A Chapter in the Refractories Handbook, edited by 

Charles A. Schacht, and to be published by Marcel Dekker, Inc., 

New York, NY 10016

Corrosion of Refractories

by 

Denis A. Brosnan, PhD, PE

Clemson University

Clemson, SC 29634-0971

Introduction

Refractories are used at elevated temperatures for structural purposes 

and they are used in many cases to contain a high temperature corrosive 

environment. This corrosive environment usually contains liquid (melted) phases 

that participate in chemical reactions with the refractory at the elevated 

temperatures resulting in refractory consumption or wear. It is usually not 

immediately obvious, but the oxidation and reduction state of the environment (as 

“redox” conditions or oxygen “activity”) can participate in and influence the 

chemical reactions that take place. Along with chemical reactions during 

corrosion, physical changes occur that may be accelerated by the corrosion 

process.

Corrosion of refractories can be defined for the purposes of this discussion 

as follows:

Corrosion of Refractories – refractory wear by loss of thickness and mass from 

the exposed face of the refractory as a consequence of chemical attack by a 

corroding fluid in a process in which the refractory and the corroding fluid react 

approaching chemical equilibrium in the zone of contact between the refractory 

and the fluid.

It is an essential point that corrosion reactions proceed in a direction 

toward localized chemical equilibrium. This means that phase equilibrium 

diagrams can be used to analyze corrosion situations and to predict chemical 

strategies to minimize corrosion and wear rates. This gives persons interested in 

refractory corrosion two options. The first is to view corrosion as a chemical and 

physical process without a detailed application of phase equilibrium diagrams –

called the “phenomenological approach”. The second is to use the information in 

the phenomenological approach and to use phase equilibrium diagrams. This 

latter option is required for a full understanding of refractory corrosion.

There are many types of refractory systems – fusion cast and/or bonded 

brick in a “working” lining placed in front of a backup or “safety” lining (or simply 

placed against a “shell”), thick-wall applications of concretes (and other 

monolithics) where the lining thickness is usually on the order of 75 mm or 
greater, and thin-wall applications of monolithics where the lining thickness is 

typically less than about 25 mm. It is beyond the scope of this chapter to cover 

every type and application possible. There are fortunately a few fundamental 

principles that the investigator can apply to any corrosion situation, and it is a 

purpose of this chapter to provide those fundamental principles.

First Fundamental Principal on Refractory and Slag Compatibility

The first fundamental principles is that “acid” refractories tend to resist 

“acid” slogs better than “basic” slogs and, conversely, “basic” refractories tend to 

resist “basic” slogs better than “acid” slogs. The definitions of acidity and basicity

in room temperature solution chemistry and of refractory chemistry at elevated 

temperature have a key difference:

Definition of Acidity and Basicity in Solution Chemistry at Room Temperature –

an acid contains an excess of hydrogen ions (H+) over hydroxyl ions (OH-) 

considering a “baseline” defined as neutrality (a “pH” of 7.0). An acidic substance 

contributes hydrogen ions to a chemical solution to make it more acidic whereas 

a basic substance contributes hydroxyl ions to make it more basic.

Definition of Acidity and Basicity in Corrosion Chemistry at Elevated Temperature

– an acidic material contains an excess of silica content (SiO2) over basic 

materials (usually CaO) considering a “baseline” defined as neutrality (a 

CaO/SiO2 ratio of 1.0). An acidic material contributes SiO2 in a corrosion 

reaction whereas a basic material contributes CaO or MgO in a corrosion 

reaction. 

This leads to the observation that acid refractories are more “compatible” 

with acid slogs, i.e. acid materials experience less corrosion loss against acid 

slogs as compared to basic slogs. In a like manner, basic refractories are more 

compatible with basic slogs than with acid slogs. 

In a strict definition of compatibility, mineral phases will not react at 

elevated temperature if they are compatible. This means on a microscopic basis, 

they will “stand beside” one another or coexist at equilibrium without reacting to 

form new substances. In general usage, the term “more compatible” just means 

“less reactive”.

There are more complex definitions of acidity and basicity for high 

temperature chemistry than those given above. For example, in steel refining the 

term “V” ratio is used where a ratio exceeding 1.0 implies a basic chemistry 

(usually a basic slag chemistry) and a ratio below 1.0 implies an acid chemistry. 

In definitions of this sort, the “V-ratio” can be expressed in terms such as the 

following:
SiO P O Al O Fe O Mn O ...

CaO MgO FeO MnO ...

V 

2 2 5 2 3 2 3 2 3

In many cases, it is convenient to use three-component phase equilibrium 

diagrams in analyzing corrosion situations. The major refractory component can 

be visualized as one apex (corner) of the triangle with CaO and SiO2 as the other 

components (apices). In these analyses, it is convenient to use the simple 

CaO/SiO2 ratio. In more complex analyses using advanced thermodynamic 

software, it may be more accurate to use the V-ratio to define “equilibrium” 

conditions.

Corrosion reactions should be viewed as attempts by the system to 

achieve compatibility by progressing toward equilibrium. Refractories are rarely 

at chemical equilibrium on a microscopic scale since they are typically made from 

mixtures of different minerals. However, at the immediate corrosion interface 

between the refractory and the slag, the localized volume elements may be at or 

close to chemical equilibrium.

Second Fundamental Principle on Porosity and Corrosion Rates

Most refractories contain void space or porosity. This porosity may be 

open pores that can be penetrated by a fluid media (i.e. “apparent” porosity) 

and/or it may be closed porosity that is not easily penetrated by fluid media. If a 

refractory contained no porosity (or brick joints, expansion joints, or construction 

joints), the corrosion reaction is limited only to the face exposed to the corrosive 

media (called the hot face). When porosity is present, particularly when open 

porosity is present, the corrosive media can penetrate the refractory causing 

destructive reactions behind the hot face.

Most investigators have found that slag corrosion rates increase linearly

with the percentage of apparent porosity within the refractory. This is usually true 

within a limited range of apparent porosity – for example in the range 12-16% 

apparent porosity – but it is not necessarily true at high apparent porosities 

(>20%). 

It is for this reason that attention is directed toward achieving higher 

densities in refractories, i.e. obtaining the lowest possible apparent porosities. 

Higher density refractories usually exhibit lower wear rates. Understandably, very 

low-density refractories may exhibit less spalling resistance leading to higher 

wear rates by another process such as thermal shock. This points to the fact that 

corrosion loss is one part of many possible contributors to an overall wear rate 

for a furnace lining. Usually one wear process, such as corrosion, is the dominant 

wear process.
Because corrosive fluids can penetrate the refractory, such penetration 

usually results in disruption and even destruction of the “matrix” of the refractory. 

The matrix is the area of “sintered fines” that hold together “bonded” bricks and 

monolithic refractories. Usually, the matrix contains more impurities and more 

porosity than aggregate particles. Therefore, corrosion affects the “weakest” 

component of the refractory at a higher rate than more dense or higher purity 

particles. In extreme cases of matrix attack, rapid erosion of coarse refractory 

aggregate can occur.

Phenomenological Description of Corrosion

Introduction

The phenomena that occur during refractory corrosion are well known to 

anyone that has been around a furnace containing reactive process components 

such as slogs or glasses. It is obvious that the reactive component, hereafter 

called the “slag”, must contact the refractory‟s exposed surface or hot face at 

elevated temperature. It is at this interface that corrosion begins and continues 

throughout the life of the lining. 

The situation is illustrated in considering a cross section of the slag and 

refractory interface (Figure 1). The refractory has a temperature gradient from the 

hot to cold face – meaning that the highest temperature is at the hot face and 

temperature declines across the refractory thickness toward its back or cold side. 

Slag is shown penetrating the refractory up until a “freeze plane” is reached. The 

freeze plane is the location in the refractory where the temperature is sufficiently 

low to cause the slag to solidify.

Figure 1: Cross Section of the Slag/Refractory Interfacial Area
The hot face of the refractory is coated with slag. In most metallurgical 

applications, this coating is relatively thin (~2-5 mm thickness). In some vertical 

wall refractory construction where the refractory is not in contact with a metal 

bath or slag pool, slag can accumulate in thick sections (>25 mm).

In some cases, two distinct slag layers can be observed on the outside of 

the refractory after a used material is sectioned. There may be a fluid “outer 

layer” over a “viscous appearing” inner layer with the latter directly contacting the 

surface of the refractory. The inner layer is typically influenced by solution of the 

refractory in the slag with the added components increasing the viscosity of the 

slag. There may be eroded particles of the refractory in the inner layer, which can 

be seen with the aid of a microscope. The two-layer situation is usually seen 

where liquid slag runs down a vertical refractory wall.

Two-dimensional representations in drawings or photomicrographs of the 

hot face zone cannot reveal the three dimensional structure underlying the area 

of the section. For this reason, pores appear that appear to be isolated, but they 

may be connected to the hot face through a channel under the observed plane. 

For this reason, pores that appear as isolated are many times filled with slag.

Slag reacts with the refractory forming new phases at the immediate hot 

face. Reaction also takes place behind the hot face where slag contacts the 

refractory at the pore walls. In metallurgical applications, metal can sometimes 

be seen penetrating the refractory along with slag. 

Third Fundamental Principle on Reactions and Temperature Gradients

A key concept is that the temperature gradient affects the extent of 

phenomena seen in slag corrosion. In a very steep temperature gradient, very 

little penetration of slag is seen with corrosion reactions more or less restricted to 

the immediate slag and refractory interface. Mobility of fresh slag (reactant) to the 

slag/refractory interface may be limited, and transport of reaction products away 

from the slag/refractory interface may be slow. Steep gradients are seen in thinwall refractory linings such as those used in utility boilers featuring water or 

steam cooling at the furnace shell.

In conventional refractory designs, the lining is at least one brick thickness 

(>225 mm), and the lining typically features a safety lining for a total thickness of 

at least 450 mm. The slag freeze plane may be located in a zone of 40-75 mm 

behind the hot face. In some cases, slag may penetrate up to 150 mm behind the 

hot face. 
Thin Wall (Steep Temperature Gradient)

The corrosion situation in a thin wall refractory lining is illustrated in Figure 

2. Here the refractory corrosion reactions occur primarily at the immediate hot 

face and thee is little or no slag penetration. Microscopic examinations usually 

show that penetration is confined to a depth of less than 100 microns (0.1 mm) 

behind the exposed hot face. This situation may be called “Stage I” of slag attack 

where reactions occur at the immediate hot face.

Figure 2: Stage I of Slag Attack

(Showing a Bonded Refractory)

It is found that the hot face temperature primarily affects the rate of 

corrosion reactions. If the hot face temperature is held just below the point that 

the products of corrosion become liquid (melt), corrosion will be very slow or nonexistent. Most authorities believe that if the hot face is maintained at no more 

than 20oC above this melting temperature (called a lowest eutectic temperature), 
reasonable corrosion rates will be observed. However, when the hot face 

temperature is more than 20oC above the eutectic, corrosion is rapid.

In some applications like utility boilers, “button fusion tests” are run to 

determine the slag melting temperature. A much more accurate test is to 

determine the eutectic or “solidus” temperature using thermal analysis. The 

solidus is recognized on heating with the appearance of an endothermic reaction 

(melting). The reaction temperature between slag and refractory can be 

determined by conducting a thermal analysis test of a mixture of ground 

refractory and ground slag.

Thin wall refractory designs rarely exhibit more extensive corrosion 

phenomena than are described as “Stage I” of corrosion. Therefore, the primary 

process variable affecting corrosion is hot face temperature. Secondary variables 

influencing corrosion rates include slag impingement velocity and slag chemistry. 

With respect to chemistry, the refractory will be soluble in the slag up to a certain 

extent (or percentage). If the slag is “satisfied” with respect to the chemistry of 

the refractory, slag corrosion should be minimal. The term “satisfied” means that 

the slag has reached the chemical solubility limit of major refractory constituents 

in the slag. Usually, the slag is not satisfied, i.e. it is corrosive. Attempts to make 

additions to the slag to obtain a satisfied chemistry (to minimize corrosion) are 

usually not economical.

There are a few cases of thick wall designs where the refractory is so 

resistant to slag corrosion that corrosion reactions are restricted to the hot face 

region. Usually these involve fusion cast refractories employed in furnaces that 

remain hot during their service life (minimal thermal cycling from hot to cold).

Thick Wall (Relatively Broad Temperature Gradient)

Thick wall refractory designs begin to exhibit corrosion on their initial 

coating of slag at elevated temperature. The initial process is Stage I of corrosion 

(Figure 2). Because of the broad temperature gradient, the refractory is 

penetrated by slag. Penetration is aided by capillary suction as the smallest 

pores in the refractory (diameter <10 microns) draw the liquid slag behind the hot 

face.

In time, extensive corrosion of the refractory takes place so that the 

refractory is at “Stage II” of corrosion (Figure 3). Stage II is characterized by two 

phenomena: (1) full penetration of the refractory and (2) extensive disruption by 

corrosion of the hot face region. Stage II follows Stage I only if there is a 

sufficiently broad temperature gradient to allow penetration.

In Stage II, the coarse aggregate in a bonded refractory exhibits 

penetration – particularly along grain boundaries (boundaries between crystals 
making up polycrystalline aggregate particles). The direct bonding between the 

matrix and the aggregate particles is disrupted, but this bonding still exists.

Slag penetration in Stage II can result in densification spalling. This type of 

spalling occurs because the thermal expansion coefficient in the slag penetrated 

zone is different than that in the unpenetrated cold face region. On continued 

thermal excursions (cooling and heating), spalling can occur at the line of 

demarcation between penetrated and unpenetrated areas. The residual lining, 

after spalling, then begins the corrosion process anew progressing from Stage I 

to Stage II again.

Figure 3: Stage II of the Corrosion Process

(Showing a Bonded Refractory)

Toward the end of the life of the refractory lining or in cases of relatively 

slow corrosion rates and where densification spalling has not taken place, the 

refractory hot face zone may progress to a final stage of corrosion that may be 

called “Stage III” (Figure 4). In this case, bonding in the hot face region and up to 

2-4 mm behind the hot face is minimal. The slag itself “appears” to be the only 

phase holding the residual aggregate particles in place. This is probably a result 

of the higher viscosity of the slag in the hot face region created as a 

consequence of dissolution of the refractory in the slag.

Because of the influence of slag viscosity maintaining some coherence in 

the hot face zone in Stages II and III with thick wall designs, a key process 

variable affecting corrosion rate is hot face temperature. Many investigators 
recommend a hot face temperature not more than 20oC above the solidus 

temperature between the slag and the refractory. Such temperature restrictions 

are impractical in many thick wall refractory designs.

Figure 4: Stage III of the Corrosion Process

Equilibrium Considerations and Phase Diagrams

A phase diagram is a map showing equilibrium phases present as a 

function of composition and temperature. The phase diagram essentially shows 

the melting relationships in a given chemical system. It is beyond the scope of 

this chapter to provide a detailed description of phase diagrams and their 

interpretation (Burgeron and Risbud).

Even without a complete understanding of phase diagrams, they can be 

used in a “simplified method” to analyze corrosion. This method provides a “first 

approximation” of the corrosion potential of the system. This technique will be 

illustrated using the Na2O-Al2O3-SiO2 system (Figure 5). In the chapter on 

Alumina-Silica Brick, two consequences of Na2O exposure were mentioned for 

alumina-silica brick including “glazing” (corrosion to form melted phase) and 

expansions. This discussion will only consider corrosion of the refractory to form 

liquids during corrosion.

The first step is to locate the point on the diagram representing the 

refractory main constituents. For this discussion, we will use a 40% Al2O3

superduty brick in contact with a corrosive phase consisting of 20% Na2O and 

80% SiO2. Both of these compositions are indicated in Figure 5 with a blue dot 
indicating the approximate composition of the slag and a green dot indicating the 

approximate composition of the brick (on the Al2O3-SiO2 composition line or 

“join”).

Figure 5: Na2O-Al2O3-SiO2 Phase Diagram Used in A Corrosion Example

The second step is to draw a line between the slag composition and the 

brick composition (indicated in red in Figure 5). It is essential to recognize that all 

potential reaction products between the refractory and the slag at equilibrium 

must be located on the red line in Figure 5.

The analysis of the corrosion potential is as follows:

1. At the slag/refractory interface, the composition must move from that of 

the refractory in a direction toward the composition of the slag, i.e. the 

local composition moves in a “northwesterly” direction.

2. The compound at equilibrium between the slag composition and the 

refractory composition is “albite” or Na2O

.Al2O3.

6SiO2, melting at 1104

oC. 

Therefore, at equilibrium, the localized melting point of corrosion products 

at the hot face is 1104oC. Note that albite is contained in the 

“compatibility” triangle between SiO2, mullite (3Al2O3

.

2SiO2), and albite. 
The compatibility triangle establishes the phases that may be present at 

equilibrium. Mullite and SiO2 (as glass or vitrified phase) are intrinsic to the 

refractory. Albite is formed as a consequence of corrosion reactions.

3. The red line joining the refractory and the slag crosses temperatures as 

low as about 1000oC. This implies that non-equilibrium liquids could be 

formed at this low temperature during corrosion. In Stage I of corrosion, 

such non-equilibrium liquids are expected. Only in Stage II of corrosion will 

sufficient refractory be dissolved in the slag to allow observation of the 

equilibrium corrosion product.

4. The red line joining the refractory and the slag composition crosses the 

Na2O

.

2SiO2 – SiO2 - Na2O

.Al2O3.

6SiO2 compatibility triangle near the 

ternary (3-component) eutectic or lowest melting temperature of 740oC. 

This implies liquids melting a low as 740oC could be formed during nonequilibrium conditions.

Considerable qualification is needed when performing this type of 

analysis. One is that the refractory contains other components than Al2O3 and 

SiO2 – notably fluxes like K2O and Fe2O3. These components can lower the 

melting points predicted by the diagram. Additionally, the slag may contain other 

components changing the melting relationships. The analysis still has 

considerable value for use in a first approximation.

Selected Models for Slag Attack on Refractories

Konig‟s Model

In 1971, Konig related the corrosion rate of refractories in bosh area of 

blast furnace walls to the temperature gradient in the wall. This analysis includes 

an assumption that the temperature gradient controls the temperature at the 

slag/refractory interface (See Figure 6 where the interface temperature is labeled 

Ts). Because the shell of a blast furnace is water cooled in the area of the bosh, 

Konig‟s analysis is appropriate for thin wall refractory with a steep temperature 

gradient fixed by the process (furnace) temperature and cold face temperature

Konig created a heat balance by assuming that the heat flux into the 

refractory surface must equal the heat flux through the furnace shell. If the heat 

flux into the refractory surface exceeds the heat flux out of the shell, the interface 

temperature (Tr) must increase causing increased corrosion rates until the heat 

balance is maintained once again. Terms used in Konig‟s analysis are given in 

Table 1.

The slag/refractory interface temperature becomes Tc (critical 

temperature) when the heat flux into the refractory equals the heat flux out of the 

shell. In other words, Tc is an “equilibrium thickness” in the lining. If the actual 
temperature is above Tc, rapid corrosion (recession) of the refractory is expected 

until a dynamic equilibrium is once again achieved.

Figure 6: Konig’s Thermal Model

Term Definition

Q1 Heat flux into the refractory surface.

Q2 Heat flux through the furnace shell.

A Area of the exposed refractory in the furnace.

t Time at elevated temperature.

Ti Temperature where i=slag surface(s), 

refractory/sag interface(r), insulation hot 

face(i), or shell(sh).

xi Thickness of component where i=slag 

layer(sl), refractory working lining(r), 

insulation(i), or shel(sh).

ki Thermal conductivity where i=refractory 

working lining(r), insulation(i), or shell(sh).

Tc Critical temperature at the slag/refractory 

interface.

fi A film coefficient for convective heat transfer 

where i= furnace environment (F) or room 

environment (E).
Table 1: Terms Used in Konig’s Analysis

The heat flow into the refractory surface is described by the equation:

Q f At(T T ) 1 F F s

The heat flux out of the furnace shell is:

Q k At(T T ) 2 ef f r E

where keff is the effective (overall) thermal conductivity through the refractory, 

insulation, shell, and film contribution in the room environment.

From heat transfer theory, the relationship for keff is:

sh

sh

i

i

r

r

E

eff

k

x

k

x

k

x

f

1

k

At equilibrium, Q1 = Q2 and Tr = Tc, so that algebraic manipulation 

provides a relationship for the “equilibrium thickness” (xr) of the working lining as:

sh

sh

i

i

E

F sl

c E

F

r r

k

x

k

x

f

T T

T T

f

x k

( )

1 ( )

Konig‟s relationship for the equilibrium thickness provides the following 

insight:

The equilibrium thickness xrincreases linearly as the thermal conductivity 

of the refractory increases.

The equilibrium thickness xrincreases at the thermal conductivity of the 

insulation increases. This is why conductive ramming materials are used 

behind high thermal conductivity basic refractories to insure the bets 

possible cooling effect of the shell.

The equilibrium thickness xrincreases as the Tc of the refractory/slag 

system increases, i.e. as more slag resistant refractories are used. This 

usually means use of higher purity refractories, i.e. higher alumina content 

or higher MgO purity.
The equilibrium thickness xr changes as the heat transfer into the lining or 

out of the shell changes, i.e. as the film coefficients change.

In attempts to apply Konig‟s analysis to the blast furnace, Herron and 

Beechan found that corrections were necessary if the refractory‟s thermal 

conductivity changes dramatically with increasing temperature. The model was 

found to describe the behavior of refractories in cyclone burners and in boiler 

applications where the initial wall thickness of refractory was ~25mm.

The Model of Endell, Fehling, and Kley

In a classic reference, Endell, Fehling, and Kley developed an empirical 

relationship for slag corrosion in a thick-walled refractory lined vessel with slag 

flowing down a vertical refractory wall. While this work was performed in 1939, it 

remains the classic reference on slag corrosion despite its empirical nature. In 

the corrosion studies, the authors used an arrangement where a solid fuel 

burned within a reaction chamber (liberating heat) with resulting ash (slag) 

impingement on walls. Terms used in the analysis are given in Table 2.

Term Definition

R Slag attack rate, i.e. refractory 

recession, in cm/s.

Lo Solubility of refractory in slag in kg of 

refractory/kg of slag.

T Absolute temperature of the refractory 

hot face.

Slag viscosity in poises.

f Fraction of ash adhering to the 

refractory wall – interpreted as slag 

impingement.

H Heat liberated in the furnace chamber in 

kcal/m3-hr.

A Ash content of the fuel in grams pf ash.

C A constant for the furnace geometry.

Table 2: Terms in the Analysis of Endell, Fehling, and Kley

The mathematical relationship for slag attack rate is:

9

1

fHA TCL 8

3

2

o

9

R
The empirical model of Endell, Fehling, and Kley for thick wall refractory 

linings provides the following information:

There is a very strong dependence of refractory corrosion rates on hot 

face temperature. In fact, temperature is the most important process 

variable that can be considered in furnace design or process control.

The slag attack rate increases linearly with the solubility of the refractory in 

the slag. While modern high Al2O3 or high MgO products were not 

included in the original research, modern practices have shown that 

corrosion rates can be reduced if the slag chemistry is adjusted to reduce 

solubility of the refractory in the slag.

In essence, the corrosion rate is inversely proportional to slag viscosity. 

This is probably why refractories exhibit some coherence in Stage III of 

slag attack because solution of refractory in the slag has increased 

viscosity at the slag/refractory interface.

In their concluding remarks, Endell, Fehling, and Kley recommend 

maintaining a hot face temperature of no more than about 30oC above the 

solidus or eutectic temperature between the refractory and slag. The authors 

further recommend that a temperature gradient in thick wall linings should not 

exceed 10oC/cm or 50oF/in. While these criteria are difficult to maintain given the 

other purposes for furnaces in processing, they establish criteria for reducing 

corrosion rates.

Case Studies of Corrosion

Alumina-Silica Brick in Ferrous Foundry Applications

Introduction

Slag used in ferrous foundry applications is found to contain lime (CaO), 

silica (SiO2), iron oxide (FeO), and other oxide constituents. As a first 

approximation, the slag can be assumed as a binary mixture of alumina and 

silica as they are usually the major components of the slag. The major 

components of the brick can be represented using the binary Al2O3-SiO2 phase 

diagram, and this allows the corrosion situation to be analyzed using the threecomponent CaO-Al2O3-SiO2 phase diagram (Figure 7).

Cases of corrosion of 70% Al2O3 and 90% Al2O3 brick are analyzed below 

using Figure 7. The composition of the brick is plotted on the Al2O3-SiO2 binary 

(2-component) side of the diagram. Then a line is drawn to the expected slag 
composition using the average CaO/SiO2 ratio of the slag. For this discussion, 

the slag is assumed to have a CaO/SiO2 ratio of ~1.0.

Figure 7: Representation of Corrosion of Alumina-Silica 

Brick In Foundry/Steel making Slag

A few things are very clear on looking at the lines joining the brick 

composition and the slag composition in Figure 7:

Both “joins” between brick and slag pass very close to the ternary eutectic 

at 1265oC implying that the slag and refractory compositions will produce 

liquid phases, i.e. exhibit corrosion, whenever they are in intimate contact 

70% Alumina

90% Alumina

Slag Composition

Refractory 

Composition
and when this temperature is reached. Other fluxing oxides typically 

reduce this temperature by as much as 50oC.

The effect of lime (CaO) – rich phases penetrating the brick is to render 

the mullite “binder” in the refractory as a non-equilibrium phase, i.e. the 

binder is dissolved into the penetrating slag phases and corrosion 

products. In 70% Al2O3 brick, the mullite is stable until the local 

composition reaches ~20% CaO, while in 90% Al2O3 brick the mullite 

disappears when the local composition reaches ~8% CaO.

The phase diagram has provided a first approximation of the reaction 

temperature between the slag and refractory for corrosion to begin, i.e. when 

liquid phase is produced as a product of corrosion reactions. This is analogous to 

the critical temperature in Konig‟s analysis or the „baseline” temperature in the 

analysis of Endell, Fehling, and Kley.

The joins between refractory composition and slag composition in Figure 7 

could have been constructed to analyze the situation with higher slag basicity, 

i.e. CaO/SiO2 of 1.5 or 2.0. In this situation, the lines would cross close to the 

ternary eutectic at 1380oC. This eutectic mixture between anorthite and gehlenite 

has the special name of “melilite”. The binary join between anorthite and 

gehlenite is shown in Figure 8, where it is easier to see the solidus temperature, 

i.e. the temperature above which corrosion begins. Thus, higher CaO-SiO2 ratios 

increase the initial reaction temperature between the slag and refractory with 

respect to formation of liquid phases and the beginning of corrosion.

Figure 8: Binary Phase Diagram Between 

Gehlenite (2CaO.Al2O3

.SiO2) and Anorthite (CaO.Al2O3.

2SiO2)

Eutectic or “lowest 

melting” liquids.

Melilite

Composition
Case Study: 70% Al2O3 Brick In An Electric Arc Furnace Roof

A microscopic study of 70% alumina brick after service in the roof of an 

electric arc furnace was performed to determine the reason for a premature 

failure of the brick in the roof. Several pieces of brick after service exhibited 

minimal thickness (Figure 9). In brick selected from other areas of the roof, a 

greater residual thickness of brick was observed. Cracks were found on saw-cut 

surfaces of most used brick from this furnace campaign, and cracks were 

arranged approximately parallel to the hot face.

Figure 9: 70% Alumina Brick After A Short Arc Furnace Campaign

The bricks were sectioned, and polished sections were produced for 

reflected light examination on an optical microscope. In the area of the immediate 

hot face, i.e. at the slag/brick interface, there was ample evidence of corrosion 

reactions (Figure 10). Recognizing that the “native” phases in the brick are 

bauxite, clay (grog), and mullite, there is no mullite visible at the immediate hot 

face. Phases in the hot face in Figure 10 include hercynite or FeAl2O4 (labeled 

“S”), calcium hexaluminate or CaO.6Al2O3 (light gray needles), lath-like corundum 

crystals exsolved from the liquid on cooling, corundum crystals (Al2O3) in a relic 

bauxite grain (lower right field), and a dark gray background of calcium aluminosilicate glass.
Figure 10: Microstructure at the immediate

Slag/Brick Interface (130X)

At a position of about 2-4 mm behind the hot face, slag penetration has 

taken place through the pore structure of the brick, and the position where mullite 

crystals appear was easily noticed because of their extreme crystal growth 

(Figure 11). The mullite crystals are seen as large light gray laths at a pore wall 

surrounded by relic bauxite grains. Small calcium hexaluminate crystals are 

located in the glass phase between mullite crystals. 

The appearance of mullite signifies the position where the CaO content of 

the penetrating phase is below about 20%. This is a consequence of the 

“dilution” of the slag by the corrosion of the brick and the lesser amount of slag 

as distance is traversed away from the hot face.

Figure 11: Position Where Mullite Crystals

Appear Behind The hot Face

At a position of about 10 mm behind the hot face of the brick, slag 

penetration stopped. Near the cold face of the brick, the microstructure of the 

brick was normal for a 70% Al2O3 brick product made using a mixture of calcined 

bauxite and clay. The conclusion from this study was that the normally expected 

corrosion reactions were found to take place, and the accelerated wear was likely 

due to poor thermal shock resistance resulting in enhanced spalling that 

ultimately shortened the campaign of the electric furnace roof. 
In summary, the corrosion processes for the 70% Al2O3 brick were:

1. Slag coated the surface of the hot refractory, and corrosion began as the 

surface temperature exceeded about 1265-1380oC. The reaction can be 

expressed in words as: “calcium (iron) alumino-silicate slag reacts with 

70% alumina refractory forming a gehlenite-anorthite glass with solution 

of mullite and bauxite in the slag”.

2. Slag penetrated the refractory filling the pores with liquid and dissolving 

the bond phase (mullite) until the local CaO content of the slag dropped 

below about 20%. Behind this area, mullite reappeared and exhibited 

crystal growth to a progressively lower extent until the freeze plane for 

slag in the brick was attained.

Case Study: 90% Al2O3 Brick From A Metal Contact Zone

In a Ferrous Foundry Furnace

In another case study involving a different foundry, 90% Al2O3 brick were 

examined for evidence of corrosion after service in a ferrous foundry furnace 

hearth/lower sidewall application. The premature failure was in a “skewback” 

brick at the furnace hearth that was supporting the lower furnace sidewalls. A 

photograph of a brick after service is shown in Figure 12. Because of the 

apparent wear of the brick, it was suspected that corrosion was a cause of 

failure.

Figure 12: 90% Al2O3 Brick After

Ferrous Foundry Service
The 90% Al2O3 brick was made using fused alumina aggregate bonded in 

a matrix of mullite and glass. The calcium-iron-alumino-silicate slag dissolved all 

of the mullite in the immediate hot face of the refractory and showed evidence of 

solution of the fused alumina aggregate particles (Figure 13). Phases other than 

corundum (from the fused alumina – labeled “A”) were hercynite (labeled “H”), 

calcium hexaluminate, and glass. The fused alumina aggregate appears to 

dissolve in the slag and form hercynite (iron aluminate spinel) at its periphery.

Figure 13: Immediate Slag/Brick Interface

Of The Used 90% Alumina Brick (130X)

The slag liquid penetrated the refractory to a depth of 10-20 mm behind 

the hot face. At a depth of about 8-10 mm behind the hot face, mullite appeared 

in the microstructure indicating that the CaO content of the penetrating slag 

dropped below about 8% (Figure 14). In the upper field of this photomicrograph, 

glass or slag (labeled “G”) penetrated between fused alumina aggregate particles 

(labeled “A”). Needle-like crystals of calcium hexaluminate are dispersed in the 

glass phase. In the lower field, mullite (labeled “M”) appears as clusters of 

individual crystals dispersed in a glass phase somewhat altered in chemistry by 

the slag penetration.

Figure 14: Area Where Mullite Reappears

Behind the Hot Face

In this case, it was determined that the corrosion reactions were those 

normally expected in a ferrous foundry, i.e. chemical solution of the bond phase 
of the brick with penetration behind the hot face. Large “veins” of solidified metal 

were found within the used brick indicating mechanical failure in service followed 

by metal penetration.

Basic Brick in Steel Making Applications

Introduction

A 90% MgO class fired, pitch impregnated brick was examined after a 

short campaign in a basic oxygen furnace producing carbon steel. The slag 

composition can be viewed as a high CaO/SiO2 ratio liquid containing a 

substantial amount of iron oxide. Because the major component of the brick was 

MgO, the ternary system CaO-MgO-SiO2 can be used to view the corrosion 

processes as a first approximation (Figure 15).

Figure 15: CaO-MgO-SiO2 Phase Diagram With Illustrations of “Joins” 

Between Brick (MgO) and Slag (On the CaO-SiO2 Join)

In the case of the 90% MgO class refractory, the brick contained 95.5% 

MgO, 2.7% CaO, and 1.4% SiO2. This implies that the “native” silicate in the brick 

is dicalcium silicate (2CaO.SiO2) - a very refractory second phase providing for a 

Slag

Compositions

Brick

Composition
part of the bonding in a product where direct sintered bonds between adjacent 

MgO aggregate particles predominate. In service, a transient period exists in the 

steel refining cycle when silicon is being oxidized in the metal bath providing for 

temporary exposure of the refractory hot face to less basic silicates. This is also 

a period of enhanced iron oxide contact with the refractory.

Since the refractory composition can be represented by the “MgO” apex 

on the phase diagram (Figure 15), joins can be drawn to different CaO/SiO2

ratios illustrating the corrosion process:

For a CaO/SiO2 ratio of 1.0, the refractory will react with the slag forming 

monticellite, CaO.MgO.SiO2, which melts at 1485oC. An iron-containing 

analog of monticellite called “iron monticellite” (FeO.MgO.SiO2) has a 

melting point of 1230oC. Therefore, corrosion starts above about 1230oC 

when liquid phases are formed. Slag penetration can result in the 

disruption of bond phases that are native to the brick lowering the 

effectiveness of the “bonding silicate phase” in the immediate hot face 

area.

For a CaO/SiO2 ratio of 1.5, merwinite or 3CaO.MgO.2SiO2 is the reaction 

product formed in the hot face. Merwinite melts “incongruently” at 1575oC 

in a process given as:

3CaO.MgO.2SiO2 = 3CaO.2SiO2 + liquid phase

It is interesting that as the slag basicity increases (i.e. as the CaO/SiO2

increases), corrosion reactions begin at higher temperature. This is one 

reason that steel makers wish to have solution of limestone slag additions 

as early as possible in a refining cycle to retard slag corrosion.

For a CaO/SiO2 ratio ≥ 2.0, the lowest eutectic temperature is at ~1790oC. 

With out considering the iron oxide and other components of the slag, this 

implies that the reaction product between the slag and the brick is above 

most steel making temperatures. However, the iron oxide plays an 

important role significantly reducing the temperature of the reaction 

products.

The Role of Iron Oxide in Corrosion

The role of iron oxide in corrosion in steel making slag and corrosion is to 

significantly reduce the melting temperatures predicted by only considering the 

CaO-MgO-SiO2 phase diagram. Some of the reaction products commonly found 

between basic refractories and steel making slag are given in Table 3. 

It is immediately obvious that most of the iron containing compounds melt 

in the area of 1200-1350oC while most of the calcium silicate phases melt 
above1485oC. It is also true that melting can typically begin below these 

indicated temperatures by as much as 50oC due to the effect of other impurities 

and due to the presence of multi-component eutectics not seen on three 

component diagrams.

System CaO-MgO-Al2O3-SiO2 System FeO-MgO-Al2O3-SiO2

Mineral Melting 

Temperature, 

oC 

Mineral Melting 

Temperature, oC 

2MgO.SiO2

(Forsterite)

1900 FeO.SiO2

(Fayalite)

1205

CaO.MgO.SiO2 

(Monticellite)

1485 FeO.MgO.SiO2 

(Iron monticellite)

1230

3CaO.MgO.2SiO2 

(Merwinite)

1575 2CaO.(Fe2O3, 

Al2O3)

.SiO2

(Iron gehlenite)

1285

2CaO.SiO2 

(Dicalcium silicate)

2130 FeO.Al2O3

.

5SiO2

(Iron cordierite)

1210

3CaO.SiO2 

(Tricalcium silicate)

2070 2CaO.Fe2O3

(Dicalcium ferrite)

1489

CaO.Al2O3

.

2SiO2

(Anorthite)

1553 4CaO.Al2O3

.

Fe2O3

(Brownmillerite)

1395

2CaO.Al2O3

.SiO2

(Gehlenite)

1593 6CaO.2Al2O3

.

Fe2O3 1365

CaO.Al2O3

.

2SiO2and 

2CaO-Al2O3-SiO2

(Melilite)

1380

Table 3: Reaction Products Between Steel Making Slag And Basic 

Refractories with Melting Points for Full Melting or Partial Melting 

(Peritectic Reactions)

The corrosion process between calcium-iron-silicate slag and magnesia 

containing refractories can depend on the mobility and concentration of 

reactants. A photomicrograph of the immediate slag/brick interface region in a 

brick from a steel making furnace location above the metal line is shown in 

Figure 16. Here the reaction products are “mixed” iron spinel or (MgO, 

FeO).(Al2O3,Fe2O3), dicalcium silicate (2CaO.SiO2), and dicalcium ferrite 

(2CaO.Fe2O3). Because of the high concentration of iron oxide dust in the 

furnace atmosphere and the lack of persistent slag attack (or flow), the spinel 

phase formed a sort of “barrier coating” over the face of the brick.

By contrast, a brick from a slag line application exhibited a very different 

corrosion process (Figure 17). Here the products of corrosion appear to be 

dicalcium ferrite (2CaO.Fe2O3) penetrating around individual magnesium oxide 

crystals (rounded gray crystals) with isolated pockets of black appearing 

dicalcium silicate (2CaO.SiO2) present in the slag layer. Magnesium oxide 
crystals appear to be eroded from the refractory surface in response to the 

extremely high temperatures and the intensity of the slag contact.

Figure 16: Corrosion Interface in an MgO Brick

From an Upper Sidewall Location

Figure 17: Corrosion Interface in an MgO Brick

From a Slag line Location

In these cases, the microstructures at the reaction interface are extremely 

different despite the fact that the brick were in the same furnace. In the upper 

sidewall, reaction with iron oxide and lime containing dust contributed to the 

corrosion process. By contrast, in the slag line, persistent contact with flowing 

slag caused corrosion and erosion processes to take place.

Case Study: 90% MgO Burned-Impregnated Brick 

After Service in the Basic Oxygen Steel making Furnace

A typical used brick from a “bottom tuyere pedestal” (also called a “bottom 

plug”) is shown in Figure 18. The bottom plug contains tuyeres where oxygen is 

injected into the molten metal bath to facilitate the refining process. The brick are 

exposed to the localized high temperatures from the exothermic reaction 

between the oxygen and the metal. It is notable that this brick exhibits a layer 

apparently in the process of spalling away from the brick hot face.
Figure 18: 90% MgO Burned Impregnated Brick

After Service in a BOF Bottom Plug

The hot face region of the brick is shown in Figure 19 at low magnification. 

The immediate slag/brick interface is in the upper field of the photomicrograph. 

The slag phase penetrating among individual MgO crystals is dicalcium ferrite 

forming a thin “dense layer” in the outer zone of the brick. Behind the dense 

layer, a zone of porosity is noted as „black areas” – areas where the impregnated 

carbon phase in the brick was removed by oxidization.

Further behind the hot face in Figure 19, a large void is seen in the lower 

left of the photomicrograph. As will be seen in the following, this void resulted 

from a defect in the formulation of the brick. With prolonged exposure to heat, the 

defect in a precursor to the “pealing” away of the surface by spalling as seen in 

Figure 18.

The corrosion process is seen in a field near the immediate hot face 

(Figure 20). Here dicalcium ferrite (bright phase) and dicalcium silicate penetrate 

around individual MgO crystals from aggregate particles. Several smaller 

rounded MgO crystals are in the final phase of dissolution. Even though

dicalcium silicate is a “compatible” phase with MgO, it will still dissolve MgO – as 

is evident from the photomicrograph. Thus, the corrosion process can, in part, be 

represented as solution of MgO in 2CaO.SiO2.
Figure 19: Hot Face Region of The MgO Brick (50X)

Figure 20: Field in the Hot Face Zone (130X)

Further behind the hot face, a zone was discovered where the “matrix” of 

fines (and carbon) shrunk away from the aggregate particles (Figure 21). This 

created voids at the periphery of the coarse aggregate particles in the brick, and 

these voids eventually consolidated leading to the larger void of Figure 19. These 

voids lead to the pealing away of layers during service (spalling of thin layers) 

Void
causing excessive wear. When confronted with the evidence, the developers of 

the brick were amazed that this separation took place because they had not seen 

it in laboratory “coking box” experiments. Eventually, the idea was advanced that 

the coking box temperature of about 1000oC was quite a bit different than the 

service temperature of the brick (≥ 1600oC).

Figure 21:

Magnesia-Chrome Brick in AOD Service

In other case studies, magnesia-chrome and dolomite refractories were 

examined after service in an argon-oxygen-decarburization (AOD) reactor 

producing stainless steel. In this instance, slag specimens were collected during 

typical refining cycles, and bricks were examined after the end of the service 

campaign.

The AOD refining cycle of slightly over one hour includes an initial period 

(decarburization cycle) where silicon is oxidized in the metal producing a very 

viscous or “chunky” slag. A photomicrograph of this early slag shows clusters of 

spinel phases or (MgO, FeO).(Cr2O3

.Al2O3.

Fe2O3) and bright metal droplets interdispersed in a calcium silicate glass (Figure 22).

Figure 22: Early AOD Slag (50X)
The angular spinel crystals and metal droplets are shown at higher 

magnification in Figure 23. The “chunky” nature of the slag is due to the viscosity 

of the slag created by the high concentration of spinel crystals.

Figure 23: Detail of Early AOD Slag

The early AOD slag also contained eroded masses of MgO crystals. One 

such fragment can be seen in the lower left field of Figure 22 and in a 

photomicrograph taken using a scanning electron microscope (Figure 24). In this 

latter picture, spinel crystals in a continuum of calcium silicate glass surround a 

dark gray mass of MgO in the center of the photomicrograph.

Figure 24: Eroded MgO Mass in Early AOD Slag

The AOD refining cycle also includes a short “reduction period” where 

argon is injected into the metal to recover chromium from the slag. In effect, the 

spinel phases are substantially reduced. This cycle also employs an increased 

slag CaO/SiO2 ratio to facilitate chromium recovery. 

A photomicrograph of a reduction slag is shown in Figure 25. Only a few 

spinel crystals protrude from the glass indicating high efficiency of reduction. The 

slag shows sign of water etching during polishing consistent with the solubility of 

merwinite and/or gehlinite (See Table 3).
Figure 25: AOD reduction Slag (130X)

Fired magnesia-chrome brick were used in early AOD linings. Despite a 

more recent trend to use magnesia-carbon products in these furnaces, the 

examination of magnesia-chrome brick illustrates several principles of corrosion 

processes. One brick from a mid-barrel location in a 50-ton AOD reactor is 

shown in Figure 26. The continuous glassy slag coating on the hot face is typical 

for AOD service.

Figure 26: Used Magnesia Chrome Brick

After AOD Service

The immediate how face area of the magnesia-chrome brick is shown in 

Figure 27. A chromite grain (labeled “Cr”) is in the center of the photomicrograps 

with a coarse magnesia aggregate particle to the left of the field. The slag layer is 

at the top of the photomicrograph. It is apparent from the “jagged” nature of the 

brick surface that the corrosion process has preferentially removed the matrix of 

the refractory leaving coarse grains to protrude into the slag layer. What appears 

to be “corrosion” of the periphery of the chromite grains can also be seen at this 

low magnification.
Figure 27: Slag/Brick Interface (50X)

A detailed view of the hot face region is shown in Figure 28 (slag layer 

toward the top of the field). Here, merwinite or 3CaO.MgO.2SiO2 can be seen 

toward the top of the field (labeled “m”). The other slag silicate is monticellite or 

CaO.MgO.SiO2. Both of these are expected as compatible phases for slag in the 

CaO/SiO2 range of 1.0-1.5 in contact with MgO containing refractories (See 

Figure 15). 

The corroded appearing edges of the chromite grains are “secondary” 

spinels that did not exist at the service temperature of the brick in the hot face 

region. The process was dissolution of the constituents of the chrome ore (Al2O3, 

Cr2O3, and Fe2O3) in the penetrating slag followed by preferential precipitation 

around the chromite grains on cooling of the lining for the last time. These spinels 

did not form on the “top” of the chromite grains in the immediate slag/brick 

interface because of the rapid quenching after the last heat of steel was made. 

Metal droplets are also seen in Figure 28. These may have formed, in part, due 

to reduction of oxides in the chromite to the metallic state.

Figure 28: Detail of the Hot Face Region (130X)
The corrosion process can be summarized as follows:

Solution of MgO in the slag phases, penetration of slag into the brick, and 

disruption of the direct bonding behind the hot face.

Solution of chromite in the slag phases and partial chemical reduction of 

oxides to the metallic state.

In magnesia-chrome brick and other bonded refractories, slag penetration 

can occur to a depth of 75 mm or more behind the hot face. While significant 

chemical corrosion is usually limited to a depth of less than about 25 mm behind 

the hot face, the presence of the slag changes the physical properties of the brick 

due to shrinkage of the refractory and other processes. If the thermal expansion 

coefficient in the penetrated zone becomes different than that in the cold face 

region, spalling can occur near the interface between the penetrated and “unpenetrated” or cold face zones. Used brick can exhibit cracks in this zone and 

even larger voids where cracks have apparently joined. This phenomenon is 

called “densification spalling”, and it can result in spalling loss of sections of 

about 75 mm in thickness.

Dolomite Brick in AOD Service

Dolomite refractories were used in AOD service in the early stages of 

development of the AOD process. Dolomite brick contain “islands” of MgO 

crystals in a continuum of CaO. The hot face region of a used dolomite brick after 

AOD service is shown in Figure 29. In this photomicrograph, the immediate hot 

face is to the right of the photomicrograph. The central area of the 

photomicrograph is the affected zone where corrosion has resulted in reaction of 

slag with the CaO continuum in the brick forming dicalcium silicate and tricalcium 

silicate (the black phase etched with water in polishing). Metal droplets are 

apparent in this region. Aside from chemical corrosion, spalling is also known to 

occur in AOD service.

Figure 29: Dolomite Brick After AOD Service (130X)
Gas Phase Reactions

The gas phase in a furnace may be the partial or even sole contributor to 

corrosion reactions or other deterioration in refractories. Some of the processes 

that have been recognized include:

1. Alkali transport to refractory surfaces resulting in surface reaction and 

penetration.

2. Reduction of the refractory and gas phase corrosion.

3. Alternating oxidation and reduction causing sequential dimensional 

changes in refractories that may enhance spalling.

The potential for alkali to cause “glazing” on the surface has been 

previously discussed (See Figure 5). It is less well recognized that alkali 

compounds can volatilize in furnace atmospheres resulting to transport of the 

volatile species to the surface of the refractory where vapor phase penetration 

may also occur.

The melting points of sodium and potassium reaction products with 

alumino-silicate refractories are given in Table 4. The melting points of the 

phases themselves are not indicative of the very low eutectic temperatures found 

in the system. For example, the lowest eutectic in the system Na2O-Al2O3-SiO2

(Figure 5) is 732oC. 

Phase Melting Point, oC

Na2O

.Al2O3.

6SiO2

(albite)

1100

Na2O

.Al2O3.

2SiO2

(carnegite)

1526

K2O

.Al2O3.

6SiO2

(carnegite)

1220

K2O

.Al2O3.

4SiO2

(leucite)

1686

K2O

.Al2O3.

2SiO2

(kaliophilite)

>1700

Table 4: Alkali Phases Found in Used Refractories
The sidewall of a sawdust fired tunnel kiln producing facing brick is shown 

in Figure 30. In the central section of the picture, the effect of the plume of 

combustion from the top burner position nearest the wall is seen (kiln draft from 

left to right). A black slag is seen coating the fireclay brick with white salt deposits 

visible in this picture taken after seven years of service. There was only surface 

glazing of the brick in the sidewall of the kiln. Glazing is also shown on the 

insulating firebrick surface in the crown of the kiln (Figure 31).

Figure 30: Sidewall of Sawdust Fired Tunnel

Kiln After Seven Years of Service

Figure 31: Crown of Sawdust Fired Tunnel Kiln After Seven Years 

Showing Surface Glazing and Sheet Spalling

The sidewall brick had expanded due to alkali attack to the point that they 

were replaced. In addition, the insulating firebrick in the crown exhibited both 

surface glazing and “sheet” spalling. Sheet spalling is well known, and it usually 

results from a subsurface expansion reaction caused by penetration of an agent 

in the gas phase that subsequently reacts with the refractory.
Alkali halides are volatilized in kiln atmospheres. They undergo hydrolysis 

reactions in a moist high temperature environment forming alkali oxide species. 

These oxides can cause glazing or expansion. The expansion reaction with 

potassium is particularly deleterious for refractories with the reaction given as:

Al2O3

.

2SiO2 + K2O = K2O

.Al2O3.

2SiO2

where Al2O3

.

2SiO2 is metakaolin (a decomposition product of fireclay) and the 

reaction product is kaliophyllite. The volume expansion due to formation of 

kaliophyllite is on the order of 15% (linear expansion of ~5%). As potassium is a 

product of combustion of sawdust, reactions of this type are expected as the 

cause of the problems in the sidewalls of tunnel kilns fired with sawdust and as 

the cause of the sheet spalling (Figure 31).

Gas phase corrosion is also known in refractories containing SiO2. It has 

long been recognized that products containing SiO2 can be “reduced” by 

hydrogen or by penetration of aluminum metal. The gas phase reduction by 

hydrogen can be represented as:

SiO2(s) + H2(g) = SiO(g) + H2O(g)

where (s) signifies a solid phase and (g) signifies a gaseous phase.

Refractories exposed to hydrogen (or carbon monoxide) can experience 

gas phase removal of SiO2 through the process of forming silicon monoxide. For 

example, at 1000oC a CO/CO2 ratio of about10-5 will be sufficient to reduce SiO2. 

Mullite reduction and silica loss at temperatures exceeding 1500oC has been 

seen in petroleum calcining kilns.

Laboratory Slag Tests

Laboratory slag tests have been developed over the years for evaluating 

corrosion potential. These slag tests range from a “static” cup test to the 

“dynamic rotary slag test. The terms “static” and “dynamic” refer to whether the 

slag is stationary or whether it moves over the refractory surface respectively. 

Some of the test methods are summarized in Table 5. The rotary slag apparatus 

is shown during a test in Figure 32. The slag is running out of the apparatus into 

a well while a technician is checking the temperature using an optical pyrometer.

Microscopic studies have shown that the same types of reactions occur in 

the rotary slag test as in brick examined after service in metallurgical furnaces. 

The slag/brick interface area in a magnesia-chrome brick from a rotary slag test 

using a steel making slag is shown in Figure 33.
Test Type Reference or Comment

Cup Static The test usually consists of a ~50mm diameter 

hole core drilled or formed into the refractory that 

is packed with slag and heated to a selected test 

temperature for a set time. The test only reflects 

the isothermal reaction and penetration potential 

between the slag and the refractory.

Drip Dynamic Slag drips over a surface with evaluation of the 

depth of a “notch” cut in the surface by corrosion. 

The test is described in ASTM C768. The test 

reflects the isothermal reaction potential between 

the slag and refractory when there is a constant 

supply of unreacted or “fresh” slag.

Rotating 

Spindle

Dynamic Refractory bars or rods ~25X25X200mm are 

partially immersed in a crucible containing slag 

while the specimen is rotating. The test reflects 

isothermal reaction potential of a flowing slag 

without a constant supply of “fresh” slag.

Rotary Dynamic Beveled refractory shapes (formed by saw 

cutting) or monolithic materials constitute a lining 

of a small rotary furnace that is constantly fed 

with slag. After the test, the specimens are 

examined for corrosion loss and slag penetration. 

This test is described in ASTM C768. The test 

reflects corrosion under a temperature gradient 

with a constant supply of “fresh” slag.

Table 5: Laboratory Slag Tests

Figure 32: Rotary Slag Test for Refractories
Figure 33: Slag/Brick interface in a Magnesia-Chrome

Brick After A Rotary Slag Test

The corrosion process in Figure 33 includes solution of the MgO in the 

slag consisting primarily of monticellite. The periphery of chromite grains appears 

to be attacked, but the “porous” spinel phases were formed on recrystallization 

from the melt when the refractories cooled after the test. The elongated nature of 

MgO crystals is typical under the influence of the steep temperature gradient in 

the rotary slag test.

Concluding Comments

Slag corrosion by liquids occurs whenever a threshold temperature is 

exceeded which is usually when melting occurs between the refractory and the 

slag. Corrosion results in solution of refractory constituents in the liquid phase 

resulting in loss of thickness of the refractory lining. The rate of corrosion is 

dependent on the chemical environment and on the hot face temperature of the 

refractory.

Models for the corrosion process point to the fact that hot face 

temperature is the most important variable in controlling refractory life. To reduce 

corrosion loss, the usual approach is to employ higher purity refractories or lower 

porosity refractories (such as fusion cast brick) because precise limits on hot face 

temperature are not practical given process goals. In new furnace designs, 

consideration should be given to hot face temperature control as a strategy to 

attain the most economical refractory cost.

Corrosion is primarily a chemical process, and the potential for corrosion 

can be estimated by reference to phase equilibrium diagrams. These diagrams 

can allow prediction of the “threshold temperature” for liquid formation. 

Microscopic techniques allow identification of particular corrosion reactions. 

Corrosion can also take place through gas phase reactions. The most 

fundamental way to limit corrosion is to understand the chemical and physical 

processes involved in corrosion and to formulate strategies to minimize these 

processes.
Dedication 

This chapter is dedicated to the memory of Mr. Walter S. Treffner, the long-time 

Research Director of the former General Refractories Company (USA). 

Mr. Treffner was an accomplished microscopist and a dignified researcher in the 

field of refractories. Mr. Joseph L. Stein, former President of General Refractories 

(USA), is acknowledged for substantial inspiration on the author‟s work.

References and Further Reading

C. Burgeron and S. Risbud, Introduction to Phase Equilibria in Ceramics, The 

American Ceramic Society, Inc., Columbus, OH (1984).

V. G. Konig, Stahl u. Eisen, 91 (2), pp. 63-69 (1971).

R. Herron and C. Beechan, Ceramic Engr. & Sci Proceedings, 2 (11-12), pp. 

1126-1140.

K. Endell, R. Fehling, and R. Kley, J. Am. Ceram. Soc., 22, pp. 105-116 (1939).