Hydrocarbon Process., Dec 1987, p 43 5. G.L. Swales, in Behavior of High Temperature Alloys in Aggressive Environments, Proc. 1979 Petten International Conference, I. Kirman et al., Ed., The Metals Society, London, 1980, p 45 6. R.A. Perkins, W.C. Coons, and F.J. Radd, in Properties of High Temperature Alloys, Proc. 1976 Fall Meeting of the Electrochemical Society, The Electrochemical Society 7. R.C. Schueler, Hydrocarbon Process., Aug 1972, p 73

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8. F.A. Prange, Corrosion, Vol 15 (No. 12), Dec 1959, p 619t 9. G.Y. Lai, M.F. Rothman, and D.E. Fluck, Paper No. 14, Corrosion/85, NACE, 1985 10. O.J. Dunmore, Proc. UK Corrosion Conf. 1982, Institution of Corrosion Science and Technology, Birmingham, UK, 1982, p 101 11. F.N. Mazandarany and G.Y. Lai, Nucl. Technol., Vol 43, 1979, p 349 12. K. Natesan, Nucl. Technol., Vol 28, 1976, p 441 13. K. Natesan and T.F. Kassner, Metall. Trans., Vol 4, 1973, p 2557 14. HSC, Chemistry for Windows, Version 6.0, A. Roine, Outokumpu Technology, Finland, www.outokumputechnology.com, accessed Dec 2006 15. ChemSage, Version 4.16, GTT-Technologies, Aachen, 1998 16. P.L. Hemmings and R.A. Perkins, “Thermodynamic Phase Stability Diagrams for the Analysis of Corrosion Reactions in Coal Gasification/Combustion Atmospheres,” Report FP-539, EPRI, Palo Alto, CA, 1977 17. R. Hultgren, P. Desai, D. Hawkins, M. Gleiser, and K.K. Kelley, Selected Values of the Thermodynamic Properties of Elements and Binary Alloys, ASM, 1973 18. L.C. Browning, T.W. Dewitt, and P.H. Emmett, J. Am. Chem. Soc., Vol 72, 1950, p 4211 19. K.H. Jack, J. Iron Steel Inst., Vol 169, 1951, p 26 20. H.J. Christ, Experimental Characterization and Computer-Based Description of the Carburization Behaviour of the Austenitic Stainless Steel AISI 304L, Mater. Corros., Vol 49, 1998, p 258 21. R.G. Coltters, Mater. Sci. Eng., Vol 76, 1985, p 1 22. S.R. Shatynski, Oxid. Met., Vol 13 (No. 2), 1979, p 105 23. W.F. Chu and A. Rahmel, Oxid. Met., Vol 15, 1981, p 331 24. Y. Nishiyama, N. Otsuka, and T. Nishizawa, Carburization Resistance of Austenitic Alloys in CH4 -CO2 -H2 Gas Mixtures at Elevated Temperatures, Corrosion, Vol 59 (No. 8), 2003, p 688 25. H.J. Grabke, “Carburization: A High Temperature Corrosion Phenomenon,’’ MTI Publications No. 52, Materials Technology Institute of the Chemical Process Industries, St. Louis, Missouri, 1998

26. S. Forseth and P. Kofstad, Carburization of Fe-Ni-Cr Steels in CH4-H2 Mixtures at 850–1000 °C, Mater. Corros., Vol 49, 1998, p 266 27. D. Jakobi and R. Gommans, Typical Failures in Pyrolysis Coils for Ethylene Cracking, Mater. Corros., Vol 54 (No. 11), 2003, p 881 28. R.H. Krikke, J. Horing, and K. Smit, Mater. Perform., Aug 1976, p 9 29. J. Klower and U. Heubner, Carburization of Nickel-Base Alloys and its Effects on the Mechanical Properties, Mater. Corros., Vol 49, 1998, p 237 30. H.J. Grabke, Metallic Corrosion, Proc. Eighth Int. Cong. Metallic Corrosion, Vol III, Mainz, Germany, Sept 6–11, 1981 31. I. Wolf and H.J. Grabke, Solid State Commun., Vol 54, 1985, p 5 32. C.M. Schillmoller, Chem. Eng., Jan 6, 1986, p 87 33. D.J. Hall, M.K. Hossain, and J.J. Jones, Mater. Perform., Jan 1985, p 25 34. J.A. Thuillier, Mater. Perform., Nov 1976, p9 35. J.F. Norton and J. Barnes, in Corrosion in Fossil Fuel Systems, I.G. Wright, Ed., The Electrochemical Society, 1983, p 277 36. O. Van der Biest, J.M. Harrison, and J.F. Norton, in Behavior of High Temperature Alloys in Aggressive Environments, Proc. International Conference (Patten, The Netherlands), Oct 15–18, 1979, The Metal Society, London, 1980, p 681 37. J.M. Harrison, J.F. Norton, R.T. Derricott, and J.B. Marriott, Werkst. Korros., Vol 30, 1979, p 785 38. O. Demel, E. Keil, and P. Kostecki, SGAW Report No. 2538, Osterreichische, Studiengesellschaft fur Atomenergie, Lenaugasse 10, A-1082 Wien, Forschungszentrum Seibersdorf, Institut fur Metallurgie 39. R.P. Smith, Trans. Met. AIME, Vol 224, 1962, p 10 40. H.J. Grabke, U. Gravenhorst, and W. Steinkusch, Werkst. Korros., Vol 27, 1976, p 291 41. S.K. Bose and H.J. Grabke, Z. Metallkde., Vol 69, 1978, p 8 42. C. Steel and W. Engel, AFS Int. Cast Metals J., Sept 1981, p 28 43. L.H. Wolfe, Laboratory Investigation of High Temperature Alloy Failure Mechanisms, Paper No. 12, Corrosion/77, NACE, 1977 44. I.Y. Khandros, R.G. Bayer, and C.A. Smith, Paper No. 10, Corrosion/84, NACE, 1984

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45. W. Steinkrsch, Werkst. Korros., Vol 30, 1979, p 837 46. L.H. Wolfe, Mater. Perform., April 1978, p 38 47. U. Van den Bruck and C.M. Schillmoller, Paper No. 23, Corrosion/85, NACE, 1985 48. R.H. Kane, Paper No. 266, Corrosion/83, NACE, 1983 49. J.F. Mason, J.J. Moran, and E.N. Skinner, Corrosion, Vol 16, 1960, p 593t 50. R.H. Kane and J.C. Hosier, “Carburization Resistance of Some Wrought NickelContaining Alloys in Simulated Industrial Environments,” Inco Alloys International Technical Report, Inco Alloys International, Huntington, WV, 1985 51. D.R.G. Mitchell, D.J. Young, and W. Kleemann, Carburization of Heat-Resistant Steels, Mater. Corros., Vol 49, 1998, p 231 52. G.Y. Lai, in High Temperature Corrosion in Energy Systems, Proc. TMS-AIME Symposium, M.F. Rothman, Ed., The Metallurgical Society of AIME, 1985, p 551 53. G.Y. Lai, B. Li, B. Gleeson, and H.L. Craig, Proposed Standard Carburization Test Method, Paper No. 3473, Corrosion/ 2003, NACE International, 2003 54. R.H. Kane, G.M. McColvin, T.J. Kelly, and J.M. Davison, Paper No. 12, Corrosion/84, NACE, 1984 55. R. Kirchheiner, D.J. Young, P. Becker, and R.N. Durham, Improved Oxidation and Coking Resistance of a New Alumina Forming Alloy 60HT for the Petrochemical Industry, Paper No. 5428, Corrosion/2005, 2005 56. F. Pons and M. Hugo, Paper No. 272, Corrosion/81, NACE, 1981 57. F.N. Mazandarany and G.Y. Lai, “Corrosion Behavior of Selected Structural Materials in a Simulated Steam-Cycle HTGR Helium Environment,” GA-A14446, General Atomic Company, San Diego, CA, Oct 1977 58. T.A. Ramanarayanan, R.A. Petkovic, J.D. Mumford, and A. Ozekcin, Carburization of High Chromium Alloys, Mater. Corros., Vol 49, 1998, p 226 59. H.J. Grabke, R. Moller, and A. Schnaas, Werkst. Korr., Vol 30, 1979, p 794 60. H.J. Grabke, E.M. Peterson, and S.R. Srinivasan, Surf. Sci., Vol 67, 1977, p 501 61. T.A. Ramanarayanan and D.J. Srolovitz, Carburization Mechanisms of High Chromium Alloys, J. Electrochem. Soc.: Solid-

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77. S.R. Shatynski and H.J. Grabke, Arch. Eisenhüttenwes., Vol 49, 1978, p 129 78. F. Bonnet, F. Ropital, Y. Berthier, and P. Marcus, Filamentous Carbon Formation Caused by Catalytic Metal Particles from Iron Oxide, Mater. Corros., Vol 54, 2003, p 870 79. E. Pippel, J. Woltersdorf, and H.J. Grabke, Microprocesses of Metal Dusting on IronNickel Alloys and Their Dependence on the Alloy Composition, Mater. Corros., Vol 54, 2003, p 747 80. F. Eberle and R.D. Wylie, Corrosion, Vol 15 (No. 12), 1959, p 622t 81. W.B. Hoyt and R.H. Caughey, Corrosion, Vol 15 (No. 12), 1959, p 627t 82. O.J. Dunmore, A Case History of a Metal Dusting Problem Which Led to a Boiler Failure, presented at UK Corrosion/82 (London), Nov 16–18, 1982 83. H.J. Grabke and M. Spiegel, Mater. Corros., Vol 54, 2003, p 799 84. G.Y. Lai, J. Met., Vol 37 (No. 7), 1985, p 14 85. E. Fitzer, W. Frohs, and M. Heine, Carbon, Vol 24 (No. 4), 1986, p 387 86. M. Maier and J.F. Norton, Studies Concerned with the Metal Dusting of Fe-Cr-Ni Materials, Paper No. 75, Corrosion/99, NACE International, 1999 87. R.A. Perkins, R.A. Padgett, and N.K. Tunali, Met. Trans. AIME, Vol 4, 1973, p 2535 88. P.J. Albery and C.W. Haworth, Met. Sci., Vol 8, 1974, p 407 89. A.F. Smith, Met. Sci., Vol 9, 1975, p 375, 425 90. C.H. Toh, P.R. Munroe, and D.J. Young, Metal Dusting of Fe-Cr and Fe-Ni-Cr Alloys under Cyclic Conditions, Oxid. Met., Vol 58 (No. 1/2), 2002, p 1 91. H.J. Grabke, E.M. Muller-Lorenz, S. Strauss, E. Pippel, and J. Woltersdorf, Effects of Grain Size, Cold Working, and Surface Finish on the Metal-Dusting Resistance of Steels, Oxid. Met., Vol 50 (No 3/4), 1998, p 241 92. A.S. Fabiszewski, W.R. Warkins, J.J. Hoffman, and S.W. Dean, The Effect of Temperature and Gas Composition on the Metal Dusting Susceptibility of Various

Alloys, Paper No. 532, Corrosion/2000, NACE International, 2000 93. S. Strauss and H.J. Grabke, Mater. Corros., Vol 49, 1998, p 321 94. J. Klower, H.J. Grabke, E.M. MullerLorenz, and D.C. Agarwal, Metal Dusting and Carburization Resistance of NickelBase Alloys, Paper No. 139, Corrosion/97, NACE International, 1997 95. K. Natesan and Z. Zeng, Metal Dusting Performance of Structural Alloys, Paper No. 5409, Corrosion/2005, NACE International, 2005 96. D.L. Klarstrom and H.J. Grabke, The Metal Dusting Behavior of Several High Temperature Alloys, Paper No. 1379, Corrosion/ 2001, NACE International, 2001 97. B.A. Baker and G.D. Smith, Metal Dusting Behavior of High-Temperature Alloys, Paper No. 54, Corrosion/99, NACE International, 1999 98. B.A. Baker and G.D. Smith, Alloy Selection for Environments Which Promote Metal Dusting, Paper No. 257, Corrosion/2000, NACE International, 2000 99. B.A. Baker, G.D. Smith, V.W. Hartmann, L. E. Shoemaker, and S.A. McCoy, NickelBase Material Solutions to Metal Dusting Problems, Paper No. 2394, Corrosion/2002, NACE International, 2002 100. H.J. Grabke, H.P. Martinz, and E.M. Muller-Lorenz, Metal Dusting Resistance of High Chromium Alloys, Mater. Corros., Vol 54, 2003, p 860 101. H.J. Grabke, E.M. Muller-Lorenz, and M. Zinke, Metal Dusting Behaviour of Welded Ni-Base Alloys with Different Surface Finish, Mater. Corros., Vol 34, 2003, p 785 102. H.J. Grabke and E.M. Muller-Lorenz, Steel Res., Vol 66, 1995, p 252 103. A. Schneider, H. Viefhaus, G. Inden, H.J. Grabke, and E.M. Muller-Lorenz, Mater. Corros., Vol 49, 1998, p 330 104. A. Schneider, H. Viefhaus, and G. Inden, Surface Analytical Studies of Metal Dusting of Iron in CH4-H2-H2S Mixtures, Mater. Corros., Vol 51, 2000, p 338 105. A. Schneider and H.J. Grabke, Effect of H2S on Metal Dusting of Iron, Mater. Corros., Vol 54, 2003, p 793

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CHAPTER 6

Corrosion by Halogen and Hydrogen Halides 6.1 Introduction Many metals react readily with halogen gases at elevated temperatures to form volatile metal halides. Some metal halides also exhibit low melting points, and some even sublime at relatively low temperatures. As a result, alloys containing elements that form volatile or lowmelting-point halides may suffer severe hightemperature corrosion. Industrial environments often contain halogen gases. Because of high vapor pressures of many metal chlorides, the chlorination process is an important step in processing metallurgical ores for production of titanium, zirconium, tantalum, niobium, and tungsten (Ref 1–3). Chlorination is also used for extraction of nickel from iron laterites (Ref 4) and for detinning of tin plate (Ref 5). Production of TiO2 and SiO2 involves processing environments containing Cl2 and/or HCl, along with O2 and other combustion products. Calcining operations in the production of (a) lanthanum, cerium, and neodymium for electronic and magnetic materials and (b) ceramic ferrites for permanent magnets, frequently generate environments contaminated with chlorine. In the chemical process industry, many processing streams also contain chlorine. Manufacturing of ethylene dichloride (EDC), which is an intermediate for the production of vinyl chloride monomer, generates chlorine-bearing environments. The reactor vessels, calciners, and other processing equipment for the above operations require alloys that are resistant to high-temperature chloridation attack. In the manufacture of fluorine-containing compounds, such as fluorocarbon plastics, refrigerants, and fire-extinguishing agents, the processing equipment requires alloys with good resistance to corrosion by fluorine and hydrogen fluorides at elevated temperatures. During the

refining operation in the production of uranium, UO2 is fluorinated at elevated temperatures (e.g., 500 to 600 °C) with HF to produce UF4 or UF6 for separation of U235 (Ref 6). The materials of construction for this processing equipment must resist corrosion by HF at both high and low temperatures. This chapter reviews data on primarily gaseous corrosion by halogen and hydrogen halides. In many high-temperature industrial processes where fuels and/or feedstocks are often contaminated with impurities, such as alkaline metals, halogen may react readily with these metals to form halide salts. Corrosion reactions under these conditions are to be discussed in later chapters dealing with high-temperature corrosion in gas turbines, coal-fired boilers, oil-fired boilers, waste-to-energy boilers, black liquor recovery boilers, and so forth.

6.2 Thermodynamic Considerations The relative stabilities of various chlorides, fluorides, bromides, and iodides are presented in Fig. 6.1 to 6.3, in terms of standard free energies of formation (∆G°) versus temperature (Ref 7). The figures also include the information about the melting points and boiling points of halides. Some of the halides exhibit low melting and boiling points. Tables 6.1 to 6.4 list melting and boiling points of some important halides and oxyhalides (Ref 8, 9). Since industrial environments typically have finite oxygen activities, the thermodynamic phase stability diagram in terms of the M-O-Cl (M represents metal, O represents oxygen, and Cl represents chlorine) system is quite useful in describing the possible corrosion products that may form on the metal. These diagrams can be constructed for major alloying

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elements of high-temperature alloys (e.g., Fe, Ni, Co, Cr, Mo, W, Al, Si, etc.). Commercial computer programs, such as HSC (Ref 10) and Chemsage (Ref 11), are available for construction of these phase-stability diagrams. Figure 6.4 shows a Ni-O-Cl stability diagram in terms of pO2 and pCl2 at 723 °C (1333 °F) (Ref 12). The

Fig. 6.1

diagram defines the boundary ( pO2 10715 atm) between Ni and NiO, the boundary (pCl2 1078 ) between Ni and NiCl2, and the boundary between NiO and NiCl2. This means that if the equilibrium gas mixture of the environment exhibits an oxygen potential ( pO2 ) higher than about 10−15 atm, NiO may form. Similarly,

Standard free energies of formation for chlorides. Source: Ref 7

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Chapter 6: Corrosion by Halogen and Hydrogen Halides / 149

NiCl2 may form if the environment exhibits a chlorine potential ( pCl2 ) higher than about 10−8 atm when the oxygen potential ( pO2 ) is below about 10−15 atm. With phase-stability diagrams, one can predict the possible phases that are likely to form on

Fig. 6.2

a metal. Let’s consider an example where nickel is exposed to an environment consisting of air with 0.1% Cl2 at 723 °C (1333 °F). The environment would be at the location that identifies pO2 being very close to 100 and pCl2 being 10−3 in the 723 °C Ni-O-Cl stability diagram (Fig. 6.4).

Standard free energies of formation for fluorides. Source: Ref 7

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Based on the stability diagram shown in Fig. 6.4, the environment is in the NiO regime. NiO oxide is expected to form on the surface of nickel when exposed to this environment. If NiO oxide scale formed on nickel is defect free, the pCl2 at the interface between NiO and Ni would

Fig. 6.3

be below the value needed to form NiCl2, thus preventing the formation of NiCl2. However, when defects and cracks develop in the NiO scale, Cl2 can permeate through the oxide scale and reach the nickel with high enough pCl2 to form NiCl2, initiating chloridation attack. The

Standard free energies of formation for bromides and iodides. Source: Ref 7

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pg 151


NiCl2 formed at 723 °C (1333 °F) would be a solid phase. As shown in Table 6.1, NiCl2 melts at relatively high temperature (1030 °C) compared with other metal chlorides, such as CoCl2 (740 °C), FeCl2 (676 °C), and FeCl3 (303 °C). The phase-stability diagram for Co-O-Cl at 723 °C (1333 °F) is shown in Fig. 6.5 (Ref 12), and those for Cr-O-Cl and Fe-O-Cl at 600 °C (1112 °F) are shown in Fig. 6.6 and 6.7, respectively (Ref 13). In addition to low melting points, some chlorides have low boiling points also. In Fig. 6.7, FeCl3 is in a gaseous state at 600 °C (1112 °F). Figure 6.8 shows WCl4 and WO2Cl2 in a gaseous state at 900 °C (1650 °F). Some fluorides, bromides, and iodides also have either low melting points or low boiling points (see Tables 6.2–6.4). Corrosion reactions can be significantly increased when the corrosion product is in either a liquid state or a gaseous state. Furthermore, many halides, although in a solid state, may exhibit high vapor pressures. When the corrosion product in a solid state exhibits a

Table 6.2 Melting points, temperatures at which fluoride vapor pressure reaches 10−4 atm, and boiling points of various fluorides Fluorides


Temperature at 10−4 atm, °C (°F)

Boiling point, °C (°F)

FeF2 FeF3 NiF2 CoF2 CrF2 CrF3 CuF MoF5 MoF6 WF6 TiF3 TiF4 AlF3 SiF4 MnF2 ZrF4 NbF5 HfF4 TaF5 NaF KF LiF MgF2 CaF2 BaF2 ZnF2 PbF2

1020 (1868) 1027 (1881) 1450 (2642) 1250 (2282) 894 (1641) 1404 (2559) 908 (1666) 64 (147) 17 (63) 2 (36) 1200 (2192) … … −90 (−130) 920 (1688) 932 (1710) 79 (174) … 97 (207) 992 (1818) 857 (1575) 848 (1558) 1263 (2305) 1418 (2584) 1290 (2354) 875 (1607) 822 (1512)

906 (1663) 673 (1243) 939 (1722) 962 (1764) 928 (1702) 855 (1571) … 24 (75) −82 (−116) −91 (−132) … 108 (226) 825 (1517) −160 (−256) 992 (1818) 583 (1081) … 615 (1139) 37 (99) 928 (1702) 788 (1450) 908 (1666) 1257 (2295) 1429 (2604) 1581 (2878) 806 (1483) 664 (1227)

1800 (3272) … … … … … … … 36 (97) 17 (63) … 283 (541) 1270 (2318) −95 (−139) … … 233 (451) … … 1704 (3099) 1502 (2736) 1681 (3058) 2230 (4046) 2500 (4532) 2215 (4019) 1500 (2732) 1293 (2359)

Source: Ref 8, Ref 9

Table 6.1 Melting points, temperatures at which chloride vapor pressure reaches 10−4 atm, and boiling points of various chlorides Chlorides


Temperature at 10 −4 atm, °C (°F)


FeCl2 FeCl3 NiCl2 CoCl2 CrCl2 CrCl3 CrO2Cl2 CuCl MoCl5 WCl5 WCl6 TiCl2 TiCl3 TiCl4 AlCl3 SiCl4 MnCl2 ZrCl4 NbCl5 NbCl4 TaCl5 HfCl4 CCl4 NaCl KCl LiCl MgCl2 CaCl2 BaCl2 ZnCl2 PbCl2

676 (1249) 303 (577) 1030 (1886) 740 (1364) 820 (1508) 1150 (2102) −95 (−139) 430 (806) 194 (381) 240 (464) 280 (536) 1025 (1877) 730 (1346) −23 (−9.4) 193 (379) −70 (−94) 652 (1206) 483 (901) 205 (401) … 216 (421) 434 (813) −24 (−11) 801 (1474) 772 (1422) 610 (1130) 714 (1317) 772 (1422) 962 (1764) 318 (604) 498 (928)

536 (997) 167 (333) 607 (1125) 587 (1089) 741 (1366) 611 (1132) … 387 (729) 58 (136) 72 (162) 11 (52) 921 (1690) 454 (849) −38 (−36) 76 (169) −87 (−125) 607 (1125) 146 (295) … 239 (462) 80 (176) 132 (297) −80 (−112) 742 (1368) 706 (1303) 665 (1229) 663 (1225) 1039 (1902) … 349 (660) 484 (903)

1026 (1879) 319 (606) 987 (1809) 1025 (1877) 1300 (2372) 945 (1733) 117 (243) 1690 (3074) 268 (514) … 337 (639) … 750 (1382) 137 (279) … 58 (136) 1190 (2174) … 250 (482) 455 (851) 240 (464) … 77 (171) 1465 (2669) 1407 (2565) 1382 (2520) 1418 (2584) 2000 (3632) 1830 (3326) 732 (1350) 954 (1749)


Table 6.3 Melting points, temperatures at which bromide vapor pressure reaches 10−4 atm, and boiling points of various bromides Bromides




FeBr2 FeBr3 NiBr2 CoBr2 CrBr2 CrBr3 CrBr4 CuBr WBr5 WBr6 AlBr3 SiBr4 MnBr2 ZrBr4 NbBr5 HfBr4 TiBr4 TaBr5 NaBr KBr LiBr MgBr2 CaBr2 BaBr2 ZnBr2 PbBr2

689 (1272) … 965 (1769) 678 (1252) 842 (1548) >800 (1472) … 488 (910) 276 (529) 309 (588) 97 (207) 5 (41) 695 (1283) 450 (842) 267 (513) 424 (795) … 267 (513) 750 (1382) 740 (1364) 550 (1022) 710 (1310) 742 (1368) 854 (1569) 398 (748) 373 (703)

509 (948) 156 (313) 580 (1076) … 716 (1321) 615 (1139) 516 (961) 435 (815) … … 53 (127) … … 169 (336) … 137 (279) … 145 (293) 690 (1274) 671 (1240) 630 (1166) 626 (1159) … … 320 (608) 432 (810)

974 (1785) … 919 (1686) … … … … 1318 (2404) … … 255 (491) 153 (307) … … 361 (682) … 232 (450) 347 (657) 1393 (2539) 1383 (2521) 1310 (2390) 1230 (2246) 1800 (3272) … 650 (1202) 914 (1677)



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high vapor pressure, the corrosion reaction can increase as well. It is generally considered that when the vapor pressure of a corrosion product reaches to 10−4 atm, the corrosion reaction can become significant. Tables 6.1 to 6.4 also include the temperature at which a halide’s vapor pressure reaches to 10−4 atm. The chlorine partial pressure (pCl2 ) needed to form a chloride corrosion product with 10−4 atm will be lower than that needed to form either solid or liquid chloride. This is illustrated in Fig. 6.9 (Ref 14). Bender and Schutze (Ref 15)

Table 6.4 Melting points, temperatures at which iodide vapor pressure reaches 10−4 atm, and boiling points of various iodides Iodides




FeI2 NiI2 CoI2 CrI2 CrI3 CuI AlI3 SiI4 MnI2 ZrI4 NbI4 HfI4 TaI5 NaI KI LiI MgI2 CaI2 BaI2 ZnI2 PbI2

594 (1101) 780 (1436) 515 (959) 869 (1596) >600 (1112) 588 (1090) 191 (376) 122 (252) 613 (1135) 499 (930) 503 (937) 449 (840) 496 (925) 660 (1220) 685 (1265) 469 (876) 650 (1202) 740 (1364) 712 (1314) 446 (835) 412 (774)

476 (889) … … 702 (1296) … 529 (984) 144 (291) 55 (131) … 227 (441) … 244 (471) 208 (406) 651 (1204) 629 (1164) 621 (1150) 425 (797) … … 316 (601) 397 (747)

935 (1715) … … … … 1207 (2205) 385 (725) 301 (574) … … … … 545 (1013) 1304 (2379) 1330 (2426) 1170 (2138) … … … 730 (1346) 872 (1602)


Fig. 6.4

Phase stability diagram for Ni-O-Cl system at 723 °C (1333 °F). Both corrosion products (NiO and NiCl2) are solid phases at this temperature. Source: Ref 12

Fig. 6.5

Phase stability diagram for Co-O-Cl system at 723 °C (1333 °F). All the corrosion products (i.e., CoO, Co3O4, and CoCl2) are solid phases at this temperature. Source: Ref 12

Fig. 6.6

Phase stability diagram for Cr-O-Cl system at 600 °C (1112 °F). All the corrosion products (i.e., Cr2O3, CrCl2, and CrCl3) are solid phases at this temperature. Source: Ref 13

Fig. 6.7

Phase stability diagram for Fe-O-Cl system at 600 °C (1112 °F). All the corrosion products are solid phases except FeCl3 at this temperature. Source: Ref 13

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Fig. 6.8

Phase stability diagram for the W-O-Cl system at 900 °C (1650 °F), showing tungsten chloride (WCl4) and tungsten oxychloride (WO2Cl2) in a gaseous state.

Fig. 6.9

Phase stability diagram for Fe-O-Cl system at 700 °C (1292 °F). The solid line represents the boundary for forming solid FeCl2, while the dotted line represents the boundary for forming FeCl2 with 10−4 atm pressure. Source: Ref 14

extensively examined the phase-stability diagrams involving the vapor pressures of chlorides reaching 10−4 atm for various alloying elements. The authors termed these phase-stability diagrams “quasi-stability diagrams.” Figure 6.10 shows the pCl2 values that are needed to form NiCl2 with 10−4 atm at various temperatures (Ref 15). Also included in the figure are two environments (air + 0.1% Cl2, and air + 2% Cl2) for illustration purpose. For both of these environments, NiO, not NiCl2 with vapor pressures of 10−4 atm or higher, is likely to form on nickel at 500 and 650 °C. However, at 800 °C

and higher, NiCl2 with vapor pressures of 10−4 atm and higher (not NiO) is to form on nickel. Molybdenum oxychlorides also exhibit high vapor pressures. Figure 6.11 shows the quasistability diagram for Mo-O-Cl system at 500 °C (Ref 15). The figure also shows that MoO2Cl2 with vapor pressures of 10−4 atm and higher is to form in the environments of air + 0.1% Cl2, and air + 2% Cl2. The metal-halogen reaction differs from other reactions, such as oxidation, in that most reaction products are characteristic of high vapor pressures and, in some cases, low melting points. The volatile halides (reaction products) formed on the metal surface can no longer provide protection against further corrosion. This is in contrast to most oxides, which generally exhibit very low vapor pressures and high melting points. Furthermore, many halides exhibit low melting points. Once the reaction products become molten, the alloy loses all protection against further corrosion, leading to rapid attack.

6.3 Corrosion in Cl2- and HCl-Bearing Environments 6.3.1 Corrosion in Cl2 Environments (No O2) This section focuses on the corrosion behavior of alloys in environments containing essentially Cl2 with no oxygen (O2) present. An excellent review of halogen corrosion data up to the mid-1970s was presented by Daniel and Rapp

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(Ref 8). They summarized the test results obtained by various authors (Ref 16–19) on chloridation of iron (Table 6.5). It should be noted that the data in Table 6.5 were obtained from short-duration tests, ranging from a few minutes to hours. Use of these test results for extrapolation to 1 year could result in significant errors. They should not be used to estimate the service life of equipment, but should instead be used for comparison purposes. As illustrated in Table 6.5, iron exhibited little corrosion attack in Cl2 at temperatures up to 250 °C (480 °F).

0

Above 250 °C, corrosion rates abruptly increased. Iron forms two types of chlorides: FeCl2 and FeCl3. The melting and boiling points of FeCl2 are 676 and 1026 °C (1249 and 1879 °F), respectively. FeCl3, on the other hand, is extremely unstable. Its melting and boiling points are 303 and 319 °C (577 and 606 °F), respectively. Bohlken et al. (Ref 20, 21) suggested that the abrupt increase in the corrosion rate of iron in Cl2 at temperatures above 250 °C (482 °F) was related to the formation of FeCl3.

p(NiCl2) ≥ 10–4 bar

500 °C

–2 650 °C

Log p(Cl2), bar

–4 NiCl2

850 °C

–6

1000 °C

–8 –10

NiO –12

Ni

–14 –30

–25

–20

–15

–10

–5

0

Log p(O2), bar

Fig. 6.10

Quasi-stability diagram for Ni-O-Cl system for NiCl2 with vapor pressures of 10−4 atm (bar) and higher at temperatures from 500 to 1000 °C (932 to 1832 °F). Source: Ref 15

0

p(MoOxCly) ≥ 10–4 bar MoOCl4

Log p(Cl2), bar

–2 –4

MoOCl3

–6 MoCl4 –8

MoO2Cl2

–10

Mo

–12 –40

–35

–30

–25

MoO2 –20

–15

MoO3 –10

–5

0

Log p(O2), bar

Fig. 6.11

Quasi-stability diagram for Mo-O-Cl system for vapor pressures of chlorides and oxychlorides being 10−4 atm (bar) and higher at 800 °C (1472 °F). Source: Ref 15

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Adding chromium and/or nickel to iron improves the alloy’s resistance to chloridation attack. Ferritic and austenitic stainless steels can resist chloridation attack at higher temperatures than cast iron and carbon steels. Brown et al. (Ref 22) reported a corrosion rate of about 600 mpy at 232 °C (450 °F) for both carbon steel and cast iron. A ferritic stainless steel (Fe-17Cr) showed a corrosion rate of about 79 mpy at 360 °C (680 °F), and a titaniumstabilized austenitic stainless steel showed a rate of 24 mpy at 418 °C (784 °F) (Ref 8). The results of studies on several stainless steels by Tseitlin and Strunkin (Ref 16) and Brown et al. (Ref 22) were summarized by Daniel and Rapp (Ref 8) and are presented in Table 6.6. Adequate aluminum when added to iron to form aluminum oxide scales is also beneficial in improving the chloridation resistance. Han and Cho (Ref 23) studied corrosion behavior of Fe3Al (Fe-12.11%Al) in Ar-1%Cl2 at 750, 800, and 900 °C using a thermogravimetric method. The alloy behaved similarly at three different temperatures, showing an initial stage of an “incubation” time before the breakaway corrosion showing a drastic weight loss, as shown in Fig. 6.12. The authors observed a thin protective Al2O3 scale during the initial “incubation” stage, and nonprotective oxide scales (Al2O3, Fe2O3) Table 6.5 Temperature, °C (°F)

77 (170) 166 (330) 198 (388) 200 (392) 230 (446) 240 (464) 240 (464) 247 (477) 251 (484) 255 (491) 260 (500) 268 (514) 279 (534) 285 (545) 285 (545) 302 (576) 304 (579) 310 (590) 323 (613) 327 (621) 381 (718) 381 (718) 540 (1004) 540 (1004) 595 (1103) 599 (1110)

and small amounts of FeCl3 and FeCl2 formed at the “breakaway” stage. Also, at the breakaway stage, the specimen showed aluminum depletion at the metal/scale interface. The thin aluminum oxide scale was observed to form on the specimen during heating to the test temperature with argon gas flowing through the test chamber (approximately 2 h) prior to switching to the test gas. The test gas (i.e., Ar-1%Cl2) was found to contain 1 ppm O2. Thus, under the test condition, Al2O3 could form on the metal as well as FeCl2, as shown in Fig. 6.13. Nickel and nickel-base alloys are widely used in chlorine-bearing environments. The corrosion behavior of nickel in chlorine at various temperatures was analyzed by Daniel and Rapp (Ref 8), using the test results of Downey et al. (Ref 24), Tseitlin and Strunkin (Ref 16), and McKinley and Shuler (Ref 25) (see Table 6.7). At temperatures up to 500 °C (930 °F), nickel showed relatively low corrosion rates. Corrosion rates became suddenly and significantly higher at temperatures over 500 °C (930 °F). Nickel reacts with chlorine to form NiCl2, which exhibits relatively high melting point (1030 °C) compared to FeCl2 and FeCl3 (676 and 303 °C, respectively). This may be an important factor, making nickel much more resistant to chloridation attack than iron. Brown et al.

Corrosion of iron in chlorine Flow rate, (L/min)

100 100 100 15 100 100 15 100 100 120 15 120 120 15 15 120 120 15 120 120 120 120 15 15 120 120

pCl2, atm

1 1 1 1 1 1 1 I 1 (c) 1 (c) (c) 1 1 (c) (c) 1 (c) (c) (c) (c) (c) (c) (c) (c)

Diluent gas

None None None None None None None None None Ar None He Ar None None He Ar None Ar He Ar He N2 N2 Ar He

Test duration, min

480 0–15, 15–480 0–15, 15–480 360 0–15, 15–480 0–15, 15–480 360 0–15, 15–480 480 … 360 … … 360 60 … … 60 … … … … … … … …

Linear rate constant(a), μm/min −4

3 × 10 3.3 × 10−3, 3.8 × 10−4 5.2 × 10−3, 3.7 × 10−4 2 × 10−4 8.5 × 10−3, 2.2 × 10−4 9.5 × 10−3, 2.4 × 10−4 2 × 10-4 1 × 10−2, 1.9 ×10−4 1.55 0.94 −4 2 × 10 1.78 2.45 −4 4 × 10 20.4 4.04 4.17 3.9 8.94 11.4 40.4 65.3 6.6 1.5 187 624

Corrosion rate(b), mm/yr (mpy)

0.16 (6.3) 0.20 (7.9) 0.19 (7.5) 0.11 (4.3) 0.12 (4.7) 0.13 (5.1) 0.11 (4.3) 0.10 (3.9) 820 (32 in.) 490 (19 in.) 0.11 (4.3 mils) 940 (37 in.) 1,300 (51 in.) 0.21 (8 mils) 11,000 (433 in.) 2,100 (83 in.) 2,200 (87 in.) 1,700 (67 in.) 4,700 (185 in.) 6,000 (236 in.) 21,000 (827 in.) 34,000 (1,338 in.) 3,500 (138 in.) 830 (33 in.) 98,000 (3,858 in.) 328,000 (13,000 in.)

(a) Rate constants are given as metal loss rates, μm/min. (b) Estimated metal loss after 1 yr of exposure. (c) Tests were conducted with chlorine partial pressures up to 0.3 atm and total pressures of 1.0 atm. However, the metal loss rates were extrapolated to 1.0 atm chlorine pressure. Source: Ref 8


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Table 6.6

Corrosion of stainless steels in chlorine(a)

Alloy

Ferritic stainless (Fe-17Cr)

Austenitic stainless (Fe-18Cr-9Ni-Ti)

Austenitic stainless (Fe-18Cr-8Ni-Mo)

Austenitic stainless (Fe-18Cr-8Ni)


Flow rate, L/min

Linear rate constant(b), µm/min

Corrosion rate(c), mm/yr (mpy)

300 (572) 360 (680) 440 (824) 540 (1,004) 418 (784) 450 (842) 480 (896) 535 (995) 640 (1,184) 315 (599) 340 (644) 400 (752) 450 (842) 480 (896) 290 (554) 315 (599) 340 (644) 400 (752) 450 (842)

15 15 15 15 15 15 15 15 15 28 28 28 28 28 28 28 28 28 28

4 × 10−4 3.8 × 10−3 6.7 × 10−2 1.35 1.1 × 10−3 4.3 × 10−2 0.13 0.47 46 1.4 × 10−3 2.9 × 10−3 5.9 × 10−3 2.9 × 10−2 5.9 × 10−2 1.5 × 10−3 2.9 × 10−3 5.9 × 10−3 2.9 × 10−2 5.9 × 10−2

0.2 (7.9) 2 (79) 40 (1.6 in.) 700 (28 in.) 0.6 (24) 20 (787) 70 (2.8 in.) 200 (7.9 in.) 20,000 (787 in.) 0.8 (31) 1.5 (59) 3 (118) 15 (590) 30 (1.2 in.) 0.8 (31) 1.5 (59) 3 (118) 15 (590) 30 (1.2 in.)

(a) Chlorine pressure was approximately 1.0 atm. (b) Duration of these tests was 60–360 min for the first two alloys and 120–1200 min for the last two alloys. (c) Estimated metal loss after one year of exposure. Source: Rcf 8

5

10 1% Cl2/Ar

0

FeCl3

0 –5

Log p Cl , g 2


750 °C

5

–10

–5

FeCl2

–10 AlCl3

–15

–15 900 °C

800 °C

750 °C

Al2O3 Fe3O4

Fe

–20

Fe2O3

Al

–20

–25 –25 0

10

20

30

40

50

60

70

–30 –60

Exposure time, h

Fig. 6.12

Thermogravimetric results for Fe3Al (Fe-12.11%Al) tested in Ar-1%Cl2 at 750, 800, and 900 °C. Source: Ref 23

(Ref 22) conducted short-term laboratory tests in chlorine on various commercial alloys. The results (see Table 6.8) suggested that, in an environment of 100% Cl2, carbon steel and cast iron are useful at temperatures up to 150 to 200 °C (300 to 400 °F) only. The 18-8 stainless steels can be used at higher temperatures—up to 320 to 430 °C (600 to 800 °F). Nickel and nickel-base alloys (e.g., Ni-Cr-Fe, Ni-Mo, and Ni-Cr-Mo alloys) were most resistant. The beneficial effect of nickel on the resistance of chloridation attack in Cl2 environments is

FeO –50

–40

–30

–20

–10

0

10

Log p O , g 2

Fig. 6.13

Phase stability for Al-O-Cl system at 750 °C. The solid circle indicates the test environment. Source:

Ref 23

illustrated in Fig. 6.14 (Ref 26). This trend is also reflected in long-term tests (Table 6.9). Alloy 600 is the most commonly used alloy for high-temperature services in Cl2 environments. Figure 6.15 shows the corrosion rates of alloy 600 in dry chlorine gas as a function of temperature (Ref 22). MTI Publication MS-3 (Ref 27) suggests corrosion guidelines for Ni200, alloy 600, alloy 400, Type 304, and steel in dry chlorine gas applications, as shown in Fig. 6.16. Tu et al. (Ref 28) performed phase analysis using x-ray diffraction on the external corrosion


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Table 6.7

Corrosion of nickel in chlorine pCl2, atm


350 (662) 375 (707) 400 (752) 407 (765) 433 (811) 465 (869) 485 (905) 500 (932) 500 (932) 525 (977) 540 (1004) 550 (1022) 550 (1022) 660 (1220) 740 (1364) 770 (1418)

~1 ~1

~1 ~1 ~1

Test duration, min

0.13 0.13 0.13 0.13 0.13 0.13 0.13 0.13 atm at 16 cm3/min 0.13 atm at 16 cm3/min 0.13 0.13 atm at 16 cm3/min atm at 16 cm3/min atm at 16 cm3/min

3000–5000 … … … … … … … 3600 … 3600 … … 3600 60 60

Corrosion rate(a), mm/yr (mpy)

0.0012 (0.05)(b) 0.0014 (0.06)(b) 0.0036 (0.14)(b) 0.005 (0.2)(b) 0.0058 (0.2)(b) 0.107 (4.2) 0.142 (5.6) 0.319 (12.6) 0.492 (19.4) 2.41 (95) 2.07 (82) 6.02 (237) 5.85 (230) 40.3 (1,587) 120 (4,724) 2,150 (84,646)

(a) Estimated metal loss after 1 year of exposure to chlorine. (b) These estimates are probably low. Source: Ref 8

Table 6.8 Corrosion of selected alloys in chlorine Approximate temperature, °C (°F), at which given corrosion rate is exceeded Alloy

0.8 mm/yr (30 mpy)

1.5 mm/yr (60 mpy)

3.0 mm/yr (120 mpy)

15 mm/yr (600 mpy)

Nickel Alloy 600 Alloy B Alloy C Chromel A Alloy 400 18-8 Mo 18-8 Carbon steel Cast iron

510 (950) 510 (950) 510 (950) 480 (900) 425 (800) 400 (750) 315 (600) 288 (550) 120 (250) 93 (200)

538 (1000) 538 (1000) 538 (1000) 538 (1000) 480 (900) 455 (850) 345 (650) 315 (600) 175 (350) 120 (250)

593 (1100) 565 (1050) 593 (1100) 565 (1050) 538 (1000) 480 (900) 400 (750) 345 (650) 205 (400) 175 (350)

650 (1200) 650 (1200) 650 (1200) 650 (1200) 620 (1150) 538 (1000) 455 (850) 400 (750) 230 (450) 230 (450)

Source: Ref 22

products formed on Ni-4Cr alloy when exposed in 105 Pa (1 atm) Cl2 at 575 and 700 °C. The authors found that the scales consisted of mainly NiCl2, CrCl3, and CrCl2. The deposits on the quartz test assembly during testing of Ni-4Cr alloy were also analyzed. These deposits were mainly NiCl2 and CrCl3 with very little CrCl2, indicating both NiCl2 and CrCl3 are major vapor phases during testing Ni-4Cr alloy. 6.3.2 Corrosion in O2-Cl2 Environments Many industrial environments may contain both chlorine and oxygen. Metals generally follow a parabolic rate law by forming condensed phases of oxides, if the environment is free of chlorine. With the presence of both oxygen and chlorine, corrosion of metals then involves a combination of condensed oxides and volatile chlorides. Depending on the relative amounts of oxides and chlorides formed, corrosion can

follow either a paralinear rate law (a combination of weight gain due to oxidation and weight loss due to chlorination) or a linear rate law due to chlorination. This is illustrated by the results of Maloney and McNallan (Ref 29) on corrosion of cobalt in Ar-50O2-Cl2 mixtures (Fig. 6.17). As shown in the figure, at high Cl2 levels, the corrosion products are primarily cobalt chloride vapor, causing the weight loss to follow a linear rate law. Because of volatile corrosion products, the reaction rate can be highly dependent on the gas flow rate. McNallan and Liang (Ref 30) showed that CoO specimens exhibited increased linear weight loss rates with increasing gas velocity when exposed in the Ar-O2-1Cl2 mixtures with pO2 being 0.01 and 0.15 atm pressures at 723 °C (1000 K). Furthermore, the oxygen partial pressure (pO2 ) of 0.01 atm resulted in higher weight loss rates than that of 0.15 atm. This is also illustrated in Fig. 6.18 (Ref 31), showing the corrosion of cobalt in Ar-O2-1Cl2 mixtures with three different concentrations (1, 10, and 50% O2) of oxygen at 650 °C (1200 °F). When the environment contained 10 and 50% O2, the corrosion reaction involved mainly the formation of cobalt oxide, thus following an approximate parabolic rate. When the oxygen concentration reduced to 1%, the corrosion reaction involved mainly volatile CoCl2, thus following an approximate linear weight loss with time. The figure also shows that the linear weight loss agreed very well with the volatilization of CoCl2. It should be noted that in the above test environments containing 10 and 50% O2, the corrosion reaction, which followed a parabolic rate due to formation of condensed cobalt oxides (Fig. 6.18), involved only a very


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mg/cm2


690

Fig. 6.14

Effect of nickel on the corrosion resistance of alloys in Ar-30Cl2 at 704 °C (1300 °F) for 24 h. Source: Ref 26

Table 6.9 Corrosion of several alloys in Ar-30Cl2 after 500 h at 400 to 704 °C (750 to 1300 °F)(a) Descaled weight loss, mg/cm2 Alloy

400 °C (750 °F)

500 °C (930 °F)

600 °C (1110 °F)

704 °C (1300 °F)(a)

Ni-201 600 601 625 617 800 310 304 347

0.2 0.02 0.3 0.7 0.6 6 28 108 215

0.3 5 3 7 7 13 370 1100 Total

47–101 127–180 85–200 … … 200–270 … … …

97 160 215 180 190 890 820 >1000 Total

(a) 24 h test period. Source: Ref 26

short-term test (2 h). It is extremely likely that upon longer exposure times the cobalt oxide scales may eventually crack (or spall), thus allowing chlorine gas to reach the underlying metal and causing formation of volatile CoCl2 and resulting in a linear corrosion rate. Corrosion of pure nickel in O2-Cl2 environments was found to be dominated by the formation of volatile NiCl2 corrosion product, and the weight loss essentially followed a linear rate law (Ref 32, 33). Figure 6.19 shows thermogravimetric results at 927 °C in Ar-50O2 containing various amounts of Cl2 from 0 to 5% (Ref 32). The figure also shows that the experimental data were in good agreement with the

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/month


theoretical calculation based on vapor pressures of NiCl2. In their testing of nickel in Ar-50O2-Cl2 containing 0, 0.5, 1.0, 1.75, 3, and 5% Cl2 at 627 to 927 °C, Lee and McNallan (Ref 32) observed a very rapid reaction referred to as “ignition” at 727 and 827 °C in the environments containing higher concentrations of Cl2. McNallan (Ref 33) suggested that this rapid corrosion reaction was caused by the reaction of chlorine with nickel at the metal/oxide scale interface to form nickel chloride vapor, which then diffuses out and is converted into powdery nickel oxide. This rapid corrosion reaction causes metal temperature to increase. He further indicated that this rapid reaction (ignition) can be prevented at temperatures below 727 °C due to low vapor pressures of nickel chloride and can also be prevented at temperatures higher than 827 °C due to formation of a protective oxide scale by rapid oxidation of nickel to reduce the ingress of chlorine through the oxide scale and thus the formation of nickel chloride vapor at the metal/oxide scale interface. In an O2-Cl2 containing environment, oxidation and chloridation can take place. As discussed in section 6.2, a protective oxide scale can lower the pCl2 below the value for forming metal chlorides at the metal/oxide scale interface. However, cracking and spalling of this oxide scale resulting from, for example, thermal

Fig. 6.15

Corrosion rate of alloy 600 in Cl2 as a function of temperature. Source: Ref 22

Fig. 6.16

MTI corrosion guidelines for Ni200, alloy 600, alloy 400, Type 304SS and steel in dry chlorine (Cl2) as a function of temperature. Source: Ref 27. Courtesy of Materials Technology Institute

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Fig. 6.17

Corrosion of cobalt in Ar-50O2-Cl2 at 927 °C (1700 °F). Source: Ref 29

Fig. 6.18

Thermogravimetric results of corrosion of cobalt in Ar-O2-1Cl2 with 1, 10, and 50% O2 at 650 °C (1200 °F). Source: Ref 31

cycling can allow permeation of chlorine to reach the metal underneath to initiate chloridation attack. Figure 6.20 illustrates the effect of thermal cycling on the initiation of accelerated chloridation attack for Fe-20Cr at 927 °C in Ar-20O2-0.5Cl2 (Ref 34). The figure shows the thermogravimetric results for three separate test

Fig. 6.19 Thermogravimetric results of corrosion of nickel at 927 °C in Ar-50O2-Cl2 (0, 0.5, 1.0, 1.75, 3, and 5% Cl2). Source: Ref 32 runs, showing similar behavior in initiating an accelerated chloridation attack right after a thermal cycle, which involved cooling the specimen for 30 min by lowering it from the test temperature (927 °C) to 100 °C in 5 min. As soon

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as the test was resumed after the thermal cycle, specimens (three separate test runs) showed accelerated chloridation attack. A significant amount of data have been generated for commercial alloys in environments containing both oxygen and chlorine. Table 6.10 summarizes short-term test results generated in Ar-20O2-2Cl2 at 900 °C (1650 °F) for 8 h (Ref 35). The results revealed several interesting trends. The best-performing alloy was an aluminum-containing alloy 214 (Fe-16Cr-3Fe4.5Al) with a small amount iron. Two worstperforming alloys were cobalt-base alloys (alloys 188 and 6B) containing large amounts of tungsten. Molybdenum-containing nickel-base alloys also did not perform well. Oh et al. (Ref 36) attributed this to the formation of oxychlorides of molybdenum and tungsten, which have very high vapor pressures. The partial pressures of WO2Cl2 and MoO2Cl2 in equilibrium with the oxides (WO3 and MoO3, respectively) and the test environment (Ar-20O2-2Cl2) at 900 °C (1650 °F) were 7.52 ×10−2 and 2.1× 100 atm, respectively. Accordingly, alloy 188 (14% W), alloy C-276 (16% Mo, 4% W), alloy 6B (4.5% W, 1.5% Mo), alloy X (9% Mo), and alloy S (14.5% Mo) suffered relatively high rates of corrosion attack. Simple Fe-Ni-Cr (Type 310 SS) and Ni-Cr-Fe (alloy 600) performed better than molybdenum- or tungsten-containing alloys. The 100

Mass change, mg/cm2

80 36 h 60

18 h

40

gravimetric results for representative alloys are summarized in Fig. 6.21 (Ref 36). Thermodynamic phase stability diagrams showing high vapor pressures of oxychlorides of tungsten and molybdenum were presented earlier in Fig. 6.8 and 6.11, respectively. In order to determine whether molybdenum oxychlorides would contribute to high corrosion rates in high-molybdenum-containing nickel alloys, such as alloy S (Ni-16Cr-14.5Mo) in O2-Cl2 environments, Jacobson et al. (Ref 37) used a high-pressure sampling mass spectrometer to measure volatile species produced from the preoxidized specimen of alloy S with 14.5% Mo in comparison with alloy 600 (Ni-16Cr-9Fe) with no Mo during the exposure of Ar-50O21Cl2. Thermogravimetric data for these two preoxidized alloys under the test condition are shown in Fig. 6.22, showing a significantly higher weight loss rate for alloy S (14.5Mo) than alloy 600 (no Mo) during the exposure of the preoxidized specimens to Ar-50O2-1Cl2 at 900 °C (1650 °F). The mass spectrometer results indicated that MoO2Cl2 along with NiCl2 and CrO2Cl2 were major vapor phases in the case of alloy S. For alloy 600, NiCl2 and CrO2Cl2 were detected. Alloy R-41 (nickel-base alloy with 1.5% Al, 3% Ti and 10% Mo) suffered less chloridation attack than other nickel-base alloys containing molybdenum despite high molybdenum content in a short-term test presented in Table 6.10. However, the results of long-term tests in Ar-20O2-0.25Cl2 by Rhee et al. (Ref 38) and McNallan et al. (Ref 39) showed that these nickel-base alloys with molybdenum, such as R-41 (Ni-19Cr-11Co-10Mo-1.5Al-3Ti) and alloy 263 (Ni-20Cr-20Co-5.8Mo-0.5Al-2.2Ti), eventually suffered severe attack despite the presence of aluminum and titanium. Figure 6.23

24 h

Table 6.10 Corrosion of selected alloys in Ar-20O2-2Cl2 at 900 °C (1650 °F) for 8 h

20 Pure oxygen 0

Alloy

4

8

12

16

20

24

28

32

36

Time, h

Fig. 6.20

Thermogravimetric results of three test runs for Fe20Cr alloy tested in Ar-20O2-0.5Cl2 at 927 °C isothermally for the first 12 h, followed by a thermal cycle by cooling the specimen to 100 °C for 30 min and raising the specimen to the test temperature to resume testing. Note that the thermal cycle resulted in the initiation of an accelerated chloridation attack. Source: Ref 34

214 R-41 600 310SS S X C-276 6B 188


Average metal affected(a), mm (mils)

0 0.004 (0.16) 0.012 (0.48) 0.012 (0.48) 0.053 (2.08) 0.020 (0.80) 0.079 (3.12) 0.014 (0.56) 0.014 (0.56)

0.012 (0.48) 0.028 (1.12) 0.035 (1.36) 0.041 (1.60) 0.063 (2.48) 0.071 (2.80) 0.079 (3.12) 0.098 (3.84) 0.116 (4.56)

(a) Metal loss + average internal penetration. Source: Ref 35


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310SS

Fig. 6.21

Gravimetric results for selected Fe-, Ni-, and Co-base alloys in Ar-20O2-2Cl2 at 900 °C (1650 °F). Source: Ref 36

Fig. 6.22

Thermogravimetric data showing weight loss of two preoxidized specimens during the exposure in Ar-50O2-1Cl2 at 900 °C (1650 °F). Source: Ref 37

shows corrosion test results for aluminumcontaining nickel-base alloys with and without molybdenum, such as 214, 601, R-41, and 263,

tested at 900 °C (1650 °F) (Ref 39). Test results for all the alloys tested at 900 and 1000 °C (1650 and 1830 °F) are summarized in Tables 6.11 and 6.12 (Ref 39). The beneficial effect of aluminum, as well as the detrimental effect of molybdenum and tungsten, on the resistance to chloridation attack in oxidizing environments is further substantiated by the results of long-term tests in another environment with a higher concentration of Cl2, as shown in Fig. 6.24 (Ref 35). Similar results were obtained by Elliott et al. (Ref 40) from tests conducted in air-2Cl2 at 900 °C (1650 °F) for 50 h (Fig. 6.25). Chloridation attack in O2-Cl2 environments at these high temperatures primarily consisted of metal wastage and internal penetration for most alloys, with the exception of Ni-Cr-Mo alloys containing high levels of molybdenum, such as alloys S and C-276, which showed mainly metal wastage with little or no internal penetration. In general, the scales formed on the alloy surface were loose when the test specimens were cooled to room temperature after the exposure test, as illustrated in Fig. 6.26. Scales were basically


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oxides, as shown in Fig. 6.27 and 6.28. Fe-Cr oxides were found to form on Type 310 SS, while aluminum-rich oxides were the major oxide phase along with little nickel-rich oxides that formed on alloy 214. Internal attack consisted of voids for some alloys and of oxides for other alloys. Some alloys appeared to contain both internal voids and oxides. Figure 6.29 shows internal void formation in alloy R-41 (Ni-Cr-CoMo-Al-Ti) and alloy 25 (Co-Cr-Ni-W) after exposure for 50 h at 900 °C (1650 °F) in air2Cl2. It was suggested that the internal penetration may involve halogen-carbide reactions and simultaneous void formation (Ref 25). Figure 6.29 shows some evidence of carbides being converted into voids during chloridation attack. Some alloys, however, showed internal oxides instead of voids, as illustrated in Fig. 6.30. Elliott et al. (Ref 40) have identified volatile species of the condensed products removed from the exit end of the test apparatus during their investigation in air-2Cl2. Their results are shown in Table 6.13. No oxychlorides were detected. Since the analysis of the volatile species was performed on the condensed phases collected at the exit end of the test apparatus, oxychlorides were likely to be in a gaseous state at the exit end,

and thus were not collected. As discussed earlier, Jacobson et al. (Ref 37) used a high-pressure sampling mass spectrometer to measure volatile species from alloy S with 14.5% Mo and alloy 600 (Ni-16Cr-9Fe) with no Mo during the exposure of Ar-50O2-1Cl2 at 900 °C (1650 °F) and found that MoO2Cl2 along with NiCl2 and CrO2Cl2 were major vapor phases for alloy S and NiCl2 and CrO2Cl2 for alloy 600. McNallan et al. (Ref 39) reported corrosion behavior in Ar-20O2-0.25Cl2 at 900 and 1000 °C (1650 and 1830 °F). This was followed by a study (Ref 41) using the same environment to investigate the same alloys at lower temperatures (i.e., 700, 800, and 850 °C). The results of the tests at 700, 800, and 850 °C (1290, 1470, and 1560 °F) are summarized in Table 6.14 Table 6.11 Corrosion of various alloys in Ar-20O2-0.25Cl2 for 400 h at 900 and 1000 °C (1650 and 1830 °F) Weight loss, mg/cm2 Alloy

214 601 600 800H 310SS 556 X 625 R-41 263 188 S C-276

900 °C (1650 °F)

1000 °C (1830 °F)

4.28 20.67 72.08 26.91 47.15 40.29 54.41 99.07 63.83 82.57 139.77 228.21 132.05

9.05 124.99 254.96 87.05 97.40 82.74 153.49 220.09 207.32 229.53 156.30 248.98 298.85

Source: Ref 39

Table 6.12 Depth of attack after 400 h at 900 and 1000 °C (1650 and 1830 °F) in Ar-20O2-0.25Cl2 900 °C (1650 °F) Alloy

Fig. 6.23

Corrosion of several aluminum-containing nickelbase alloys with and without molybdenum in Ar-20O2-0.25Cl2 at 900 °C (1650 °F). Source: Ref 39

214 601 600 800H 310SS 556 X 625 R-41 263 188 S C-276


Total depth(a), mm (mils)

0.023 (0.9) 0.150 (5.9) 0.061 (2.4) 0.264 (10.4) 0.127 (5.0) 0.252 (9.9) 0.043 (1.7) 0.191 (7.5) 0.086 (3.4) 0.152 (6.0) 0.046 (1.8) 0.152 (6.0) 0.099 (3.9) 0.218 (8.6) 0.208 (8.2) 0.272 (10.7) 0.114 (4.5) 0.244 (9.6) 0.130 (5.1) 0.193 (7.6) 0.216 (8.5) >0.356 (14.0) 0.315 (12.4) 0.353 (13.9) 0.300 (11.8) 0.320 (12.6)

1000 °C (1830 °F) Metal loss, mm (mils)

Total depth(a), mm (mils)

0.013 (0.5) 0.203 (8.0) 0.330 (13.0) 0.203 (8.0) 0.191 (7.5) 0.152 (6.0) 0.318 (12.5) 0.356 (14.0) 0.381 (15.0) 0.368 (14.5) 0.254 (10.0) 0.419 (16.5) 0.419 (16.5)

0.051 (2.0) 0.295 (11.6) 0.386 (15.2) 0.424 (16.7) 0.246 (9.7) 0.300 (11.8) 0.434 (17.1) 0.437 (17.2) 0.457 (18.0) 0.424 (16.7) 0.417 (16.4) 0.472 (18.6) 0.450 (17.7)

(a) Metal loss + internal penetration. Source: Ref 39


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Fig. 6.24

Corrosion of several nickel- and cobalt-base alloys in Ar-20O2-1Cl2 at 900 °C (1650 °F). Source: Ref 35

Fig. 6.25

Corrosion of several iron- and nickel-base alloys in air-2Cl2 at 900 and 1000 °C (1650 and 1830 °F) for 50 h. Source: Ref 40

(Ref 41). The corrosion behavior of alloys as a function of temperature from 700 to 1000 °C (1290 to 1830 °F) can best be summarized in Fig. 6.31 using three different alloy systems (Ni-Cr-Mo alloys S, Fe-Ni-Cr alloy 800H, and Ni-Cr-Al alloy 214). As discussed earlier, refractory metals, such as molybdenum and tungsten, are detrimental to chloridation resistance

in oxidizing environments at high temperatures. Alloy S was found to be less corrosion resistant than alloy 800H. However, both alloys S and 800H suffered increasing corrosion attack with increasing temperatures. This represents a typical trend for most alloys in oxidizing environments. One exception is the Ni-Cr-Al system. As illustrated in Fig. 6.31, alloy 214 showed a sudden decrease in corrosion attack as the test temperature was increased from 900 to 1000 °C (1650 to 1830 °F). This sharp reduction in corrosion attack at 1000 °C (1830 °F) was attributed to the formation of a protective Al2O3 scale. At lower temperatures, such as 900 °C or less, the kinetics of Al2O3 formation was not fast enough to form a protective oxide scale in O2-Cl2 environments. Formation of a protective Al2O3 scale is favored at higher temperatures (e.g., 1000 °C or higher). Most of the data presented so far were generated at fairly high temperatures (i.e., 900 °C and higher). Chloridation was quite aggressive at those high temperatures. Schwalm and Schutze (Ref 42–44) investigated a large number of commercial alloys at significantly lower temperatures, varying from 300 to 800 °C (572 to 1472 °F) for times up to 300 h in air-2Cl2. The corrosion behavior of various alloys were generated in terms of the decrease in specimen crosssection thickness (i.e., metal loss) and the depth of internal attack. Extensive characterization of the corrosion products formed on the alloys was

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investigated using SEM/EDX analyses, which showed distribution of elements including oxygen and chlorine in the corrosion products. Their results are briefly summarized below. Figure 6.32(a) shows the decrease in metal cross-section thickness for 2.25Cr-1Mo steel (10CrMo9 10), alloy 800H, alloy AC66, alloy 45TM, and alloy 690 after testing for 300 h as

Fig. 6.26

a function of temperature. The data on the depth of internal corrosion attack are presented in Fig. 6.32(b). The corrosion attack at 300 °C (572 °F) was quite negligible after 300 h for the alloys tested including 2.25Cr-1Mo steel. Nevertheless, the Cr-Mo steel was found to exhibit a fragile oxide scale, which contained Fe, O, and Cl. All other alloys including two

Loose scales on samples of several nickel-base alloys after testing at 900 °C (1650 °F) in Ar-20O2-1Cl2 for 100 h

Semiquantitative EDX analysis, wt% Area

1 2

Fig. 6.27

Fe

Cr

Si

85.1 69.5

14.9 29.9

… 0.6

Scanning electron micrograph showing oxide scale formed on Type 310SS sample exposed at 900 °C (1650 °F) for 400 h in Ar-20O2-0.25Cl2. The results of the EDX analysis of the corrosion products on the areas, as marked No. 1 and No. 2, are tabulated. Magnification bar represents 33.3 µm

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nickel alloys (alloys 45TM and 690) showed pitting type of attack. The phase in the pit was heavily enriched in Cl and O with some Cr, Fe, and Ni. Figure 6.33 shows the morphology of the pit along with an x-ray map for Cl. At 500 °C (932 °F), 2.25Cr-1Mo steel suffered both significant thickness reduction and internal attack, while alloys AC66, 800H, 45TM, and 690 showed little attack. At 650 °C, alloy AC66 suffered significantly more thickness loss than alloys 800H, 45TM, and 690. The elemental distribution in the corrosion products for alloy AC66 (worst alloy in this group) is shown in Fig. 6.34. The chromium oxides with a layer of iron-rich oxide that formed on AC66 became convoluted. Chlorine was detected at the metal/ oxide scale interface and in the metal underneath the metal/oxide scale interface, where internal attack was observed. Alloy 690, on the other hand, showed a continuous chromium oxide scale. No chlorine was detected. Some internal

attack was observed underneath the oxide scale, and these internal particles are believed to be aluminum oxides. The morphology of the corrosion products formed in alloy 690 is shown in Fig. 6.35. Corrosion behavior for alloys 59, C-2000, and HR160 after testing for 300 h as a function of temperature is shown in Fig. 6.36. All three alloys exhibited little corrosion attack at 300 and 500 °C. At 650 °C, both alloys 59 and HR160 continued to exhibit little corrosion attack, while alloy C-2000 suffered much more corrosion attack. Both C-2000 (Ni-23Cr-16Mo-1.6Cu) and 59 (Ni-23Cr-16Mo-0.3Al) exhibit similar chemical compositions except C-2000 contains additional 1.6% Cu and alloy 59 contains additional 0.3Al. It is not clear whether Cu in alloy C-2000 was responsible. At 650 °C, both alloys 59 and HR160 were found to perform better than alloys 690 and 45TM (Fig. 6.32). Alloy 59 was found to exhibit a thin, continuous chromium-rich oxide

Semiquantitative EDX analysis, wt%

Fig. 6.28

Area

Ni

Cr

Al

Fe

1 2

72 6

20 5

4 88

4 1

Scanning electron micrograph showing oxide scale formed on alloy 214 sample tested at 900 °C (1650 °F) for 400 h in Ar-20O2-0.25Cl2. The results of the EDX analysis of the corrosion products on the areas, as marked No. 1 and No. 2, are tabulated. Magnification bar represents 33.3 µm

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scale after exposure at 800 °C for 300 h in air2Cl2. as shown in Fig. 6.37. Also observed was an outer Fe-, Ni-, and Cr-rich oxide scale. Also observed were a small amount of chlorine and slight molybdenum enrichment at the metal/ chromium-oxide scale interface. At 800 °C, all three alloys showed higher corrosion attack. Alloy 59 continued to perform the best. The

Fig. 6.29

high concentration of molybdenum (16%) showed no detrimental effect on the alloy’s corrosion resistance in this oxidizing environment containing 2% Cl2. Thermodynamically, this environment at 800 °C, molybdenum oxychloride (MoO2Cl2) would be stable, as shown in Fig. 6.11. Molybdenum oxychloride was believed to contribute to high corrosion rates for

Scanning electron micrographs showing internal void formation in (a) alloy R-41 and (b) alloy 25 after exposure for 50 h at 900 °C (1650 °F) in air-2Cl2

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high-molybdenum-containing alloys, such as alloys C-276, S, and so forth when exposed to O2-Cl2 environments at 900 °C (1650 °F) and higher (Ref 35, 39, 40). The chloride phases exhibiting 10−4 atm pressure that are predicted thermodynamically under this test condition are NiCl2, CoCl2, CrO2Cl2, FeCl3, and MoO2Cl2 (Table 6.1), while the oxides are Al2O3, SiO2, Cr2O3, and Fe2O3. HR160 alloy showed an inner

layer of silicon-rich oxide scale underneath the chromium-rich oxide scale with slight internal silicon oxides after exposure at 800 °C for 300 h in air-2Cl2, as shown in Fig. 6.38. The figure also shows that silicon oxides contained what appeared to be tiny internal voids. Schwalm and Schutze (Ref 44) also tested alumina former alloy 214 (Ni-16Cr-3Fe-4.5AlY-Zr), intermetallic Fe3Al (Fe-5.5Cr-15.9Al0.2Zr), and intermetallic TiAl (Ti-36Al). The corrosion data are shown in terms of crosssection thickness loss (Fig. 6.39a) and internal attack (Fig. 6.39b) after exposure for 300 h at temperatures from 300 to 800 °C (572 to 1472 °F) in air-2Cl2. Alumina former alloy 214 performed very well, comparable to alloy 59 (Fig. 6.36). Fe3Al, while showing little corrosion attack at 300 and 500 °C, suffered severe corrosion at 650 °C. At 800 °C, Fe3Al showed very little corrosion again. The oxide scales formed on Fe3Al at 650 °C consisted of nonprotective, multilayers of Cr-containing Fe2O3, Al2O3, and Fe2O3. At 800 °C, a thin, continuous aluminum-rich oxide scale was found to form on Fe3Al when tested at 800 °C. It is clear that

Semiquantitative EDX analysis, wt%

Table 6.13 Volatile species of condensed products(a) after testing at 900 °C (1650 °F) for 50 h in air-2Cl2

Area

Ni

Cr

Fe

Al

Ti

Alloy

1 2 3 4 5

1 19 25 29 56

62 66 60 45 26

1 7 12 13 18

27 6 2 10 …

9 2 1 3 …

Alloy 214 Alloy 601(b) Type 310SS Alloy 800H(b) Alloy 25 Alloy 625 Alloy 617 Alloy 263 Alloy C-276

Fig. 6.30

Scanning electron micrograph showing oxide scales and internal oxides for alloy 601 exposed at 900 °C (1650 °F) for 400 h in Ar-20O2-0.25Cl2. The results of the EDX analysis of the corrosion products on the areas, as marked No. 1, No. 2, No. 3, No. 4, and No. 5, are listed.

Major constituents

… NiCl2, AlCl3 FeCl3·2H·2O, NiCl2·6H2O FeCl3·2H2O, NiCl2·6H2O CoCl2, NiCl2·6H2O, WCl6 NiCl2·6H2O, FeCl3·2H2O, MoCl5 NiCl2·6H2O, FeCl3·2H2O, CoCl2 NiCl2·6H2O, CoCl2, MoCl5, FeCl3·2H2O NiCl2·6H2O, MoCl5

(a) Collected at the downstream, cooler section of the test apparatus. (b) Tested at 1000 °C (1830 °F) for 50 h. Source: Ref 40

Table 6.14 Depth of attack for various alloys after 400 h at 700, 800, and 850 °C (1290, 1470, and 1560 °F) in Ar-20O2-0.25Cl2 700 °C (1290 °F) Alloy

214 600 800H 310SS 556 S C-276 188 Source: Ref 41

800 °C (1470 °F)

850 °C (1560 °F)


Total depth, mm (mils)



Metal loss. mm (mils)


0.010 (0.4) … 0.025 (1.0) … … 0.079 (3.1) 0.033 (1.3) …

0.010 (0.4) … 0.033 (1.3) … … 0.081 (3.2) 0.046 (1.8) …

0.018 (0.7) 0.020 (0.8) 0.023 (0.9) 0.036 (1.4) 0.020 (0.8) 0.145 (5.7) 0.066 (2.6) 0.058 (2.3)

0.061 (2.4) 0.086 (3.4) 0.046 (1.8) 0.053 (2.1) 0.051 (2.0) 0.150 (5.9) 0.071 (2.8) 0.074 (2.9)

0.018 (0.7) 0.038 (1.5) 0.031 (1.2) 0.031 (1.2) 0.020 (0.8) 0.224 (8.8) 0.163 (6.4) 0.025 (1.0)

0.066 (2.6) 0.132 (5.2) 0.097 (3.8) 0.061 (2.4) 0.079 (3.1) 0.257 (10.1) 0.175 (6.9) 0.264 (10.4)


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higher temperatures are favored for forming an aluminum oxide scale. Unlike Fe3Al, TiAl suffered severe corrosion attack at 500, 650, and 800 °C. In many combustion processes involving burning fuels or feedstock containing hydrocarbon, combustion environments under substoichiometric combustion conditions most often would contain CO2. McNallan et al. (Ref 45) examined the effect of CO2 on the chloridation resistance of Fe-Cr and Fe-Ni-Cr alloys and found that CO2 significantly increased the alloy’s metal wastage and internal attack. The authors compared the environments between Ar-20O2-2500 ppm (0.25%) Cl2 and Ar-20CO22500 ppm Cl2 for Fe-20Cr and 800H at 927 °C (1700 °F). Figure 6.40 shows metal loss data for both environments as a function of exposure time. For both alloys, the CO2-Cl2 environment caused more metal wastage than the O2-Cl2 environment with Fe-20Cr being more severely affected than alloy 800H. The authors did not offer an explanation in the paper. It was likely that when the environment was switched from Ar-20O2-Cl2 to Ar-20CO2-Cl2 the environment changed from oxidizing to reducing with its oxygen potential being reduced to closer to, or in, the CrCl3 and FeCl2 regimes in the Cr(Fe)-O-C diagrams, resulting in formation of more volatile

chlorides, thus more metal loss. In addition to metal loss, the internal attack for alloy 800H was more severe in the CO2-Cl2 environment than the O2-Cl2 environment (Fig. 6.41). The authors hypothesized the mechanism of the internal corrosion by forming internal chromium carbides, which are then converted to chromium chlorides with carbon reacting with chlorine to form more chromium carbides. Thus, internal corrosion is the result of formation of internal voids and pores in the metal. Fe-20Cr alloy, on the other hand, suffered no internal corrosion. The authors attributed this lack of internal corrosion to the alloy’s low carbon content and ferritic structure. In another paper by McNallan et al. (Ref 46), the CO2-containing environments were further examined. The authors observed internal carburization of Type 310SS in Ar-20CO2 at 800 °C for 24 h. When the test was conducted in Ar-20CO2-Cl2 at 800 °C for 24 h, Type 310SS suffered internal corrosion attack in forms of voids and carbides. Similar internal attack was observed for alloy 800. The corrosion, suggested by the authors, proceeded in two stages with carburization preceding the chlorine-accelerated oxidation. In combustion environments, H2O is invariably present among the combustion products produced. The oxygen potential is dictated by partial pressures of oxygen and hydrogen. Hydrogen reacts with Cl2 to form more stable HCl molecule. The chloridation behavior in HClcontaining environments is discussed in the next two sections. 6.3.3 Corrosion in O2-HCl Environments In some combustion environments, chlorine is present as HCl instead of Cl2. This section covers the corrosion data in oxidizing environments containing HCl. In an oxidizing environment containing HCl, chlorine partial pressure, pCl2 , can be calculated from the equilibrium condition of the reaction below. The oxide and chloride phases that are likely to form thermodynamically can then be determined from the thermodynamic phase stability diagrams discussed earlier in section 6.2. 4HCl(g)+O2 (g)=2Cl2 (g)+2H2 O(g)

Fig. 6.31

Corrosion behavior of alloy 214 (Ni-Cr-Al-Y), alloy S (Ni-Cr-Mo) and alloy 800H (Fe-Ni-Cr) in Ar20O2-0.25Cl2 for 400 h at 700–1000 °C (1290–1830 °F). Source: Ref 39 and Ref 41

The effects of adding O2 to HCl environment were extensively studied by Ihara et al. (Ref 47) on iron, by Ihara et al. (Ref 48) on nickel, and by Ihara et al. (Ref 49) on chromium. Adding O2

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to HCl significantly accelerates the corrosion of iron (Ref 47), as shown in Fig. 6.42. The increased corrosion was attributed to the formation of highly volatile FeCl3 (Ref 47). For nickel, adding O2 to HCl does not significantly affect the metal’s corrosion rate, as shown in Fig. 6.43 (Ref 48). This is due to the formation of primarily NiCl2 whether it is 100% HCl or O2-HCl mixtures (Ref 48). Adding O2 to HCl suppresses the corrosion rate of chromium at lower temperatures (400–600 °C) by forming Cr2O3 and accelerates the corrosion rate at higher temperatures (700 and 800 °C) by forming highly volatile CrCl3 (Ref 49). The corrosion rate of chromium as a function of O2-HCl

mixtures at different temperatures is shown in Fig. 6.44 (Ref 49). Devisme et al. (Ref 50) investigated the effect of O2 and CO2 in HCl-bearing environments on the corrosion behavior of four nickel-base alloys, alloys C-276, 600, 601, and 214. Corrosion data in terms of metal loss, which was determined by weight change and metallographic examination of the test specimen, are summarized in Tables 6.15 and 6.16. Adding 2% O2 to Ar20HCl significantly increased corrosion attack for alloy C-276 (Ni-Cr-Mo alloy), but significantly decreased corrosion attack for alloy 214 (Ni-Cr-Al alloy). Both alloys 600 and 601 were not significantly affected. Nevertheless, adding

1600 1575

Alloy AC66

Decrease, µm

500 400

Alloy 800H

10CrMo 9 10*

300 Alloy 45TM 200 Alloy 690

100 0 300

350

400

450

500

550

600

650

700

750

800

Temperature, °C

(a) 225

Alloy AC66

200

10CrMo 9 10*

175

Depth, µm

150 125 Alloy 800H

100 75 50

Alloy 45TM Alloy 690

25 0 300 (b)

350

400

450

500

550

600

650

700

750

800

Temperature, °C

Fig. 6.32 Corrosion behavior of 2.25Cr-1Mo steel (10CrMo9 10), alloy 800H, alloy AC66, alloy 45TM and alloy 690 tested for 300 h at temperatures from 300 to 800 °C (572 to 1472 °F) in air-2Cl2; (a) decrease in thicknesses as a function of temperature, and (b) depth of internal corrosion attack as a function of temperature. Source: Ref 42

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10O2 to Ar-5HCl significantly increased corrosion attack for alloys C-276, 600, and 601. For Ar-5HCl, adding 0.5% H2O to the environment caused increased corrosion for all four alloys, particularly severe for alloys 600 and 601. The test results obtained by Devisme et al. (Ref 50) showed that overall, alloy C-276 (Ni-Cr-Mo) performed better than alumina former alloy 214 in reducing environments, such Ar-HCl and ArHCl-H2 (Table 6.15). Addition of 10% CO2 was found have less effect on corrosion attack for four alloys in Ar-20HCl environment at 600 and 700 °C, as shown in Table 6.16. Ganesan et al. (Ref 51) investigated the corrosion behavior of nickel-base alloys (625, 825, and 600) and iron-base alloys (800HT, 316SS, and 347SS) in a combustion environment consisting of N2, 4 and 9O2, 12CO2, 500 ppm SO2 with two levels of HCl (1% HCl and 4% HCl). For both levels of HCl, three nickelbase alloys were significantly more resistant to chloridation attack than iron-base alloys. For three nickel-base alloys, alloys 625 and 600

Fig. 6.33

Scanning electron backscattered image (a) and an x-ray map for Cl (b), showing a typical pit on alloy 800H tested for 300 h at 300 °C (572 °F) in air-2Cl2. Source: Ref 42

were better than alloy 825. Their test results are shown in Fig. 6.45 to 6.47. Smith and Ganesan (Ref 52) conducted further extensive studies on the corrosion behavior of iron-base alloys (Type 316SS, Type 347SS, and alloy 800HT) and nickel-base alloys (alloys 825, 600, and 625) in simulated combustion environments consisting of N2, O2, SO2 and various amounts of HCl at 426, 593, and 704 °C (800, 1100, and 1300 °F). Also included in their studies was the effect of H2O in the environment on the alloys’ corrosion behavior. Table 6.17 summarizes the test results generated from tests conducted in N2-10O2-50 ppm SO2-500 ppm HCl at 426, 482, and 593 °C (800, 900, and 1100 °F) for 1008 h. HCl is known to be more corrosive than SO2 in high-temperature corrosion. Thus, in this environment, corrosion attack is primarily from HCl. The test results show that the environment that contained about 500 ppm HCl was not corrosive at all for Type 316SS, Type 347SS, 800HT, 825, 600, and 625 at temperatures up to 593 °C (1100 °F). When the HCl content was increased to 4%, both Types 316 and 347 showed much higher corrosion rates at 593 °C (1100 °F), while alloys 800HT, 825, 600, and 625 showed little corrosion attack at 593 °C (1100 °F) after 1008 h, as shown in Table 6.18. However, when the temperature was increased to 704 °C (1300 °F), only high-nickel alloys, such as alloys 600 and 625 exhibited good corrosion resistance in the environment containing 4% HCl (Table 6.19). The environment containing 10% HCl became very corrosive to nickel-base alloys 825, 600, and 625 even at 593 °C (1100 °F) (Table 6.18). At high temperatures, nickel- and cobalt-base alloys, while exhibiting low metal loss, suffered more internal attack in O2-HCl environments. Elliott et al. (Ref 53) examined the corrosion attack in terms of the metal loss and internal penetration for nickel-base alloys (alloys 214, 600, and 601) and cobalt-base alloys (alloys 25 and 188) along with Fe-Ni-Co-Cr alloy (alloy 556), Fe-Ni-Cr alloy (alloy 800H), and Type 310SS after testing in Ar-5.5O2-1HCl-1SO2 at 900 °C (1650 °F) for 800 h with thermal cycling to 200 °C (390 °F) every 100 h. Although the environment contained SO2, hydrogen chloride (HCl), which is known to be more corrosive than SO2, was the primary corrodent causing the corrosion attack. All alloys, while exhibiting low metal losses except Type 310SS, suffered significant internal penetration attack. These test results are shown in Fig. 6.48.

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6.3.4 Corrosion in HCl and HCl-Bearing Reducing Environments Brown et al. (Ref 22) reported corrosion rates of various commercial alloys in a dry HCl environment (Table 6.20). Test duration varied

from 2 to 20 h. Thus, extrapolation to a year would yield an unreliable corrosion rate. Hossain et al. (Ref 54) performed long-term tests in HCl on several nickel alloys and one stainless steel (Type 310SS). Their results are summarized in Table 6.21 and Fig. 6.49. Type 310SS was found

(a)

(d)

(b)

(e)

(c)

(f)

Fig. 6.34

(a) Scanning electron backscattered image of the corrosion products formed on alloy AC66 tested at 800 °C for 300 h in air2Cl2 and the x-ray maps showing elemental distribution for (b) chlorine, (c) chromium, (d) oxygen, (e) iron, and (f) nickel. Source: Ref 42

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to be the worst among the alloys tested. All nickel alloys (Ni-201, 601, 625, C-4, B-2, and 600) were much better than Type 310SS. In this HCl test environment containing no O2, the molybdenum-containing nickel-base alloys, such as alloys 625 (Ni-22Cr-9Mo-3.5Nb) and C-4 (Ni-16Cr-15.5Mo), were found to be the

best performers. In a study by Hossain et al. (Ref 54), nickel performed reasonably well in HCl until the temperature reached 700 °C (1290 °F). At 700 °C, nickel was inferior to many nickel-base alloys, such as alloys 600, 625, and C-4 (Table 6.21). Alloy 400 was found to be very susceptible to chloridation

(a)

(d)

(b)

(e)

(c)

(f)

Fig. 6.35

(a) Scanning electron backscattered image of the corrosion products formed on alloy 690 tested at 800 °C for 300 h in air2Cl2 and the x-ray maps showing elemental distribution for (b) chlorine, (c) chromium, (d) iron, (e) oxygen, and (f ) aluminum. Source: Ref 42

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attack in HCl (Ref 54). Alloy 400 specimens (6 mm diam × 12 mm length) were completely destroyed after exposure for 100 h at 400 °C (750 °F) in HCl. Alloy 600 is commonly used for high-temperature service in HCl environments. Figure 6.50 shows the corrosion rates of alloy 600 in HCl gas as a function of temperature (Ref 22). MTI Publication MS-3 (Ref 27) suggests corrosion guidelines for Ni200, alloy 600, alloy 400, Type 304, and steel in HCl environments, as shown in Fig. 6.51. In reducing environments, such as Ar-4H24HCl, investigated by Baranow et al. (Ref 35), Ni-Cr-Mo alloys, such as alloys C-276 and S,

were significantly better than alloys 600, 625, 188, and X. Their test results generated from the 8 h tests at 900 °C (1650 °F) in Ar-4H2-4HCl are shown in Fig. 6.52. In general, alloys suffered very little metal losses, but suffered significant internal penetration attack. Figure 6.53 shows the cross sections of the specimens in various iron- and nickel-base alloys after testing in Ar4H2-4HCl at 900 °C (1650 °F) for 8 h (Ref 55). A study was performed by Brill et al. (Ref 56) investigating the chlorination resistance of nickel and nickel-base alloys in H2-10HCl. Pure nickel, Ni-Mo, and Ni-Cr-Mo alloys containing little or no iron were found to be more resistant

Alloy C-2000

500

Decrease, µm

400 Alloy HR160

300

40 30 20

Alloy 59

10 0 300

350

400

450

500

550

600

650

700

750

800

Temperature, °C

(a) 22 20

Alloy C-2000 Alloy 59

18 16

Depth, µm

14

Alloy HR160

12 10 8 6 4 2 0 300

(b)

Fig. 6.36

350

400

450

500

550

600

650

700

750

800

Temperature, °C

Corrosion behavior of alloys 59, C-2000, and HR160; tested for 300 h at temperatures from 300 to 800 °C (572 to 1472 °F) in air-2Cl2; (a) decrease in thicknesses as a function of temperature, and (b) depth of internal corrosion attack as a function of temperature. Source: Ref 43

pg 174


Fig. 6.37

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(a) Scanning electron backscattered image of the corrosion products formed on alloy 59 tested at 650 °C for 300 h in air-2Cl2 and the x-ray maps showing elemental distribution for (b) chromium, (c) chlorine, (d) molybdenum, (e) iron, (f ) nickel, and (g) oxygen. Source: Ref 43

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than some nickel-base alloys with higher iron contents. This is illustrated in Fig. 6.54, showing alloy 205 (pure nickel), B-2 (Ni-28Mo), C-4 (Ni-16Cr-16Mo), and 59 (Ni-23Cr-16Mo) being much more resistant than alloys 625 (Ni-22Cr9Mo-3Fe) and 600H (Ni-16Cr-9Fe) (Ref 56).

Fig. 6.38

In another study by Devisme et al. (Ref 50) on chlorination of Ni-Cr-Mo alloy C-276, NiCr-Fe alloy 600, Ni-Cr-Fe-Al alloy 601 (1.4Al), and Ni-Cr-Al-Fe alloy 214 in Ar-HCl environments. Their test results conducted at 600 °C in Ar-5HCl, Ar-10HCl, and Ar-20HCl are

(a) Scanning electron backscattered image of the corrosion products formed on alloy HR160 tested at 800 °C for 300 h in air-2Cl2, and the x-ray maps showing elemental distribution for (b) chromium, (c) chlorine, (d) silicon, (e) oxygen, and (f ) nickel. Source: Ref 43

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shown in Fig. 6.55. Barnes (Ref 57) investigated nickel-base alloys in Ar-HCl mixtures with concentrations of HCl varying from 13 to 100% and found that Ni201 (pure nickel) was slightly more resistant than alloy 600 (8% Fe) for all the concentrations tested except 100% HCl, as shown in Table 6.22. Other nickel-base alloys were also included in his test program, with the test results summarized in Table 6.23. Except for pure nickel (Ni201), nickel-base alloys suffered internal attack by formation of internal voids. This surface zone with extensive internal voids was found to be highly depleted in chromium in the matrix near voids. For example, alloy 600 showed about 1.7% Cr and 4.7% Fe in the matrix in the porous zone after testing at 735 °C (1355 °F) for 100 h in Ar-33HCl (Ref 57). Barnes’s test environments were strictly HCl (inlet gas), thus making pCl2 in the environment significantly higher than that if the initial inlet gas contain both H2 and HCl. This is shown in Fig. 6.56, where pCl2 was plotted as a function of temperature for 100% HCl, Ar-33HCl, and H2-30HCl along with several metal chlorides

(Ref 57). The figure shows that there is no significant difference in pCl2 between the 100HCl and Ar-33HCl environments. However, with the

Decrease, µm

400 350 300 250 200 150 Fe3Al 50 TiAl 40 30 20 10 Alloy 214 0 300 350 400 450 500 550 600 650 700 750 800 Temperature, °C

(a) 105 100

Depth, µm

95 30 25 20 15 10 5 0

Fe3Al

Alloy 214

TiAl 300 350 400 450 500 550 600 650 700 750 800

(b)

Temperature, °C

Fig. 6.39

Corrosion behavior of alumina former alloy 214 and intermetallics Fe3Al and TiAl; tested for 300 h at temperatures from 300 to 800 °C (572 to 1472 °F) in air-2Cl2; (a) decrease in thicknesses as a function of temperature, and (b) depth of internal corrosion attack as a function of temperature. Source: Ref 44

Fig. 6.40

Metal loss as a function of time for alloys Fe-20Cr and 800H tested at 927 °C (1700 °F) in (a) Ar20O2-2500 ppm (0.25%) Cl2 and (b) Ar-20CO2-2500 ppm Cl2. Source: Ref 45

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Fig. 6.41

Depth of internal penetration as a function of time (h0.5) for alloy 800H tested at 927 °C (1700 °F) in (a) argon 20% O2, 0.25% Cl2 and (b) argon 20% CO2, 0.25% Cl2. Source: Ref 45

introduction of H2 in the H2-HCl gas mixture, pCl2 becomes significantly lower, as shown in H2-30HCl. Even though chlorine partial pressure (pCl2 ) stays relatively unchanged (Fig. 6.56) thermodynamically, the kinetic of the HCl-metal reaction (corrosion rate) in Ar-HCl mixtures was found to increase with increasing HCl concentration (HCl partial pressure). This is shown in Fig. 6.57 for both Ni201 and alloy 600 tested at 735 °C (1355 °F). It should also be noted that Barnes’s test environments involving mainly HCl were extremely corrosive for nickel alloys tested. In addition to high pCl2 values, the environments were so reducing that no oxides were thermodynamically stable. Thus, in those test environments, the gas-metal reaction mainly involved formation of chlorides (highly volatile corrosion products) and no oxides (nonvolatile solid phases). Figure 6.58 shows nickel chlorides that formed on a Ni201 coupon after testing at 735 °C (1355 °F) in Ar-33HCl (Ref 57). In a simulated waste incineration environment consisting of N2, 12% CO2, 500 ppm SO2, and

1

1

Fig. 6.42

Effect of oxygen in O2-HCl mixtures on the corrosion rate of iron at 300–700 °C (570–1290 °F). Source: Ref 47

Fig. 6.43

Effect of oxygen in O2-HCl mixtures on the corrosion rate of nickel at 400–700 °C (750–1290 °F). Source: Ref 48

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pg 179


1% HCl, Ganesan et al. (Ref 51) observed that nickel-base alloys, particularly alloy 625, performed significantly better than iron-base alloys, such as alloy 800HT, Type 316SS, and Type 347SS, as shown in Fig. 6.59 and 6.60. Although not discussed in the paper, it is believed that the oxygen potential (pO2 ) was high enough to form some oxides, such as Cr2O3. Thus, little corrosion attack was detected for nickel-base alloys even when the test temperature was increased to 704 °C (1300 °F) (Ref 51). Formation of chromium oxides can significantly reduce the corrosion attack in reducing, HCl-bearing

1

800HT

316SS 347SS

Fig. 6.44

Effect of oxygen in O2-HCl mixtures on the corrosion rate of chromium at 400–800 °C (750–1470 °F). Source: Ref 49

Table 6.15 Corrosion of nickel-base alloys in terms of metal loss µm (mils) after 500 h at 600 °C (1112 °F) in the indicated test environments

Fig. 6.45

Weight change as a function of exposure time for nickel-base alloys (alloys 625, 600, and 825) and iron-base alloys (alloy 800HT, 316SS, and 347SS) in N24O2-12CO2-1HCl-500 ppm SO2. Testing was initially performed at 649 °C, then increased to 704 °C, and finally to 760 °C as indicated. Source: Ref 51

Metal loss, μm (mils) Environment

Ar-20HCl Ar-20HCl-2O2 Ar-5HCl Ar-5HCl-10O2 Ar-5HCl-0.5H2O Ar-5HCl-3H2

C-276

600

601

214

60 (2.4) 330 (13.0) 35 (1.4) 120 (4.7) 90 (3.5) 5 (0.2)

150 (5.9) 185 (7.3) 50 (2.0) 140 (5.5) 240 (9.5) 15 (0.6)

150 (5.9) 120 (4.7) 90 (3.5) 160 (6.3) 255 (10.0) 15 (0.6)

260 (10.2) 65 (2.6) 30 (1.2) 55 (2.2) 80 (3.2) 20 (0.8)

800HT

Source: Ref 50

347SS

Table 6.16 Metal loss after 500 h at 600 and 700 °C in the indicated test environments

316SS

Metal loss, μm Environment

Ar-20HCl Ar-20HCl-10CO2 Ar-20HCl Ar-20HCl-10CO2 Source: Ref 50

Temperature, °C

C-276

600

601

214

600 600 700 700

60 110 200 280

150 100 280 195

150 120 225 250

260 230 360 190

Fig. 6.46

Weight change as a function of exposure time for nickel-base alloys (alloys 625, 600, and 825) and iron-base alloys (alloy 800HT, 316SS, and 347SS) in N29O2-12CO2-1HCl-500 ppm SO2. Testing was initially performed at 649 °C, then increased to 704 °C, and finally to 760 °C as indicated. Source: Ref 51


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pg 180


800HT 316SS

347SS

704

Fig. 6.48

Fig. 6.47

Weight change as a function of exposure time for nickel-base alloys (alloys 625, 600, and 825) and iron-base alloys (alloys 800HT, 316SS, and 347SS) in N29O2-12CO2-4HCl-100 ppm SO2. Testing was initially performed at 593 °C, then increased to 704 °C, and to 816 °C, and finally to 927 °C as indicated. Source: Ref 51

Table 6.17 Corrosion rates in terms of metal loss for iron- and nickel-base alloys at 426, 482 and 593, 900 °C (800, 900, and 1300 °F) for 1008 h in N2-10O2-50 ppm SO2-500 ppm HCl

The metal loss and internal penetration for nickelbase alloys (alloys 214, 600, and 601) and cobaltbase alloys (alloys 25 and 188) along with Fe-Ni-Co-Cr alloy (alloy 556), Fe-Ni-Cr alloy (alloy 800H), and Type 310SS tested in Ar-5.5O2-1HCl-1SO2 at 900 °C (1650 °F) for 800 h with thermal cycling to 200 °C (390 °F) every 100 h. Source: Ref 53

Table 6.19 Corrosion rates in terms of metal loss for iron- and nickel-base alloys at 593 and 704 °C (1100 and 1300 °F) in N2-10O250 ppm SO2-4HCl Metal loss 593 °C(a)

Metal loss 426 °C

482 °C

593 °C

Alloy

µm/yr

mpy

µm/yr

mpy

µm/yr

mpy

Type 316 Type 347 Alloy 800HT Alloy 825 Alloy 600 Alloy 625

0.07 0.09 0.07 0.02 0.04 0.02

0.003 0.004 0.003 0.0008 0.002 0.0008

2.54 2.29 1.02 1.27 2.03 1.78

0.1 0.09 0.04 0.05 0.08 0.07

5.08 7.11 5.33 2.54 3.05 2.79

0.2 0.28 0.21 0.1 0.12 0.11

704 °C(b)

Alloy

µm/yr

mpy

µm/yr

mpy


914 1245 74 20 25 16

36 49 2.9 0.8 1.0 0.6

3,810 … 12,039 2,083 41 152

149 … 470 81 1.6 5.9

(a) Test duration: 72 h. (b) Test duration: 192 h. 1.0 µm = 0.001 mm = 0.0394 mil. Source: Ref 52

1.0 µm = 0.001 mm = 0.0394 mil. Source: Ref 52

Table 6.20 Corrosion of alloys in dry HCl(a) Approximate temperature, °C (°F), at which given corrosion rate is exceeded

Table 6.18 Corrosion rates in terms of metal loss for iron- and nickel-base alloys at 593 °C (1100 °F) for 72 h in N2-9O2-12CO2100 ppm SO2-4 and 10HCl Metal loss 4% HCl

10% HCl

Alloy

µm/yr

mpy

µm/yr

mpy


914 1245 74 20 25 16

36 49 2.9 0.8 1.0 0.6

… … … 1066 1219 1549

… … … 42 48 60

1.0 µm = 0.001 mm = 0.0394 mil. Source: Ref 52

Alloy

0.8 mm/yr (30 mpy)

1.5 mm/yr (60 mpy)

3.0 mm/yr (120 mpy)

15 mm/yr (600 mpy)

Nickel 600 B C D l8-8Mo 25-12Cb 18-8 Carbon steel Ni-resist 400 Cast iron Copper

455 (850) 425 (800) 370 (700) 370 (700) 288 (550) 370 (700) 345 (650) 345 (650) 260 (500) 260 (500) 230 (450) 205 (400) 93 (200)

510 (950) 480 (900) 425 (800) 425 (800) 370 (700) 370 (700) 400 (750) 400 (750) 315 (600) 315 (600) 260 (500) 260 (500) 148 (300)

565 (1050) 538 (1000) 480 (900) 480 (900) 455 (850) 480 (900) 455 (850) 455 (850) 400 (750) 370 (700) 345 (650) 315 (600) 205 (400)

675 (1250) 675 (1250) 650 (1200) 620 (1150) 650 (1200) 593 (1100) 565 (1050) 593 (1100) 565 (1050) 538 (1000) 480 (900) 455 (850) 315 (600)

(a) Based on short-term laboratory tests. Source: Ref 22

environments. This is illustrated by the test results generated by Strafford et al. (Ref 58). The authors tested Fe-Cr alloys containing 2, 5, 9, 14, and 25% Cr at 1000 °C (1832 °F) in a

H2-H2O-HCl mixture giving pO2 value of 1.5 × 10−16 atm and pCl2 value of 10−8 atm at the test temperature. Based on the M-O-Cl stability


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diagram at the test temperature, the environment was in the location where Cr2O3, but not FeO, was to form, and FeCl2 was to form. Test results show that Fe-25Cr alloy exhibited no weight loss for the exposure time up to 100 h. This suggests that once the alloy contains sufficient chromium to form a continuous Cr2O3 scale, as in the case of Fe-25Cr, corrosion can be significantly

reduced or prevented. The results also indicate that Fe-Cr alloys with low chromium contents, such as 2%, 5%, and 9%, suffered significant corrosion attack (Fig. 6.61). In an investigation into possible candidate alloys for a process developed by the Bureau of Mines for extracting alumina from Kaolinitic clay, Carter et al. (Ref 59) tested various

Table 6.21 Corrosion of selected alloys in HCl at 400, 500, 600, and 700 °C (750, 930, 1110, and 1290 °F) Metal loss mg/cm2 400 °C (750 °F)

500 °C (930 °F)

600 °C (1110 °F)

700 °C (1290 °F)

Alloy

300 h

1000 h

100 h

300 h

1000 h

100 h

300 h

96 h

Ni-201 601 310 625 C-4 B-2 600

1.19 1.58 3.26 0.74 0.55 0.75 0.93

0.91 1.47 5.16 1.1 1.12 0.76 0.81

1.60 2.57 6.74 2.42 2.09 2.10 1.69

2.89 4.14 13.65 3.78 3.36 2.65 3.31

4.86 9.38 46.60 8.64 7.24 5.87 7.81

11.46 9.01 15.65 6.79 7.31 12.93 7.67

37.7 19.46 32.6 14.6 19.14 62.3 17.3

377 102.5 1025 26.5 34.9 126.4 49.6

Source: Ref 54

Fig. 6.49

Corrosion rates of several iron- and nickel-base alloys in HCl at 400 to 700 °C (750 to 1290 °F). Source: Ref 54


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alloy. Long-term tests (1584 h) at 260 to 380 °C (500 to 720 °F) showed corrosion rates of 0.3 to 8 µm/yr for alloy 625 and 4 to 10 µm/yr for alloy R-41. The results were in agreement with those obtained from short-term tests. Titanium and Ti-0.2Pd were also tested in the same environment for 15 days. The corrosion rates for titanium were found to be 246, 28, and 13 µm/yr at 400, 300, and 200 °C, respectively. Ti-0.2Pd was found to corrode at 1900, 108, 0, and 2 µm/yr at 500, 400, 300, and 200 °C, respectively. In a study by Reeve (Ref 60) involving 80% HCl and 20% H2O, corrosion rates were obtained for steel and stainless steels (Fig. 6.62). Mild steel was found to suffer corrosion rates of less than 0.76 mm/yr (30 mpy) up to 400 °C (752 °F), and the 18-8 stainless steel, less than 1.0 mm/yr (40 mpy) up to 500 °C (930 °F). These values were significantly lower than those predicted by Brown et al. (Ref 22) in dry HCl (Table 6.20). It appears that addition of significant amount of H2O into HCl environment makes HCl less corrosive.

month

commercial alloys in an environment of 40% HCl and 60% H2O at various temperatures up to 500 °C (930 °F). The results of iron- and nickelbase alloys are summarized in Tables 6.24 and 6.25. Corrosion data at temperatures below 200 °C (390 °F) are not included because of dew point corrosion. Stainless steels and nickel alloys tested showed low corrosion rates at all test temperatures in this environment. A cobalt-base alloy (alloy 188) and an Fe-Ni-Co-Cr alloy (Multimet alloy) were also tested in the temperature range of 315 to 375 °C (600 to 710 °F), showing no measurable attack for alloy 188 and a corrosion rate of only 0 to 3 µm/yr for Multimet

6.4 Corrosion in F2- and HF-Bearing Environments 6.4.1 Corrosion in F2 Environments Fig. 6.50

Early studies of corrosion in fluorine gas for various metals and alloys were carried out by

Corrosion rate of alloy 600 in HCl as a function of temperature. Source: Ref 22

Corrosion rate, mm/yr 0.025 700

0.05

0.10

0.25

Tubes / Internals

0.50

1.00

2.5 1290

Vessels / Pipes 1110

600 Nickel 200 and Alloy 600

400

1020 930

Type 304

300

750

Alloy 400

572

Carbon steel

200

Temperature, °F

Temperature, °C

500

100

390

0 1

2

4

6

10

20

40

60

100

Corrosion rate, mpy

Fig. 6.51

MTI corrosion guidelines for Ni200, alloy 600, alloy 400, Type 304SS and steel in HCl as a function of temperature. Source: Ref 27. Courtesy of Materials Technology Institute

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Metal loss

214 Continuous internal penetration

C-276

S

601

R-41

600

625

188

X 0

40

80

120 160 200 240 280 320

360

Corrosion rate, mils/month

Fig. 6.52

Corrosion rates in terms of metal loss and internal penetration for nickel- and cobalt-base alloys at 900 °C (1650 °F) in Ar-4HCl-4H2. Data was based on 8 h tests. Source: Ref 35

Myers and DeLong (Ref 61). Their test duration was short, normally 4 h. The longest test run was 15 h. The corrosion rates they obtained are summarized in Table 6.26. These rates, extrapolated from short-term tests, are not recommended for estimating the service life of process equipment. Rather, the data are useful for making performance comparisons among various alloys. The results obtained by Myers and DeLong (Ref 61) indicated that nickel had good corrosion resistance in fluorine gas at temperatures up to 500 °C (930 °F). Even at temperatures higher than 500 °C (930 °F), corrosion rates for nickel were significantly lower than those of other alloys. Nickel is commonly used for plant equipment handling fluorine at temperatures up to 500 °C (930 °F) (Ref 62). The resistance of nickel to fluorine gas was attributed to the formation of an adherent nickel fluoride scale (Ref 62). The reaction of nickel and fluorine was found to follow a parabolic rate law (Ref 63).

Hauffe (Ref 64) summarized the results of corrosion tests by Myers and DeLong (Ref 61), Jarry et al. (Ref 62), and Lukyanchev et al. (Ref 65) in Table 6.27. The tests were very short in duration, from 30 min to 32 h. The corrosion rates obtained by Myers and DeLong were two orders of magnitude higher than those observed by Jarry et al. and Lukyanchev et al. Based on the corrosion rates observed by Jarry et al. and Lukyanchev et al., nickel exhibited good resistance in fluorine at temperatures up to 810 °C (1490 °F) (about 12 mpy). Hale et al. (Ref 66) tested six different types of nickel with various degrees of purity at 590 °C (1100 °F) for 95 h in fluorine and found no measurable corrosion for all the samples, except one with a slightly higher silicon content (about 8 mils of general corrosion). The results by Hale et al. (Ref 66) are shown in Table 6.28. They also found that the material’s purity played an important role when tested at 700 °C (1290 °F) for 210 h (Ref 66). Both high-purity vacuum-melted nickel and electrolytic nickel showed negligible attack. Low-carbon nickel I (about 0.015% Si), carbonyl nickel, and “A” nickel (commercial grade pure nickel) began to show intergranular void formation at 700 °C (1290 °F). At the same temperature, low-carbon nickel II (about 0.05% Si) suffered significant general corrosion and intergranular attack (void formation). Silicon appears to be a detrimental impurity in pure nickel in resisting corrosion attack by fluorine. Results obtained by Steindler and Vogel (Ref 67) showed that “A” nickel (commercial pure nickel) suffered significant corrosion at 750 °C (1380 °F), as shown in Table 6.29. A general review on the kinetic aspects of nickel-fluorine reactions can be found in Ref 8. Additions of alloying elements to nickel generally are detrimental to fluorine corrosion resistance. Many nickel-base alloys were found to be significantly more susceptible than nickel to fluorine corrosion (Ref 61, 66, 67). Alloy 600 (Ni-15.5Cr-8Fe), Inco 61 weld wire (Ni-1.5Al-3Ti), Ni-O-Nel (Ni-21Cr-31Fe-3 Mo-1.75Cu), INOR-1 (Ni-20Mo), INOR-2 (Ni5Cr-16Mo), INOR-3 (Ni-16Mo-1Al-1.6Ti), INOR-4 (Ni-17Mo-1.7Fe-2Al-1.7Ti), INOR-5 (Ni-13Mo-2.7W-2.2Cb+Ta-1Mn), Hastelloy W (Ni-25Mo-5.5Fe-2.5Co), Hastelloy B (Ni25Mo-6Fe), HyMa 80 (Ni-16Fe-4Mo), and Monel (Ni-30Cu) suffered significantly more corrosion than nickel in fluorine at 590 °C (1100 °F), as shown in Table 6.28 (Ref 66).

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pg 184


Fig. 6.53

Internal penetration in terms of voids for various iron- and nickel-base alloys after testing in 900 °C (1650 °F) for 8 h in Ar-4H2-4HCl. Source: Ref 55

0

Metal loss, µm

–50

–100

–150

–200

Fig. 6.54

Corrosion of nickel-base alloys in H2-10HCl at 850 °C with 24 h cycles. Source: Ref 56

–250

5% HCl 10% HCl 20% HCl

–300 C276

Dura-Nickel (Ni-3Ti-1.5Al), 70Cu-30Ni, and Ni-10Co suffered slightly more corrosion than nickel at 590 °C (1095 °F) (Ref 66). Fluorides of molybdenum, tungsten, titanium, and other elements have low melting points and/or high vapor pressures (Table 6.2). Iron is significantly less resistant to fluorine attack than nickel. Myers and DeLong (Ref 61) observed that Armco Iron is resistant to fluorine

600

601

214

Alloy

Fig. 6.55

Corrosion in terms of loss of sound metal (µm) of alloys C-276, 600, 601, and 214 in Ar-5HCl, Ar10HCl, and Ar-20HCl at 600 °C for 500 h. Source: Ref 50

at temperatures up to 250 °C (480 °F). Steels were found to be less resistant than Armco Iron (Ref 61). The concentration of silicon appears


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Table 6.22 Metal loss rate(a) in Ar-HCl and 100HCl environments at 735 °C (1355 °F)(b) Metal loss Alloy

Ni201 Alloy 600 Ni201 Alloy 600 Ni201 Alloy 600 Ni201 Alloy 600

Environment

mm/yr

mpy

Ar-13HCl Ar-13HCl Ar-33HCl Ar-33HCl Ar-52HCl Ar-52HCl 100HCl 100HCl

2.8 3.5 5.3 7.9 9.6 11.4 18.4 14.2

110 138 209 311 378 449 724 559

(a) Metal loss rate did not include internal void penetration. Internal void penetration was a small portion of the total metal loss. (b) Test durations from 15–100 h. Source: Ref 57

to be important in affecting the steel’s fluoridation resistance, as shown in Table 6.26 (Ref 61). SiF4 has a very low melting point and high vapor pressure (Table 6.2). Ferritic and austenitic stainless steels, except Type 347SS, showed negligible corrosion at 200 and 250 °C (390 and 480 °F) (Table 6.26). At higher temperatures, corrosion of these alloys became significant. Jackson (Ref 68) also reported significant corrosion rates for several austenitic stainless steels at 370 °C (700 °F), as

Table 6.23 Estimated metal loss rates for pure nickel and nickel-base alloys tested at 685, 735, and 785 °C (1265, 1355, and 1445 °F) in 100HCl Metal loss Alloy

Ni201 Alloy 600 Alloy HR160 Alloy 214 Alloy 602CA Alloy HR160 Ni201 Alloy 600 Alloy HR160 Source: Ref 57


mm/yr

mpy

685 (1265) 685 (1265) 685 (1265) 686 (1265) 686 (1265) 735 (1355) 785 (1445) 785 (1445) 785 (1445)

4.2 7.5 7.0 5.7 28.6 45.6 16.7 26.4 96.0

165 295 276 224 1130 1800 657 1040 3780

Fig. 6.57

Metal loss rate (mm/yr) as a function of HCl concentrations (pHCl) in Ar-HCl mixtures at 735 °C (1355 °F). Source: Ref 57

H2-30% HCI

Fig. 6.56

Thermodynamic equilibrium chlorine partial pressure (pCl2 ) as a function of temperature for several environments (100HCl, Ar-33HCl, and H2-30HCl) and several chlorides. Source: Ref 57

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pg 186


Fig. 6.58

Nickel chlorides formed on Ni201 after testing at 735 °C (1355 °F) for 15 h in Ar-33HCl. Source: Ref 57

347SS

316SS 316SS 347SS

Fig. 6.59

Mass change as a function of time for nickeland iron-base alloys tested initially at 593 °C (1100 °F), then increased to 649 °C (1200 °F), and finally to 704 °C (1300 °F) in N2-12%CO2-500ppm SO2-1%HCl. Source: Ref 51

Mass change as a function of time for nickel- and iron-base alloys tested at 593 °C (1100 °F) in N212%CO2-500ppmSO2-1%HCl. Source: Ref 51

shown in Table 6.30. Cobalt and cobalt-base alloys are not as resistant to fluorine as nickel (Ref 66) (Table 6.28). Limited data for other metals, such as copper, aluminum, and magnesium, are shown in Tables 6.26, 6.29, and 6.31. Aluminum was resistant to fluorine at temperatures up to approximately 500 °C (930 °F) (Table 6.26 and 6.31). AlF3 has a relatively high melting point (1197 °C or 2187 °F) (Ref 64). Corrosion of nickel, alloy 400, and alloy 600 by various volatile metallic fluorides was investigated by Vogel et al. (Ref 69) in their studies on

volatile fission product fluorides, which were associated with the development of a process to recover uranium and plutonium from partially spent nuclear reactor fuels. Most of these fission product fluorides were much less corrosive than fluorine gas. However, whenever fluorine gas was present along with the fluoride, the corrosion rate was generally more aggressive. Nickel and alloy 400 exhibited relatively low corrosion rates for all the volatile fluorides at 500 °C (930 °F), except TeF6, as shown in Table 6.32. Alloy 600, on the other hand, suffered extremely high corrosion rates.

Fig. 6.60


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6.4.2 Corrosion in HF Environments Hydrogen fluoride is generally less corrosive than fluorine for most metals and alloys. However, it is still very corrosive at elevated temperatures. Corrosion rates of various metals and

Fig. 6.61

Weight change as a function of exposure time at 1000 °C (1832 °F) in a H2-H2O-HCl mixture giving pO2 value of 1.5 × 10−16 atm and pCl2 value of 10−8 atm for Fe-Cr alloys containing various amounts of chromium. Source: Ref 58

Table 6.24 Corrosion of iron-base alloys in 40HCl-60H2O Corrosion rate(a) μm/yr Temperature, °C (°F)

500 (930) 400 (750) 300 (570) 210 (410) 200 (390)

316LSS

29-4SS

430SS

E-Brite 26-1

1020 steel

4130 steel

483 15 3 … 7

… … 10 3 …

… … … 8 …

363 5 2 … 3

2100 326 700 … 46

1700 406 870 … 51

(a) Linearly extrapolated from 15 d (360 h) laboratory tests. Note: mpy = (µm) × 0.0394. Source: Ref 59

Table 6.25 Corrosion of nickel-base alloys in 40HCl-60H2O

alloys in HF based on short-term tests were reported by Myers and DeLong (Ref 61), as shown in Table 6.33. Nickel was the most resistant among the alloys tested. Copper, alloy 400, and alloy 600 were slightly worse than nickel. Carbon steels and stainless steels showed poor resistance. Tyreman and Elliott (Ref 70) found that nickel, cobalt, copper, and molybdenum were more resistant to HF corrosion than iron, chromium, niobium (columbium), and tantalum (Fig. 6.63). Chromium was found to be detrimental to resistance to corrosion by HF for Ni-Cr alloys (Fig. 6.64) (Ref 71). Chromium fluorides were the major corrosion products for Ni-Cr alloys tested in HF (Ref 71). Marsh (Ref 72) did an extensive investigation for his Ph.D. thesis on the resistance of several nickel-base alloys to corrosion attack in HF environments at elevated temperatures. In NiCr alloys tested in HF at 650 °C (1200 °F), chromium fluorides were found to form on the metal surface as well as internally as internal phases. This is illustrated in Fig. 6.65, showing elemental distribution for Ni-40Cr alloy tested at 650 °C (1200 °F) in anhydrous HF environment (Ref 72). Surface corrosion products were found to be enriched in chromium and fluorine, but depleted in nickel, and similar elemental distribution was observed for the internal phases. Several commercial nickel-base alloys were also tested. Figure 6.66 shows corrosion attack of alloy 600 in HF after 92 h at 650 °C (1200 °F). In this case, both the corrosion products formed on the metal surface and the internal phases were found to consist of both CrF3 and FeF2, as determined by x-ray diffraction analysis (Ref 72). In Ni-Cr alloys, it appears that NiCr-Mo alloy 625 was more resistant to HF corrosion attack than Ni-Cr alloys X750 and

Corrosion rate(a), μm/yr Temperature, °C (°F)

500 (930) 380 (720) 380 (720) 380 (720) 375 (710) 375 (710) 365 (690) 350 (660) 350 (660) 315 (600) 310 (590) 290 (550) 260 (500)

Ni-200

600

625

R-41

B-2

… 13 15 15 25 13 20 3 18 13 … … 30

… 58 46 58 48 48 46 30 33 15 … … 13

132 … 13 23 10 15 13 8 15 8 0 … 8

… 3 3 3 5 0 3 3 5 0 0 … 3

… 15 18 20 13 28 33 … 36 … … … 28

(a) Linearly extrapolated from 15 d (360 h) laboratory tests. Note: mpy = (µm) × 0.0394. Source: Ref 59

Fig. 6.62

Corrosion rates of mild steel, Type 304SS, and alloy 800 in 80HCl-20H2O at 300 to 600 °C (570 to 1110 °F). Source: Ref 60

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Table 6.26

Corrosion of various metals and alloys in fluorine Corrosion rates(a), mpy

Materials

200 °C

250 °C

300 °C

350 °C

400 °C

450 °C

500 °C

600 °C

650 °C

700 °C

Nickel 400 600 Copper Aluminum Magnesium 430SS 347SS 309CbSS 310SS Armco Fe Steel (0.007% Si) 1020 (0.22% Si) 1030 (trace Si) 1030 (0.18% Si) 1015 (0.07% Si) Music wire (0.13% Si)

… … … … … Nil 8.4 Nil Nil Nil Nil Nil 456 24 … … …

… … … … … Nil Nil 1,740 Nil Nil 24 192 5,760 96 … … …

… … … … … Nil 3,060 2,556 900 372 108 48 7,920 108 9,000 9,960 4,800

… … … … … … 936 6,204 5,544 4,248 96 2.4 1,764 Nil … … …

8 6 456 1,920 Nil … 936 9,540 7,980 6,732 288 144 6,480 180 … … …

23 18 1,152 … Nil … … … … … 3600 … 18 in. 6,480 … … …

61 24 744 1,440 156 … … … … … 139 in. … … 238 in. … … …

348 720 2,040 11,880 216 … … … … … … … … … … … …

192 960 1,560 … … … … … … … … … … … … … …

408 1,800 6,120 … … … … … … … … … … … … … …

(a) Note: mpy × 0.0254 = mm/yr. Source: Ref 61

Table 6.27 Corrosion of nickel in fluorine at various temperatures Temperature, °C (°F)

300 (570) 300 (570) 400 (750) 400 (750) 400 (750) 500 (930) 500 (930) 500 (930) 550 (1020) 600 (1110) 660 (1220) 720 (1330) 810 (1490)

Table 6.28 Results of corrosion tests in fluorine at 590 °C (1100 °F) for 95 h

Pressure, atm

Test duration, h

Corrosion rate, mm/yr (mpy)

Ref

0.9 0.9 0.9 0.9 0.99 0.9 0.9 0.99 0.99 0.9 0.99 0.99 0.99

8 32 0.5 28 … 0.25 30 … 2 32 1.5 2 1.7

0.005 (0.2) 0.00073 (0.03) 0.015 (0.59) 0.0012 (0.05) 0.21 (8.3) 0.018 (0.7) 0.003 (0.12) 1.5 (59.1) 0.036 (1.4) 0.017 (0.7) 0.150 (5.9) 0.240 (9.4) 0.310 (12.2)

62 62 62 62 61 62 62 61 65 62 65 65 65

Source: Ref 64

601 containing no significant amounts of molybdenum, as illustrated in Fig. 6.67 (Ref 71). Corrosion attack in HF for alloys 601, 625, and N is shown in Fig. 6.68 to 6.71. Alloy N (Ni-5Cr16Mo) was found to be significantly more resistant to HF corrosion than nickel-base alloys 600, 601, and 625. This suggests that nickel-base alloys containing low chromium and high molybdenum would be more resistant to HF corrosion attack than Ni-Cr alloys. Molybdenum showed little corrosion attack in HF after 10 h at 850 °C (1560 °F), as shown in Fig. 6.63 (Ref 70). The results of Zotikov and Semenyuk (Ref 73) indicated that both molybdenum and tungsten were quite resistant to corrosion attack by HF at temperatures from 300 to 800 °C (up to 700 °C for tungsten), as illustrated in Table 6.34. Molybdenum appears

Alloy

Depth of corrosion attack(a), mm (mils)

High-purity nickel sheet High-purity nickel rod Low-carbon nickel (II) Low-carbon nickel (I) Carbonyl nickel Electrolytic nickel “A” nickel 70-30 Cupronickel Inconel (low carbon) Inco “61” weld wire Duranickel Ni-O-NEL

NM NM 0.20 (8) NM NM NM NM 0.06 (2.5) 0.64 (25) 0.64 (25) 0.10 (4) >0.42 (16.5)

INOR-1

>0.83 (32.5)

INOR-2

>0.83 (32.5)

INOR-3

>0.83 (32.5)

INOR-4

>0.81 (32)

INOR-5 Hastelloy alloy B

0.76 (30) >0.48 (19)

HyMa 80 90Ni-10Co 80Ni-20Co Monel

0.37 (14.5) 0.06 (2.5) >0.13 (5) >0.80 (31.5)

Hastelloy alloy W

>0.79 (31)

310SS

>0.60 (23.5)

Carpenter 20

>0.66 (26)

Haynes alloy No. 25

>0.34 (13.5)

Cobalt

>1.52 (60)

(a) NM, not measurable. Source: Ref 66

Comments

… … General attack … … … … General attack General attack General attack General attack Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides General attack Completely converted to fluorides General attack General attack General attack Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides Completely converted to fluorides


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to improve the corrosion resistance of nickelbase alloys in HF. The results published by International Nickel Company (Ref 74) showed that alloy N (Ni-16Mo-5Cr) and alloy B (Ni25Mo) were slightly better than nickel in HF (Table 6.35). Field testing in an HF-bearing environment showed that alloy N and alloy S (Ni-16Cr-15Mo) were better than alloys 625 and C-22 at 900 °C (1650 °F) (Ref 75). Barnes (Ref 76) investigated the corrosion behavior of a wide range of materials in HF environments at 1000 °C (1832 °F). Tests were Table 6.29 Corrosion of several metals and alloys in fluorine(a) Metal

Test temperature, °C (°F)

Time, h

Weight gain, mg/cm2

Corrosion rate, mm/y (mils/y)

Ni(b) 600 Cu Ni(b) 400 600 Cu Ni(b) 400 600 600 Cu

550 (1,020) 550 (1,020) 550 (1,020) 650 (1,200) 650 (1,200) 650 (1,200) 650 (1,200) 750 (1,380) 750 (1,380) 750 (1,380) 750 (1,380) 750 (1,380)

6.2 5.0 2.42 5.28 6.0 5.6 4.8 4.1 6.1 4.7 4.7 5.8

1.9 706 17.4 21.5 16.0 1,743 134 100 −1, 831 4, 907 −12, 220 278

0.11 (4.4) 81 (3,200) 2.8 (110) 1.5 (59) 1.0 (41) 180 (7,100) 10.9 (430) 9.0 (353) 74 (2,900) 610 (24,000) 660 (26,000) 19 (750)

(a) Tests were conducted in flowing gas (30 to 130 cc/min). (b) Nickel “A” (commercial grade pure nickel). Source: Ref 67

Table 6.30 Corrosion of several alloys in fluorine(a) Corrosion rate, mm/yr (mpy) Alloy

Exposure time, h

400

Ni-200

304 304L 347 Illium “R” 600

5 24 120 5 24 120 5 24 120 5 5 5

200 °C (400 °F)

370 °C (700 °F)

540 °C (1000 °F)

0.013 (0.5) 0.048 (1.9) 0.76 (29.8) 0.013 (0.5) 0.043 (1.7) 0.29 (11.3) 0.003 (0.1) 0.031 (1.2) 0.18 (7.2) 0.084 (3.3) 0.043 (1.7) 0.62 (24.5) 0.013 (0.5) 0.031 (1.2) 0.41 (16.1) 0.003 (0.1) 0.010 (0.4) 0.35 (13.8) 0.155 (6.1) 40 (1565) … 0.191 (7.5) 153 (6018) … 0.65 (25.4) … … 0.102 (4.0) 108 (4248) … 0.152 (6.0) 0.32 (12.7) 103 (4038) 0.015 (0.6) 2.0 (78.0) 88 (3451)

(a) Tests were conducted in flowing fluorine. Source: Ref 68

Table 6.31

Corrosion of aluminum in fluorine

Test temperature, °C (°F)

26 (79) 201 (394) 356 (673) 543 (1009) Source: Ref 68


0.00087 (0.03) 0.0003 (0.01) 0.57 (22) 2.2 (87)

conducted in Ar-5HF, Ar-15HF, and Ar-35HF mixtures. In nickel-base alloys, alloys B-3 (Ni28Mo-1.5Cr) and 242 (Ni-25Cr-8Mo) were found to be significantly better than alloys 600, 617, and 602CA. His test results are summarized in Table 6.36 and Fig. 6.72. Both Ni-Mo alloys containing low Cr formed thin surface scales with very little internal fluoride penetration. Surprisingly, alloy 188, a cobalt-base alloy with high chromium (22%) and high tungsten (14%), also exhibited good corrosion resistance with a thin surface scale and little internal fluoride penetration. Nickel aluminide intermetallic (Ni-8Al-8Cr-1.4Mo-1.7Zr) was found to exhibit good corrosion resistance in Ar-5HF. However, when tested in Ar-35HF, the nickel-aluminide Table 6.32 Corrosion of nickel and nickel-base alloys by various volatile fluorides at 500 °C (930 °F) Corrosive environment

Test duration, h

GeF4 + F2 AsF5 AsF5 + F2 SeF6 SeF6 SeF6 + F2 MoF6 MoF6 + F2 MoF6 + F2 TeF6 TeF6 TeF6 + F2 SF6 UF6 F2 F2

8.4 7.0 7.2 7.0 29.5 6.6 9.2 6.0 7.0 18.9 5.7 7.8 28.7 28.8 7.0 8.6

Corrosion rate(a), mm/yr (mpy) Ni-200

Alloy 400

Nil(b) 0.22 (8.8) 0.22 (8.8) Nil 0.44 (17.5) 0.67 (26.3) 0.22 (8.8) 0.22 (8.8) (c) (c) 0.22 (8.8) 0.22 (8.8) 0.22 (8.8) Nil 0.22 (8.8) Nil (c) (c) 8.9 (350) 1.8 (70) 135 (5326) 56 (2190) 0.22 (8.8) 0.22 (8.8) Nil Nil 0.22 (8.8) … 0.66 (26) 0.22 (8.8) … …

Alloy 600

29 (1139) Nil 45 (1770) 0.44 (17.5) … 22 (850) 0.22 (8.8) 44 (1726) 63 (2488) 2.4 (96) 5.6 (219) 40 (1568) Nil … 9.8 (385) 31 (1209)

(a) Calculated from weight loss after descaling. (b) Rates reported as nil are less than 0.001 mils/h. (c) Scale was not completely removed by descaling. Source: Ref 69

Table 6.33 Corrosion of various metals and alloys in anhydrous HF Corrosion rate, mm/yr (mpy) Material

Nickel 400 600 Copper Aluminum Magnesium Carbon steel (1020) 304 347 309Cb 310 430 Source: Ref 61

500 °C (930 °F)

550 °C (1020 °F)

600 °C (1110 °F)

0.9 (36) 1.2 (48) 1.5 (60) 1.5 (60) 4.9 (192) 12.8 (504) 15.5 (612) … 183 (7,200) 5.8 (228) 12. 2 (480) 1.5 (60)

… 1.2 (48) … … … … 14.6 (576) … 457 (18,000) 43 (1,680) 100 (3,960) 9.1 (360)

0.9 (36) 1.8 (72) 1.5 (60) 1.2 (48) 14.6 (576) … 7.6 (300) 13.4 (528) 177 (6,960) 168 (6,600) 305 (12,000) 11.6 (456)


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suffered extensive corrosion attack, while alloy 242 remained resistant under the same test condition. The results indicated that Ni-Al system was not as good as Ni-Mo system in resisting HF corrosion attack. Barnes (Ref 76) also tested various pure metals, which included gold (Au), platinum (Pt), palladium (Pd), chromium, Ni200, Ni201,

Fig. 6.63

Corrosion of various metals in HF and HF-50H2O at 850 °C (1562 °F) for 10 h. Source: Ref 70

Fig. 6.64

Effect of chromium on resistance to HF at 650 °C (1200 °F) for Ni-Cr alloys. Source: Ref 71

Ni270, and copper. The test results on pure metals are summarized in Table 6.37 (Ref 76). Precious metals (Au, Pt, and Pd) showed no weight changes and little changes in specimen surface appearance after testing. However, both gold and platinum specimens were found to exhibit surface etching on the specimen surface. When tested in a higher concentration of HF (Ar-35HF), palladium was heavily corroded with some molten corrosion product formed on the specimen surface. There was no discussion in the paper about the performance of gold and platinum in Ar-35HF mixture. Copper was also found to exhibit few changes in specimen appearance after testing in Ar-5HF. In fact, the copper specimen retained its original metallic luster after testing in Ar-5HF. Chromium was found to be quite susceptible to fluoridation attack when tested in Ar-5HF. The metal suffered extensive internal chromium fluoride penetration and extensive weight loss due to vaporization of chromium fluorides. Nickel was found to be extremely resistant to HF corrosion attack. Three different types of nickel were tested: Ni200 and Ni201 (99% pure nickel), and Ni270 (99.9% pure nickel) (Ref 76). Ni201 is a low carbon nickel (0.02% C), while Ni200 is a high carbon nickel (0.15% C). The impurities and minor elements in these three nickel specimens in the test program are: 0.15C, 0.4Fe, 0.35Si, 0.35Mn, and 0.25Cu for Ni200; 0.02C, 0.4Fe, 0.35Si, 0.35Mn, and 0.25Cu for Ni201; and 0.02C, 0.001Cu, and 0.05 max Fe for Ni270. The test results after 15 h in Ar-5HF are summarized in Table 6.37 (Ref 76). All nickel specimens retained metallic luster appearance after testing. A nickel wire that had been used for holding test specimens during testing of specimens in the Ar-5HF test gas at 1000 °C for 1600 h was sectioned for metallographic examination, showing only a penetration of about 20 µm (0.8 mils) of internal nickel fluoride precipitates. These nickel fluorides, which were randomly distributed, were found to be less than 2 µm in diameter. Some nickel fluorides were also detected on the metal surface. Ceramic materials and graphite were also tested by Barnes (Ref 76) under the same test conditions. The results are summarized in Table 6.38. Graphite was found to be extremely resistant to HF corrosion. Ceramic materials, such as alumina (polycrystalline or single crystal), sintered silicon carbide, chemical vapor deposited (CVD) silicon carbide, and silicon nitride, suffered high corrosion rates.

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Fig. 6.65

Scanning electron backscattered image (a) and x-ray maps for Cr (b), Ni (c), and F (d) for Ni-40Cr alloy tested at 650 °C (1200 °F) for 22 h in anhydrous HF environment. Source: Ref 72. Courtesy of Glyn Marsh

Fig. 6.66

Optical micrograph showing corrosion attack of alloy 600 after testing in HF for 92 h at 650 °C (1200 °F). Source: Ref 72. Courtesy of Glyn Marsh

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Adding H2O to HF appears to make HF less corrosive. Figure 6.63 shows the corrosion data generated by Tyreman and Elliott (Ref 70) by

Fig. 6.69

Optical micrograph showing corrosion attack on alloy 601 after testing in HF for 16 h at 650 °C (1200 °F). Source: Ref 72. Courtesy of Glyn Marsh

Fig. 6.67

Corrosion kinetics of alloys 625, X750, and 601 as a function of time in HF at 650 °C (1200 °F). Source:

Ref 71

Fig. 6.70

Optical micrograph showing very little corrosion attack (about 1 to 2 µm in depth) on alloy N (Ni-5Cr-16Mo) after testing in HF for 46 h at 650 °C (1200 °F). Source: Ref 72. Courtesy of Glyn Marsh

Fig. 6.71

Fig. 6.68

Optical micrograph showing corrosion attack on alloy 625 after testing in HF for 142 h at 650 °C (1200 °F). Source: Ref 72. Courtesy of Glyn Marsh

Scanning electron micrograph showing the corrosion scale (mainly NiF2 with some CrF3, as determined by x-ray diffraction analysis) formed on alloy N after testing in HF for 115 h at 650 °C (1200 °F). Source: Ref 72. Courtesy of Glyn Marsh

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comparing the 100% HF and HF-50%H2O environments for nickel, copper, cobalt, iron, chromium, molybdenum, niobium, and tantalum. All metals except molybdenum showed more corrosion attack in HF than in HF-50H2O. Corrosion data generated in HF-10H2O and HF50%H2O are summarized in Tables 6.39 and 6.40, respectively. 6.4.3 Corrosion in O2-HF Environments Oxygen appears to make HF more corrosive. During nuclear fuel reprocessing, stainless steel fuel cladding was being chemically removed by reacting it with an O2-HF mixture (40–60% HF) at 377 to 627 °C (710 to 1160 °F) (Ref 77). The Table 6.34 Corrosion of molybdenum and tungsten in HF containing 0.6% H2O Corrosion rate, mm/yr (mpy) Test temperature, °C (°F)

Molybdenum

Tungsten

300 (570) 400 (750) 500 (930) 600 (1110) 700 (1290) 800 (1470)

0.003 (0.1) 0.011 (0.4) 0.014 (0.6) 0.023 (0.9) 0.144 (5.7) 1.3 (51)

0.003 (0.1) 0.009 (0.4) 0.017 (0.7) 0.022 (0.9) 0.91 (36) …

removal rates (or corrosion rates) were more than 1 mm/h (40 mils per h (Ref 77). Macheteau et al. (Ref 78) also found that oxygen contamination accelerated fluoridation attack of iron. Marsh and Elliott (Ref 79) observed that cobalt exhibited a protective CoF2 film with a very low corrosion rate when exposed to HF at 650 °C (1200 °F). However, once air was introduced to mix with the test gas of HF (i.e., HF-O2-N2 mixture), the corrosion rate increased significantly. This is illustrated in Fig. 6.73 (Ref 79). The rapid corrosion attack of cobalt when air was mixed with HF was associated with the formation of numerous Co3O4 oxide protrusions, as shown in Fig. 6.74. At the Co3O4 oxide and the cobalt metal interface, no protective fluoride films were observed except remnant segments of CoF2 films were still present as indicated the lower line in Fig. 6.74. It is believed that as soon as the air was introduced to mix with HF gas stream, a protective CoF2 film was “broken,” prompting the formation of Co3O4 oxide protrusions. Schutze and Simon (Ref 80) tested a number of alloys in N2-5O2-3HF at 1100 °C for 75 h in

Source: Ref 73

Table 6.35 Corrosion of various nickel-base alloys in HF at 500 to 600 °C (930 to 1110 °F) Alloy


Alloy C Alloy 600 Alloy B Ni-200 Ni-201 Alloy 400 Alloy K-500 70Cu-30Ni

0.008 (0.3) 0.018 (0.7) 0.051 (2) 0.229 (9) 0.356 (14) 0.330 (13) 0.406 (16) 0.406 (16)

Tests were performed in HF (7 lb/h) for 36 h. Source: Ref 74

Fig. 6.72

Depth of internal fluoride penetration for alloys tested in Ar-5HF at 1000 °C (1832 °F) for 15 h.

Source: Ref 76

Table 6.36 Corrosion behavior of nickel-base alloys along with one cobalt-base alloy (alloy 188) and nickel aluminide Intermetallic (IC221M) in Ar-5HF at 1000 °C (1832 °F) for 15 h Alloy

Alloy 242 Alloy B-3 IC221M Alloy 188 Alloy 617 Alloy 600 Alloy 602CA

Maximum attack(a), mm/side (mils/side)

Specimen appearance

Comments

0.015 (0.6) 0.02 (0.8) 0.025 (1.0) 0.051 (2.0) 0.13 (5.3) 0.15 (5.9) 0.18 (7.1)

Surface scale Surface scale Surface scale Surface scale Nodules on surface Nodules on surface Molten corrosion product

Good resistance at 35HF … Extensive corrosion at 35HF(b) … Extensive internal fluorides Extensive internal fluorides Extensive internal fluorides

(a) Surface metal loss + internal fluoride penetration. (b) IC221M exhibited good resistance in Ar-5HF, but poor resistance in Ar-35HF. The specimen was almost completely converted to fluorides. Source: Ref 76


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Table 6.37 Materials

Corrosion behavior of pure metals in Ar-5HF at 1000 °C (1832 °F) for 15 h Weight change, mg/cm2

Corrosion rate, mpy

Specimen appearance

0 −0.2 0.2 −35.25 −0.1 −0.1 0.0

0.0 1.1 … 1454(b) 1.4 1.4 0.0

Unchanged Matte grey Unchanged Green Unchanged Unchanged Unchanged

Gold Platinum Palladium Chromium(a) Ni200 Ni201 Ni270

Comments

… … Heavily corroded in 35HF Total depth of attack = 8.3 mils No internal attack No internal attack No internal attack

(a)Tested for 50 h at 1000 °C in Ar-5HF. (b) Based on total depth of attack (maximum metal affected) of 8.3 mils/side in a 50 h test. Source: Ref 76

Table 6.38 Corrosion behavior of ceramic materials and graphite in Ar-5HF at 1000 °C (1832 °F) for 15 h Materials

Graphite (CZR-1) Graphite (DFP-2) Alumina (polycrystalline) Alumina (sapphire) CVD SiC Sintered SiC Si3N4


Corrosion rate, mpy

Comments

0 0 −2.4 −0.9 −5.7 −6.4 −2.2

0.0 0.0 141 52 411 458 155

Unchanged Unchanged General wastage Frosted appearance General wastage General wastage General wastage

Source: Ref 76

Table 6.39 Corrosion of several alloys in HF-10H2O at 850 °C (1560 °F) Alloy


304LSS 310SS 800H 600 625

100–130 (4000–5000) 100–130 (4000–5000) 100–130 (4000–5000) 23 (900) 6.7 (265)

Source: Ref 70

Table 6.40 Corrosion of nickel and Alloy 400 in 50HF-50H2O at various temperatures Corrosion rate, mm/yr (mpy)

their search for a candidate alloy for air nozzles in a circulating fluidized-bed combustor for reprocessing spent pot lining material in aluminum production. Among the metallic materials tested, alloy 242 was found to exhibit the smallest metal loss (50 µm). Although the alloy suffered extensive nitridation attack, the authors considered that nitridation attack would not affect the alloy’s performance as air nozzles. In a lowtemperature test (450 °C) in N2-8O2-10CO215H2O-5HF, Crum et al. (Ref 81) observed no measurable corrosion attack for many nickelbase alloys and a superaustenitic stainless steel after 155 h of exposure. Their test results are summarized in Table 6.41. Stress-corrosion cracking (SCC) resistance of these alloys was also included in the test program using the same test environment with U-bend test specimens. All alloys showed no cracking after 100 and 155 h except alloy 600. Authors did not explain why alloy 600 suffered SCC while other alloys including many nickel-base alloys and one superaustenitic stainless steel (alloy 25-6MO) showed no cracking.

6.5 Corrosion in Bromine and Iodine Environments Very little data have been reported on the performance of metals and alloys in bromine and

550 °C (1020 °F)

600 °C (1110 °F)

650 °C (1200 °F)

700 °C (1290 °F)

750 °C (1380 °F)

Nickel 0.79 (31) 1.83 (72) 2.74 (108) 3.66 (144) 3.05 (120) Alloy 400 … 0.61( 24) 1.52 (60) 3.96 (156) 5.18 (204) Source: Ref 61

iodine environments at elevated temperatures. Miller et al. (Ref 82) investigated the corrosion behavior of copper, nickel, and nickel-base alloys in bromine at 300 and 500 °C (470 and 930 °F) (Table 6.42). Copper suffered rapid attack by bromine at 300 °C (570 °F). The CuBr compound exhibits very high vapor pressures. The vapor pressure of CuBr reaches 1× 10−4 atm (a value considered high enough to cause hightemperature corrosion) when the temperature is as low as 435 °C (815 °F) (Table 6.3). Nickel exhibited good corrosion resistance in Br2 at 300 and 500 °C (470 and 930 °F). Bromine reacts with nickel to form NiBr2, which melts at 965 °C (1769 °F), significantly higher than the melting point of CuBr. Duranickel 301 (Ni-5Al) had a corrosion resistance similar to nickel. Alloy 400 was less resistant than nickel and Duranickel 301. Smith and Ganesan (Ref 83) investigated the corrosion behavior of various commercial alloys in a simulated combustion environment containing a very high level of HBr (about 4%) at 593 and 927 °C (1100 and 1700 °F). The test


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gas (the inlet gas) consisted of 9% O2, 12% CO2, 4% HBr, 100 ppm SO2, and balance N2. The equilibrium partial pressures of the test environment at 593 and 927 °C (1100 and 1700 °F) are listed in Table 6.43. The test results, which are summarized in Table 6.44, showed surprisingly very little corrosion attack at 593 °C (1100 °F) for all the alloys tested, which included austenitic stainless steels (Type 309, 316, and

347), alloy 800, and nickel-base alloys. At 927 °C (1700 °F), nickel-base alloys showed little weight loss, while austenitic stainless steels, Fe-Ni-Cr alloy 800, and high-iron-containing nickel alloy 825 suffered high weight losses. The authors (Ref 83) analyzed the corrosion products spalled from test specimens as well as Table 6.41 Corrosion rates generated from tests in N2-8O2-10CO2-15H2O-5HF at 450 °C (842 °F) for 100 and 155 h Alloy

Alloy 600 Alloy 400 Alloy 825 Alloy C-276 Alloy 686 Alloy 622 Alloy 25-6MO

Corrosion rate, mm/yr (mpy) from 100 h tests

Corrosion rate, mm/yr (mpy) from 155 h tests

0.01 (0.4) 0 0.01 (0.4) 0 0 0 0

0 0.002 (0.1) 0 0 0 0 0

Source: Ref 81

Table 6.42 Corrosion of several alloys in bromine at 300 and 500 °C (570 and 930 °F) Alloy

Ni-201

Fig. 6.73

Weight change of cobalt at 650 °C (1200 °F) when exposed to HF during the first 12 h, showing a very low corrosion rate. At time “X,” air was introduced into the test gas to mix with HF, rapid corrosion attack was observed. Source: Ref 79


300 (570) 300 (570) 300 (570) Duranickel 301 300 (570) 300 (570) 400 300 (570) 300 (570) Copper 300 (570)(a) 300 (570)(a) Ni-201 500 (930) 500 (930) Duranickel 301 500 (930) 500 (930)

Time, days

Weight loss, mg/cm2


11 11 11 10 10 2 2

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