Current limitations of high-temperature alloys in practical applications

Oxidation of Metals, Vol. 44, Nos. 1/2, 1995

Current Limitations of High-Temperature Alloys in Practical Applications

John Stringer* and Ian G. Wrightt:~

Received March 24, 1995

There are several engineering systems which require materials of construction to tolerate elevated temperatures, and aggressive environments of one kind or another; and where, furthermore, the performance of the system is limited by the materials capability. This paper reviews a number of these systems, drawn principally from the electric power industry, and describes both the current approaches to improving the materials capability, and potential directions for research and development for the future. Particular emphasis is given to cases where the problems related to oxidation and high-temperature corrosion are of major importance.

KEY WORDS: oxidation; corrosion; high temperatures; erosion; wear.

INTRODUCTION

There are a number of engineering systems within which there are regions of high temperature. These include, for example, metals extraction systems such as blast furnaces; materials treatment systems such as steel heat treatment furnaces or pottery kilns; heat engines such as gas turbines, and high- temperature electrochemical systems such as molten carbonate fuel cells. The chemical environment can also vary widely, from simple oxidizing environments such as air or clean combustion gases with excess air, through mixed oxidant environments, such as clean substoichiometric combustion

*Electric Power Research Institute, Palo Alto, California. tCorrosion Science and Technology, Metals and Ceramics Division, Oak Ridge National

Laboratory, Oak Ridge,Tennessee. ~At the time of the preparation of this paper, Dr. Wright was a Visiting Scientist at the Electric Power Research Institute, on leave from Battelle Columbus Laboratories, Columbus, Ohio.

265 0030-770X/95/0800~0265507.50/0 �9 1995 Plenum Publishing Corporation

266 Stringer and Wright

gases, or combustion gases containing sulfur, to carburizing atmospheres such as those in reformers, or sulfidizing atmospheres such as those in sulfur boilers. The high-temperature environment may be gaseous, as in those examples quoted in the last sentence, liquid, as in molten carbonate fuel cells or glass melters, or solid, as in solid oxide fuel cells. In addition to the potentially aggressive chemical environment, the high-temperature component may be subjected to mechanically aggressive factors, such as high- velocity particle-containing gas streams, as for example in the combustion gas in a pressurized fluidized bed entering the expansion turbine, or the fly-ash laden gas in a conventional pulverized-coal-burning boiler passing through the heat-exchanger tube banks; or low-velocity high-particle-density environments, such as those in the bed of fluidized-bed combustors.

In most of these systems, the high-temperature region is of crucial importance to the overall process, and the materials of construction for this region present significant problems, both in terms of the maximum temperature that can be attained, the maximum size of the hot region, and the lifetime of this part of the system. The critical properties required of the materials of construction may include strength, toughness, repairability, thermal conductivity, chemical stability, and coefficient of thermal expansion, among others; and it is important to recognize that the properties may be critical not only at the maximum system temperature, but at intermediate temperatures and even room temperature. A material selected for resistance to the high-temperature environment may be sensitive to aqueous corrosion or stress-corrosion cracking resulting from condensates during down times, for example; or may have unacceptable low-temperature mechanical properties, such as very low toughness. A further requirement is that the material must be stable, and not exhibit long-term aging processes which degrade either the chemical or mechanical properties. This requirement is very important, since by definition solid-state diffusion processes are relatively rapid at high temperatures, so phenomena such as approach to equilibrium, grain- boundary precipitation, grain-growth, recrystallization, and annealing, must be considered. A well-known example is the tendency in high-alloy steels and similar alloys to precipitate topologically close-packed (TCP) phases, of which the best known is o--phase (sigma-phase).

In this paper, the majority of the examples given will be drawn from the electric power industry, because this is the area of the authors' principal experience; but it is important to recognize that related problems exist in many other industries. Furthermore, most emphasis will be given to the aspects of surface stability: mechanical stability in the cases of erosion, abrasion, and wear; and chemical stability in the cases of oxidation and high-temperature corrosion.

Current Limitations of High-Temperature Alloys in Practical Applications 267

The General Characteristics of the Problem

A designer has a number of alternatives for a system which contains a high-temperature region. These include:

�9 Cool the critical components. This may be done actively, by circulat- ing a coolant through the component at risk, or passively, by lining the structural member with a low-conductivity layer that separates it from the hot zone.

�9 Limit the temperature in the hot zone to a level that can be tolerated by the available materials of construction. This usually carries thermodynamic penalties, but the costs of these must be balanced against the costs of using higher-priced materials.

�9 Moving the most critical components at risk to a lower-temperature part of the system. This can sometimes be done in an initial design, but is usually difficult to do as a remedial action if the problem appears while the system is in service.

�9 If the problem is related to the stability of the structural material in the high-temperature environment, it may be possible to modify the environment. This may mean altering the chemistry or (for example) removing erosive particles.

�9 In the case of surface stability, it may also be possible to use a coating or other surface modification.

�9 If these approaches do not work, or carry unacceptable economic penalties, the designer may select an alternative material from the standard well-understood materials available (replacing a low-alloy steel by a higher-alloy steel, for example).

�9 FinMly, if none of the preceding approaches are effective, it may be necessary to develop a new structural material. In most cases, designers wish to avoid this, because it takes a long time, not only to develop the material, but to acquire sufficient engineering data, to learn about the fabrication methods, and to learn about the main, tainability aspects. However, this approach has been commonly used in the development of gas turbines over the years.

In this paper, the following problems will be addressed:

�9 Hot-section materials in gas turbines. This presents one of the most severe situations for materials in terms of the strength and toughness requirements coupled with the corrosive environment. The approaches combine several elements:

Active cooling of the hot-section components


A passive cooling component, in the use of thermal barrier coatings Coating for resistance to oxidation and corrosion The development of new structural alloys

Materials for large coal-fired boilers. The approaches include

Identification of the root causes of tube failures Selection of superheater tube materials for advanced steam conditions Coextruded superheater tubes for boilers experiencing fireside corrosion Coatings for water-wall tubes experiencing severe corrosion Treatments to reduce steam-side oxide growth and exfoliation

Materials for gas-cooler heat exchangers in coal gasifiers. The approaches include

Moderating the component temperature Use of coatings and claddings New materials

Other possibilities include

Modifying the environment chemistry Modifying the cycle to enable the component to be moved to a location where the conditions are less severe

Hot-Section Materials in Gas Turbines

The gas (or combustion) turbine uses the Brayton cycle. In this, air is first compressed; then heated. The heated gas is then exhausted through an expansion turbine: the turbine drives the compressor, and the excess energy above that required for compression is available to do work. In a utility gas turbine, the shaft drives the generator producing electricity. The gas leaving the turbine is still quite hot, and the efficiency of the system can be improved by making use of this heat in some way. The heating of the compressed air may be direct or indirect. Essentially all utility gas turbines use direct heating, in which fuel is mixed with the compressed air and burnt in a combustion chamber. However, indirectly-fired turbine systems using fluidized bed combustors or high temperature nuclear reactors as sources of heat have been studied: the working fluid is a clean gas which is heated by means of a heat exchanger. The indirectly-fired gas turbine may be open cycle, in which the exhaust from the gas turbine is vented, or (more commonly) closed cycle, in which the exhaust gas is returned to the low-pressure side of the compressor.

As with any heat engine, the efficiency is related to the difference between the maximum and minimum (absolute) temperatures in the working

Current Limitations of High-Temperatore Alloys in Practical Applications 269

cycle, in this case the inlet temperature to the turbine and the exit temperature, divided by the maximum temperature. This is the Carnot efficiency and real cycles achieve only a fraction of this. For example, in a modern gas turbine the turbine inlet temperature may be 1288~ (2350~ and the exhaust temperature may be 593~ (1100~ corresponding to a Carnot efficiency of 44.5%; the actual efficiency of the engine is closer to 34%.

Improvements in efficiency can obviously be made by either increasing the turbine inlet temperature or reducing the exhaust temperature at the back end. The advantage of the gas turbine is that there is considerable range to increase the inlet temperature, since the upper limit is the stoichiometric flame temperature: the practical limitations are in part set by the available materials of construction and the effectiveness of the hot-section materials cooling. Reduction of the exhaust temperature is much more difficult in the simple cycle, and the best approach is to use the sensible heat in the exhaust gas in some other way. In recent years, this has been done by adding a steam cycle (Rankine cycle) by using the exhaust gas to boil water in a heat-recovery boiler; the steam is then expanded through a steam turbine in the usual way. The overall system is then referred to as a combined cycle, and efficiencies as high as 58% have been attained.

The increases in the turbine inlet temperature in the early days of the aviation gas turbine were paced by the capabilities of the alloys available for the hot-section components. These components are, simply, the combustion chamber; the duct that directs the hot gas from the combustion chamber onto the turbine airfoils (this is called the transition piece in the case of a can-type combustor design, and the inlet volute in the case of a silo-type combustor design) ; and the first stages of the turbine itself. A stage consists of a stationary airfoil followed by a rotating airfoil. The stationary airfoil directs the gas flow onto the rotating stage and accelerates it, and in the case of the first stage it is called the inlet nozzle guide vane.

Alloy Developments

The alloys that were developed for the nozzle and the rotor airfoils are generally called superalloys, and may be either cobalt-based or nickel-based. Historically, the progress of the development of these alloys is well-known: a comprehensive description is contained in Sims, Stoloff and Hagel, 1 and another by Gell et al.2; and a recent paper by Molloy 3 summarizes the development and current status of investment-cast superalloys. For both classes of superalloys, the maximum temperature at which acceptable strength could be retained leveled out at approximately 900~ (1652~ which is of the order of 65% of the absolute melting point of the pure metals. As this limit was recognized, materials specialists started to examine the


possibility of using metals with higher melting points: the preferred candi- dates were molybdenum and niobium. Alloys based on these materials did not prove very successful, largely based on the poor oxidation resistance. However, before this problem became crucial, the concept of cooling the airfoils was developed, and a large part of the increase in the turbine inlet temperature to the present levels is attributable to increasing sophistication in cooling. This has carried with it a requirement for increasing sophistication in the casting of the airfoils, particularly the first-stage rotor blades, to form the internal cooling passages.

Although cooling accounted for most of the increase in the turbine inlet temperature over the last 20 years, there has also been some progress in the capabilities of the alloys. 4 A large part of this has addressed the issue of the grain boundaries, since the sites of initial failure were grain boundaries normal to the long axis of the airfoils. The initial progress was the addition of small amounts of zirconium and boron to modify the grain boundaries. The beneficial effects had been recognized for a long time: in 1957 Decker, Rowe and Freeman 5 showed that 0.01% Zr+0.009% B added to Udimet 500 (the nominal compositions of the alloys mentioned in the text are listed in Table I) could increase the creep rupture life [172 MPa, 870~ (25 ksi, 1600~ by a factor of 13, and elongation by a factor of 7. Other changes in both composition and heat treatment have been aimed at modifying the composition and distribution of carbides. However, again a radical change has been made. This was the use of directional solidification to attempt to produce a microstructure in which all the grain boundaries were parallel to the long axis of the blade. There is some controversy about the originator of this approach (see Ref. 4) but generally it is believed that VerSnyder 6 and others at Pratt and Whitney Aircraft in the early 1960s were largely responsible. Some changes in the composition of alloys was required, but generally the directionally solidifed (DS) alloys are modifications of existing superalloys. However, it was soon discovered that the addition of 2% Hf coupled with directional solidification made it possible to use high-strength eutectic compositions. DS alloys are widely used in both aircraft-derivative engines, and in advanced industrial "Frame"-type engines; they would be more widely used but for the need to optimize fabrication which leads to greater costs.

The logical extension of this approach was to grow single-crystal (SC) airfoils. Single crystal alloys offer futher flexibility in alloy design and associated increase in temperature capability for the following reasons7:

�9 The directional solidification process results in large MC-type carbides in some alloys, and these carbides are often precracked and initiate matrix fatigue cracks under cyclic loading conditions. 8'9


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Fig. 1. 10,000 h rupture strength of the single-crystal superalloys CM SX-2 and -3, Nasair 100 and SRR 99 compared to conventionally cast (CC) and directionally solidified (DS) Mar-M 247. These data were taken from various Cannon Muskegon brochures and the Metals Handbook.

�9 If some grain boundaries are not perfectly aligned with the solidification direction, creep cracks may initiate at these boundaries where they intersect a free surface.

�9 The removal of grain boundary strengthening elements, including C, B, Zr, and Hf, might improve the fatigue properties by elimination of MC-type carbides and could increase the incipient melting temperature and therefore the creep resistance, because these elements are melting-point depressants. The increase in the melting point permits solutioning of large amounts of coarse ),' which later precipitates as fine 7', enabling realization of the full strength potential. Addition of Re has also proved a boon in improving the creep strength by partitioning to )/. The improved creep strength realizable from single crystal alloys over DS and conventionally-cast materials is illustrated in Fig. 1. The alloy Nasair 100, which is a single crystal derivative of Mar M 247, shows a 22~ (40~ temperature increase capability with respect to the conventionally-cast Mar M 247, and about 17~ (30~ with respect to DS Mar M 247.

�9 In the absence of grain boundaries, more flexibility in alloying might be achieved that would result in an optimum balance of creep-rupture strength and oxidation and hot corrosion resistance. Development of leaner alloys also improves castability.

Single crystal blades (alloy PWA 1480) were initially put into commercial service in 1982 in Pratt and Whitney's engine in a Boeing 767 aircraft and


Airbus A310. Howmet, the market leader in SC casting for aircraft components, has shipped more than 10,000 single-crystal castings. All major aircraft engine manufacturers have developed and applied new SC alloys in recent years. The only use of SC alloys in land-based turbines to date has been the use of CMSX-4 alloy in Solar Industrial Turbine's Mars T-14000 engine in 1990. l~ The application of single crystals to large land-based utility turbines has, however, lagged behind for several reasons.

First, if SC alloys are to be used for industrial gas turbines, more corrosion-resistant alloys must be developed. The alloys reported in the literature so far are low-chromium compositions with adequate corrosion resistance for aircraft turbines, but not for industrial turbines.

A second and more severe problem for industrial turbine application is the scale-up to larger sizes. Problems will be encountered in two areas:

(i) Single-crystal casting technology relies on the existence of a thermal gradient which is typically 7~ (320~ twice as large as in DS casting. H The larger the blade size the more difficult it becomes to establish such a high gradient. If the gradient is too small, casting defects such as freckles and equiaxed grains occur. Material cast under lower gradients will have inferior properties, even if the casting defects mentioned above are avoided. One reason is that the dendrite spacing depends on the gradient. Low gradients result in large dendrite spacings which are difficult to solution- ize, and cause inferior ),' structures. 12 Another reason is that a higher level of porosity is to be expected which will affect particularly fatigue strengh. ~3 Some of the problems encountered with low gradients can be overcome by using lower growth rates. This would lead, however, to lower throughput rates and higher cost.

(ii) The combination of the high superheat temperatures of above 1500~ (2732~ required to obtain suitable gradients, the long periods of contact between metal and mold wall, and the high concentration of reactive elements can lead to considerable metal/mold interaction. At the same time, the mold material is required to have sufficient strength to support the weight of the melt. Thermal expansion coefficients of the mold and the metal have to be matched in order to avoid problems due to thermal stresses during cooling. All this makes the mold technology rather critical. The problems will become more severe in industrial castings, since the weight of the melt is larger and the furnace times are longer.

Nevertheless, it is expected that single-crystal airfoils will become a standard feature in large industrial turbines within the next few years. It is worth noting that although the idea was reported as early as 1968, the benefits were generally appreciated, and both Pratt and Whitney and General Electric


engaged on a considerable in-house research effort, it is only within the last few years that DS components have become widely used; and SC components are still not "bill-of-materials" on many engines. There has been much discussion recently of the accelerating pace of the transfer from research to practice, and in some cases periods like three years have been suggested: it has to be appreciated that in the field of advanced materials the interval between research and practice is still quite long.

Oxidation and Corrosion

In the early days, the superalloys used were all high in chromium, and formed protective Cr203 scales. However, the suitability of chromia scales diminishes as the temperature rises, particularly in high-velocity gas streams, because of the further oxidation to the volatile species CrO3. However, the alloy developments for nickel-base superalloys resulted in an increase in the A1 content, and several of the earlier high-strength alloys were capable of forming A1203 protective scales, which are much more suitable for high- temperature applications. Associated with this was a decrease in the chromium content, and these alloys proved to be sensitive to a form of accelerated oxidation which was associated with the deposition of an alkali metal sulfate on the airfoil surface. This type of attack is called hot corrosion. Because of the presence of sulfur in the deposit, and the presence of sulfides in the corrosion product, the attack was also referred to as sulfidation, but this term should be kept for those reactions where reaction with sulfur is the major process. The first type of hot corrosion reported involved only sodium sulfate, and existed over a range of temperatures between the melting point of the deposit and its dew point. This is now called Type I hot corrosion. Later, another form was observed, initially in marine applications at part load, and it was originally termed low-temperature hot corrosion. This turns out to be associated with the presence of a molten sulfate which involves a transition metal (usually cobalt) as well as sodium. This complex sulfate is not very stable, and its stabilization requires a significant partial pressure of SO3. The range over which the corrosion appears is between the melting point of the salt and its dissociation temperature. This form of attack is now called Type II hot corrosion.

The three processes: oxidation, Type I hot corrosion, and Type II hot corrosion, represent significant problems for the durability of the gas turbine. The relative importance of the processes depends on the application and the siting of the turbine. In many cases the local concentrations of alkali-containing species in the fuel or the intake air are sufficiently low that only oxidation need be considered, and the alumina-forming superalloys may be adequate. For other circumstances, the salt deposition is not sufficient to cause severe


hot corrosion, but must be considered. Alloys have been developed which will perform well under these conditions: the cobalt-base alloys X-40 and FSX-414; and the nickel-base alloy IN 738 are of this class, and have been the materials of choice for industrial gas turbines since 1970 (there are some other alloys in this group). However, there are other conditions where the salt intake is so great, even if only transiently, that even these materials are inadequate.

The design of corrosion-resistant superalloys involves a sacrifice of some strength, and some turbine designers have preferred to treat the strength and the corrosion resistance as separable entities, using a high-strength alloy for the main body of the airfoil, and then applying a corrosion-resistant coating to the surface. Additional protection for the more corrosion-resistant materials is conferred by using a coating, and thus most high-temperature vanes and blades have been coated for over 20 years.

Coatings

The earliest coating was aimed at increasing the surface content of aluminum, since even the alumina-forming superalloys contain barely enough aluminum to maintain a protective scale. Alumina scales tend to spall on cooling, and then hopefully a replacement scale is formed on reheat- ing; the lifetime is thus related to the number of heating and cooling cycles and the magnitude of the aluminum reservoir in the substrate. Processes to form an aluminum-rich layer on a metal surface existed well before the development of the modern gas turbine. The technique used for much of the early period was pack aluminizing (sometimes called pack cementation), in which the component to be aluminized was placed within a bed consisting of an inert powder material (usually alumina) mixed with a relatively small amount of aluminum or an aluminum-containing alloy, and a constituent called the activator, which is a halide such as sodium chloride. The pack is contained within a metal box (retort) and heated for a fairly long time at a fairly elevated temperature, depending on the alloy, the thickness of the aluminized layer required, and so forth. Until quite recently, the majority of components in service were pack aluminized. This process is an example of chemical vapor deposition (CVD), and some recent developments have been made by addressing it as a CVD process, and in some cases physically separating the step of forming the vapor species, which transports the aluminum to the component surface, from the surface itself. This permits much better chemical control.

In 1972, there was a major advance, 14 in which platinum was incorpor- ated in a diffusion aluminide. There are various ways of doing this, but the most usual method is to electroplate a thin layer of platinum onto the


component, and then aluminize in the usual way. Currently, the most frequently used modified aluminides are RT-22 (this is a Chromalloy designa- tion; other companies have similar coatings with different designations) and RT-44, which contains rhodium as well as platinum and is used for cobalt- base superalloys.

Pratt and Whitney initiated a program to develop a completely different approach to coatings, in which the coating was overlayed onto the substrate by evaporating the metal components of the coating from a source using an electron beam, and allowing the vapor to condense on the substrate. This is called electron beam physical vapor deposition, EB-PVD. Working with Airco Temescal, by the late 1960s they had developed a range of MCrA1Y overlay coatings, where M is Fe, Co, Ni, or a combination of these. 15 A typical composition is Co-15 to 20% Cr 12% A1 0.1 to 0.9% Y; the role of the yttrium is to restrict the spallation of the protective oxide on thermal cycling; the chromium offers resistance to hot corrosion.

The effort devoted to the development of coating compositions has been as energetic as that for the base alloys. A wide range of coating compositions is available from vendors: a recent EPRI report briefly lists a number of the more common 16 coating systems. The most recent development has been the introduction of rhenium, akin to its introduction into the base alloys. It is believed to improve the performance, but as yet there appears to have been no explanation of the mechanism of the effect.

As mentioned before, single-crystal airfoils present a new set of problems for coatings, and it is hoped that alloy compositions can be developed which do not require coating, at least for the less severe conditions.

The most recent development in coating technology is the use of coatings with low thermal conductivity (thermal barrier coatings, or TBC) to maintain a relatively low metal temperature at high inlet gas temperatures, and to reduce the amount of cooling air required. The most successful coating of this type is zirconia (ZrO2), where the cubic form of zirconia is partially stabilized by yttria (Y203), or magnesia (MgO). The coating is applied, usually by plasma spraying but also by EB-PVD, over a bond coat which is usually an MCrA1Y. Currently, TBCs are used on combustors, transition pieces, parts of the nozzle guide vanes, and the platforms of the blades: some Pratt and Whitney JT9-D aircraft turbines have been flying since 1989 with TBCs on the rotating blade airfoils for evaluation purposes.

An important problem with TBCs concerns the bond coat, and in particular the oxidation and hot corrosion of the bond layer. The TBC itself is generally not impermeable: ZrO2 supports anion (oxygen) transport, and it is usually arranged for the ceramic itself to contain a network of microcracks as a way of overcoming the problems associated with the differences in the coefficients of thermal expansion of the ceramic and the substrate. This


means that oxygen or other corrodent species may reach the bond coat- TBC interface under conditions which may easily lead to low oxygen activity environments, and which can result in continued oxidation of the bond coat, and even accelerated local corrosion. Much attention is currently focused on understanding how changes at the bond coat-TBC interface influence the mode of degradation of the coating system, and on the possible incorpor- ation of oxygen diffusion barriers in the TBC.

MATERIALS IN LARGE COAL-FIRED BOILERS

Much of the subsequent material on the structure of boilers is a very simplified extraction from "Steam: Its Generation and Use. ''17

Approximately 65% of the electricity generated in the United States is produced by generating stations based on the Rankine cycle, in which steam is generated in a boiler, expanded through a steam turbine which drives a generator, and condensed before being returned to the boiler. This is thus an indirectly-fired closed-cycle system; the compression is (in effect) produced by the boiling. In detail, the water is preheated as it returns from the condenser by steam tapped from the turbine, and enters a horizontal tubular heat exchanger in a pass of the boiler where the combustion gas is passing downwards prior to exiting the system (this is called the economizer). From the economizer, the preheated water enters a (usually) vertical tubular heat exchanger which forms the walls of the combustion chamber, and boils; the relative fraction of steam-to-water rises as the fluid rises through the walls. In drum-type boilers, the tubes are manifolded at the top of the walls into a large drum, which separates the steam from the water. The water returns to the bottom of the wall and recirculates, and the steam passes into a horizontal tubular heat exchanger which is located above the economizer. This is called the primary superheater, and is heated by convection from the flue gas. From this, the steam passes into a vertical tubular heat exchanger hanging in the roof of the combustion chamber: these tubes are arranged in a flat array called a platen, and the heat exchanger consists of a number of platens in parallel which are manifolded into a header. This heat exchanger, the secondary or finishing superheater, receives heat by convection from the flue gas and, in some designs, by radiation from the combustion zone. From the header, a main steam line carries the superheated steam to the turbine. The first part of the turbine is called the high-pressure (HP) turbine, and the exhaust from this returns to the boiler where it is reheated in a sequence of heat exchangers. The finishing reheater is again manifolded into a header, from which a reheat system line carries steam back to the second section of the turbine, which is called the intermediate-pressure (IP) turbine. Generally, the exhaust from this enters directly into the third section


of the turbine, which is called the low-pressure (LP) turbine. However, sometimes there may be a second stage of reheat between the IP and the LP.

On the fireside of the boiler system, coal is pulverized to 70% passing through a 200 U.S. standard sieve (75/lm) or better. This is transported by primary air at 54-93~ (130-200~ to the burners, which are in the lower part of the combustion chamber, from perhaps 10 to 20 m (35 to 65 ft) above the bottom of the unit. Secondary air preheated to about 316~ (600~ is supplied by the forced draft fans and mixed with the pulverized coal in the throat of the burner. A single burner is typically in the range 15-90 MWth, so that a 500-MWe boiler could have from 20 to 100 burners. The objective is to form a fireball in the combustion chamber enclosure separated from the walls. Heat is transferred from this fireball to the waterwalls largely by radiation, with a maximum flux to the walls of the order of 0.3-0.4 MW/m 2 at a level just above the top of the burners, falling to about 0.1-0.2 MW/m 2 at the top of the combustion chamber.

The gas temperature immediately above the fireball is typically of the order of 1538~ (2800~ and that impacting the leading tubes of the pend- ant finishing superheater is some 111 ~ (200~ lower. At this location, the gas stream turns into a horizontal pass, containing a number of vertically oriented heat exchangers. The gas temperature entering this pass is approximately 1093~ (2000~ and the mode of heat exchange is increasingly convective.

All of these temperatures are approximate, and depend on the furnace design, the fuel type, and so forth; and there may be large variations across the face of the superheater, for example.

After the horizontal pass, the flue gas turns into a downward convection pass, which contains the primary superheater and the economizer, both of which are typically horizontal tube banks. From there, the gas enters the air heater, which preheats the secondary combustion air, and then passes to the environmental control systems and finally to the stack.

This brief description is of a recirculating, single reheat, subcritical boiler which represents the majority of the units in the United States. Typ- ically, the maximum steam pressure, that of the steam entering the high- pressure (HP) turbine, is 16.6 MPa (166 bar; 2400 psi), and the temperature is 538~ (1000~ The pressure in the reheater is close to 3.8 MPa (38 bar; 550 psi) and the exit temperature is 538~ (1000~ again. The boiling point of water at 16.6 MPa (2400 psi) is 349~ (660~

An alternative method is to operate the system above 22.1 MPa (221.2 bar, 3208.2 psi) which is the critical point for the water/steam system. At pressures above this, water does not form a two-phase steam/water mixture on boiling: the system forms a single-phase, turning continuously


from a fluid resembling water to a fluid resembling steam: in a glass tube supercritical boiling is indicated by the disappearance of the interface between the liquid and the vapor. In this case, there is no need for a steam drum, and there is no recirculation. Supercritical once-through units in the United States operate at 23.8 MPa (238.5 bar; 3500 psi). The superheating process is the same, and the main steam temperature is again 538~ (1000~ the reheat is also essentially the same, but more frequently there are two stages of reheat. Since a number of references in the subsequent part of this paper are to work performed in the U.K., it should be pointed out that the steam temperatures there are a little higher, typically 565~ (1049~

The efficiency of a large steam plant depends essentially on the difference between the temperature of the steam on entry to the HP turbine and the temperature of the steam-water mixture exciting the low-pressure (LP) turbine. This last temperature is determined by the back-end pressure, and is a function of the condenser. For a modern boiler the back-end pressure is typically 400 Pa (4 x 10 . 3 bar; 58 x 10 .3 psi) or so, corresponding to a saturation temperature of 30~ (86~ assuming of course that adequately cool water is available to cool the condenser. There is little flexibility at this end of the system, so that further improvements in efficiency require an increase in the maximum steam temperature. The steam pressure does not in itself affect the efficiency directly, but higher pressures allow the system to be smaller for a given power, and permit more efficient steam turbines to be built. From 1905 to 1957, the main stream temperatures and pressures, the size of the units, and the overall power plant efficiency, all increased progressively. The maximum in the steam conditions was reached with the double-reheat, supercritical units at Philo [30MPa, 621/565/538~ (4500 psi, 1150/1050/1000~ and at Eddystone [34.5 MPa, 649/565/ 565~ (5000 psi, 1200/1050/1050~ the latter unit attained a coal pile- to-bus bar efficiency of 42~ (although, it should be remembered, without flue gas desulfurization).

However, the industry has retreated from these advanced conditions to the values given above. The reasons are not altogether clear: the greater expense of the plant, because of the larger amount of high-alloy austenitic steel that had to be used, a perception of increased unreliability, through other causes such as heat checking in the main steam valves, attributable to having to make these heavy section components from austenitic steels with relatively poor thermal conductivities; an increased incidence of tube failures; and a decrease in unit flexibility have all been suggested. In part, the increase in tube failures was attributed to fireside corrosion. In the last few years, there has been a renewal of interest in advanced steam conditions in the United States of America, Japan, and Europe. EPRI sponsored two


system studies 18'19 to identify the most desirable configuration, the possible benefits, and those areas which presented problems or would require additional development; and following these a development plan was generated. 2~ Later, a workshop 21 was held to report on the progress of this study 22 and of efforts elsewhere in the world. At more or less the same time, the U.S. Department of Energy assessed the materials requirements for advanced steam cycle systems 23 and conducted research into materials for superheaters and reheaters for advanced steam conditions. 24

Following these initial studies, EPRI selected intermediate conditions for its "state-of-the-art power plant" (SOAPP) 25 which became the subject of much more detailed development, aimed at defining a realistic option for utilities in the relatively near future. An important alloy development for these advanced steam concepts was the improvement in the Fe-9%Cr-1%Mo steels, which led to the T-91 and P-91 steels. These high-strength ferritics allowed heavy-section components to be designed using solely ferritic materials, with their higher thermal conductivity.

Materials Issues in Pulverized Coal Fired Boilers

The typical utility pulverized coal fired boiler and the associated systems is a very large structure indeed, and there are a large number of materials issues. In fact, there is a limited range of materials that can be used for boiler tubes, and all such alloys must be qualified in accordance with the ASME codes. For the highest-temperature components in the boiler, which are the finishing superheater tubes, the finishing reheater tubes, and the pipes which carry the steam to the turbine, the issues are related to creep and to oxidation and corrosion. The relatively cooler evaporator tubes which are in the waterwalls facing the combustion zone nevertheless may also suffer corrosion damage. The still-cooler economizer tubes in the boiler downpass do not experience significant corrosion, but may experience erosion from the fly ash; under some circumstances the superheater and reheater tubes may also experience erosion. Oxidation can be an issue on both the fireside and the steam-side of tubes. In this paper, only superheater and reheater fireside corrosion and waterwall fireside corrosion will be discussed because of limitations on space; but it should not be concluded that these are the only, or even the most important materials issues in pulverized coal-fired boilers.

To help in putting these issues in context, the general issue of boiler tube failures will be briefly mentioned.

Causes of Boiler Tube Failures

Boiler tube failures (BTF) are the leading cause of forced outages and availability losses in fossil-fueled power plants. 26 Recent statistics compiled


by the North American Electric Reliability Council 27 indicate that BTFs are responsible for approximately 70% of the forced outages and 50% of the lost availability of fossil-fueled power plants. Twenty-two distinct BTF mechanisms have been identified as a result of cooperative research between the U.S. utility industry through EPRI in cooperation with the major boiler manufacturers (RP- 1890-1). The major failure mechanisms of boiler tubing are

A. Stress rupture (short-term overheating; high-temperature creep; dis- similar metal welds)

B. Water-side corrosion [caustic corrosion; hydrogen damage; pitting (localized corrosion); stress corrosion cracking]

C. Fireside corrosion (low temperature; waterwall; coal ash; oil ash) D. Erosion (fly ash; falling slag; sootblower; coal particle) E. Fatigue (vibration; thermal; corrosion) F. Lack of quality control (maintenance cleaning damage; chemical

excursion damage; material defects, welding defects)

The two failure mechanisms of most relevance to this symposium are waterwall fireside corrosion, and fireside corrosion of the superheaters and reheaters ("coal ash" and "oil ash" corrosion in the above listing). Of these, fireside corrosion of the superheaters and reheaters has been relatively uncommon in U.S. utility experience in recent times, but it has been studied in very great detail over the years.

Waterwall Fireside Corrosion in Coal-Fired Boilers

Rapid wastage of the waterwalls is sometimes experienced in coal-fired boilers; it is normally associated with a failure to establish the combustion zone properly in the center of the furnace ("flame impingement on the walls") with the result that a region of local low oxygen activity is present at the waterwall surface. The high heat flux means that a slag layer develops with a molten outer surface, although the metal surface temperature is normally quite low. Reid 28 comments that the first real attempts to investigate external corrosion of wall tubes began in the 1940s. Severe metal loss was occurring on wall tubes beneath slag deposits related to the flame pattern in slag-tap furnaces. Reid notes that at the time of writing (1971) there was still no general agreement as to the mechanism of the reaction, and in spite of a considerable amount of work in the intervening 20 years that is still true. Better control of the fuel/air distribution in the furnace, coupled with accumulated experience with burner design and location, and the trend in recent years to dry-bottom furnaces, have eased the problem considerably. Nevertheless, Manny and Natanson 29 reported recently that metal wastage


rates of 30 35 nm/h (10-12 mil/yr) are still occurring, compared to the maximum of 8 nm/h (3 mil/yr) that would be expected based on normal oxidation. The early studies, reported by Reid, Corey and Cross, 3~ made it clear that the metal surface temperature in regions of severe attack did not exceed 375~ (707~ and that there was no difference in temperature between severely attacked tubes and other nearby tubes which appeared to suffer no attack. Corrosion may occur at metal temperatures as low as 315~ (600 ~ F).

The slag layer on the wall contains particles of unburnt carbon and particles of unoxidized pyrite, FeS2. In the earliest work, it had been postula- ted that the local "reducing conditions" might be a factor because volatile iron carbonyls might be formed. This was quickly discounted, and the early American work described by Reid concluded that the low-oxygen activities were unimportant and that the alkali iron trisulfates were the important corrosive species: the model presented in the section on superheater corrosion in coal fired boilers was, indeed, originally developed to explain wall tube corrosion. The low local oxygen partial pressure would appear to mili- tate against the development of high local SO3 partial pressures and thus reduce the probability of formation of the trisulfate, but the stabilities of complex sulfate phases under these conditions have not been studied in detail. In addition, the lowest reported melting points of these phases are of the order of 550~ (1022~ and it seems unlikely that other components could depress them below 500~ (932~ Alkali iron pyrosulfate has been suggested as the responsible molten species, because of its lower melting point, but this is even less likely since the partial pressure of SO3 required to stabilize it is even higher: all in all, it seems unlikely that these complex sulfate species could be stable in the low-oxygen-activity environments.

Because of extensive furnace wall corrosion experienced in some large pulverized coal boilers in the United Kingdom, this problem has been extensively studied by what was then the U.K. Central Electricity Generating Board (CEGB). The earlier work has been summarized by Cutler, Flatley and Hay31; and a series of papers was presented in the conference edited by Meadowcroft and Manning. 32 The initial view developed by the CEGB investigators had been that chlorides play a key role in the reaction. The low oxygen potential conditions were believed to result in the surface release of "volatile chlorine and sulfur gases." The active chemical species responsible for the high corrosion rates were believed to be HC1, H2S/SO2 and CO. Laboratory experiments in gaseous environments suggested that HC1 reacts with the outer oxide grain boundaries of previously formed protective scale, creating microchannels through which (eventually) HC1 gains access to the metal-scale interface. There, a volatile iron chloride is formed which escapes into the bulk furnace gases, where it is oxidized, reforming HC1.


Metal (

~ ~ ~ / ~ / / - - FeS lamellae

~ , f - Fly ash spheres

~,~'~g Unburnt coal ,2/--oa lc'os

Fe 3~ matrix

Fig. 2. The general structure of the corrosion scale formed on furnace wall tubes: intermediate case. The scale thickness is 200 300/ira; the corrosion rate is approximately 300 nm/h, In somewhat milder corrosion conditions there is also intergranular penetration into the substrate. (After Cutler, Flatley and Hay31).

However, it is clear from boiler experience that an important reaction product is iron sulfide, FeS. The CEGB differentiated three different structures corresponding to three reaction regimes: in the first, corresponding to a corrosion loss of 200 nm/h (69 mil/yr), a thick scale was formed: the outer layer was porous Fe203 mixed with ash; the inner layer had an Fe304 matrix with FeS islands embedded in it; the metal exhibited intergranular internal attack. The second involved the formation of a thinner scale: the inner layer had extensive sulfide lamellae, and the outer layer was less thick. There was no intergranular attack of the metal: the rate of metal recession was approximately 300 nm/h (103 mil/yr) (Fig. 2). 31 The most rapid corrosion (600 nm/h ; 207 mil/yr) corresponded to the formation of a thin scale which was almost entirely FeS, with ash particles embedded in it. There was a thin inner layer of Fe304, which may be porous or laminated; there was no internal attack of the metal.

More recent evaluation of the corrosion in practice has resulted in a discounting, to some extent, of the role of chlorine except in the most severe cases, since it is now believed that the low-oxygen-activity environment in the presence of sulfur (and, perhaps, carbon) is sufficient to account for most of the corrosion.

Japanese investigators note that some corroded wall tubes exhibit deep corrosion penetrations in the form of parallel grooves normal to the tube axis with a spacing of the order of 1 ram: sulfur prints show clearly the existence of sulfides associated with these penetrations. French 33 has also reported these features. Their origin is believed to be related to the interaction of thermally induced alternating stresses and the corrosive environment. This phenomenon has been discussed recently by Dooley, 34 who comments that this is an issue primarily with supercritical units: most of the


155 supercritical units in the United States suffer from this failure mechanism. The cracking originates from the fireside surface in regions of high heat flux, 35 and has been referred to by a variety of names: circumferential cracking, horizontal cracking, transverse cracking, craze cracking, elephant hiding and alligator-skin cracking. Unless failure occurs, cracks are very often not detected unless a careful and extensive inspection is performed. Very often the susceptible areas are overheated and suffering from excessive fireside corrosion.

A comprehensive EPRI project has been undertaken for the last three years to monitor these areas for temperature and strain (fireside and coldside), cycle chemistry and flue gas environment to determine the critical and overriding parameters. 36 The root cause of the cracking and overheating is due to a combination of the internal deposits raising the peak tube temperature and the fireside slag shedding, which may be driven by sootblowing or may occur on a random basis, depending on location in the boiler. The internal deposits (Fe304) are often "rippled." Similar overheating and failures have occurred in other countries, 37 where rippled magnetite formation on the waterside surfaces has been directly related to elevated tube temperatures in the high heat flux areas and to boiler pressure drop loss problems. The influence of internal oxide deposits on the fireside metal surface temperature is an important factor, which the early work completely neglected. Obviously, a fireside deposit reduces the metal surface temperature toward the water/steam temperature inside the tube; but adequate heat transfer can only be obtained if this effect is small, or if the wall deposits frequently detach.

The solutions to the waterwall corrosion problem include: finer grinding of the coal to accelerate combustion, improved burner design and maintenance; and bleeding air along the walls to try to keep the local environment oxidizing. For the circumferential cracking problem, oxygenated water treatment can help by greatly decreasing the magnetite deposition on the waterside surface, and methods of reducing the peak heat flux locally, such as the maintenance of slag cover, or the application of refractory coatings, may be useful in some cases.

However, where severe corrosion is encountered, a materials solution appears to be the most attractive. The CEGB has conducted extensive tests with coextruded tubes, 32'38 in particular Type 304 stainless steel outside a carbon steel inner layer, or Type 310 over a carbon steel inner layer. In both cases the outer layer was 1.75 mm (69 rail) thick and the inner layer 6.55 mm (0.258 in) thick. Both materials had, at the time of the report, accumulated 27,000 h in a 500-MWe boiler burning a 0.3% C1 coal. The corrosion rates were between 4 and 10 times lower than those of carbon steel in the same boiler: the maximum rates were 77 nm/h (27 mil/yr) for the Type 304

Current Limitations of High-Temperature Alloys in Practical Appfications 285

combination and 55 nm/h (19 mil/yr) for the Type 310. However, it should be pointed out that with coated systems the important quantity is not the corrosion rate, but the life of the coating: since the coatings were only 1.75 mm thick, these corrosion rates would remove them after 2.5 years in the first case and 3.6 in the second. It was noted that, because of the low temperature, the full protection that one might expect from a developed Cr203 layer was not obtained, and a suggested improvement was to increase the chromium content still further by using an IN 671 outer layer. Diffusion aluminizing has been successful in some cases. A three-layer plasma spray coating, involving an initial exothermic layer to develop a good bond with the substrate, an intermediate layer to provide an aluminum (and perhaps also a chromium) reservoir, and a final layer to seal the porosity, has had some success. The durability of coatings, and their freedom from local defects, is still a matter of doubt, however.

Fireside Corrosion of Superheaters and Reheaters

Fireside corrosion is associated with the formation of ash deposits on the tubes. Crossley 39 reported that work by the U.K. Fuel Research Station in the 1940s had indicated that the external fouling of boilers in the U.K. was connected with the alkali and SO3 contents of the deposits. It was believed that alkalis existed in coals in two forms: chlorides and complex silicates (this is in fact an oversimplification: alkali chlorides are not common in coals, but for this argument the point is not important). The chlorides could be expected to volatilize readily during combustion, but the silicates were expected to retain alkali in the ashes. Thus, it was found convenient to use the amount of chlorine in the coal as the indicator of the volatile alkalis, to distinguish the alkali available to participate in fouling from that in the complex silicates. As a consequence, the determination of coal-chlorine content has been an important item in the ultimate analysis of coals in the United Kingdom since that time. The most common superheater deposits that led to outages were typically alkali-bonded, and so a correlation was sought between the rate of fouling by alkali-bonded deposits and the amount of chlorine in the coals. A survey indicated that there appeared to be a good correlation between the rate of fouling and the chlorine content in the coal for 72 power stations with stoker-fired boilers, whereas for 14 stations with pulverized coal-fired boilers, there was little fouling due to alkali-bonded deposits, and no relationship of fouling with chlorine in the coal. The relative immunity to bonded deposits of pulverized coal (PC)-fired boilers was attributed mainly to the very large amount of fly ash from PC firing in excess of the amount of binding agent required to produce a bonded deposit.

Although PC-fired boilers appeared to be relatively immune to fouling with coals containing up to 0.5 wt% chlorine, the rate of fouling when this


immunity had broken down was appreciably faster than in boilers fired by mechanical stokers. Fouling trials carried out on coals with various levels of chlorine (up to 1.0 wt%) indicated that no boilers fouled when burning coals of very low chlorine content, but all boilers fouled when burning coals of very high chlorine content. Between the extremes of chlorine content, fouling was accelerated or promoted by increasing the amount of sulfur in the coal.

By the time the hot flue gas reaches the radiant superheater it has cooled into the range 1200 to 1000~ (2192 to 1832~ and the entrained ash would be expected to be largely solid. The proposed effect of sulfur was that relatively low amounts of SO3 lead to the formation of normal alkali sulfates as binding agents, whereas increasing the amount of SOs leads to the produc- tion of the corresponding acid sulfates which are considerably more fusible than the normal sulfates. The amount of ash from the coal also was of considerable importance, since it appeared that the ratio of ash to binding agent was a major factor in determining whether bonded fouling deposits could form.

There have been several theories for the buildup of the deposit. The early theories have been summarized by Bishop4~

(i) Alkali metal salts in the vapor phase condense on the tubes to form a sticky layer ("flypaper") which collects impacting particles of fly ash.

(ii) The initial deposit consists of powdery fly ash; the insulating effect of this ash results in the outer layers of the deposit becoming hotter than the inside, and this temperature gradient allows a partial decomposition of sulfates in the hotter parts, with the SO3 so formed diffusing toward the colder metal surface. An inner dense alkali sulfate-rich layer forms.

(iii) Alkal metal or alkaline-earth oxides are deposited on the surface, and are then converted to low-melting point sulfates and pyrosulfates by reaction with the SO3 in the bulk gas phase, and perhaps with other elements in the deposit.

(iv) Alkali chloride vapors from high-chlorine coals, or alkali hydrox- ide vapors from low-chlorine coals condense upon the tubes, and are subsequently converted to sulfates.

(v) Silicon compounds such as SiS, SiSz, or SiO are evolved as vapors from silicate minerals during combustion, are then deposited as fine aerosol particles on the tube surface, and rapidly oxidized to silica.

(vi) Molten and partially molten ash particles in the hot gas stream impact the cold metal surface, and freeze in place.


Flue gas stream ~, Friable ash ~ deposit

~ - 1 ~ [ Iron-rich, hard ~ : . ~ , k , ~ " r e d slag"

~ W h i t e , alkali-

~ Hard black Fe304 scale

Fig." 3. Deposits and corrosion area for a. 41typical reheater tube. (After Nelson and Cam ).

Bishop also identified a number of factors which could contribute to fouling. These include

(a) The presence of relatively minor amounts of molten material, usually alkali-metal-salt-rich, on the surfaces of the fly ash particles themselves, or the tubes

(b) The presence of very fine particles (aerosols) of alkali salts, usually water-soluble, which encourage sintering of the fly ash

(c) The presence of acicular crystals of oxide which project from the oxidized tube surface, trap ash particles, and encourage the buildup of deposits

(d) An irregular tube surface, which is likely to influence collection during the first few hours of deposition

Generally, it appears that a white deposit is formed first, on both the leading tube surface (that facing the gas stream) and the trailing surface. This white deposit is largely alkali sulfate. Later, an ash deposit forms, largely on the leading surface. Not all deposits, by any means, result in enhanced fireside corrosion; but a deposit is essential for accelerated corrosion. The deposit on a corroding tube has the general appearance shown in Fig. 3, which is essentially that of Nelson and Cain. 41 The black inner layer is iron oxide, essentially Fe304 ; the white layer contains significant amounts of alkali sulfates, together with many of the ash constituents but relatively little iron; the "red slag" layer is high in iron, low in alkalis, but with sulfate


FLUE GAS DIRECTION F'LUE GAS DIRECTION

/ ' ~ F R I A E , LE LAYER C / ~

/ l,OOO. ISOT ERM / ~NIN~h,.~--_MOLTEN , E D ~ \ LAYER B 10:~5 "r ISOTHERM,-~ ~ ~ / ~ ~ i l ~ L ~ . MOLTEN WHITE /.~ ~ ~ " ~ L O O S E L Y r ~ ~ " I " - LAYER D

l~l;~ V / ~'/'-/; / ,~//."~'~,~ BONDED L AY ER A J l ; ~ J V ' / / / / / / / / / / / / / ~ IB~l

~~.,i \ . v /

Fig. 4. The effect of steam temperature on deposit structure. (After Borio, Plumley and Sylvester42).

content intermediate between that of the ash and that of the white layer. When dissolved in water, both the white layer and the red slag layer give acid solutions; the solution derived from the outer ash is essentially neutral. Borio, Plumley and Sylvester 42 have described the structure on tubes with two different steam temperatures, as shown in Fig. 4, and concluded that the alkali-rich material was molten in the range of 590-700~ (1094 1292~ Figure 5, taken from "Steam, Its Generation and Use," issued by Babcock and Wilcox, 17 shows the analysis of different regions of the deposit.

Figure 6, from Cutler and Raask, 43 shows the temperature distribution through a superheater tube for a typical set of circumstances. This diagram may be somewhat oversimplified. The overall heat flux to the superheater is indeed of the order of 0.2 MW/m 2, but as is clear from Fig. 3, the thickness of the insulating ash layer on a leading tube is greater at the front than at the back. As a result, the heat flux will vary round the tube, being least where the ash layer is thickest. Wastage takes place at the edges of the deposit--the "five o'clock and seven o'clock positions"--where presumably the heat flux is higher, and the situation may approach that shown in Fig. 4. In the limit, the metal temperature beneath the deposit will be closer to the steam temperature as the heat flux falls, and the temperature of the "alkali sulfate" layer shown in Fig. 4 will also be lower, by as much as 100~ (180~ However, the local SO3 partial pressure can be expected to diminish from the thick part of the deposit to the thin part, and the interaction of the two variables, temperature and SO3 partial pressure in the


Gas Flow 2100F (1149C)

%

Outer Intermediate Inner Layer Layer Layer by wt %by wt % b y w t

SiO2 23.5 23.3 7.6 A1203 14.0 11.5 1.7 Fe203 36.0 11.0 70.5 TiO2 0.9 < 0.1 < 0.1 CaO 1.3 < 0.1 < 0.1 MgO 1.3 1.1 <0.1 Na20 0.3 1.7 0.15 K20 2.9 13.5 1.3 NiO <0.1 <0.1 0.3 Cr203 < 0.1 < 0.1 7.0 SO3 7.3 27.5 10.0 CI 0.02 < 0.01 < 0.01

Water Soluble, % 9.0 45.4 9,0 pH 3.0 2.2 4.3 Excess SO3, % 0.5 11.2 11.8

Fig. 5. Analyses of a typical ash deposit from an 18Cr 8Ni superheater tube. (From Stultz and Kitto17).

vicinity of the alkali sulfate-rich layer, will determine the location of the maximum attack.

As has been pointed out in the discussion of hot-corrosion in gas turbines, the alkali sulfates themselves are not molten in this temperature range. Na2SO4 melts at 884~ (1623~ and K2SO4 at 1069~ (1956~ the two form a continuous range of solid solutions with a minimum melting point of 823~ (1530~ As pointed out earlier in the analogous situation for waterwall corrosion, in the search for low-melting-point phases, early investigators suggested the pyrosulfate (Na,K)2 $207 which melts at temperatures of the order of 400~ (752~ but is unstable: it dissociates to its constituent sulfates unless the partial pressure of SO3 is high. Corey, Cross and Reid 44 as early as 1945 had identified a phase-- the alkali iron trisulfates,


Metal

652(

Alkali sulfate }

Iron- Oxide rich l Ash

ash I

I I I / /

I,( 77

, - / l / I , I

Flue gas (-~ 1000~

977

Temperature (~

1 Heat flux = 0.2 MW/m2

Fig. 6. The general temperature distribution through a superheater tube, the oxide, and the deposit. (After Cutler and Raask43).

(Na,K)3 Fe(SO4)3--which several later investigations suggested were important in the corrosion. Cain and Nelson 45 showed that there was a minimum melting point of 554~ (1029~ at a Na : K ratio of 2 : 3 (Fig. 7). In addition, there are analogous alkali aluminum trisulfates: the pure salts melt at approximately 650~ (1202~ but again there is a minimum melting mixture at approximately 540~ (1004~ these melting points may be further depressed by the presence of other constituents such as LizSO4. The alkali iron trisulfates, like the pyrosulfates, are unstable and will dissociate to the constituent sulfates unless there is a relatively high-local SO3 partial pressure; Reid 28 remarked that they will dissociate at 540~ (1004~ unless the SO3 partial pressure is at least 25 Pa (25 x 10 -5 bar/3.6 x 10 -3 psi); this is, however, considerably less than would be required to stabilize the pyrosulfates.

The problem is then to show how the SO3 partial pressure can be enhanced over the levels normally present in the furnace, which seldom exceed 5 Pa (5 x 10 -5 bar; 7 x 10 -4 psi). Krause, Levy and Reid 46 were able to show experimentally that the SO3 levels immediately above oxidized metal surfaces in the gas were considerably enhanced, presumably due to catalytic oxidation of SO2. The more successful laboratory tests, such as those of Kihara, Isozaki and Ohtomo 47 and of Rehn 48 have found it necessary to use an atmosphere containing both SO2 and 02 passed over a catalyst. Rehn used a gas containing 3.6% O2 and 0.25% SO2 ; the catalyst was Fe203, and the equilibrated gas contained 50-80 Pa (50 80 x 10 -5 bar; 7 - 12 x 10 3 psi) S03.


1200 I I I I l

J 1150 h

I IO0

Q. E

Joso

~ooo I I I I 0 2 4 6 8 - -~

No- - 8 6 4 2 0

Molor Rotio of K 3 Fe(S04) 3 to No 3 Fe(S04) 3

Fig. 7. Melting points as a function of composition in the system Na3Fe(SO4)3-K3Fe(SO4)3. (After Cain and Nelson45).

The temperature dependence of the corrosion is usually described in terms of a "bell-shaped curve." Borio, Plumley and Sylvester 42 show the general form of the corrosion for austenitic stainless steels in laboratory tests in Fig. 8, and also showed that the corrosion range corresponded fairly closely to the range within which a 1:1 N a : K alkali-iron trisulfate was molten or sticky; the maximum corrosion rate corresponded to the maximum temperature at which the salt was wholly molten. Cutler and Raask 43 also related the decline to the decreasing stability of the low-melting-point complex, and included the contribution of the heat flux. For a given metal temperature, the corrosion rate increases with increasing gas temperature. 49 This could be related to the migration of SO3 or the liquid sulfate species in the temperature gradient.

Although the "bell-shaped curve" is generally regarded as describing the temperature dependence of the corrosion, Rehn also commented that the rapid decrease in corrosion rate at temperatures above that of maximum attack, and the actual maximum itself, may well be artefacts of the experi- mental technique used to determine them. In particular, the location of the maximum will depend on the local SO2/SO3 concentration; and that if a molten alkali iron trisulfate is formed at the lower temperature, the temperature can be raised and rapid corrosion will continue, whereas the corrosive liquid will not spontaneously form at the higher temperature.

It has been suggested that the presence of carbon in the deposit may enhance corrosion. This can happen either by the transport of unburnt coal


72.0 ' ' C ~ I:I Na.K COMPLEX

,,60.0 ~ "

g Na COMPLEX "" " \ \

(,,,}

36.0 Z 0 DRY , ~

24.0 . . . . MoL N W \

850 950 1050 11N 1250 TF..M PERATURE, ~

-7:

! 0.19

0

Fig. 8. The physical state of complex alkali sulfates as a function of temperature compared to the corrosion behavior of austenitic alloys. (After Borio, Plumley and Sylvester42).

with the fly ash, or as the result of sooting during the oil-fired startup of a coal-fired boiler. Rehn examined this issue, and found little effect at 675~ (1247~ however, at 790~ (1454~ carbon-free ash did not form the alkali-iron trisulfate, and there was thus little enhanced attack due to the synthetic ash deposit; but the carbon-containing deposit did, apparently, produce enhanced corrosion resembling the liquid salt attack (Table II). Further studies to identify the details of the corrosion process were not undertaken.

Table II. Maximum Penetration in Laboratory Tests Using a Synthetic Ash With and Without Carbon, and a Simulated Combustion Gas, at

785~ From Rehn 48

Alloy

Maximum penetration (pm)

Without carbon 15% carbon added

Type 304 No attack a 457 b Type 347 No attack Edge pitted Type 310 No attack 305 21-6-9 No pits 254 b IN 800H No attack 457 b IN 811 No attack 533 b In 671 No attack - -

"Ash not fused; molten alkali iron trisulfate apparently not formed. bFused ash visible on all specimens showing attack; largely pitting.


The effect of chlorine has been a matter of some controversy for many years. The subject was discussed in some detail in a recent Conference 5~ and further research is currently in progress to determine the effect of chlorine in Illinois Basin coals on fireside corrosion. The experience of the CEGB (mentioned earlier) has been that the corrosion of superheater tubes was a linear function of coal chlorine content. However, chlorine is not detected in the deposit, and the possible chloride species would not appear to be thermodynamically stable at these temperatures. Holmes and Meadowcroft 5~ note that it is now believed that the hydrogen chloride (which is the form in which all the chlorine is present in the combustion gas) acts as a release agent for both the sodium and the potassium in the coal, although they also speculate that perhaps there is a direct effect of chlorine after all. Cutler and Raask 43 concluded that the deposition of alkali sulfates is not a linear function of chlorine content, but approximately 0.3% C1 is required in the coal before rapid deposition occurs, and this has led to a criterion for coal accept- ability if fireside corrosion is to be eliminated; however, it is almost certain that this is an oversimplification. Rehn 48 examined the effect of adding NaC1 to his synthetic ash. At 675~ (1247~ there was perhaps some enhancement of attack, but the effects were small. At 790~ (1454~ however, there was evidence that the ash had fused (as was the case with carbon additions) and there was a considerable pitting attack. However, between the pits there was no attack at all. Mayer, Manolescu and Thorpe 52 have also studied the effect of chlorine in a laboratory test: their approach was to add HC1 to the synthetic combustion gas. Their gas was 80 vol% N2, 16% CO2, 4% 02, 0.2% CO, 0.18% SO2, together with 0, 0.1, 0.2, 0.8 and 2% HC1. The gas was at 1060~ (1940~ and the metal temperatures were 560~ 700~ and 1000~ (1040, 1292, and 1832~ The exposure time was between 5 and 450 h. Sulfur prints indicated the presence of sulfide near the scale-metal interface in specimens oxidized in the presence of HC1, which may suggest some role in the corrosion. Brooks, quoted by Holmes and Meadowcroft ~ proposed the following expression for the rate of fireside corrosion (C, in nm/h) for 580~176176176 and 800~ 1472~ < Tg < 1400~176

C = aL([CI] - 0.1) (Tin - 500) Tg 2 ( 1 )

where Tm is the metal surface temperature and Tg is the gas temperature, both in ~ [C1] is the chlorine content of the coal in wt%; a is a constant (=4.08 • 10 -13) and L = 1.0 for nonleading tubes and 2.2 for the leading tubes, where "leading tube" means a tube in the bundle within about three tube diameters of clear gas in front of it.

Borio, Plumley and Sylvester 42 suggested a procedure for assessing the corrosiveness of coal. Bituminous coals were obtained from six geographic areas in the United States, and a total of eighteen 300-h tests were conducted


in the Solid Fuel Burning Test Furnace of Combustion Engineering. The gas temperature was 1093~ (2000~ and Type 321 stainless steel was used as the standard test material; the specimen temperature was close to 593~ (1100~ It appeared that the most significant factors were sulfur, iron, the alkaline earths (Ca and Mg), and the alkali metals (Na and K). The correlation was much better with the acid-soluble alkali than with the total alkali. The pyritic sulfur correlated well with the iron content. Iron is present in the coal largely as pyrite, FeS2 ; and in most U.S. bituminous coals the ratio of pyritic and organic sulfur is fairly constant. A nomogram was constructed from these data, and from this a corrosion index was calculated; there appeared to be a reasonably linear correlation between this index and the fireside corrosion experienced in the test furnace. In addition, they inserted probes into the combustion gas stream and determined the acid soluble Na20/K20 ratio. They noted that this value was different from that calculated on the basis of the coal analysis, but that it was probably a more accurate measure of the corrosivity, because it would account for the different release of the two alkalis during combustion. In their work, the effect of chlorine was not examined, but clearly it would have an effect on the alkali release rate.

Materials Behavior in Superheater Corrosion

There have been several tests of the relative corrosion resistance of different alloys. Plumley, Accortt and Roczniak 53 have presented a brief report of an extensive study undertaken with U.S. Department of Energy support to assess boiler tube alloys for advanced power cycles, using a similar approach to that described above, coupled with longer-term field studies. In these tests, the Appalachian coal was the least aggressive, and the Illinois coal the most aggressive of those studied; its ash fusion temperature was also the lowest, interestingly.

Both Kihara, Isozaki and Ohtomo, 47 and Rehn q8 have used a laboratory system to study the relative corrosion resistance of materials. Figure 9 summarizes the results of Kihara et al. Sumitomo Metal Industries have also studied the fireside corrosion of candidate boiler materials using a similar laboratory test, and the data have been presented by Fujikawa and Makiura54: the results generally were in agreement with Kihara et al. Alloy 17-14 CuMo, which is a high-strength steel which was used at the Eddystone high-temperature boiler and has been considered as a candidate for advanced steam cycle units, was considerably worse than the others. Of the more conventional materials Type 316 stainless steel was significantly worse than Type 304" the stabilized steels were better, with Type 347 somewhat better than Type 321. These results suggest that Mo, and to a lesser extent Ti, had


500

400 -- / 17-14CuMo / \

Esshete 1 300 - - g

347 ~ \ ~_ 200 -- ~~~/.///', ~k .m 800H Q)

�9 (2-"

600 650 700 750 Temperature (~

Fig. 9. Corrosion rates of some alloys in a laboratory test using a synthetic ash (37.5 mole% NR2S04, 37.5 mole% K2804, 25 mole% Fe203) in a synthetic flue gas (80% N2, 15% CO2, 4% 02, 1% S Q , saturated with H20). Exposure time 50 h. (After Kihara, Isozaki and Ohtomo47).

adverse effects on the corrosion resistance of a Fe-18Cr-8Ni steel, whereas Nb appeared to be beneficial. Latham, Flatley and Morris 55 have reported the relative corrosion rates in coal-fired boilers of Types 316, 321 and 347 stainless steels, Esshete 1250, Incoloy 800; and Type 310 stainless steel and IN 671 as claddings. Samples of Esshete 1250, Type 316 and 347, and IN 671/IN 800 coextruded tubes were installed into leading tube positions of the hottest reheater of a 550-MWe boiler; the metal temperature was 650~ (1202~ the gas temperature was approximately 1150~ (2102~ and the coal fired contained an average of 0.45% chlorine. After 15,400 h, specimens were removed for examination. In addition, corrosion probe studies were conducted for 2000 h in the same boiler, and also in a boiler burning 0.3% C1 coal at a location where the gas temperature was 950~ (1742~ Because of its high scaling rate, the Esshete tube was bandaged after only 6000 h to avoid failure. After completion of the test, wastage flats were clearly visible on the Types 316 and 347 stainless steels; the IN 671- clad tube appeared to be unattacked. Interestingly, the corrosion on the Type 316 tube was 2.5 mm in 15,400 h, compared to 3.2 mm on the Type 347 tube, the reverse of the behavior reported by other investigators. Type


310-clad material exhibited a loss of 1.1 mm after 18,000 h in the same boiler. The probe tests did not show a strong dependence on temperature in the range of 600~176 (1112-1292~ with little evidence of the "bell- shaped curve" reported in the laboratory tests (it will be remembered that Rehn also suggested that this curve might perhaps be an artefact of the laboratory testing procedure, and not be duplicated in practice). However, Latham et al. suggested that perhaps the problem was that the temperatures fluctuated over the probe, so that all sections saw a significant range of temperatures during the exposure. Under the more aggressive conditions, Types 316 and 347 were severely corroded, with maximum rates of more than 100 nm/h on Type 316.

Under the less severe condition, Type 316 and Esshete 1250 were the most heavily corroded materials. With the 0.3% C1 coal, there was intergranular penetration of Type 316, Type 347, and Esshete 1250, but none on Type 310; but with the 0.45% C1 coal there was extensive intergranular attack; the grain boundary penetrations were sulfides. Both the Type 347 and the Type 310 formed some deep oxide-filled pits with intergranular penetration in the metal beneath them in the higher chlorine coal environment. Overall, the authors concluded that the corrosion resistance of the alloys was in the order IN671 (best) > Type 310>Type 347~Type 316 ~Eshete 1250. Note that it was difficult to distinguish between the last three in spite of the lower chromium content of the Esshete 1250. If Type 310 is to be used as a cladding, it was suggested that the chromium content be specified to be at least 25%, and that modifying it with approximately 1% Nb decreases the intergranular penetration; a minimum of 1% Si is also beneficial.

"Bandages" formed by wrapping a strip of a corrosion-resistant material around a corrosion-prone section of a tube have been extensively used by the CEGB. The corrosion-resistant material is usually Sicromal. Mortimer and Latham 56 have discussed the performance of bandage shields on Type 347 reheater tubes at the Thorpe Marsh coal-fired power station. The lifetime of these tubes was less than four years, and several causes were suggested: rather high operating temperatures (680~176 1256- 1292~ the occasional occurrence of a reducing atmosphere; and the use of substantial periods of oil firing. In addition, starting before their report (1972), a high-chlorine coal (0.3 C1) was introduced. Four bandage materials were tried: Sicromal; a Fecralloy; an austenitic steel (Fe-22% Cr-16% Ni); and a nickel-base alloy (Ni-16% Cr). The two latter alloys are not described any more completely in this report. The Sicromal behaved best; the general condition of the Fecralloy was quite good, but there was some embrittlement and local internal oxidation and sulfidation. The remaining two alloys exhibited more extensive attack. Bandaging depends on relatively poor thermal contact between the bandage and the tube, so the bandage operates at


a higher temperature than the tube, catching the ash preferentially. As a consequence, the heat transfer to the heat exchanger is reduced. This may not be a problem if only small regions of the leading tubes have to be bandaged, but for more severe corrosion extending over a greater area of the exchanger, the reduction in heat transfer would not be acceptable. Nowa- days, this approach is being abandoned in favor of co-extruded tubes.

Coating has been less successful: Rehn 48 has reviewed the available information and reported on a study of a number of possible coating systems using the test procedure described earlier. Diffusion coatings of aluminum, chromium or silicon on various substrates were examined; the aluminized surfaces appeared satisfactory, whereas the chromized surfaces suffered severe local pitting; in one case, the attack was reported as being much deeper than the uncoated substrate. In the case of siliconized materials, there appeared to be no pitting, but the protection was not very considerable. A variety of plasma-sprayed and metallized coatings was studied; coating porosity was a problem. Plasma sprayed IN 671 was generally satisfactory, although the bulk alloy showed deep local pitting. A coating consisting of plasma sprayed MgZrO3 over an initial bond coat of Ni-AI on stainless steel (Type 304) appeared excellent in the short-term tests. Longer-term testing 57 confirmed these good results, and indicated that the coating should be less than approximately 300 pm to obtain good thermal cycling resistance.

In summary, it seems likely that the variations in corrosion resistance among the conventional stainless steels are not sufficiently great to be a reason for selecting one over another; local fluctuations in conditions may produce larger changes than the differences observed in laboratory tests. In severe conditions, coextruded tubes may be the most convenient solution, in spite of the additional cost. It is necessary to optimize Type 310 if that is to be used as a cladding, and longer-term exposures under a variety of conditions are needed. Although some examples of rapid attack of IN 671 have been reported, it appears to be the most resistant material of those currently available as a cladding.

MATERIALS IN COAL GASIFICATION SYSTEMS

General Features of a Coal Gasification System

The general principle of a coal gasifier for utility application is to produce a gas from coal while retaining as much of the chemical energy in the product as possible. Gasifiers may also produce gas as a chemical feedstock, and in some cases the desired product is methane, as a substitute natural gas (SNG); however, in this last case there is a significant energy penalty. Possible routes include partial oxidation, so that the product is largely carbon monoxide; or the water-gas reaction, in which case the product is a


mixture of carbon monoxide and hydrogen:

C 4- H20 = H2 + CO (2)

However, this reaction is endothermic, so that part of the coal must be oxidized to carbon dioxide to provide the heat, and the product gas is a mixture of hydrogen, carbon monoxide, carbon dioxide, and water; if the gasifier is air-blown as opposed to oxygen blown, there will also be a considerable amount of nitrogen. The relative proportions of the different gases vary from process to process.

There are three basic ways in which the gasification reaction can be accomplished: in a fixed bed, a fluidized bed, and an entrained gasifier:

Thefixed bed is best known as the Lurgi gasifier, which has been used for many years. Coal is loaded into the top of the gasifier vessel, and air (or, less commonly, oxygen) together with steam is blown in from the bottom. A combustion zone is established at the bottom of the bed, and the hot gas rises upward and drives the water-gas reaction (Eq. (2)). The still-hot product gas rises through the coal bed, preheating it, first driving off the volatiles, and at the top of the bed driving off the moisture. The incombustible coal material (the ash) is withdrawn from the bottom of the gasifier. The product gas is relatively cool: the exact temperature depends on the moisture and volatile content of the feedstock. The tars and volatiles must be condensed out, and the gas is usually water-quenched, thus also removing the ammonia which is the product of the fuel-bound nitrogen. The sulfur in the coal is present as H2S, and this can be removed by conventional techniques such as the Selexol process. The problems are the losses associated with the tars and volatiles, and the tendency for the process to blow off the coal fines at the top of the bed without their being gasified. A high-fines feedstock is thus undesirable.

British Gas has developed the Lurgi process to operate at higher pressures and temperatures, so the ash actually melts and is tapped out of the bottom. The effect is to increase the throughput. Techniques have been developed to blow the recovered tars and the fines through the oxidant tuyeres, increasing the overall efficiency.

Fluidized bed gasifiers use a bed of particles which is fluidized by air or oxygen blown upwards through it. The coal is injected into the bed, which operates so that the oxidant flow is less than the stoichiometric quantity: partial combustion heats the bed, and the remainder of the coal is gasified. The temperature of the bed is typicaly 850~ (1562~ and the release of tars and volatiles is significantly reduced or eliminated. Fines are a problem with this technology also, but proponents claim to have developed techniques to overcome it. The product gas is at the bed temperature, so the sensible heat is significant, but not really enough to warrant a heat


exchanger. The gas is normally quenched, with some recovery system for the heat in the quench water.

The entrained gasifier resembles a burner: the coal, typically in a slurry (although dry feed systems are available) is injected through the center of the burner, and mixes with the oxidant which is injected in the outer con- centric part of the burner. The temperature in the combustion zone is very high, and all the volatiles and tars are completely gasifed. There is also no problem with fines: in fact, the coal is typically pulverized before gasification. The incombustible material forms a slag, which runs down the walls of the gasifier vessel and into a quench pond at the bottom, from which it is lock- hoppered out. The product gas temperature is very high, and it is essential to recover the sensible heat, which is done with a radiant heat exchanger in a vessel directly above the burner, followed by a convection heat exchanger in an adjacent vessel.

The two principal designs of entrained gasifier are those of Texaco and of Shell. The Texaco gasifier was the basis of the Cool Water integrated gasifier combined cycle (IGCC) on the Southern California Edison system which was partially funded by EPR158; the Shell gasifier is the basis of the IGCC at Buggenum in The Netherlands. 59 Both these systems are oxygen- blown. A major materials issue is the radiant heat exchanger (syngas cooler), and the problem reduces to one of determining the maximum permissible metal temperature. Much of the subsequent section is drawn from a recent paper by Bakker. 6~

Materials for Radiant Heat Exchangers in IGCC Systems

Metal Temperatures

Most of the heat exchange surface in syngas coolers is used to evaporate water. The metal temperature of the tubes is largely determined by the water temperature, which depends on the desired saturated steam pressure. Presently, the usual pressure is 10 MPa (1500 psi) which will produce 320~ (608~ saturated steam and corresponding metal temperatures of 340- 400~ (644-752~ depending on the syngas temperature. It is expected that the steam pressure will be raised in the future to that presently used in subcritical boilers, i.e., 18 MPa (2586 psi), producing 350~ (662~ saturated steam and metal temperatures in the 380-450~ (716-842~ range. This is the primary temperature range of interest, which most metal alloys will experience. 6~ Some coal gasification systems will require superheating of at least part of the steam in the syngas coolers to 500-550~ (932- 1022~ with corresponding metal temperatures in the 550-600~ (1022- 1112~ range. This is the highest temperature that alloys will experience in present generation gasifiers, at least in significant volume; there may be


small components, such as instrumentation wells, where the temperature may be higher, but these situations will not be discussed here.

Gas Composition

The composition of raw syngas is rarely measured before quenching the gas after its passage through the syngas cooler. Gas compositions given by developers are often on a water-free basis, after desulfurization. An extensive compilation of gas compositions was given earlier by Natesan. 61 As an approximation, the composition of the gas from an oxygen-blown entrained slurry-fed gasifier like the Texaco is (in vol%) 35-45% CO, 10-15% CO2, 27-30% H2, 15-25% H20, and 0.2-1.2% H2S; for a dry-fed gasifier such as the Shell the corresponding compositions are 62 64% CO, 2-4% CO2, 27- 30% H2, 0-3% H20, 0.2-1.2% HzS. These compositions are for the gas at equilibrium at the gasifier operating temperature. Bakker 6~ states that the cooling is sufficiently rapid that these compositions are quenched in at the metal temperature of the heat exchanger. This presents a problem: for ana- lyzing a corrosion reaction in a mixed gas of this kind it is normal to describe the gas in terms of the chemical activities of the oxidizing species, Po~, psi, and ac. For a gas at equilibrium, these are easily calculated from the ratios of the partial pressures of the molecular species, Pco2/pco, etc. However, it is arguable whether it makes any sense to talk about thermodynamic activities for species in a non-equilibrium mixture, although Bakker refers to a method proposed by Perkins. 62

It is, however, clear that the oxygen activity will be low, and much lower in the dry-feed gas. The sulfur activity is also low, but as it turns out, higher than the oxygen activity: it appears to be very much the same in the two gases. In gasifier systems the carbon activity is always high, close to unity.

Corrosion Mode

The corrosion in a gasifier environment is mixed oxidant corrosion. In this mode, the ability of the system to form a stable protective scale is disturbed by the second ( third, . . . ) oxidant. Generally, the only stable protective scale is an oxide, and the protective nature of the oxide scale can be disturbed by the sulfur and/or carbon in the gas. This is the case, although in coal gasifier environments the bulk of the reaction product is a sulfide; the protective oxide is formed beneath this outer sulfide layer. There are various ways that the continuity of the protective oxide scale can be disturbed, and it is beyond the scope of this paper to describe them in any detail.


The Behavior of Materials in the Gasifier Radiant Cooler

As pointed out earlier, the materials that can be used for boiler tubes are limited by the ASME codes. For an evaporator heat exchanger, only ferritic steels are allowable, at least for the surface in contact with water. Generally, for the pressure and temperature conditions characteristic of a gasifier radiant cooler, a plain carbon steel would be sufficient. However, for such tubes, the only stable corrosion product is iron sulfide; the corrosion rate is linear, and the rate would be unacceptable.

For high-chromium austenitic steels, the stable specie would be chromium oxide, which is capable of forming a protective scale with an acceptably low rate of growth. The oxygen and sulfur partial pressures of most syngas compositions fall in the Cr203 and Fe(Ni)S phase stability fields when plot- ted on a S-O-M phase stability diagram (where M is Fe, Ni, or Cr). Within this field there is generally a "kinetic boundary" as defined by Perkins 63 and Natesan 61 to the right of which the growth rate of Cr203 is fast enough to form a protective, sulfur-free, Cr203-rich scale. Unfortunately, almost all gas compositions of commercial processes fall to the left of the kinetic boundary, and form less protective mixed oxide/sulfide scales, although Cr203 is still the thermodynamically stable species. Intuitively one would expect that the corrosion rate would increase with decreasing Po2/ps2 ratio, although it is not clear how the O/S ratio in the corrosion product will affect the corrosion rate. Isothermal corrosion tests of 600 h duration were, therefore, run at 550~ (1022~ in a syngas to which increasing amounts of water were added. 6~ Water additions up to 5% appeared to increase the corrosion loss for the 20% Cr alloy, but not that of the 25% Cr alloy. All alloys showed an abrupt decrease in corrosion rate when the water content was increased from 10-15%, despite the fact that the change in oxygen pressure was small. Bakker remarks that an elegant theoretical explanation for the observed data eludes us so far.

The corrosion rates observed in syngas with higher water contents are quite low for most alloys and lead to reasonably low corrosion losses, generally smaller than 23 nm/h (8 mil/yr) even when extrapolated linearly. The corrosion losses in syngas with a low water content are unacceptably high when extrapolated linearly. Consultation with various process developers indicated that linear extrapolation is probably too conservative, and that parabolic extrapolation is more appropriate. To check this, a few isothermal tests of different lengths were performed in a water-free syngas. The results indicated that this is a fair assumption for the 25% Cr alloy (Type 310), but may not be valid for the alloys containing only 20% Cr. 6~

Aluminum and silicon are the traditional alloying additions to improve oxidation and sulfidation resistance, presumably because they are able to


form protective silica- and alumina-rich subscales. Laboratory data at 550~ (1022~ confirmed the beneficial effect of A1 and Si on corrosion resistance in syngas, 58 although the presence of a protective oxide layer could not be confirmed, and iron and nickel sulfides were present in the outer scale. The performance of A1- and Si-containing alloys in actual service is less straightforward. Diffused aluminum coatings containing 10-40% A1 on low- alloy steels have corroded rapidly in all plants gasifying coal with a chlorine content above 0.1-0.15%. It is believed that the HC1 in the gas removes t h e A1 in the coating as A1CI3, which decomposes to A1203 in the scale, thus providing additional HC1 for further attack. 64 A recent, very preliminary corrosion test in a pressurized, dry syngas, using MA 956, suggested that same mode of material loss may happen here in the form of pitting.

Experience with Si-containing alloys is more limited. Only one alloy containing 3.7% Si was extensively tested in a pilot and a demonstration plant producing syngas containing 5 25% H20. 65'66 In the more severe environment of the pilot plant, where high-chlorine fuels were gasified, extensive intergranular corrosion and cracking was noted. Under the more moder- ate operating conditions of the demonstration plant, intergranular corrosion occurred. Because of these experiences, the investigation of Si-containing alloys was not continued in gasifiers producing syngas with a high water content.

The Effect of Other Oxidants: Chlorine and Carbon

The possible effect of chlorine on coal-fired boiler waterwall corrosion, where the local environment may not differ greatly from that in a gasifier, has been mentioned above. Corrosion tests in pressurized syngas have showed that HC1 may increase the corrosion rate somewhat, although the scatter is great. So far, there is no clear-cut evidence that HC1 has an effect on the corrosion rate of stainless steels, except for those containing aluminum. 6~

The main detrimental effect of HC1 in the syngas is its effect on aqueous corrosion during downtime. 67 Deposits on heat exchanger surfaces in plants gasifying coals with as little as 0.1% chlorine quite often contain soluble chlorides, such as FeC12, NaC1 and NH4C1, up to several percent. When such deposits are allowed to get wet during downtime, the condensate can become quite acidic, especially when some of the iron sulfide oxidizes to iron sulfate. Polythionic acids can also form. There is ample evidence from laboratory studies, as well as plant exposures, that rapid pitting and even stress corrosion cracking of common stainless steels can occur under such conditions. The best defense against downtime corrosion is the prevention of condensates by keeping the heat exchangers above the dewpoint of the air, or lowering the dewpoint by blanketing the syngas coolers with dry


nitrogen or air. If condensation cannot be prevented, alloys containing Mo and V will usually resist pitting.

Carbon monoxide is the major constituent in syngas. Protective, pure Cr203-rich scales usually do not form in syngas, so that the potential of carburization is certainly present. Evidence for carburization in plant exposure was summarized in a recent paper. 68 It was shown that carburization readily occurs at temperatures above 600~ (1112~ in alloys such as Type 310 and Incoloy 800. Bakker 6~ commented that carburization followed by chromium depletion and accelerated sulfidation/oxidation may be the failure mode which will limit service life at superheater temperatures. Longer-term exposure tests are needed to confirm this.

SUMMARY AND CONCLUDING REMARKS

The examples of components where high-temperature oxidation and corrosion resistance are significant issues for the designer and operator that have been instanced above are intended to show that research and development in this general area is still needed, emphasizing the relevance of this symposium. However, it is also intended to show that the areas of practical importance are far from simple, and the research to optimize the corrosion resistance must be consistent with the other requirements of the system: the need for strength and/or toughness in the materials; the need to ensure that the resistance to environmental degradation must be maintained throughout the working range, and not simply at the maximum temperature; the need to ensure that the solutions are consistent both with the manufacturability of the component and with the integrity of the rest of the system; and that cost is always an important consideration.

There are traps in conducting research hoping to address situations of practical concern. Often, the lifetime required of the engineering system in practice is very long, perhaps as long as 30 years (262,800 h). Of course, it may be sufficient for individual components to have shorter lives: major overhauls which would allow components to be replaced may take place at three-year intervals (26,280 h) ; but in this case a very high degree of confidence in the ability of the component to last this time will be required. During this period, the equipment may cycle, and may even have extended periods when it is idle. It may be difficult or inconvenient to exclude the ambient atmosphere during these times, and the possibility of the initiation of damage by "downtime corrosion" at low temperatures must be considered. The fuel chemistry may change for all sorts of reasons, and the corrosion resistance must either be tolerant of such changes, or the sensitivities must be carefully defined. Similarly, the operating conditions may change:


a plant which was designed for steady-state operation may change to inter- mittent operation later in its life, and again the corrosion solution must either be tolerant, or the degree of intolerance needs to be defined.

A solution to the corrosion problem which is acceptable under laboratory conditions may not be acceptable at industrial scale: for example, fluctuations in composition within the normal commercial specification limits must be tolerable; the solution must be tolerant of the sort of surface finish encountered in practice, the minor scratches that are part of real-world installation, and so forth. Often, structures in practice are very large, and there may be a significant amount of elastic strain in operation. For high- temperature components, there will also be a degree of time-dependent plas- tic strain due to creep: the solution must also be tolerant of these.

The extrapolation of short-term data is risky. Most laboratory investigators will fit (often force-fit!) their results to parabolic rate curves, and use the parabolic rate constants so determined to extrapolate to practical life- times. First, real parabolic growth is unusual; and although the error may not be important for short extrapolations, it can produce quite important errors over long times. Second, in many cases the real problem is "break- away," when the corrosion rate accelerates abruptly. To some extent, testing may be accelerated by increasing the temperature, but this too carries considerable risks. For example, below approximately 570~ (1058~ the oxide grown on iron in oxygen or air at atmospheric pressure is principally magnetite (Fe304), with a thinner outer layer of hematite (Fe203). Above this temperature, the oxide is principally wustite, with thin outer layers of magnetite and hematite. Small amounts of chromium in the steel increase this transition temperature somewhat. Reducing the partial pressure of oxygen in the atmosphere can alter these results: in low-oxygen steam, for example, hematite may be absent. Obviously, attempting to estimate the oxidation behavior of low-alloy steels in boilers, which typically will experience temperatures below approximately 540~ (1004~ by increasing the test temperature is unlikely to be successful. In the case of Type I hot corrosion, the initiation of the reaction depends on a number of factors which are difficult to model with confidence, but certainly some condensed alkali sulfate on the corroding surface is essential: if the temperature is raised the sulfate may evaporate. In the case of Type II hot corrosion or molten salt- accelerated corrosion of superheaters the problem is even more specific, because the complex alkali-containing sulfates will dissociate if the temperature is increased significantly.

Nevertheless, it is obvious that 50,000 h testing is out of the question on a regular basis, and certainly as a method of qualifying new materials. Probably, 5000 h testing is required, with perhaps modest acceleration, coupled with a more careful empirical determination of the form of the kinetics


for extrapolation purposes. In addition, every opportunity should be taken to study and characterize specimens from operating units with long histories.

A further problem is that, in practice, many of the important applications involve the materials being subjected to a thermal gradient. For example, in a boiler the heat-exchanger tubes have steam at a maximum temperature of (say) 600~ (1112~ on the inside, and combustion gas at a temperature of the order of 1300~ (2372~ on the outside. The metal surface temperatures on both the inside and the outside will differ from these fluid temperatures, and will vary with the oxide or deposit thicknesses. Some investigators believe that the thermal flux itself may have an effect on the corrosion.

For corrosion in mixed gases there is a further complication. The resi- dence time of the hot gases in the vicinity of the heat exchanger or indeed in the system itself may be quite short in comparison to the times required for the different gas phase reactions to approach equilibrium. The metal surface itself may catalyze some or all of these reactions, but nevertheless the reaction will be taking place in a nonequilibrium environment. Few investigators have considered this; and indeed the analysis is difficult and quite different to the methods of analysis customary in the community: it is also difficult to model in the laboratory.

In laboratory studies of oxidation, it is often assumed that the aim is to reduce the oxidation rate by considerable amounts--two or three orders of magnitude, for example. This is seldom the case in practice. If anything, the rate of growth of alumina scales is too slow! If the rate of growth were a little faster, it would be possible to extend the beneficial effects of alumina to lower temperatures. For practical situations, an improvement in the lifetime by a factor of 3 is often not only of great value, but may be all that can be used. There is no point in extending the life of a component beyond the life of the rest of the system.

From these brief remarks, it can be concluded that, first, there is much useful work to be done in the field of high-temperature oxidation and corrosion; and, second, the most useful research requires the abandoning of many traditional assumptions. This is a major intellectual challenge for the next few years.

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Current limitations of high-temperature alloys in practical applications