2 korosi pada keramik

5 Corrosion of Hard Materials K. G. Nickel and Y. G. Gogotsi

5.1 Introduction

Corrosion may be defined as the physical and chemical alteration of a material due to its interaction with the environment of interest. It must be emphasized that corrosion resistance is not a material property but a system property and real environments have a high variability in both chemical and physical boundary conditions.

Corrosive alteration leads to a change in the desired original properties and so corrosion usually, but not exclusively, has adverse effects and causes a decrease in strength, toughness, hardness, and increases wear. There are many reasons for the degradation of properties by corrosion. Decrease of the cross-section carrying the load, pitting, stress induced by phase changes, and grain separation are some of the common features.

Some hard materials, namely Sic and Si3N4, are designed to serve at high temperatures, where corrosion resistance is an essential feature. This is the reason why a large proportion of the ceramic literature on corrosion, including recent textbooks [ 1-51, deals with these materials.

In order to react with the environment, the material must not be in equilibrium with it. The hard materials covered in this chapter are non-oxides and hence this condition is fulfilled for all of them in almost every application environment, including ambient air. The applicability of hard materials thus rests on the kinetics of the processes.

Before discussing individual materials we will introduce important terms and concepts in corrosion like attack modes, media classes, location of attack, kinetics, and their parameters. We feel that this is necessary for understanding of the discussion of the materials, because there is no generally accepted standard procedure that allows us to compare the corrosion behavior of materials as easily as other properties like mechanical strength or hardness.

5.2 Corrosive Media

The compositional variability of environments is so vast that only very few attempts have been made to classify corrosion behavior from the point-of-view of the discrete environment at hand. A noteworthy exception is the Dechema handbook [6],in which valuable information may be found if the corrosive agent is simple and known (e.g. acetates, chlorine, steam). However, its main thrust is on the corrosion of metals.

Handbook of Ceramic Hard Materials Edited by Ralf Riedel

Copyright Q WILEY-VCH Verlag GmbH, D-69469 Weinheim (Federal Republic of Germany), 2000

5.3 Corrosion Modes 141

The more common classification scheme is to divide the corrosive media into their state of aggregation, that is to subdivide into corrosion by solids, liquids and gases. While solid state corrosion is rarely dealt with, we have vast amount on literature on hot gas corrosion. The case of corrosion by liquids is commonly further subdivided into more specific cases, such as aqueous corrosion (e.g. acids and water), corrosion by glasses, metal melts and salt melts. The last case is for historic reasons known in the form of a rather misleading expression: hot corrosion. A special case, which spans from the liquid into the gaseous state is given by the corrosion in hot water systems: hydrothermal corrosion.

At first glance this may not make too much sense, but the experience is that the different names mentioned above correspond to main forms of the attack mode and often refer to different application temperature levels.

5.3 Corrosion Modes

5.3.1 Active and Passive Corrosion

A fundamental distinction between corrosion modes is the division into active and passive corrosion. The active mode is characterized by the loss of material to the environment and results in the decrease of size and weight of the specimen. The loss may be in form of gaseous or dissolved species.

The passive mode denotes a process, where the material reacts with the environment to yield a new condensed phase, often in form of a layer or scale on the surface. In most cases it is associated with a weight and/or component size gain.

In the strict sense, the passive mode implies that the newly formed phase is retard- ing the process, that is corrosion is slowed down with time (protective scale), but there are cases in which the scale is nonprotective: a scale with cracks, low viscosity, or foamy texture may not hinder the access of the corrosive agent to the substrate.

5.3.2 Homogeneity and Location of Attack: Internal, External and Localized Corrosion

If there is a perfectly homogeneous surface attack and a removal of material or a scale growth is taking place solely and uniformly on the surface of the material, the process is denoted external corrosion.

This is most likely if we have a perfect single-phase material. As soon as grain boundaries in general or secondary phases appear as discrete particles or grain- boundary phases the process is likely to become more complex, because each phase of the assemblage will react differently to the environment at hand.

An easily envisaged form of the problem is the attack of oxygen on a S ic containing discrete free carbon particles. At high temperature the carbon reacts in an active mode (formation of CO,T), while the Sic forms condensed Si02. Inhomogeneous or

142 5 Corrosion of Hard Materials

local corrosion with bubble and pit formation is the result, at least in the early stages of this oxidation [7].

If the secondary phase is a grain-boundary phase, which allows a relatively fast transport of the corroding agent into the material via the grain boundaries, we will have an attack of the main material and/or the grain-boundary phase not only on the surface, but in a zone stretching out into the bulk of the specimen. This is known as internal corrosion.

Despite its localized character on a microscopic scale, intergranular corrosion may be quite homogenous on the macroscopic scale. Additionally, the process of internal corrosion is also capable of slowing down with time, so it may offer some protective quality.

It must be emphasized that those processes do not exclude each other. Internal corrosion may occur while a scale is formed on the outside. We will discuss the problem in the context of the individual materials.

5.4 Corrosion Kinetics

5.4.1 Physical Boundary Conditions

In modeling the kinetics of corrosive processes and quoting numbers of kinetic constants it is usually inherently implied that we have steady-state conditions with a constant activity of the corrosive agent in the environment and an infinite reservoir of the material to be corroded.

The first condition is met in many technical applications, where the corrosive medium is flowing, the second one may be true or not. But even the first condition is violated for some corrosion conditions, such as when a substance is deposited on the surface of the material by a physical process and can hence be treated as an attack by a substance which was implanted at a given amount without replenishing.

Exact corrosion kinetics must be modeled by solving the second law of Fick for the geometry of the case at hand. However, in some cases a net effect may be calculated from simple thermodynamics, as for closed system conditions in active corrosion [8]. For the case of diffusion through scales it has been demonstrated that quasi-steady-state modeling is often a good approximation for an exact solution, at least for conditions tD /x2 > 2 [9] (where t = time, D = diffusivity, x = layer thickness). Some basic solutions for situations with instant singular corrosion can also be found in the literature [lo].

But even under steady-state conditions there is a profound influence of physical boundary conditions on corrosion behavior. The most widely known example of this is the boundary between active and passive oxidation of silica-formers. The classic modeling has been done by Wagner (1 11 for silicon.

The line of argument is as follows. Active oxidation of silicon occurs by the formation of gaseous SiO and passive oxidation by the formation of silica:

Si + 402 H SiOt (1)

5.4 Corrosion Kinetics 143

and

Si + O2 H Si02.

Calculations of the oxygen partial pressure P(0,) for the boundary condition both by simply taking the values from the JANAF-tables [12] or by using a modern thermochemical program [ 131 yield similar results with an equilibrium P(0,) as low as bar at 1000°C. The prediction from the equilibrium would thus be, that at conditions with P ( 0 2 ) > lop2* bar we should find the formation of a silica layer on a bare Si surface. Experimental observation shows that this only takes place at P(02)-levels more then 20 orders of magnitude higher.

Wagner’s solution was to transform the problem into a physically controlled one. Silica is formed at conditions, which allow its condensation from the gas phase in equilibrium with the substrate:

4 Si + Si02 * SiO. (3)

If reaction (1) is fast, a concentration gradient of O2 inwards and SiO outwards must exist. Then the interdiffusion coefficients and the effective thickness of the boundary layer will control whether the critical P(Si0) of the condensation reaction (3) is reached or not. From tables and estimations of those physical constants we can relate back to the oxygen pressure of the bulk gas necessary to induce the critical P(Si0) on the surface. The calculation yielded reasonable results for the active- passive boundary of Si in streaming atmospheres with low oxygen contents and accordingly the theory was later applied to other silica-formers [ 14,151.

This success story demonstrates that equilibrium calculations disregarding physical boundary conditions can be misleading, and that thermochemical calculation is capable of giving good results if those boundaries are taken into account.

5.4.2 Active Corrosion Kinetics

Under steady-state conditions the loss of material to the environment by decomposition into gaseous species or dissolution is usually a reaction or a diffusional problem. For reaction control we have the flux J as

J = krci, (4)

where J = mass flux, k , = reaction coefficient, ci = concentration of agent at interface, and for first-order reactions a linear law follows under steady-state conditions.

The basic shape of the kinetics for diffusion control can be seen by looking at Fick’s first law

dc dl ’

J = -D-

where I = diffusion path length. With constant conditions, the gradient from the surface to the bulk of the corrosive medium becomes constant, 1 stays constant and accordingly the flux, J , becomes constant with time. Therefore the change in


10

h - 5

Y f c 0 In In a 0 0 2

.-

v

- 8 -5

0 1 2 3 4 5 6 7 8

log (time [s])

Figure 1. Predicted behavior of a model case for active corrosion under different external conditions.

size or mass (Ax) is also linear with time t:

AX = kit. ( 6 )

This expected linearity is usually experimentally confirmed. A theoretical prediction of the size of the h e a r rate constant kl in E q . (6) for a given temperature and composition of medium and material is possible, when thermodynamic data are available to calculate the reaction, and boundary parameters such as the geometry of the sample, viscosity, stream velocity, Schmidt and Reynold’s numbers are known or can be estimated [16].

Other approaches to active corrosion prediction utilizing thermochemical calculations [8,17,18] require the experimental determination of effective parameters. They show both the importance of physical boundary conditions and the extremely low level of partial pressures at which active corrosion is potentially dangerous.

In Fig. 1 a model case applying such simplified kinetic equations [17] is illustrated. Thermodynamic calculations for Sic in H2 yield partial pressures of about lo-’ bar for SiH4 and CH4 even at moderate temperatures of 1 100°C, so that the reaction

Sic + 4H2 w SiH4 + CH4 (7)

is used as the dominant reaction. Geometric boundary conditions (a tube of 1 m length with an inner diameter of 30 cm) allow us to calculate a volume flow through such a pipe from a given stream velocity. The calculated material recession after a given time ranges from less than a monolayer in the closed system to more then 70mm per year for ideal, and hence gas velocity independent, evaporation. The model for intermediate cases indicates that both mild and conservative applicability limits for high speed civil transport systems [18] are not met even at low gas velocities.


It is interesting to note that different approaches [17,18] result in applicability limits, which specify partial pressure limits of <lop7 bar to be tolerable in applications with strong security requirements in the active corrosion mode unless reaction control retards the process.

5.4.3 Basic Passive Corrosion Kinetics

Passive corrosion with scale formation on the substrate is often more complicated. The simplest form is that of reaction control. If a reaction is slow compared to the delivery of the agent to the reaction site, the concentration of the agent at the interface is as high as the external concentration of the corrosive medium. For a given temperature this reaction rate constant, k , dictates then the linear rate constant k~ in Eq. (6) and linear kinetics prevail. Only this time Ax is positive and denotes a layer growth or a mass gain.

Likewise Eq. (6) is correct, if the chemical reaction rate constant is high (= fast reaction) and the growing scale is not hindering the agent from access to the substrate. This situation is most easily envisaged in a scale with plenty of short- circuits like cracks or open porosity. Here the concentration gradient develops from the substrate into the bulk of the environment. Accordingly the analysis is identical to that for active corrosion kinetics above, which gave the linear kinetics.

Another simple case is a homogeneously growing single-phase layer, which allows only a slow diffusional transport of corrosive medium to the substrate. Here the steady state means for Eq. (5) that a layer, and hence the diffusion path length 1, is growing with time while the absolute concentration difference Ac stays constant. At the substrate-oxide interface, the equilibrium with the material is achieved or the concentration of the agent can be approximated to zero, while the concentration at the surface of the scale is constant at the solubility limit of the scale material.

The law appropriate for this situation as derived from Fick's first law is called the parabolic law due to diusional control and introduces the parabolic rate constant k,, which has units of m2 s-' or kg2 m-4 s-' , depending on the definition of Ax as scale growth or mass gain:

(ax)2 = k,t (8)

or

ax = fi. (9)

The derivation of the parabolic law above is only a good approximation for a situation with a sufficiently thick pre-existing scale. It has been shown in a classic paper by Deal and Grove [19] on the oxidation of silicon, that the correct parabolic law should recognize the chemical reaction rate and the gas transport coefficient along with the diffusion coefficient. Their analysis demonstrated that the parabolic rate law is rather

(ax)2 + axA = B( t + .) (10)


with a linear ‘short-term’-solution for t << T and a parabolic ‘long-term’ solution identical to Eq. (8) for t >> T ( A , B, T = constants). It should be noted that all those constants vary with temperature and that the so called ‘help time’, T , can be of appreciable length: For silicon at 1700°C we have a time of at least several tens of hours, during which the kinetics are basically linear! A nice review of the complexity of the ‘simple’ parabolic oxidation of Si, particularly in its initial stages, has been given by Irene [20].

A third basic law for passive corrosion comes from the analysis of the situation, where the reaction product formed on the material is completely blocking the further access of the corrosive agent. This law was derived from the study of metals at moderate temperatures. It is easily envisaged: If every particle formed protects the substrate completely, we have a situation, in which the surface area available for reactions is diminished rapidly. The mathematical form following is the logarithmic law

with the logarithmic rate constant klog. The same equation is valid for a more complex situation, pertinent to ceramic systems. If the diffusion through a scale is rate controlling and the scale crystallizes a phase, which is either completely blocking for the agent or has a diffusion coefficient much smaller than the original scale, the effective cross section for the agent transport is likewise diminished. For a constant crystallization rate we will end up with Eq. (1 1).

However, natural nucleation and crystal growth is complex, dependent on many variables and usually not constant with time. The energy requirements are often so strongly temperature dependent that we see either none or full crystallization. Thus in pure systems we will often have kinetic breaks with time or temperature rather

h - u)

c ._

=I 300 P 2 e c 2. 200 u) ln Q) c Y 0 .-

100 tJ >. m -

0 0 50 100 150 200 250 3

Time (arbitrary units)

Figure 2. Basic passive laws.

I0


then a nice curve following Eq. (11) exactly. This may be different for systems in which phase formation or crystallization is linked to chemical changes of the scale, induced by diffusion of components from the interior.

The shapes of curves with behavior according to the basic laws is illustrated in Fig. 2. It is obvious from Fig. 2 that data from a single point (here: t = 28 and x = 100) would not allow prediction of the behavior at any other time unless the law is specified.

5.4.4 Kinetic Breaks

The basic laws of Section 5.4.3 are for constant growth under constant conditions. However, a scale growing according to the parabolic law may change to a different behavior only after some time has elapsed.

There are a number of reasons for such kinetic breaks, shown in Fig. 3 . Again the most easily envisaged case is a crystallizing scale. Crystallization is often accompanied by a volume expansion, which may induce stresses in the scale and lead to spalling. The renewed fast access of the medium to the substrate accelerates reaction rates. If spalling is connected to a critical thickness, we find pseudo-linear growth built from individual parabolic segments (Fig. 3) . If the first break is not repaired by the following growth, for example if the scale is not spalling but permanently cracks, we find break-away behavior, where parabolic behavior changes to linear.

We may also find breaks which slow down corrosion, such as if a scale crystallizes only after some critical thickness or time is reached without disintegration. A parabolic-logarithmic behavior may follow. Note that this break is not apparent to the naked eye unless very long times of observation are involved.

Figure 3. Common kinetic break types.


*Oo0 I n 1500-

cu m

=! Y

Y P 1000- 0 .- 5

- 5 500-

0 20 40 60 80 100 1

Time (a.u.) 0

Figure 4. Growth change due to developing diffusion coefficients, D.

5.4.5 Complex Kinetics

The real corrosion behavior may be far more complex then the ideal laws suggest, because of changes in properties of the scale with growth and the operation of differing types of corrosion at the same time.

Changes in the properties may be microstuctural changes like the development of bubbles, pits, or internal porosity, where effective surfaces or effective diffusion path lengths are changed. Other common changes are those of the effective diffusion coefficient. If a scale is not built from a phase of constant composition like the SiOz formed on Si, but contains elements from external impurities or internal secondary phases, which are concentrated or diluted with time, the diffusion coefficient will change with these changes. Accordingly the overall behavior must deviate from parabolic behavior.

Such changes are illustrated in Fig. 4. Note that the increase of D with time makes the shape more linear with time, while the decrease in D brings it to a curve more like the logarithmic law.

The development of scales on highly porous media has been reported to show a typical break of growth speed with increasing temperature. The model case shown in Fig. 5 involves indeed no true kinetic break, the parabolic rate constant increases regularly with temperature and is identical for short and long times. The accepted reason for the shape of the 'high-T' curve in Fig. 5 is the pore closure with the associated breakdown of the effective surface. A typical factor is of the order of 300 if a sample with a geometric surface of 1 cm2 has a specific surface of 0.1 m2/g-' in the porous state.

A relatively simple analytical treatment for complex growth patterns has been suggested by Nickel [21]. It works on the assumption that the basic laws of


30

Time (a.u.)

Figure 5. Typical development of layer growth with increasing temperature in porous media with pore closure due to passive corrosion.

growth can operate simultaneously. Growth according to

Ax = k,t + + klog Iog(t)

is evaluated by stepwise multiple linear regression. The stepwise procedure was adopted to give fit constants with more physical meaning. The fit allows quantification of individual contributions of the basic laws and is illustrated in Fig. 6.

Other attempts to model complex kinetics have been put forward by Ka11 et al. [22]. Both complex modeling methods do not incorporate true kinetic breaks and thus need improvement in order to be used as a general growth law.

0 500 I000 1500

Time (a.u.)

Figure 6. Model for complex growths kinetics after Eq. (1 1) [21] with individual contributions.


6

- 4 ? m Y

% 2 c m c 0 0 E

3 -2

0) a .-

-4

0 20 40 60 80

Time (a.u.)

Figure 7. Weight change curves obtained from the simultaneous action of parabolic weight gain and two different linear weight loss rates (‘para-linear behavior’).

An important special case of complex kinetics is the simultaneous action of basic passive laws together with active corrosion. A typical case for this is when a scale grows but is consumed at the same time by evaporation. The resulting shape of curves is shown in Fig. 7 and has been described as ‘para-linear’ [23] behavior. It may be analyzed with Eq. (12) omitting the logarithmic term.

It should be noted that small variations in the linear loss rate constant give macroscopically different behavior: the curve labeled ‘data 1’ in Fig. 7 might wrongly be interpreted as ‘hardly any change after a short period and ‘data 2’ as ‘active corrosion after an incubation period’.

5.5 Corrosion Measurement

5.5.1 Experimental Methods

The most common type of corrosion experiment is the after-treatment test, that is investigations on property changes are done after the end of the test by any of a variety of means (weight, dimensions, penetration, phase change, retained strength etc.). The limitations are discussed in the section on corrosion data (5.5.2).

For the corrosion of ceramics by ordinary liquid media, the testing is usually done by immersion tests: the sample is placed in a heated retort or autoclave inside a stirred excess of corrosive medium. Common test conditions are at the boiling point of the medium (e.g. 10% H2S04, 30% NaOH) for a week.

The corrosion by melts is classically studied by the methods of the refractory industry like drop-, crucible-, and finger tests [24]. While those methods allow an

5.5 Corrosion Measurement 151

easy ranking they are hardly capable of producing quantitative data outside the narrow conditions set by the standardized experimental set-up [25].

For hot corrosion testing burner rigs can be applied [26]. The agent is injected into the flame of the burner and the sample is held or agitated at a distance, which allows deposition of the corrosive agent from the gas phase at a constant rate.

In the ordinary furnace test for gas corrosion the samples are placed in a furnace for a given time in the atmosphere of interest. The advantage of the furnace test is that many different samples can be tested in one set of conditions, the drawback that it gives only one data point in time per experiment. The pitfalls are apparatus related: Temperature and atmosphere constancy have to be guaranteed, cross- contamination must be avoided, and interaction with furnace linings and crucible material minimized. The danger of those points is often underestimated; there are reports of different results from a mere change in the furnace size [27].

Contamination by furnace linings is a source of error for all corrosion experiments. There is ample evidence that the commonly used linings of A1203-ceramics do contaminate samples with A1 and alkalis via gas phase transport, even if high- purity A1,03 is used. The contamination is strongly enhanced in wet atmospheres

In situ measuring of corrosion is mostly done by thermogravimetric measurement, that is the monitoring of a weight change during a isothermal period of time. The advantage here is that the true shape of changes is recorded, which allows proper kinetic modeling. The limitations of the method come firstly from the apparatus: Resolution, accuracy, precision, and the often unchecked long- term stability are major points of concern. The pitfalls of contamination by furnace interior and crucibles, and local atmosphere changes are identical to those mentioned above.

For coarse estimations of danger levels in corrosion nonisothermal thermogravimetry or differential thermal analysis (DTA) can be used.

[281.

5.5.2 Corrosion Data

From the descriptions above the reader will accept that there are plenty of problems of defining data and parameters for the characterization and comparison of corrosion resistance. Standardization is under discussion [29], but is not generally accepted. Only for very specific corrosion cases are the first norms and recommen- dations available [30-341.

This is the reason why even excellent data collections on advanced ceramics [35], become very fuzzy when corrosion is concerned. We will discuss here briefly the different data types.

For a general case we first look for the changes in dimensions (Fig. 8). There are three principal changes: a thickness of a scale growing on the substrate, a change in the overall size of the component, and penetration depths. The most interesting parameter from the point-of-view of an engineer would be the penetration depth, because it would allow evaluation of the effective cross section of the component after corrosive attack and hence prediction of the remaining strength.


layer thickness

corroded size

Figure 8. Principal dimensional changes during corrosive attack.

For homogenous surface attack by active corrosion the case simplifies to the recession as there is no scale growing and the remaining cross section is of virgin material.

The parameter of interest is then the recession rate, which is identical to kl of Eq. (6) for linear kinetics. In corrosion science the SI-units ms-' or the more vivid unit mmyear-' of engineering practice are useful. The relation to mass loss is obtained by multiplying by the density of the lost material to yield kgmp2 s-' or equivalent units.

It should be emphasized again that this parameter is useless unless the physical boundary conditions are specified. The quotation of partial pressures of outgasing species is often more meaningful.

A problem arises when the attack is localized, because the attack is then different for different phases and the outer surface is no longer identical to the virgin material. The evaluation of penetration depth and rate may be performed in some cases using dye tests. If penetration goes with formation of porosity, dyes like methylene-blue or phosphorescing agents can be applied to show the depth of corrosion in cross sections.

More complex situations, such as with chemical modification of the grain- boundary phase, may be measured in optical thin sections, but they need usually more careful evaluation by chemical profiling with microanalytical techniques. Note that the kinetics of active corrosion may then deviate strongly from linear kinetics because leaching by liquid media or evaporation out of channels often involves diffusional problems.

A k1 is then not the parameter for ranking. A better choice would be the reference to critical effects, such as the 'time needed for 0.1 mm penetration at 0.7 T,' [36]. Regrettably, such parameters are extremely scarce in the literature.

As the experimental determination and the investigations afterwards are highly time-consuming and expensive, an often followed method of characterization is by the effect of corrosion on mechanical properties, usually strength or wear properties [37,38], in addition to thermogravimetric data.

5.5 Corrosion Measurement 153

\

I

Changes (a.u.)

\ \

0.0 -

\

Moving to the more complex cases with scale growth, Fig. 9 shows the range of kinetic curves one may find. Figure 9 shows a parametric study for a single case of para-linear behavior. Not even the sign of the curves is constant. We must therefore discuss the individual contributions.

For linear scale growth the discussion above is valid. For processes of retarded scale growth with time we describe the layer growth by k, (units of m2 s-l or equiva- lents) for simple parabolic behavior (Eq. (8)).

The calculation of the appropriate mass change unit is not as straightforward as in the linear case. This is because the mass change recorded by weighing or thermo- balance is measuring the difference in masses and not the mass of the layer.

The latter may be inferred if the chemical reaction and the properties of the scale material are known, introducing a stoichiometric factor u. As an example the reaction

Sic + 1 4 0 2 = SiO2 + CO (13)

means that I mole of the condensed product Si02 (= 1Msio2 = 60.09g with a density psio2 of 2 . 2 g ~ m - ~ in the amorphous state) is produced from 1 mole of Sic (= 40.09 g). Thus a recorded mass change of Ax = nMproduct - nkfmaterial

(= 19.99g) in Eq. (12) represents the production of a full mole of SO2. In general it follows that the factor u is defined from the reaction equation as

(14) (nM)condensed product

(nM)condensed product - (nM)material ' U =

which is 3 for reaction (13). Because of the form of Eq. (8) the two k, are hence related by

3 [ y 1 Pproduct 2 [ $1 k,(layer-thickness) - = kp(mass-change)

U 2


Neglecting the factor u has lead to occasional inconsistencies in the literature. The discussion of it seems trivial but a close look at Eq. (14) shows that for other cases, for example

SiBCN3 + 2.25 O2 H SO2 + 4 B203 + CO + 1.5 N2, (16)

u can become very large with the result that substantial scale production may occur almost invisible to thermogravimetry.

The layer thickness derived from any of the kinetic equations may be recalculated to give a solution for homogeneous recession:

From this discussion one might conclude that layer thickness measurements are the better choice for monitoring passive corrosion. While this is true from the point of resolution (ellipsometry allows the determination of scales down to the nm scale) it is difficult for high-temperature in situ recording and it implies dense transparent layers, which may not be the case.

Thus a safe characterization will ask for both quantities to be measured indepen- dently. Inconsistencies between measured and calculated values from Eq. (1 5) will then immediately point to problems concerning scale structure and/or internal corrosion.

Additional recording of mechanical property changes has the advantage of including information on potentially dangerous new flaw populations derived by corrosion. The overlay of several parameters (flaw geometry, cross-section changes, stress field changes) will not allow us to use mechanical testing as the only criterion.

Thus, at the present time we prefer to compare and rank materials in terms of time needed for penetration to a certain depth at the temperature of interest to make any type of kinetic equation accessible. The fixed ‘critical condition’ is taken to be 0.1 mm, following Ashby and Jones [36], the time unit is in years to give a measure for engineering practice. An engineer may then be able to decide that a material is not useful for long-term applications if it drops to log t, < 0 at the temperature of interest where tc is the critical time in years for a penetration of 0.1 mm. In the following sections on the materials this is calculated but the uncertainties have to be kept in mind.

5.6 Materials

5.6.1 Diamond and Diamond-like Carbons

Diamond is known to resist all acids and most other chemical substances in the low-temperature range to 100°C. However, transformation of diamond into other carbon phases may occur under local contact stress at room temperature [39,401

5.6 Materiuls 155

Diamond will be etched by caustic alkalis and some oxidizing salts, for example by NaClO and KOCl at 380°C [41], and K and Na nitrates above 400°C. Diamond reacts with metals, which form carbides (e.g. W, Ta, Ti, and Zr) or which dissolve carbon (e.g. Fe, Co, Mn, Ni, and Cr) [42]. Detailed knowledge of the temperatures of reactions with metals is reviewed in the context of experiments with metals for the catalysis of diamond synthesis [43].

Diamond corrosion is usually a problem of oxidation and graphitization. Graphi- tization in inert atmospheres was observed at about 1500°C and becomes very fast at 2100°, where sizable diamonds are converted to graphite within minutes [44]. An extrapolation of the data gives a T M 1450°C for t, = 0. It has been reported that graphitization is prevented in pure H2 up to 2200°C [45].

Surface graphitization associated with concurrent oxidation, which may be found at T > 600°C, is not thought to be a true conversion but an autocatalytic surface reaction involving CO, species.

The oxidation of diamond is clearly an active corrosion process. At least up to 700°C diamond has a fast reacting { 11 1)-plane, an intermediate { 1 lo} and a slow { 100)-plane, which indicates reaction control. At higher temperatures and/or lower oxygen pressures gas diffusion becomes rate determining in analogy with graphite [46],and this is indicated by a more even attack [47]. Hence corrosion rates are faster or start at lower temperatures for fine powders compared to films and the corrosion in air is faster than in low-oxygen environments [48].

At 720°C diamond will burn in an oxygen jet. It is often stated in the literature that oxidation starts at 600°C. The theoretical onset of oxidation should be well below 600"C, because from the etching rates of Evans [44] for the { 11 1)-plane in 1 bar oxygen a t, well below zero ( ~ 0 . 0 2 ) is calculated. Linear extrapolation would give T M 340°C for tc = 0. However, in analogy with graphite there may be an accumulation of C-0 complexes at the surface, which is faster then their removal between 300°C and 500°C [49]. Surface graphitization at T > 600°C points in this direction.

Also in analogy with other carbon forms, impurities may act as catalysts for the oxidation, which can decrease oxidation temperatures by more then 100°C and increase rates by orders of magnitude [50].

To a first approximation the temperature sensitivity of the oxidation rate at T < 800°C is less dependent on the structure of diamond (films) then on crystallo- graphic orientation, the data for the diamond faces bracket those of thin layers [47].

As described in more detail in the contribution on hydrothermal synthesis of diamond in this book, hydrothermal reactions start for fine powders at 7OO0C, but a significant etching and dissolution of large diamond grains is observed only at T 2 800°C [51].

5.6.2 Carbides

The carbides considered as hard materials include B4C, WC, and Sic. In oxidizing environments, in particular, their behavior is so different that they can not be compared directly.


5.6.2.1 Silicon Carbide

Compilations of corrosion data on Sic-based materials have been given in some hand- and textbooks [2,35,52,53], many papers on it are included in the proceedings of specific meetings [3,5,54].

5.6.2.1.1 Corrosion by Liquid Media There is only limited information about the behavior of pure Sic in low-temperature water, acids, and bases. A relatively pure sintered Sic is applied in many pumping systems for the chemical industry [%].This provides evidence that it is hardly affected by chemical attack.

Significant attacks can occur via a grain-boundary phase in other Sic types, and this is usually accompanied by a decrease in strength and Weibull modulus. Thus it is clear that liquid-phase sintered Sic will show less corrosion resistance with a behavior depending on the chemistry of the additives. For Y/Al-sintered Sic significant corrosion was reported for a treatment in boiling 20% HC1 and 50% NaOH; concentrated acids (HZSO4, H3P04) and bases (KOH) had much less effect [56].

Sic with free Si (Si-Sic) is affected in particular by hot strong bases (10M NaOH) and hot diluted acids (H2SO4 at pH = 3) [57].

5.6.2.1.2 Hydrothermal Corrosion In contrast to the good behavior of Sic at lower temperatures a strong attack can be obtained in supercritical water. At pressures of 100 MPa there is some indication from powder experiments that initial reactions start at 300°C [58, 591, strong corrosion is observed at T > 500°C [60,61]. The lower temperature onset is supported by tribological studies [62].

The main mechanism is the formation of silica from Sic along with CH4, COX f C. The silica is then dissolved in HzO. The dissolution rate of silica will play a vital role in the kinetics of the process. Basically the attack should have active corrosion character (Eq. (6)).

Chemical modification of the water to increase the pH will enhance the dissolution of silica [63] and hence further increase corrosion of Sic.

5.6.2.1.3 Corrosion by Metals Sic can be attacked by a number of metals, because often silicides and in some cases other carbides are stable phases and eutectic temperatures are low. From [52] we have evidence of reaction with Bi (600"C), Li (815"C), Mg (8OO0C), A1 (750"C), Ti (575"C), Ta (12OO0C), Nb (1300"C), Cr (lOOO"C), Mo (1200"C), W (1500"C), Ni (500"C), Co (1150"C), Fe (lOOO"C), Cu (950"C), Pd (600"C), and Au (950°C). Resistance to attack was noted for Na (350"C), Ca (1 180"C), Cd (500"C), Sn (600"C), Pb (815"C), and Ag (962°C).

These data have to be treated with some caution, because many results are obtained from treatment in vacuum without a protective scale present. The kinetics of the reaction have been described as parabolic for a number of cases. However, they depend also on physical conditions like the wetting behavior of the liquid metal or a formed melt [64].

Thus, it is obvious that Sic with free Si will react more readily than Sic. Tempera- tures of possible reactions may be evaluated with the aid of phase diagrams for the appropriate Si-metal system.

5.6 Materials 157

From own experience we know that Pt can be used as crucible material in oxidation studies of Sic. However, in vacuum the formation of melt and silicides at the eutectic of the system (= 1400°C) is so fast that individual Pt grains drill deep holes into the material within seconds.

The use of Sic as a hard material in wear application, can thus be endangered by a reaction with the metal if the silica scale is eroded by the process and contact temperatures are high enough.

5.6.2.1.4 Corrosion by Salt Melts (Hot Corrosion) Data and reviews of the mechanisms of the hot corrosion behavior of Sic have been presented in a number of papers by Jacobson and coworkers [65-681.

A representative member of the hot corrosion agents is Na2S04. Its generation is explained in combustion environments by

2 NaCl + SO2 + 0.5 O2 + H 2 0 H Na2S04 + 2 HCl. (18)

Since the action in the liquid state is regarded as most effective, the first condition for the window of hot corrosion is given by the melting and dewpoint of Na2S04. The melting point of 884°C does not vary much, the dewpoint depends on pressure, S-content of the fuel and NaC1-concentration in the atmosphere. For conditions with NaCl > 1 p.p.m. and combustion pressures of 100 bar it is calculated to be 1100-1200°C from Eq. (18).

However, the effective corrosion action is generally believed not to be by the sulfate itself. It stems from the activity of Na20 arising from the decomposition

Na2S04 * N a 2 0 + SO3

x Si02 + Na20 H N a 2 0 - (SiO2).,.

(19)

(20)

because only the Na20 is attacking the protective scale of Si02 by

This is the reason why other compounds releasing alkali oxide on decomposition, such as Na2C03, have similar hot corrosion effects. It should be noted that complex salts, in particular those containing vanadium, can have an even stronger hot corrosion effect, as is well known from metals and oxides [69,70].

From Eq. (19) it can be calculated that only for conditions with a low P(SO3) will we have a dissociation sufficient to drive reaction (20) to the right side. Consequently the window of hot corrosion has a third axis, the P ( S 0 3 ) of the combustion, and it has been computed [71] that the window should only open for low-S fuels (<0.5%, a condition met by many modern fuels) and at comparatively high temperatures (FZ>IOOO"C).

In reality the attack already starts at the melting point of Na2S04. This is attributed to internal controlling factors. Many Sic-based materials contain free carbon, which promotes the dissociation of Na2S04 and makes the silicate melt more basic. The dramatic attack of Na2S04 on Sic with free carbon has been confirmed experimentally [72].

As is described in more detail for Si3N4, recent studies suggest that free carbon is an enhancing factor but not a necessary condition [73,74]. S ic may act as the reducing agent to promote sulfate dissociation either directly or via an oxygen gradient in the silica scale consumed by reaction (20).


It is very difficult to quantify hot corrosion in laboratory experiments with pre- loaded samples, because the salt melts propagate rapidly both by spreading and gas phase transport, resulting in uneven coverage of the sample and changes with time.

A much better constancy of conditions is achieved in burner-rig methods [75-781. However, even here a quantification is difficult as the usually highly porous scales containing sulfate, silicate and silica are liable to spalling. Furthermore, the attack is often in the form of pitting, and so there is a change in exposed surface and the attack is nonuniform. The statistical nature of pit distribution is probably not due to a distribution of secondary phases but the result of the action of surface tension effects, because two immiscible liquids (sulfate and silicate) with differing surface stresses are present [74].

Therefore, while weight changes of chemically cleaned corroded samples are good indications for the strength of attack, they remain unsatisfactory and a widely used criterion is by the influence of corrosion on mechanical properties. Even for cases with good control of chemical changes, a prediction of the effect on strength can not follow directly, because the fractures occur usually from the newly formed pits (Fig. 10 [78]).

None the less, as seen in Fig. 10, the relation between pit size and failure is evident for most Sic materials. A model for the prediction of properties from corrosion conditions would thus have to give an expression for the formation and growth of pits with time, temperature, and external properties. The author is not aware of a fully consistent model to yield this result.

Thus we are left with individual studies. A study on commercial Sic [78] yielded the typical result that a reduction in strength of 20-30% is obtained after 4 0 h exposure at 1000°C.

(ac)"', pm

500 1 1 200 100 50

2 400 z d

E 300 5

c ; 200 3 0 m c

t 100

0.75 K,c -- - ( a ~ ) " ~

0 Type A 0 Type B A Type C 0 Type D

0.1 0.2 0.3 (Pit dimensions)-'", (ac)-'", pm-"'

Figure 10. Relation of fracture stress and pit dimension from a study of four different Sic materials 1781.

5.6 Materials 159

Figure 10 also shows that the relation with pit size is not true for all Sic-materials: Type D shows no correlation and has a much stronger decrease in strength (= -50%). It was a material with free silicon. It follows that Si-Sic has less hot corrosion resistance and that the attack and its consequences are due to the distribution of the preferentially corroded Si.

To improve the hot corrosion resistance current developments aim for a protection by mullite-based coatings [79,80].

5.6.2.1.5 Corrosion by Hot Gases Halogens do not form condensed species with Si at the temperatures of interest. Thus, if they are pure (no oxygen or water present) they will automatically induce active corrosion behavior and calculated pressures of species like SiF4 or SiCI4 over Sic will be as high as the concentration of F and C1 in the gas.

In the presence of enough oxygen to form silica, these pressures will be significantly reduced for the case of attack by C12 and accordingly lower corrosion rates are measured [81]. However, pressures of SiC14 over Si02 can exceed lop6 bar at temperatures above 700°C in high-C12 environments.

Calculated pressures for an attack by H2 are likewise significant even at low temperatures [8]. The calculated main low-T species is CH4, but the kinetics are so unfavorable for its formation that in reality the beginning of significant active corrosion is at temperatures above 13OO0C, as is known from etching studies [82]. This is only true for pure Sic: grain boundaries and secondary phases in sintered S ic are attacked at temperatures as low as 1000°C [83].

For the survival of S ic it is therefore of great importance whether silica is formed or not. The boundary for the active-to-passive oxidation is described in Fig. 1 1. It shows the literature data and a Wagner-type theoretical model [15] based on the

H i m 76

Gulbranran 66

Vaughn so *

-8 I 1 I 1

1,200 1,400 1,600 1,800 2,000

Temperature ("C)

Figure 11. Transition temperatures and oxygen pressures for active and passive modes of oxidation of Sic after [15].


influence of carbon activity and the presence of condensed SiO. Even though the model might be challenged in view of new thermodynamic data, arguing against condensed SiO, the general position of the boundary between active (I) and passive oxidation is known from the experiments.

The region denoted active (11) in Fig. 11 is for a special situation, when the pressures of the produced gas species (CO) exceed 1 bar, which can lead to bubble formation and spalling, and hence quasi-active behavior, even though nominally silica is produced. Experimental evidence for this was obtained at temperatures of about 1760°C [84].

It should be noted that the oxygen pressures for the transition in Fig. 1 I should not be confused with those obtained by COX-equilibria. Despite low oxygen partial pressures of, for example CO in equilibrium with carbon, we have passive oxidation of Sic up to x 1400°C in CO, because CO is a reactive and oxidizing gas species for Sic [52,85,86]! Consequently active corrosion in COX-environments is measured only at very high temperatures [87].

In passive oxidizing environments Sic forms silica (Eq. (13)). For pure Sic a parabolic behavior is evident [88]. For SSiC there is evidence that oxygen does not penetrate deeply into the interior [89], leading to a relatively homogeneous surface oxidation.

Therefore SSiC has often parabolic oxidation rate constants nearly as good as chemical vapor deposition (CVD)-Sic and powder. From a data collection [53] plus some more recent data [88] the critical time t, for 0.1 mm recession is computed and shown in Fig. 12. We have omitted data from experimental changes, which were deliberately made with high additive contents to study their influence on oxidation.

It is evident from Fig. 12 that basically all cases meet a long-term stability criterion up to very high temperatures.

Some cautionary remarks have to be added: At very high temperatures the kinetics may not be adequately described by the parabolic law; at T > 1700°C the problem of active (11) oxidation will become important; SSiC can show bubble and pit formation at much lower temperatures, which is usually attributed to the additives B and/or A1 [90].

This bubble-and-pit formation is sometimes only present after a ‘incubation period’ of several lOOh at T > 1600°C. Most oxidation studies run only for times < 100 h. Therefore long-term studies may be necessary for true qualification of Sic.

Industrial utility of Sic is not guaranteed by a sufficient resistance to oxidation by O2 or air. The influence of other components of the atmosphere or transferred impurities has to be evaluated. Often this is difficult because strong variations in temperature and gas composition are encountered.

None the less it has been shown that SSiC can resist flue gases of aluminum remelting furnaces for up to 1000 h without much change in strength [91], although this is certainly not true for all plants of this type [92]. Similarly nitrogen-based gases of the system N2-H2-C0 did not decrease the strength of SSiC, while endothermic gases had an adverse effect (-20-50%) at 1200-1300°C [93,94] in long-term experiments. Typically B-doped Sic behaved better then Al-doped or those with free Si. A good resistance of Sic-based materials to SO2 or H2S has also been reported [95].

5.6 Materials 161

A

I " 4 *

* *

*

* ssic A powder&SC A CVD

800 1000 1200 1400 1600 1800 900 1100 1300 1500 1700

Temperature ("C) Figure 12. Calculated critical times (years) for 0.1 mm recession of S ic materials, using the data compilation of [53].

The influence of water on the corrosion of Sic seems to be threefold: It enhances the oxidation rate slightly; it transports impurities (e.g. Al and alkalis) which can strongly enhance oxidation; and it is able to evaporize the silica scale inducing para-linear behavior under appropriate flow conditions [23]. This may impair some applications.

5.6.2.2 Boron Carbide

Boron carbide is among the hardest materials yielding only to diamond and boron nitride. It is also one of the most corrosion-resistant compounds at room or moderate temperatures. When considering the corrosion resistance of boron carbide materials, it is important to remember that they are rarely stoichiometric, with the carbon content varying from 9.88 to 23.4% [96] Many of them contain free carbon or sintering aids. Thus their behavior depends on the chemical composition.

5.6.2.2.1 Corrosion by Liquid Media Pure boron carbide is insoluble in HCl, H2S04, and HN03, even when boiling. Resistance of hot-pressed boron carbide to sulfuric acid solution has been investigated [97]. The material was produced from a powder containing 77.8-79% boron, 21.6-20.4% carbon, 0.25-0.7Y0


boron oxide, and 0.25-0.7% free boron. Dense and additive-free B4C specimens did not interact with 10% sulfuric acid either at room temperature or on boiling. A small mass loss (<O. 1 mg cm-2) in the initial period of boiling can be explained by surface impurities. As has been shown [98], the corrosion resistance of boron carbide powder in water, mineral acids and alkalis is largely dependent on its purity and the ordering of the crystalline structure. The studies on the electrochemical behavior of hot-pressed boron carbide in H2SO4 and NaOH solutions [99] have demonstrated that the corrosion resistance of B4C depends to a large extent on the content of additives and impurities.

However, during heating, strong oxidizing acids and oxidant mixtures such as Cr03 + H2SO4, KIO3 + H3P04, K2Cr207 + KIO3, Cr03 + H2SO4 + H3P04, KMn04 + H2SO4, H2SO4 + HC104 + K2Cr207 [98] oxidize free carbon present in many boron carbide materials. To remove free carbon from B4C, alkaline solutions of hydrogen peroxide, bromine, and others were also used [98].

5.6.2.2.2 Hydrothermal Corrosion and Corrosion in Water Vapor The interactions of boron carbide with water vapor starts at 250°C. The following reactions are possible:

B4C + 8 H20 + 2 B203 + C02 + 8 H2,

B2O3 + H20 + 2 HBO2,

B203 + 3 H20 + 2 H3B03.

(21)

(22)

(23)

and

Gaseous boric acid removes a boron oxide film. The rates of formation and removal of the B2O3 film are equal at 550-600°C in air with a dew point of 25-70°C and at 650°C with a dew point of 88°C. At higher temperatures, B203 is formed at a higher rate than it is removed by the interaction with water vapor. Therefore, at low temperatures boron carbide is oxidized with water vapor more rapidly than with dry air, at high temperatures the situation is quite the opposite [2].

Hydrothermal corrosion of B4C has been studied theoretically and experimentally [loo]. Reaction (21) as well as reactions leading to the formation of CH4 and CO are possible. Reactions that lead to the formation of carbon oxides dominate only at low pressures and carbide:water ratios. Under all other conditions, the formation of methane has been predicted. HB02, H3B03 and H3B306 gas molecules, which were predicted by thermodynamic simulation under various temperatures and pressures, are very stable thermodynamically. They may condense to some type of boric acid on cooling.

Unlike Sic, boron carbide did not show any evidence of free carbon formation under hydrothermal conditions. Neither thermodynamic simulation, nor experiments show the formation of elemental carbon under hydrothermal conditions. Even the traces of carbon that were present in the as-received B4C powder disappeared after a hydrothermal treatment.

Tests of a hot-pressed B4C-lO%SiAlON composite demonstrated high corrosion resistance up to 500°C. Even at 500"C, almost no changes of the sample surface occurred under the experimental conditions.

5.6 Materials 163

5.6.2.2.3 Corrosion by Melts Hot-pressed boron carbide is attacked by fused alkalis. The use of boron-carbide-based materials at high temperatures is limited at present, which probably results from the fact that they have been investigated primarily at room or moderate temperature [ 1011.

5.6.2.2.4 Corrosion by Hot Gases Boron carbide is mainly used at relatively low temperatures [96], therefore its oxidation has not been studied as thoroughly as that of the Si3N4 and Sic ceramics. There are data on the oxidation of B4C powders [lo21 and several papers on oxidation of boron carbide materials produced by different methods [103-1071.

In [ 1081 different reactions in the system B4C-02 were calculated thermodynamically at oxygen pressures of 105 and 1.3 x lop3 kPa. It was shown that the oxidation of B4C may result both in condensed (B203, B, C) and gaseous (CO, COz, B202, BO, BOz) products. At high oxygen partial pressures the reaction

B4C + 402 ++ 2B203 + C02 (24)

should dominate. This reaction was confirmed by the results of experimental investigations on powders and sintered specimens [ 1091.

Oxidation of boron carbide powders starts at about 600°C. Studies on hot-pressed B4C containing iron, aluminum, magnesium, calcium, titanium, silicon, and other impurities at a level of up to 4% and oxidized in air [lo41 demonstrated that upon nonisothermal heating of specimens exothermal effects due to oxidation can be registered above 600°C. A detectable mass gain starts above 700°C. At 1000°C the oxidation process intensifies and above 1200°C the material starts oxidizing considerably, with a high thermal effect. At temperatures above 12OO0C, the specimen weight decreases due to the B203 vaporization, the rate of which increases significantly at this temperature [ 1091. The X-ray diffraction (XRD) analysis of oxidized specimens did not reveal any solid oxidation products other than B203. This is in agreement with the thermodynamic calculations [98] and the experimental investigations showing that reaction (24) dominates B4C oxidation.

The kinetic curves obtained under isothermal conditions (Fig. 13) follow the paralinear law (Fig. 7) with competing B4C oxidation and B 2 0 3 vaporization.

3, 1

2 N . 6 1

F o @ E -10

-20 < 1 2 3 4 5 -300

t, h

Figure 13. Oxidation kinetics of hot-pressed B4C at: (1) 800°C; (2) 900°C; (3) 1000°C; (4) 1100°C; and (5) 1250°C; reproduced from [2].


Thus, the oxidation kinetics due to the boron oxide vaporization at temperatures above 1200°C result in weight loss of specimens (Fig. 13), which becomes linear in time after the first 30min of oxidation. On oxidation at temperatures below 1100°C a liquid boron oxide layer (T, of amorphous B2O3 is 450°C [llO]) uniformly covers the specimen surface.

Under the conditions of active B203 vaporization the diffusion of impurities apparently does not exert great influence on the oxidation process, unlike in the case of oxidation of silica formers. One may suppose that up to 1200°C the oxidation rate of B4C is controlled by the oxygen diffusion through the B203 layer.

Differences in the thermal expansion coefficients of B203 and B4C lead to cracking of the oxide layer after cooling from high temperatures. The XRD analysis revealed B203 and H3B03 in the surface layer of oxidized specimens. However, H3B03 is not formed during the oxidation of materials, but during the exposure of specimens to air at room temperature for several days due to the hydration of boron oxide according to reaction (23). The oxidized specimens kept in air were covered with a white film due to the hydration of boron oxide.

It is important to note that the specimens of ceramics oxidized at 1400°C had a graphite layer of up to 30 pm thickness under the B203 film [107]. If the oxidation process was accompanied by friction, the graphitized layer was formed at much lower temperatures. The formation of free carbon was also observed upon oxidation of B4C-Sic composites [l 1 11.

Active oxidation of B4C at a low partial oxygen pressure leads to the formation of gaseous boron oxides and removal of them from the surface of the specimen. Similar to other materials, porosity increases the reactive surface of boron carbide specimens and weight gain on oxidation [ 1061.

All the above primarily refers to materials containing relatively small amounts of impurities. However, boron carbide ceramics usually contain 7-20% Al or A1203 to increase hardness and 6-1 3% silicon to improve mechanical characteristics [98]. Aluminum- and silicon-doped materials often exhibit structural nonuniformity. They have areas of pure boron carbide, and silicon- and aluminum-doped areas [107].

In contrast to the oxidation of the additive-free material, weight gain instead of its loss is observed above 1200"C, that is the rate of oxide formation is higher than the rate of its vaporization. Electron microscopy investigations of specimen surfaces of Si- and Al-doped materials heated to 1430°C demonstrated glassy oxide layers containing boron, oxygen, silicon, and aluminum. Thus, the oxide layer is a borosilicate glass. Crystallization of aluminum borates forming needle-like crystals also has been found in our work. The dissolution of silica and alumina in boron oxide retards the rate of its vaporization. In general, addition of silicon and aluminum to boron carbide facilitates the formation of borosilicate glass on oxidation and decreases the rate of boron oxide vaporization, thus increasing the oxidation resistance of the material [ l l 11 [2]. However, to form a continuous protective layer on the specimen surface, it is necessary to provide a uniform distribution of additives in the material excluding the existence of pure boron carbide areas.

5.6 Materials 165

B4C reacts with C 0 2 to yield B203 and CO or free carbon [96]. Boron carbide neither interacts with sulfur and phosphorus vapors, nor with nitrogen up to 1200°C. BN can be formed upon reaction with nitrogen at higher temperatures, or when ammonia is added. With chlorine it reacts above 1000°C to form BC13 and graphite. Bromine and iodine do not react with B4C [98].

5.6.2.1 Tungsten Carbide

Hard materials on the basis of WC have usually a metallic binder phase of Co and/or Cr, Fe, Ni, and additions of other carbides (Tic, TaC, NbC), which deter- mine the corrosion resistance. Accordingly, these composites are attacked by many acids and bases. The authors refer the reader to previous work [112,113] for more details.

Tungsten carbide has a low resistance to oxidation. It is the result of low protective properties of tungsten oxides. Dissolution of oxygen in the carbide structure starts at low temperatures, and W 0 2 and W03 form upon oxidation at 500°C and higher. The oxidation rate increases drastically at 800°C due to vaporization of W 0 3 [108].

According to thermodynamic simulation [loo], at a carbide:water ratio of 1 : 10, tungsten carbide should be completely oxidized producing H2W04 or W 0 3 at lower temperatures and W 0 2 at higher temperatures. Carbon is oxidized to CO and C 0 2 in the case of a low WC:H20 molar ratio. Formation of free carbon upon hydrothermal corrosion of WC was predicted to be possible at a high WC:water ratio upon hydrothermal corrosion of WC. The stability of WC increases with increasing pressure.

Raman spectra obtained from the hydrothermally treated WC powders demon- strate the presence of graphite. Hydrothermal treatment of a WC-G%Ni commercial cemented carbide demonstrated a rather low corrosion rates below 500°C. Only a very thin layer of corrosion products covered the surface of the samples after a treatment at 350°C. At 500"C, the corrosion rate increased and all WC grains from the surface layer of the sample were dissolved in the supercritical fluid. This led to a slight decrease of the sample weight. Ni does not interact or interacts very slowly with water under these conditions. Therefore, a network of Ni grain-boundary layers remained on the surface. Ni grain-boundary phase serves probably as a diffusion barrier and protects the material from further corrosion. At 700°C, the corrosion rate increases and the surface of the sample was covered with a thick layer of corrosion products.

No changes in the phase composition of the samples were found by XRD after the treatment at 350°C. WC, Ni, and traces of W 0 3 were found after hydrothermal treatments in the temperature range 400-500°C. W03 dissolves in water and a weight loss of the samples was observed. At 700"C, oxidation of Ni occurs and nickel oxide interacts with W03 (or H2W04) leading to the formation of NiW04.

Also, hydrothermal corrosion of WC-Zr02 composites has been studied, but at lower temperatures when less significant interaction of WC with water occurred [114, 1151.


5.6.3 Nitrides 5.6.3.1 Silicon Nitride

In many cases the references quoted for Sic can be used to access corrosion data on Si3N4-based materials [2,3, 5, 35, 53, 54, 116, 1171. Si3N4 ceramics are much more complicated in their behavior because sintered types come with a significant amount of grain-boundary phase, which determines the overall behavior.

5.6.3.1.1 Corrosion by Liquid Media The higher susceptibility of Si3N4 ceramics to corrosion by liquid media is already apparent from reports on hydroabrasive wear, which is much higher then the wear of S ic [55].

It is not the attack of the matrix Si3N4, which controls aqueous corrosion but that of the grain-boundary phase. Often this grain-boundary phase is a oxidic silicate glass with or without small nitrogen contents. To predict the corrosion resistance in water, acids, and bases the reader is thus referred to the literature on glass and glass corrosion [118,119].

Rules of thumb would imply that a glass with more network-modifiers is less corrosion resistant to corrosion, materials with less glass content are better and that a greater crystallinity of the grain-boundary phase is of advantage.

The kinetics of the leaching processes are often linear for HF and strong bases, while selective acid etching can be parabolic. Leaching is always associated with a loss in strength and tends to correlate linearly with degree of dissolution from the glass matrix as has been shown for HF and HCl with Y - and Al-leaching [120]. At 70°C a treatment for 240 h with 1 M HCl can reduce the bending strength by z 50%.

5.6.3.1.2 Hydrothermal Corrosion The onset of reactions of Si3N4 powder with water at 100 MPa according to

Si3N4 + 6 H 2 0 H 3 Si02 + 4 NH3 (25)

starts at temperatures below 200°C [61], probably at about 130°C [121]. At 500°C the powder is completely oxidized within a few hours.

Reaction (25) would suggest a passivating behavior but the solubility of silica is favored even at low temperatures in the alkaline water present due to the dissolution of NH3. Hence basic reaction kinetics are linear, it is a form of active corrosion.

The grain-boundary phase of Si3N4-ceramics may be even stronger affected by hydrothermal corrosion. This leads to pit formation at conditions of 300°C and 8.6MPa [122]. The shape and depths of pit formation depend on the additive type. It is interesting to note that a hot isostatically pressed (HIP) Si3N4 ‘without sintering additives’ was attacked more strongly then those with additives, because the grain boundaries with silica from the oxygen content of the powder were leached and resulted in grain dislodgment [ 1231.

It should be noted that hydrothermal corrosion should not be confused with corrosion by wet atmospheres and that the behavior changes at higher temperatures when N2 + H2 is the product gas rather then NH3 [60].

5.6 Materials 167

5.6.3.1.3 Corrosion by Metals Reactions of Si3N4 with metals and alloys are a complex issue, because the atmosphere both for oxygen and nitrogen partial pressures plays a role and silicides, nitrides, or ternary compounds can form. A good compilation of the partly conflicting observations has been provided [ 1 161. For the interaction of the related SiAlONs and Fe-based alloys see published work [124].

Predictions for reactions can be made on the basis of Ellingham-type diagrams. In noble gas atmospheres or under vacuum, Si3N4 reacts readily with most transition elements below 1000°C and with many other metals at varying temperatures.

Whether this is true for a given Si3N4 ceramic under application conditions depends on many parameters like a silica layer protection, diffusion coefficients, wetting behavior etc. The reader is here referred to the reference mentioned above [116].

5.6.3.1.4 Corrosion by Salt Melts Hot corrosion of Si3N4 has been extensively studied and reviewed by researchers at NASA labs [67,68,77, 125, 1261. The basic attack is that on the protective silica layer, which makes the behavior similar to that of Sic and basic mechanisms can be taken from the chapter on S ic above.

Recent investigations [74] gave more insight into the details of the kinetics, which is shown in Fig. 14. The surface tension differences between sulfate and silicate melts result in a fractal pattern of hillocks of silicate melt beneath the sulfate, hereby exposing unaltered silica with a strong oxygen potential gradient to the sulfate. On consumption of this silica the decomposition of the sulfate is enhanced up to a gas bubble formation. The exposed areas are the sites for stronger corrosion, which can ultimately produce pits. However, the process does slow down, when the sulfate is used up or a critical thickness of silicate layer is reached.

There is a strong indication that the grain-boundary phases modify the development. Wide shallow pits and/or exposure of individual Si3N4-grains are observed [127,128], and Na advances into the grain-boundary phase.

The resulting degradation of mechanical properties is of similar order as for S ic (20-50%) [26,127,129]. However, it seems that the degradation is not linear with time but levels out with time [129]. Furthermore, in pitted samples there is an increase in the Weibull modulus, with the result that the failure stress at low failure probabilities is decreasing much less than the average strength [ 1301.

Taking into account the higher level of initial strength of Si3N4-based ceramics the remaining strength after hot corrosion is usually higher compared to SIC.

5.6.3.1.5 Corrosion by Hot Gases Active corrosion has usually an even more pronounced effect on strength, compared to Sic, as the phase between the grains may vanish leading to grain dislodgment and inhomogeneous internal corrosion [ 13 11. All gases which produce high partial pressures of Si-bearing species (namely halogens) are therefore dangerous, particular in low-oxygen environments [ 132- 1341.

In mixed and oxidizing atmospheres Si3N4 has a superior resistance to C1- containing environments relative to Sic materials [ 135-1 371. A small water content in H2S-H2 mixtures is also capable of producing a silica scale, which slows down


of the Na-sulfate

with amounts of

immiscible at 9OOOC

1- n Ess Na,SO, Na,O * @i02 3J SiO,

(watersoluble) (not watersoluble)

Figure 14. Schematic model for the hot corrosion kinetics of silica-formers after [74].

5.6 Materials 169

~

corrosion [ 1341. As for Sic the plain influence of a water vapor content in oxidizing environments is probably moderate [ 1381 unless it transports impurities.

The active(1)-to-passive transition has been discussed with some controversy but most models [14,139,140] and experimental data [141-1431 point to very similar conditions as outlined in Fig. 11 for Sic.

The passive oxidation of pure Si3N4 has been intensively studied. Thermodynamic data suggest that a two-step oxidation

2 Si3N4 + 1; O2 H 3 Si2N20 + N2

Si2N20 + 14 O2 + 2 SiOz + N2

(26)

(27)

and

and accordingly a duplex layer in a oxygen potential gradient, developing in a scale, is favored. However, the discussion of the correct kinetics, the presence or absence of a sublayer of Si2N20, the activation energy, and other issues are still not settled [14,144-1511. Extremely low oxidation rates up to high temperatures are, however, clear.

This contrasts strongly with the behavior of sintered Si3N4. In Fig. 15 we have calculated tc for pure and sintered qualities, assuming parabolic growth and using the data compiled earlier [53].

1 -0 .. 0 .

0. . 0 .

0

-1

- 2 1 ::'"""": , I , I 0 : , , , , , , HPISSN .

-3 800 1000 1200 1400 1600 1800

900 1100 1300 1500 1700

Temperature ("C) Figure 15. Calculated critical times (years) t , for 0.1 mm homogeneous recession by parabolic silica scale growth on Si3N4-based materials using data from the collection of [53].


A number of things can be learned from Fig. 15. First, sintered Si3N4 never reaches the level of oxidation resistance of pure Si3N4. Typically a temperature difference of 250°C exists to have comparable behavior for even the best qualities.

Second, at low to moderate temperatures pure Si3N4 oxidizes much slower then Sic (Fig. 12), but sintered Si3N4 is comparable to Sic. Third, there is a very wide scattering of results for sintered Si3N4. And last, but not least, at about 1400°C there is a rapid breakdown in log t, to values below and much below zero, indicating very rapid oxidation of sintered Si3N4.

The reasons for these features are manifold and again related to the presence of grain-boundary phases. We often have a very complex passive oxidation behavior with leaching of grain-boundary constituents to the scale, crystallization, internal oxidation, bubble formation, scale cracking, and other features. All these phenomena are different for different systems and vary with the amount and exact composition of the additives, For detailed information the reader is referred to the textbooks and proceedings already mentioned and some reviews, which cite papers from a vast amount of literature [140,152, 1531.

We have to point out some facts, which should make the reader cautious in inter- preting Fig. 14 too strictly. At low temperatures some Si3N4 types (sintered with pure Y or rare earth additives) contained secondary phases, which oxidize rapidly and with large volume increases to yield the oxide phase. Since at low temperatures (<lOOO°C) the outer scale becomes ridged, it leads to cracking of the specimen [154] and breakaway or pseudolinear behavior. This is part of the reason why there is such a large spread in data from Si3N4 ceramics even at low temperatures.

Particularly at T > 1 300"C, we commonly have deviations from the parabolic law and all the forms described in Sections 4.3.2 and 4.4.3 can be observed. Since the extrapolation to t, is done on the basis of a parabolic approximation it will not give very reliable t,.

Perhaps the most disturbing characteristic of sintered Si3N4 oxidation is the appearance of internal oxidation accompanied by outer scale growth. Leaching of grain boundary constituents into the scale (Fig. 16) is observed for Y, Al, Mg, REE, and others to varying degrees. The leaching is connected to internal oxidation and is for some cases the rate limiting step [156]. In the leached zone, which extends into the sample with the same order of magnitude as the outer scale thickness, we often find increased porosity and hence decreased hardness [155]. At least this zone must be incorporated into the value for penetration but can not be gained from measuring layer thickness or mass change.

There is furthermore recent evidence from transmission electron microscopy (TEM) studies for the dissolution of matrix grains in the grain-boundary phase during oxidation at high temperatures, which widens the grain boundary films as a function of depth into the material [157,158]. This type of penetration will be even more difficult to detect and quantify.

Therefore, for most Si3N4 ceramics it will be necessary to have mechanical property measurements along with corrosion to evaluate their applicability. None the less there are rules of thumb, which may help in materials selection. The most important one is that the level at which oxidative corrosion becomes the prime problem, is connected to the eutectic of the system sintering additive(s) + Si02.

5.6 Materials 171

Figure 16. EDX map for Y and SEM picture of SSN oxidized at 1400°C for 990 min [ 1551, showing a leaching zone beneath the scale.

This may well apply also to the pure Si3N4 qualities, because here the eutectic Si3N4-Si02 in the vicinity of 1600°C may be the ultimate temperature for long- term use [ 1591.

The challenge for sintered Si3N4 to be applicable at T > 1400°C has found first success in studies which show that, for example, additions of MoSi, and surface treatments are capable of enhancing the oxidation resistance [ 160, 1611.

5.6.3.2 Boron Nitride

Hard materials based on dense crystalline modifications of boron nitride find wide application in cutting tools [162]. Since these materials are produced at high pressures, the size of samples is usually small. Corrosion has been studied primarily in view of abrasive and tool applications of BN and the data are limited.

However, there exists much information on corrosion of hexagonal graphite-like BN. Similar to the graphite and diamond phases of carbon, the reaction products are the same for all modifications of BN. However, hard modifications of BN typically show a higher corrosion resistance. Thus the available information on graphitic BN can be used for evaluation of the lowest corrosion resistance limit of hard BN.

Although materials based on the high-pressure phases possess very high hardness and strength, they are unstable at high temperatures, which imposes restrictions on their application. When they are heated in air, along with oxidation, the wurzite and sphalerite modifications can be transformed into the graphitic one [ 1621.

5.6.3.2.1 Corrosion by Liquid Media The resistance of pure boron nitride to attack by diluted acids is lower than that to attack by concentrated ones and decreases in


the order HN03 -+ HC1 + H2S04. It interacts slowly with concentrated sulfuric acid to form ammonium sulfate and boric acid [I631 and the corrosion rate of dense specimens is extremely low. Boron nitride dissolves on interaction with concentrated hydrofluoric acid (94% H F solution) by the reaction

BN + 4HF H NH4BF4 (28)

and on heating with ammonium fluoride and concentrated sulfuric acid by the reaction

BN + 3 NH4F + 2 H2S04 % 2 (NH4)2S04 + BF3. (29)

Boron nitride specimens are resistant to attack by hydrochloric, sulfuric, and phosphoric acid solutions with additions of oxidizing agents (KMn04, K~Cr207, KC104), but decompose most rapidly in a 5% H2S04 solution with KC104 addition [ 1631. The corrosion resistance of powders and bulk specimens in alkaline solutions is rather high and almost independent of the concentration.

5.6.3.2.3 Corrosion by Melts Boron nitride fused with potassium carbonate decomposes by the reaction

BN + K2C03 H KB02 + KOCN. (30)

Nitride excess leads to a certain amount of KCN. If carbon is present, the reaction proceeds as follows:

4 BN + 3 K2CO3 + 2 C H K2B407 + C02 + 4 KCN. (31)

Boron nitride reacts with sodium formate [163]. It must be emphasized that boron nitride has an exceptionally good resistance to

most molten salts and glasses, probably due to its resistance to wetting. It is inert to a number of corrosive molten salts, such as alkali halides, lithium borate and cryolites, and has a good resistance to molten glasses. However, it does react with molten alkali carbonates and hydroxides.

5.6.3.2.2 Corrosion by Hot Gases Thermodynamic calculations [ 1081 show that the oxidation of boron nitride at atmospheric and low pressures over the whole temperature range to 2000°C is dominated by

4 B N + 3 0 2 ~ 2 2 2 0 3 + 2 N 2 . (32)

The oxidation of boron nitride as well as boron carbide is distinguished by an oxide layer remaining in a liquid state and vaporizing considerably over the whole temperature range. However, the study of the mechanisms of oxidation of boron nitride materials is complicated by the existence of several BN modifications with considerable differences in their structure and properties. The oxidation of a stable hexagonal a-modification of BN with a graphite lattice has been studied most thoroughly [104].

The studies on Hexanite-R ceramics based on the wurzite modification of boron nitride [164] demonstrated that the oxidation kinetic curves (Fig. 17) were very similar to that for B4C (Fig. 13). A rather high oxidation resistance is also exhibited

5.6 Materials 173

E -1 - a

-2 -

I I

3 O 1 2 t, h

Figure 17. Oxidation kinetics of wurzite BN Hexanite-R at: ( I ) 900°C; (2) 1000°C; (3) 1 IOo"C, and (4) 1200°C; reproduced from [2].

by a cubic boron nitride-based material having sphalerite structure (Elbor-R) and materials that are a mixture of wurzite and sphalerite BN [165].

Graphitic boron nitride is hydrolyzed with humid air to form ammonia and boric acid:

BN + 3 H 2 0 @ H3B03 + NH3. (33)

Boron nitride is resistant to hydrogen and sulfur dioxide; it reacts with fluorine at room temperature by the reaction

2 BN + 3 F2 @ 2 BF3 + N2. (34)

The absence of interaction with other halogens and some halides can be used for removal of boron, boron carbide and other impurities from boron nitride [98].

5.6.4 Carbonitrides 5.6.4.1 Silico-carbonitrides

The few investigations available for the new precursor-derived silico-carbonitrides do not permit their corrosion behavior to be assessed in detail. Most studies of interest here deal with thermal stability only.

The issue of thermal stability is important for these phases, because precursor- derived ceramics are amorphous and have stoichiometries where detrimental structural and chemical changes may occur on crystallization. The highest temperatures recorded so far come from phases in the system Si-B-C-N (2000°C [166]).

For the carbonitrides, the stoichiometry of the amorphous product is Si3+.TN4Cs+p [167] and thus in terms of possible crystallization products there is always C in addition to Si3N4 and Sic. In precursor-derived Si-C-N-ceramics,


micro/nanocrystalline Si3N4 and Sic appear at M 1400-1500°C [168]. Thermal stability experiments show that in Si-GN-ceramics the reaction

Si3N4 + 3 C @ 3 Sic + 2N2 (35)

takes place at about the thermodynamically predicted temperature (M 1405°C at 1 bar N2). The effectivity of this reaction is well known from the sintering of Si3N4-SiC composites [ 1691. Its suppression by oxidation [I 701 indicates an internal pressure build-up of several bar.

In ceramics from polyborosilazanes, on the other hand, ordering effects are demonstrated only at 1700°C and strong crystallization at 1800°C [171]. The crystallization of Si3N4 at T 2 1750°C in this system was also confirmed by Baldus et al. [ 1721 who reported even higher crystallization temperatures, sometimes exceeding 1900°C.

In boron-containing systems the initial Si3N4-crystallization without decomposition may indicate that the carbon is not present in the form of a reactive free carbon: boron may keep carbon in some other (amorphous?) form to suppress or to retard the kinetics of reaction (35).

It is a very old experience that the crystallization of amorphous substances is influenced by numerous parameters [ 1731. Therefore we consider it very dangerous to take the observed crystallization behavior as a fixed value.

Despite a lack of data on active oxidation behavior some thermal stability tests give good hints. It relates to the oxygen impurities of the precursor material. The mass loss observed and clearly attributed to oxygen removal from silicocarbo- boronitrides [166] starts from T M 1300°C. Thus in an environment with low P(02) , at which no protective layer is formed, we will have a substantial loss by active oxidation.

In passive oxidation some very promising data for the new ceramics have been given [170,174, 1751. It is indicated that silicocarbonitrides behave as good as pure Si3N4 [170] and that the boron-containing types are even more resistant [175]. A t , value calculated from the k, given by [174] would be over 300 years at 1550"C!

We are not aware of hot corrosion testing of precursor-derived ceramics. A prediction rests upon the argument that for those which contain boron a top scale compositions in the Na-B-Si-0-system should form, that is a classical glass forming system with low melting points and viscosities [176]. Apart from the hot corrosion degradation by itself it is completely unknown whether this will induce crystallization of the amorphous precursor-derived ceramics.

5.6.4.2 Titanium Carbide, Nitride, and Carbonitrides

Carbides and nitrides of transition metals are usually hard and some of them are used as hard materials. Tungsten carbide, which was briefly described above, is a major component of cutting tools and wear-resistant materials. Titanium carbide, TIC, and nitride, TiN, are two other major hard materials that are widely used in tools. They also find application as hard coatings (in particular, TiN [177]) and components of ceramic matrix composites [178]. Both TIC and TiN show a very

5.6 Materials 175

wide stoichiometry range and almost never exist in the stoichiometric composition. Due to their structural identity, they also form solid solutions in the Ti-C-N system. Oxygen can be dissolved in the material until saturation of carbon or nitrogen vacancies [179] with no changes in the f.c.c. (NaC1) crystal structure. In all cases, whether we consider TIN, Tic, and TiC,N, materials, the oxidation behavior is similar and these rutile-formers will be analyzed together.

The peculiarities of the defect structure of TIC and TiN allow for oxygen dissolution on the first stage of oxidation. Titanium oxides appear later, when the material structure is saturated with oxygen. Anatase was observed at lower temperatures and/or short oxidation times ( 5 min). With increasing time and temperature, rutile is formed as the only oxidation product [108].

Oxidation of TIN powders [180], thin films [177], and dense bodies [I081 was extensively investigated. It was shown that oxidation of TiN, following a logarithmic law, starts at room temperature. Above 500°C the oxidation rate increases and the kinetics follow a parabolic law. Oxidation of a TiN powder prepared by plasma vapor phase synthesis and having the specific surface area of 46 m2 g-' is noticeable above e2250"C and results in the formation of Ti02 as anatase and brookite at lower temperatures, and as rutile at higher temperatures [181].

The scales formed at oxidation of sintered TIN in air can be divided into two major types with the following layer systems: At low temperatures (5 800°C) TiN-dense Ti02-air; at high temperatures (1 000°C) TiN-porous Ti02-dense Ti02-air.

The latter can develop into a multilayered scale during the long-term oxidation above 1000°C [181]. Growth of such rutile scales was based on the simultaneous diffusion of titanium and oxygen ions through the oxide layer in opposite directions.

We can suppose that the fine grain inner layer grows by inward diffusion of oxygen, while the coarse grain outer layer grows by outward cation diffusion.

Oxidation of TIC is similar to that of TIN. Many carbides of transition metals form carbon at the carbide-oxide interface. In the case of TIC, no carbon was formed [182]. Oxidation of titanium and carbon at 1000°C occurs with the same rate [183].

Both TiN and TIC are widely used in ceramic matrix composites for improving electrical conductivity and mechanical properties [ 178, 184,1851. The formation of a rutile scale was observed on the surfaces. The growth of such a scale should be related to the diffusion of titanium to the surface of the composite and its oxidation according to the reaction:

2TiN + 2 0 2 = 2Ti02 + N2. (36)

The protective properties of rutile are fairly low compared to those of alumina or silica. Thus, the presence of TiN or TIC limits the high-temperature applications of the above composites. At small amounts and particle size of TIN in silicon nitride ceramics, a continuous silicate film can be formed, covering TIN particles and protecting them from further oxidation [ 1781.

Thermodynamic analysis of the chemical equilibria in water-carbide systems [loo] indicated that carbon, corresponding oxides, CH4, C02, CO, and H2 are the main products of hydrothermal corrosion of metal carbides in the temperature


range 300-1000°C. There exists a possibility for selective etching of T i c and carbon layers formation. A low H 2 0 to carbide molar ratio results in the formation of carbon layers. Lower pressure is beneficial here. A high H 2 0 to carbide molar ratio means that under equilibrium conditions carbon can be oxidized to CO and C02. Experimental studies showed that no protective oxide layer was formed upon hydrothermal corrosion of carbides of transition metals. Raman spectra obtained from the hydrothermally treated TIC powders demonstrated the presence of graphite. Anatase crystals were also observed on the surface of TIC. With growing temperature, anatase was partly transformed to rutile [loo]. This is in agreement with other published data on corrosion of TIC and TiN [186].

5.6.5 Titanium Diboride

TiB2 is widely used as a high-temperature material, particularly in coatings [187] and as a component of composites [ 1881. According to thermodynamic calculations for the Ti-B-0 system under oxygen partial pressures close to the ambient one and sufficiently high partial pressures of B203, Ti02 is a stable reaction product. Lower titanium oxides are stable only at reduced oxygen pressure, and TiB only at extremely low values of PBl0,. Thus, the main products of TiB2 oxidation in air or oxygen are Ti02 and B2O3. As the oxygen partial pressure is reduced, Ti305 becomes more stable. It was experimentally detected at the scale-boride interface at an oxygen partial pressure of 14.66 kPa [189].

TiB2 oxidation above 700°C was approximated by the parabolic rate law (9) [190]. Deviations from the parabolic oxidation start at temperatures of 950-1 100°C when vaporization of B2O3 becomes noticeable. A cubic law was suggested [191,192], but we assume that a paralinear law (Fig. 7) should better describe the kinetics, similar to the case of BN and B4C.

However, unlike for boron carbide and nitride, the kinetics becomes parabolic again at higher temperatures (>12OO0C). This can be explained by the fact that when B203 is completely evaporated the process is controlled by the diffusion through the titanium dioxide film.

Impurities and metal binders were found to improve the oxidation resistance of TiB2 [193]. This was explained by the fact that they accelerate the sintering of a rutile scale, thus increasing its protective properties.

The oxidation behavior of an equimolar TiB2-A1N composite has been studied in the temperature range 700-1300°C in flowing oxygen [194]. Throughout this temperature range, the oxidation rate is governed by diffusional mechanisms, but three temperature domains can be identified. In the range 7O0-85O0C, only the oxidation of titanium diboride occurs. Calculations assuming a cubic diffusional kinetic law result in the apparent activation energy of 180 kJ mol-' . In the temperature range 900-1000°C, the oxidation of TiB2 still dominates over oxidation of AlN, but the formation of aluminum borates A14B209 and A118B4033 occurs and leads to a more complex behavior. An apparent activation energy of 80kJmol-' was determined for this stage. Above 1 050"C, aluminum nitride oxidation becomes

References 177

more significant and the oxidation kinetics can be described by a parabolic law with an apparent activation energy of 472 kJ mol-' .

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