NASA Contractor Report 174639 

Literature Survey on Oxidations and Fatigue Lives at Elevated Temperatures 

H. W. Liu and Y. Oshida 

Syracuse University 

Syracuse, New York 

March 1984 

Prepared for 

NATIONAL AERONAUTICS AND SPACE ADMINISTRATION 

Lewis Research Center 

Under Grant NAG 3-348 
CONTENTS 

1. INTRODUCTION ............................................... 1 

11. OXIDATION OF Ni AND Ni-BASE SUPERALLOYS .................... 3 

111. GRAIN BOUNDARY OXIDATION ................................... 10 

IV. STRESS GENERATION AND ITS EFFECTS ON OXIDATION ------------- . 13 

V. MICROSTRUCTURAL CHANGES DUE TO LONG-TERM HEATING ----------- 17 

VI. HIGH TEMPERATURE FATIGUE CRACK INITIATION AND PROPAGATION -- 2 0 
I, INTRODUCTION 

Fatigue lives at elevated temperatures are often shortened by creep and/or 

oxidation. Creep deformation creats voids and weakens the material. In conjunction with cyclic load, creep damage may accelerate the initiation and growth of 

fatigue cracks. Oxides are brittle solids that have very low fracture strength 

and fracture toughness. Therefore, oxidation also accelerates fatigue damage 

both in terms of fatigue crack initiation and fatigue crack propagation. There 

is a considerable amount of information on creep and fatigue interaction and 

their acceleration of fatigue damage!'' This review, however, concentrates on 

oxidation effects on fatigue damage. 

It is well known that the data scatter band of stage I1 fatigue crack growth 

is very narrow. Therefore, the wide scatter of fatigue lives might be due to the 

large scatter in fatigue crack initiation and possibly due to the scatter in 

stage I crack growth!21 The crack initiation period could be shortened considerably by the fractures of carbides and intermetallic compounds or by grain boundary 

oxidation. The oxidation accelerated fatigue crack initiation may contribute to 

the zcztL?r ~f fa'iz-e lives. 

Bricknell and ~oodford"~ observed embrittlement as a reduction of the tensile elongation of nickel and nickel alloys due to high-temperature oxidation. 

They tested three grades of nickel (Ni200, Ni270, and Spec-pure Ni) following 

0 air exposure at 1000 C. Tests in a variety of gaseous conditions have revealed 

oxygen as the damaging species and have shown that even low partial pressures 

of oxygen can result in considerable decrease in tensile elongation. 

Cyclic deformation significantly increases the oxidation rateL4] Coffin [ 5 I 

also observed the inability of a film to provide protection against oxidation 

wherever a highly localized deformation occurs. The repeated rupture of the oxide 

film during a fatigue test results in the extremely high oxidation rate in these 

localized regions. cyclic stress or strain comes from either mechanical or thermal effects. Thermal stresses arise from temperature gradient due to transient 

temperature fluctuation and the mismatch of the thermal coefficients of the 

surface oxide and the substrate. 

Coffin et a1. 16] have studied cast Udimet 500 subjected to low-cycle fatigue 

at 1500~~ (815'~) and showed the fatigue crack initiation to be the result of a 

localized ridging. Crack propagation occurs along grain boundaries accompanied 
. by substantial amounts of oxidation. Coffin et al. proposed a "stress oxidation" 

mechansim for crack initiation and propagation, especially along grain boundaries. 

The presence of oxide-filled intergranular cracks is a common observation at high 

temperatures. It is clear that oxidation may shorten the fatigue life considerably. However, a detailed quantitative relation between oxidation and fatigue 

life is still lacking. 

This literature survey relevant to high-temperature low-cycle fatigue in 

nickel-base superalloys reviews general oxidation, selective oxidation, internal 

oxidation, grain boundary oxidation, multi-layer oxide structure, accelerated 

oxidation under stress, stress-generation during oxidation, composition and substrate microstructural changes due to prolonged oxidation, fatigue crack initiation at oxidized grain boundaries and the oxidation accelerated fatigue crack 

propagation along grain boundaries. 
11. OXIDATION OF Ni AND Ni-BASE SUPERALLOYS 

Nickel-base superalloys are the most complex and the most widely used for 

high-temperature applications. The required properties of nickel-base superalloys 

for applications such as aircraft engine components are high-temperature tensile 

strength, thermomechanical fatigue resistance, low thermal expansion, as well as 

oxidation resistance. Oxidation behavior of the nickel-base superalloys is extremely complex for the following reasons. 

1). Each alloying element has different affinity for oxygen. Since the affinity in metalloxide reactions is nearly proportional to the heats of decomposition of oxides per mole of oxygen!7' one may estimate the ease of oxide formation 

of each alloying element by comparing the AH (heats of decomposition) data dec 

listed in Table 1L8] because the higher the AHdec, the more stable the oxide. 

Table 1 shows that oxides of chromium and aluminum form more readily than oxides 

of other major alloying elements. NiO is also formed because Ni is the major 

constituent of the alloy. The oxides of tantalum, silicon, and zirconium are 

seldom formed, even though they are stable oxides, because Ta, Si, and Zr are 

minor coilst it:ienta. 

Depending on the affinity of an element for oxygen and the stability of its 

oxide, one or two of the alloying constituents are selectively oxidized. Generally any alloy containing a small quantity of a metal (which has a high affinity for oxygen, forms an oxide showing resistance to diffusion of oxygen and 

the solute atoms, and has little solubility in the oxide of the solvent metal) 

should have a strong oxidation resistance if the oxide film has a strong adherence 

[9 I to the substrate. 

At a reduced partial pressure of oxygen, that is, under conditions of limited 

availability of oxygen at the surface, oxidation is selective!lol A lower partial pressure of oxygen will also develop at the protective external 
Figure 1 shows an example of the complex oxides formed on (a) ~en6 41 (19% 

Cr, 11%Co, 10%Mo, 3.1%Ti, 1.5YA1, and 0.7%Fe) and (b) Udimet 700 (15%Cr, 19%Co, 

SXMo, 3.5%Ti and 4.5%A1)!lo1 The kinetics of the reaction and the main reaction 

products change with temperature and oxidation time. Hence, at the early stage 

of oxidation, the growth rate of oxide is predominantly linear. The rate deterL mining process involves the transport of the reactants of the oxidation process. 

12' By continuing the oxidation or by increasing the thickness of the oxide film 

(or scale), diffusion becomes the rate controlling process and the growth rate 

becomes parabolic, as discussed below. 

Let us consider the complex oxide structure and composition, which is more 

common in Ni-Cr-Co-Mo-Ti-A1 alloys. Suppose the external oxide scale consists of 

NiO, Cr203, A1203, (Cr,A1)203, NiCr204, NiA1204, Ni(Cr,Al) 0 and Ti02 and the 2 4 

internal oxide subscale includes A1203 Then, as shown in Figure 2, the following steps occur. (a) At an early stage of oxidation, a compact NiO layer is formed 

because Ni is the major constituent. 

(b) As the oxidation proceeds, a porous NiO layer forms beneath 

the compact NiO layer. However, the total NiO layer thickness is increasing becsuse the diffusion rate of Mi in NiC 

is higher than that of Cr in Cr203, as will be discussed 

later. At the same time, Cr203 and Al 0 oxides form at the 2 3 

interface of the inner porous NiO layer and the substrate. 

The 'rapid diffusion of the nickel atoms toward the outer 

surface causes the porous structure. The inward diffusion 

of oxygen toward the substrate and the reaction of the oxygen with chromium result to form Cr 0 Some of these oxides 2 3' 

may be formed inside the porous NiO layer, resulting in a 

solid-state reaction to form a spinel-type oxide such as 

NiCr204, NiA1204, or a mixture of the two as Ni(Cr,Al) 0 2 4' 

(c) As the oxidation proceeds further, more spinel-type oxides 

may be formed and small particles of A1 0 will be formed 2 3 

as a subscale beneath the metal surface. 

(d) By increasing the thickness of both external and internal 

scales, Ti02 will be formed at the scale/substrate interface because Ti is enriched relatively to the other elements, which are already depleted through oxidation. 

-4- 
3). When oxygen dissloves in the alloy phase during oxidation, the less 

noble alloy component may form small oxide particles within the alloy, as shown 

in Figure 2 (c), and bands of oxide may form beneath the external scale/alloy 

interface. This type of scale is known as subscale and the process is called 

internal oxidat ion. 

4). In some cases, liquid oxide phases may also formed on a surface. This 

may occur (i) if materials are exposed to the vapor of a low-melting oxide during oxidation or (ii) in oxidation of alloys with an alloying component that 

forms a low-melting oxide. The typical low-melting oxides are MOO (m.p. : 795O~) 3 

and V205 (m.p. : 674'~) !I3] The presence of liquid oxide phases may lead to excessively fast oxidation and disintegration of the alloy. This phenomenon is 

commonly termed catastrophic oxidation, often occurs in boiler or superheater 

tubes where the combustion gas contains V 0 which condenses and reacts with 2 5' 

fly ash and oxide scales on the heat exchanger tubes. This process is also known 

as V-attack (i.e. the vanadium-attack). An analogous phenomenon may exist in the 

case of the volatile oxide CrO The CrO oxide forms in the temperature range 3 ' 3 

950'~ - 1020 C. 114y15 y161 Although the vapor pressure of CrO is usually low in 3 

this tcnperature region, it is volatile at such high temperatures. 

Selective oxidation, growth of oxides, reaction between different oxides, 

and internal oxidation are always observed in nickel-base superalloys to a certain 

extent depending on chemical compositions, temperature, oxygen pressure, and time 

of oxidat ion. 

Before discussing the kinetics of oxidation, let us summarize the alloying 

effects of individual constituents in nickel-base superalloys. 

Chromium: 1) Chromium appears to be an obvious choice for high-temperature 

alloying constituent by virtue of a natural protective layer . - 

[I71 of chromic oxide (Cr 0 ) which adheres to a surface. 2 3 

2) The "critical" minimum content of chromium to ensure the formation of a protective Cr 0 scale is approximately 15% [I81 or 

[19,20,21] 2 3 

20 - 25%. 

3) For these alloys, a thin external NiO layer is formed by rapid 

overgrowth during the early stage. Under this layer is the protective Cr 0 layer, which controls the oxide growth, as an 2 3 

intermediate spinel-type oxide layer (such as nickel chromite: 

[22,2 33 NiCr 0 ) often does. 2 4 
Cobalt: 1) Cobalt (as well as Ni, Fe, Cr, Mo, W, and V) is an austenitic 

former element and it enhances the y-phase formation. Cobalt is 

[241 a weak solid solution strengthener. 

117,251 2) improves the creep strength. 

[261 3) raises the solvus temperature of y'-phase, Ni3(A1,Ti). 

Molybdenum: 

[I71 1) improves the creep strength. 

2) Molybdenum (as well as Cr and W) is one of the best substitutio- 

[I91 nal solid solution strengtheners. 

Tantalum : 

[271 1) improves the creep resistance. 

Niobium : 

1) Niobium (as well as Ti and Ta) substitutes for Ni in Ni A1 type 

[241 3 

y' precipitate. 

2) Niobium (as well as Cr, Mo, W, V, Ta, and Ti) is an effective 

119 I carbide former (MZ3C6, M6C, MC) . Zirconium : 

1) Zirconirrn (as well as Xi, Ce, aad Ca) is effectfve t3 prc2uzz 

the protective oxide film (or scale) due to the "keying" effect 

[28,291 for anti-spalling. - 

1261 2) promotes ductility. 

Manganese : 

1) improves the strength as a solid solution metal in Ni without 

[251 appreciable reduction of ductility. 

Boron: 1) segregates to grain boundaries and improves creep properties, ( 

1241 by decreasing elongation and increasing rupture stress). 

Aluminum: 

1) increases the oxidation resistance by formation of continuous 

[201 A1203 layer. 

2) Aluminum (as well as Cr) is one of the best substitutional solid 

[I91 solution strengtheners. 

Chromium and Alminum: 

1) Simultaneous additions of Cr and A1 to Ni produce a protective 

complex oxide structure of NiO, Cr203, and A1 0 and their spinel 2 3 I..-.,-, 

type mixed oxides such as NiCr204, NiA1204, and Ni(Cr,Al)204. LJU, 

[ 20 I 31'32'33'341 This oxide gives rise to high oxidat ion resistance. 
2) The Cr/Al ratio plays an important role .in the types of oxide 

[30,311 formed . 

(i) For both low Cr and A1 contents, an external NiO scale 

forms over internal scales of Cr 0 and A1203, 2 3 

(ii) For high Cr (>15%) but low A1 (~3%) contents, both Cr20j 

and A1 0 have been observed as external oxides, and 2 2 

(iii) For relatively high Cr (>15%) and high A1 (>3%) contents, , 

an exclusive A1 0 scale is formed. 2 3 

When nickel is heated in oxygen, it forms an adherent film scale of nickel 

oxide, NiO, which controls the rate of further oxidation by acting as a diffusion 

barrier between the metal and the atmosphere. The mechanism and the rate-controlling process for the initial formation of the film and its growth at relatively 

low temperatures (i.e., below 400°c), are not at all clearly understood, but at 

higher temperature the growth follows a parabolic law. 

Wagner's mechanism[351 for the parabolic relationship is that of film thickening controlled by diffusion due to a concentration gradient. Since the oxide 

film is composed ef cations (metal ions) and anions (oxygen ions), this diffusion 

across the film must be accompanied by a simultaneous diffusion of electrons in 

the same direction as the cations, or in the opposite direction to the anions. 

Hence the oxidation of nickel is governed by the diffusion (as a rate-determinins factor) of charged vacancies in the oxides or of only one type of ion (either 

cation or anion) or of electron through the lattice defects of oxides. 

Methods of investigating the growth of oxide layers include determination of 

changes in thickness of the scale, in weight of the metal sample, or in volume 

of the surrounding gas!361 Most quantitative data have been obtained as changes 

in weight (AW) per unit surface area. Kinetic theory is concerned primarily with 

the progress of a reaction with time. Thus, the first task is to find relationship between oxidation and time t. The parabolic relationship of the oxidation 

2 2 

process (AW) = a x t represents a straight line when (AW) is plotted against 

time t. If this line does not intersect the origin of the coordinate axes, the 

2 

more general form of the parabolic equation, (AW) = a x t + C, is applicable. 

The constant a in the above equations is the parabolic rate constant. 

Figure 3 shows the self-diffusion coefficients of a metal ion (cation) in 

its single or complex oxides. Figure 4 shows the parabolic rate constant of 
several nickel-base superalloys. 

It is well known that the temperature dependence of reaction rates is governed by an exponential relation known as the Arrhenius equation. According to this 

equation, the temperature dependence of diffusion rates is generally written in 

the form D = D e -Q/RT , where D and Q are constangs. On the other hand, according 0 

to Kubaschewski and H~pkinsj~~l~the temperature-dependent parabolic rate constant 

for the oxidation is given by the expression a = aoe -Q/RT , where a. = 3.2 and 

Q = 45,000 cal/mole for pure nickel oxidation over the temperature range 750 - 

1240'~. 

One may calculate the activation energies for self-diffuison and parabolic 

oxidation from the data in Figures 3 and 4. The results are tabulated in Tables 

2 and 3, respectively. Because the magnitude of the activation energies of the 

diffusion and the oxidation processes are similar, it is likely that the diffusion 

process is rate controlling. However, because the oxidation processes are complex 

chemical reaction, it is difficult to pinpoint which diffusion process controls 

a specific oxidation process. 

Heindlhofer et a1!381 have pointed out that no binary alloy can be expected 

to oxidize cxnctly in prcportion tn the square root of the elaspcd time at constant 

temperature, because chemical composition in metals immediately adjacent to the 

scale changes during that time as a result of selective oxidation and internal 

oxidation. Although this is true in the early stage of oxidation, after the 

steady state is established, the concentrations at the oxide/substrate interface 

139 I may usually be considered constant. 

Results obtained by Yao 1401 on pure nickel, by ~ettrid~e[~~] on Nimonic 

alloys, and by Kofstad [411 on en: 41, Mar-M 200, and IN-100 are shown in Figure 

2 4 5 with a common unit of (mg /cm /sec) for a parabolic rate constant, a. The parabolic rate constant a decreases in the order of pure nickel - Nimonic series I 

----+ IN 100 -----+ Rene 41/Mar-M 200. This order agrees approximately with the 

alloying effects of Cr, Al, and Ti on oxidtaion, as discussed above. 

Under stress or strain, the oxidation reaction will be accelerated because 

(i) the protective oxide layer is cracked, particularly under a cyclic loading, 

and (ii) dislocations and surface slip steps are sites of localized high energies 

where oxidation is accelerated. The effect of applied stress has been studied 

on several nickel alloys. Oxidation proceeds at an unaffected rate with increasing stress until a critical level of stress is reached, above which oxidation 

is more rapid. This critical stress to produce accelerated attack is plotted in 

-8- 
Figure 6 14*' as a function of exposure temperature for a number of alloys. At the 

critical applied stress, which generally corresponds to the stress required to 

produce 1% extension in 100 hours, the rate of metal deformation inhibits the 

formation of protective and adherent oxide scales. Thus, the continual exposure 

of fresh metal results in a faster oxidation rate. In this regard, any stress 

that produces excessive shear at the oxide/metal interface can cause cracking, 

separating, and exfoliation of the protective oxide and exposure of fresh metal. 

/ Creep testing of 10-mil Rene 41 sheet showed that selective depletion of alloying elements increases significantly with increasing stress. 

Rhines et a1!431 have studied the effects of externally applied forces on 

oxidation rate by suspending high purity nickel tubes in a nickel-tubed furnace 

with various dead load (tensile stresses: 3050, 6100, 9150, and 12199 psi). The 

tested tubes were paired, one being stressed and the other unstressed as they 

were oxidized together in the furnace. Rhines et al. found that as the load is 

increased, the oxidation rate first diminished both inside and outside the tubes, 

af ter which it increases. Antolovich [441 attempted to that account of the effect 

of stress in terms of the stress dependent diffusion constant for the parabolic 

kinetics . 

Recently, I keda et al. [451 studied the effects of strain on the oxidation of 

stainless steels (Fe- 18@ 20Cr, Fe-20Cr-12Ni, and Fe-17Cr-6Ni) in oxygen and 

in a i r a t 900~~. The oxidtaion of Fe- 18 - 20Cr f erritic steels was not accele- 

- 6 rated by tensile loadings that produced average strain rates of 10 10-~/sec. 

However, the oxidation of the austenitic steels was accelerated at the average 

- 6 strain rates of the order of 3 x 10 /sec. The authors speculated that the absence of accelerated oxidation in the ferritic steels could be attributed to 

the rapid diffusion of chromium through these steels and the resulting rapid 

healing of defects produced in the oxide scale during straining. The low diffusion rate of chromium in the austenitic steels, according to this speculation, 

did not lead to repair of these defects, and thus, resulted in increasing oxidation during straining. 

These observations clearly indicate that engineering materials such as 

nickel-base superalloys with grain boundaries, carbides, and precipitates 

should show locally accelerated oxidation, particularly at such microstructural 

discontinuities, when subjected to an externally applied load and high-temperature oxidation. The spalling and thermal shock resistance in oxides and autogenerated stress in oxides will be discussed in chapter IV. 
111. GRAIN BOUNDARY OXIDATION 

The grain boundary is an important site for oxidtaion because (i) its composition may differ from that of the matrix, (ii) local disorder may make it chemically more active than the matrix, (iii) it may act as a site for precipitation 

of harmful or matrix-depleting phases, and (iv) it may function as a rapid diffu1461 sion path. 

Several interconnected facts on grain boundary oxidtaion appear to be esta149 I blished in the works by Gulbransen et ad4'] Rh~din!~~] Kubaschewski and Hopkins, 

and ofs st ad!^^] Figure 7(a) illustrates the following phenomena: (i) On 

a crystal surface, the orientation of the oxide film is directly related to that 

of the crystal surface. (ii) The rate of growth of such a film depends on its 

orientation. (iii) On a polycrystalline metal surface, the first layer is randomly 

oriented and the oxide grains are of different thicknesses owing to the variation 

in growth rate with orientation. (iv) The oxide films immediately above the grain 

boundaries or precipitates are thicker than these formed on metal matrix phase. 

The rapid oxidation at grain boundaries was observed by Sywestrowicz on 

1521 titaziun; thlz ffi;r,a. Reporting sz the SEX exminztizns of oxidstioz c!:arzct?- 

ristics of alloy 800H, Miley et al. [531 found that growth rates of Cr-rich oxide 

were greatest at the grain boundaries. When external scale had begun to form, the 

substrate developed with the appearance of oxide particles within the alloy matrix 

ahead of the interface between metal and external scale by the process of internal 

oxidation. These internal oxide particles appeared to a greater extent at or near 

the grain There also exists a depletion zone of alloying elements 

at or near the grain boundaries. The oxide particles at the boundaries grew to 

form continuous films, which progressively enveloped the grains. At the same time, 

oxidation proceeded within the grain by the formation of further oxide particles 

ahead of the thickening grain boundary film in much the same way that oxidation 

occured initially at the metal surface. 

Shida et studying preferential intergranular oxides in Ni-A1 (0.55 - 

4.10ZAl) alloys at 800 - 1000~~, identified the several types of intergranular 

oxide morphologies. They reported that thin and continuous oxides developed along 

grain boundaries in Ni-0.55A1 and Ni-1.5A1 and that thick oxide particles developed around and in the boundaries in Ni-2.45A1 and Ni-4.10A1 at 900 and 1000~~. 

Reddy et a~!~" have studied diffusion in AIZOj scales formed on NiCrAl 

18 alloys by the proton activation technique employing the 0 isotope as a tracer. 
They found that tile 018 profiles identified a zone of dxide penetration beneath 

the external scale and that oxygen transport was primarily via short-circuit 

paths such as grain boundaries. Bricknell and his coworkers [3,58,59,60,611 have 

done extensive work on oxygen embrittlement in Ni and its alloys. Complex oxides 

were identified on the boundaries at depths roughly an order of magnitude greater 

[58,591 than those for matrix internal oxidation. Using Auger spectroanalysis, they 

0 observed the enhanced O/Ni peak height ratios in the air-exposed (1000 C x 200 

hr) samples. [601 However, this technique has disadvantages; (i) the detected 0 

peak cannot distinguish between external oxide and internal oxide, and (ii) this 

0 peak also cannot distinguish inward diffusing oxygen from ambient and dissolved 

oxygen in the substrate matrix. Therefore, the experimental technique to measure 

the penetration depth of oxide formed along the grain boundaries has not been well 

established. The measurement of grain boundary penetration is still based on a 

conventional technique, examination of the cross section of an oxidized sample 

under an optical microscope. 

Grain size has an effect on the oxidation behavior. Armanet et al. [621 reported 

on the oxidation behavior of 80Ni-20Cr alloy as a function of average grain size. 

Depe~lding on the st ructu~r, tdo siages oE two different types of gravimetrlc Curves were noted. The first, which is nonparabolic, is due to the growth of a NiO 

layer over the entire surface of the sample except at grain boundaries, where 

Cr203 forms. The second curve is parabolic and is due to the formation of a continuous layer of chromium sesquioxide (Cr 0 ) at the grain boundaries. The first 2 3 

stage is shorter, and the parabolic stage sets in sooner as the average grain 

size is smaller. Pischitelli et al. i631 also observed that the oxidation rate in 

Fe-29Ni-17Co alloy depends on the alloy grain size and the amounts of the minor 

alloy constituents and that the rate is enhanced at grain boundaries. 

As the oxidation process continues, grain boundaries will be preferably 

oxidized by extended external oxidation or increased. internal oxidation, or both, 

resulting in the morphology shown in Figure 7(b)..The outer surface of external 

oxide scale will be cracked because of thermal and mechanical stresses to be 

discussed in the next chapter. 

Table 4 summarizes the penetration depth of oxide and its thickness along 

the grain boundary, grain size, and thickness of oxide formed on the surface of 

various superalloys. This table shows that by prolonging the oxidation process 

(i) inward oxide penetration may reach two grain sizes in depth and (ii) the 
thickness of oxide formed along grain boundaries increases to roughly twice that 

of the oxide formed on the surface. Therefore, the grain boundary oxide must be 

crucial for thermomechanical fatigue phenomena (either in crack initiation or 

in crack propagation). Since the relevant data under unstressed conditions are 

lacking, it is premature to draw conclusions here. Suffice it to say that stress 

cycling accelerates the oxidation rate particularly at grain boundaries. 

The high-temperature engineering components whose service life limits are 

controlled by fatigue have been designed to a crack initiation criterion. A component is considered to have failed and is removed from service as soon as a crack 

reaches the size of 0.031 inch (= 0.8 mm or 800 pm). r641 This critical crack size 

is comparable to the grain boundary oxide penetration depth in some of the hightemperature alloys shown in Table 4. For longer service times, the penetration 

depth could conceivably be deeper. 
. IV. STRESS GENERATION AND ITS EFFECTS ON OXIDATION 

During high-temperature oxidation of a metal, internal stresses (called an 

auto-generated stress) arise from various causes mentioned below. The oxidation 

behavior will be influenced not only by the externally applied stress but also 

[ 65 I by this auto-generated internal stress during oxide formation, 

An isothermal stress can be attributed to four main causes [66y671: (i) . . the 

epitaxial relationship between the oxide and the metal, [66y681 (ii) the volume 

ratio between the metal and the oxide formed, [66 y671 (iii) composition changes in 

either the metal or the oxide !69 ' 70y 71y 721 and (iv) the influence of vacancies 

[68,73,74,75,76] generated during oxidation. 

It has been stated that stress may arise in oxide films as a result of epitaxial interactions between oxide film and metal substrate. These interactions 

include oxidelmetal lattice mismatch, formation of pores in the oxide and metal 

[77,781 and at their interface, and the growth morphology of the oxide. 

The stress is produced by lattice expansion associated with oxygen solution 

in the surface layers of the metal, by the volume change that occurs when metal 1 

is converted +-#? its oxide at 2 coherent ir,t.erface, or by the volume change due 

[791 to the existence of a substantial gradient in stoichiometry across the oxide. 

If an oxide film is a coherent and adherent film and if lattice mismatch exists, 

internal stress will occur. Table 5 lists the volume ratio ($), defined as the 

ratio of the molecular volume of oxide (MeOx) to the atomic volume of metal 

(Me) j8 pilling and Bedworth [801 proposed that oxide films could be divided into 

two classes with regard to their protective abilities by assuming that oxide 

growth takes place through an inward migration of oxygen through the scale; (i) 

If the oxide/metal volume ratio is smaller than one (e.g., K20, BaO, or MgO), 

,a tensile internal stress will exist in a coherent film, and the oxide will fail 

to cover the entire metal surface. The oxide will be nonprotective because it is 

unable to form a coherent surface film, thus continually exposing fresh metal I 

surface. (ii) If the volume ratio is larger than 1, the oxide will be protectivb 

except for oxides with extremely high ratios (e.g., Na 0 Ta205, Moo3, WOg, ori 2 5' i V 0 ). These oxides are accompanied by the formation of cracked or porous scales, 2 5 

possibly after an initial stage during which the scale remains coherent. I 

1 

I Even if an ideal situation existed in which the oxide volume and metal vo-1 

lume were equal, it is most likely that each phase would have a different coeffi- 
cient of thermal expansion. Thus, if coherene were maintained at all times, interfacial strains would be induced upon any change in temperatures. In general, the 

thermal dilation is very significant, and the magnitude of the internal stresses 

[651 is often equal to greater than the fracture stress of the oxides. 

If transport occurs by an outward migration of metal by the vacancy mechanism, vacancies may coalesce to form cavities and pores at the metallscale interface and produce appreciable porosity in the oxide scale after more extended reaction. The porosity is probably due to the development of stresses with subsequent cracking, a phenomenon called breakaway oxidation!781 ~heref ore, the 

oxide that have medium value of volume ratio are more likely to be protective. 

In polycrystalline material, large stresses may also develop at the grain 

boundaries owing to the differences in oxidation rates (i) between neighboring 

grains and between grains and grain boundaries and (ii) between different crystal surfaces of different orientations !781 Formation of new phases (such as a 

spinel-type oxide by a solid-state reaction) during the oxidation reaction may 

[all further serve as a source of stress. 

An initial compressive stress in the oxide due to epitaxial growth is followed by a period in which the surface regions of the nelix were in tension because 

of vacancy injection into the metal. The net stress may become compressive again 

as oxidat ion proceeds. [771 Poole et a1!821 measured the stress in thin films (less 

3 2 than 500 2 thick) of CuZO and FelOg as compressive stresses of 1.6 x 10 (Kglcm ) 

3 2. and 2.7 x 10 (~g/cm ), respectively. These stresses, whatever their source, can 

have a profound effect on film structure. These stresses may cause rupture in 

the scale. 

The internal stresses do not necessarily lead to rupture. During isothermal 

oxidation at high-temperature, plastic flow and sintering of the oxide may also 

take place. The stress relief due to plastic deformation of oxide may prevent 

rupture of the scale. The plastic deformation could be due to slip or could in- 

[831 volve high-temperature creep in the oxide. Sintering of the oxide may rake 

place through various diffusion processes such as volume, grain boundary, and 

surface diffusions, the rates of which all increase exponentially with tempera- 

[781 ture. 

Nonuniform dimensional changes cause infernal stresses. These stresses are 

more significant in oxide films than in bulk metal substrates. Oxides lack ductility to relieve the stress; consequently, cracking and spalling are common in 
oxides!84 Nonuniform dimensional changes may arise iron any of the following 

184 I processes. 

(i) Oxide phase transformation 

These may be exemplified either by the reconstructive transformation of 

quartz to cristobalite of SiO or by the displacive transformation of low cris- 2 

tobalite to high cristobalite. Either change invlolves heat-induced expansion, 

which produces tensile stresses in the unaltered part of the material and compressive stress in the transformed part. 

(ii) Differential thermal expansions 

These are typified by mullite porcelains (A1 0 2Si0 ), which exhibit 2 3 2 

differential expansion in polyphase materials, because mullite has a linear coo efficient of 5 x cm/cm/ C, while the surrounding glassy matrix has an ex- 

- 6 o pansion coefficient of 7 x 10 cm/cm/ C. Nimonic has linear thermal expansion 

o 0 1851 coefficient of 18 x c/cm C in the temperature range 200 - 1000 C, 

while the spinel-type oxide has a linear thermal expansion coefficient of 7.6 - 

0 [86,871 8.8 x cm/cm/ C in the same temperature range. 

(iii) Anisotropic thermal expansion within crystals 

AF zn esarople, an alumina (41 O ) insulator can ha-re adjacent grains "2 3 

0 that expand 8.3 and 9.0 x lom6 cm/cm/ C if their c-axes are perpendicular to 

each other. 

(iv) Temperature differentials and their effect on dimensional changes 

Here the simplest example is glass, which might be cooled rather slowly so that the outside layer contracts thermally while the center is still hot 

and semielastic. A tensile stress exists in the outside layer when the inside 

is in compression. If the tensile stress in the outside layer causes nonelastic 

deformation and relieves the tensile stress, and if the glass is further cooled, 

the center will be in tension and the outside layer will be in compression. Further cooling places the center under tension (and the surface under compression) 

because the rigid surface cannot continue to contract as the center contracts. 

The ways in which intermittent cooling from the service temperature affects 

weight loss, scale fracture, buckling, and spalling have been widely recognized. [ 

88'89990'911 Oxide spalling is enhanced by temperature cycling and the thermal2 fatigue conditions. A spalling weight loss of 1.0 mg/cm corresponds approximately to a loss of in ./side of metal consumed. Oxide spalling usually exposes 

an alloy that is depleted in the protective components (Cr and/or Al) so that the 
subsequent oxidation rate is faster, with considerable metal being consumed in 

the process!421 The potential severity of cyclic versus steady temperature service is illustrated for six commercial alloys in Table 6!421 These alloys exhibit 

a wide range of behavior because of the complexity of their reaction products, 

as well as differing extents of surface and internal oxide attack. 

Rapidly forming oxides may have more porosity and larger nonstoichiometry 

than oxides that form slowly. Hence, a rapidly forming oxide would be more highly 

stressed and would rupture before reaching - the thicknesses observed with slowly 

forming oxide^!^ ' 37 9921 Wolf et a1Lg2 presented exploratory work on the phenomelogy of gaslmetal reactions, considering the anisothermal oxidation of 

nickel and titanium during heating and when subjected to a variety of controlled 

heating rates. They used two different heating patterns for anisothermal oxidation; (i) heating rate (114 OC, ~OC, and 8Oclsec.) up to Tm (maximum temperature) 

of 950°, 100oO, and 1050~~ with holding times of 1 and 4 hours, and (ii) after 

0 the various heating rates, a cooling rate of 8 Clsec., followed by noholding 

time at Tm. At a very early stage of isothermal oxidation, the weight gain of 

the rapidly heated sample (8°~/sec.) was approximately 8 times larger than that 

Ecated sio~l~ (~.25~~!sec.j. 3owevcr, after a 1-hr holding tin=, the mcics of 

weight losses are reduced to 1 - 1.5. 

The role of thermal shock in oxidation weight change in high-temperature 

alloys was investigated in research conducted at Lewis Research Center. r931 Sheet 

0 specimens were oxidized in 1-hr cycles, terminated by slow (1 Clsec.) or fast 

0 (240 C/sec.) cooling, to accumulate 100 hours at temperature. Materials were 

B-1900+Hf, IN-601 (both oxidized at 1100~~) , TD-Ni, TD-NiCrA1, IN-702, and Hoskins 

0 875 (oxidized at 1200 C). Records of the progressive weight change are shown in 

Figure 8. For TD-Ni and Hoskins 875, no oxide spalling was observed under either 

.cyclic condition. For TD-NiCrA1, the shock condition resulted in extensive cracking of the metal and increased oxide loss by spalling. For other materials, slow 

cooling resulted in greater oxide loss by spalling. 
V. MICROSTRUCTURAL CHANGES DUE TO LONG-TERM HEATING 

The important time-dependent properties such as creep strength and oxidation depend on the metallurgical structures. The metal temperatures of turbine 

blade vary between 500'~ and 950 O C. [94y951 At a high enough service temperature, 

the metallurgical structures of nickel-base superalloys may change. 

Above 0.6Tm (Tm: melting temperature), the y' phase (Ni3(A1,Ti)) ripens 

(increases in size) at a significant rate. The higher the y' particle size, the 

higher the creep resistance. Therefore, minimizing ripening will help to retain 

long-term creep resistance, [961 by (i) increasing y' percent [981 or (ii) adding 

[961 high-partitioning and slow-diffusing Nb and Ta to the alloy. 

Depending on their formability, nickel-base superalloy products are usually 

used as (i) sheet such as Waspaloy and Rene 41 wieh 25% y', (ii) wrought (the 

strongest) such as Udimet 700 with 25% y', and (iii) cast such as IN-100 and 

Mar-M 200 with 55 - 65% y' 1971 The sheet M-252 of low volume percent y' has lower flow stress than cast Inconel alloy 700 of high percent y'. The sensitivity 

of flow stress to changes in y' particle size was much greater in the more lean 

[961 alloys. '&ile a high volume fraction of y' is desirable for high yield and 

tensile stresses, the fatigue strength seems to be lowered with increasing volume 

[99 I fraction of y' particles. 

Several general rules can be derived to relate carbide types to alloy compositions after approximately 5000 hours of service. These rules are derived 

from the computer modelling studies on the formation and decomposition of carbides [ 101, [loo] and long-term (700 to 5000 hours) metallurgical experimental data. 

lo2] The general findings are as follows: (1) Below the critical temperatures of 

matrix phase stability (815'~)~ grain boundaries are relatively stable; and MC 

and M C are the principal equilibrium phases. (2) Low chromium and high colum- 23 6 

bium, tantalum, molybdenum, tungsten, and iron stabilize MC carbides and improve 

ductility. (3) At and above ~OOO~C, Cr C has been found only in a basic Ni-Cr- 7 3 

Ti-A1 superlalloy such as Nimonic 80A. Addition to the basic superalloy of cobalt 

(Nimonic go), molybdenum (Waspaloy), tungsten (D-979), or columbium (Inconel 

alloy X-750) removes this chromium carbide. (4) At a temperature of 980°c, close 

to the upper design region, MC, M C, and M C are pseudoequilibrium phases in 6 23 6 

superalloys other than Nimonic 80A. High Cr, Ti, Al, and Ta favor the formation 

of M23C6; high Mo and W favor M C; and Nb and Ta favor MC. (5) As the service 6 
temperature increases, the carbon content of nickel alloys may slowly decrease. 

After very long exposure, alloy such as Udimet 500 (0.05C) have appeared to 

lose all MC due to decarbonization. [961 

These microstructural changes (y'-ripening, carbide-precipitation, and decarbonization) due to long-term exposure at high-temperature may cause different high-temperature low-cycle fatigue or thermomechanical fatigue behaviors. 

The y'-ripening causes the decrease of an effective element (Al) for the formation of the protective oxide, and as a result, overall oxidation resistance 

decreases. Carbide-precipitation provides more crack-initiation sites, to be 

discussed in the next chapter. Therefore, both microstructural changes may 

shorten the high-temperature fatigue life in nickel-base superalloys. 

The turbine blade alloys are generally heat treated to provide a duplex y' 

morphology consisting of large cuboidal precipitates, which confer good creep 

resistance and small spherical precipitates that increase the tensile strength. 1 

96,1031 

The effects of microstructure on elevated temperature fatigue and creep 

crack growths were studied by Merrick et al. [lo4] on Waspaloy and P/N Astroloy 

0 

at 650 2. They rlyorted that iu Z;lspaioy, changes in y' size and distribution 

did not markedly affect the fatigue crack growth rate. However, they noticed 

that an increase in fatigue crack growth rate occurred at low test frequencies 

and was associated with a transition to intergranular crack propagation. In 

P/M Astroloy, a coarser grain size lowered the fatigue crack growth rate. Serrated grain boundaries are beneficial under creep loading but have no effect 

on the fatigue crack growth rate. 

There seems to be general agreement, however, that a further complicating 

factor is that the precipitate structure generally encountered in nickel-base 

alloy is not stable under stress at elevated temperatures. [491051 Phase stability, 

particularly of the precipitates is also related to mechanical stress. Gasturbine blades made of Ni-alloy El893 were investigated in the initial state and 

after 11,000 hours of operation at 7300~1~~' After service the tensile strength 

increased 10 - 22%, but the ductility and impact strength were reduced 50% because 

of the aging process. 

From the foregoing review of material degradation under the combination 

of stress and high-temperature oxidation, it is clear that these phenomena may 

have even more important effects on oxidation behavior at grain boundaries and 

that this could be crucial to HTLCF (High-Temperature Low-Cycle Fatigue [1061) 
and TMF (Thenno-Mechanical Fatigue I1071 
VI. HIGH TEMPERATURE FATIGUE CRACK INITIATION AND PROPAGATION 

Surface fatigue cracks often originate as intergranular cracks in air at 

high temperatures, especially at low frequencies. Both environment (particularly oxidation) and time-dependent flow can infuence the propensity for intergranular surface crack initiation, since cracks are often developed as a consequence of the grain boundary sliding, void coalescence, and grain boundary 

[lo81 oxidation. 

Surface oxidation is detrimental to fatigue because it can cause the formation of stress raisers (such as voids or oxide cracks, as was discussed in 

chapter IV), and therefore it can accelarate crack inititaion. Paskiet I1091 has 

shown that preoxidation of a specimen at 982'~ followed by fatigue testing at 

0 760 C produced many surface intergranular cracks, whereas, testing of specimens 

without prior oxidation produced a single intergranular crack. Thus, preoxidized 

phases or grain boundaries served as incipient cracks. ~u~ita[llO] has also 

proposed a model for accelarated oxidation of surface-connected slip bands, which 

serve as crack nuclei. Danek [1111 and Achter ['12] have explained that at high 

stresses end short fatigue lives, oxygen izdsorption at thc crack ti? zcc:lcr:t:s 

the rate of crack growth in air, while at low stresses, a thicker oxide is formed, 

which is thought to strengthen the metal. 

According to localized oxidation leads to a "notch" at the 

surface and localization of the strain, which accelerates further oxidation and 

sharpens the notch. This reinforcing cycle continues until local fracture occurs. 

This results in a crack, which they can propagate into the material. It has been 

seen that the initial region of oxidation is often broad and diffuse, but the 

notch created by the material removed through oxidation acts to concentrate the 

, deformation further, leading to a wedge-shaped notch and finally a crack. On A286 

(iron-base superalloys) , Coffin identified highly localized surface oxidation as 

the nucleation site for fatigue cracks under low-cycle fatigue. off in['l4] tested 

0 A286 and Udimet 500 at 816 C and reported that crack initiation was transgranular 

-8 in vacuum (10 torr) and intergranular in air. Therefore, oxidation strongly 

affects fatigue crack initiation, particularly along grain boundaries. 

In fact, it has been generally accepted that in engineering alloys, fatigue 

crack initiation occurs at some preexisting defect. [11591161 Defects can be micropores that are introduced in casting or atomization, inclusions (both metallic and 

nonmetallic), and microstructural features such as carbides and brittle second 
phases. ['I7 "18' In addit ion, micros t ruc tural features such as grain boundaries, 

twin boundaries, or precipitates can be crack nucleation sites. In general, defects can be any discontinuity in the microstructure which acts to raise the 

local stress level. As a result, localized plastic deformation can be generated 

in regions adjacent to the defects even at a nominal stress well below the yield 

stress of the material. This localized deformation process is important in fatigue 

crack initiation. ['16] Therefore, the greater the plastic strain concentration, 

the greater its effect on crack nucleation. 

0 In cast Udimet 500, crack initiation in low-cycle fatigue at 650 C occurs 

[I19 I preferentially at surface carbide particles, which cracked during testing. 

Subsequent crack propagation was found to be in the stage I mode at this tempe0 rature. At 815 C, crack initiation occurred preferentially at grain boundaries 

and interdendritic carbide particles, with subsequent propagation in the stage 

0 I1 mode. At 982 C, fatigue crack initiation takes place totally at grain boundaries. 

In the case of columnar and single-crystal Mar-M 200, [1201 cracks still ori0 ginated at the carbides and micropores at 760 C and 927'~. On the other hand, in 

conventionallp cast Mar-M 200, grain boundaries were the preferred sites for 

0 crack initiation and propagation at 760'~ and 927 C. The mode of propagation 

in the columnar and single-crystal materials was transgranular. The growth 

rate was slower and fatigue lives were one or two orders of magnitude higher 

than those of conventionally cast Mar-M 200, in which cracks grew intergranularly 

Wells et al. [1219122'1231 studied Udimer 7000 and found that (i) from the 

room temperature to 64g0cj localization of slip in discrete bands led to slip0 band crack initiation at free surfaces, (ii) at 760 C, surface cracking became 

predominantly intergranular, because of the reduction in grain boundary strength 

0 and (iii) at 927 C, crack initiation as well as propagation were intergranular. 

[I191 This observation, in general, agrees with that of Sabol et al. 

As temperature increases, indicating of surface intergranular crack initiation have been found in wrought Udimet 700. [1221 At 870'~ all the crack initiation in this material is at the surface and is intergranular. Often, these 

intergranular surface cracks propagate only one grain diameter before a transition to transgranular cracking occurs. 

Crack initiation has been observed at oxidized carbides at the surface 

o [124] where carbides were decomposed of Inconel 718 tested in air at 850 C. 
[I251 Fatigue crack was initiated at carbides in Rene 95. / 

Antolovich et a1 . concluded from work on ~eng 80 and Rene 77 that lowcycle fatigue damage at elevated temperatures is primarily in the form of oxidation or oxide penetration along surface-connected grain boundaries. A crack is 

said to initiate at a critical degree of boundary penetration., as was discussed 

in the previous chapter. 

Recently, Matsuoka et al. reported that fatigue cracks are initiated 

at grain boundaries of Fe-9Cr-1Mo steel tested at 593O~ and 700'~ since the 

early fracture is intergranular. 

Less direct, but nevertheless significant, is the metallographic evidence 

that crack propagation in air at high-temperature is accompanied by highly localized oxidation on the fracture surface. This, in turn, leads to oxide extrusion, or ridging, as in the case of cast Ni-base superalloys or the formation 

of complex surface oxide structures as for A 286. Dennison et al. [I29 1 

found a similar behavior for cast IN-100, and fatigue fracture occurred by the 

propagat ion of surf ace-nucleated cracks. Romanoski [1301 studied high-temperature 

low-cycle fatigue on ~eng 80 and 111-100 at 871°c and 1000~~ and observed that 

crack initiation and propagation generally occurred at grain boundaries because 

they possess poorer mechanical strength. 

Most alloys when tested at elevated temperatures under cyclic loading, 

including hold times, usually exhibit a lesser fatigue resistance in terms of 

cyclic life if the hold is in tension rather than in compression. 'I3' However, 

Teranishi et al. 11311 found that a compression hold can be more damaging than a 

tension hold for 2ziCr-1Mo steel. The reason for the effect of hold time on cyclic 

life can be related to the behavior of the oxide immediately after the hold time. 

After a tension hold, the oxide spalls to produce a new surface, which, at least 

in the early stages of the test, did not contain macroscopic cracks. On the other 

hand, after a compression hold, the oxide cracks rather than spalls, thereby 

creating localized stress and strain concentrations. These facilitate the early 

nucleation of fatigue cracks in the substrate, one of which will ultimately 

be responsible for failure of the specimen. 

Suzuki et al!t321 using 1Cr-1Mo-CV steel tested at 500 - 600°c, found that 

the repeated formation of Fe 0 oxide followed by spalling accelerates the fatigue 2 3 

crack propagation rate. Similar results were obtained by Skelton on Cr-Mo-V 

steel. 

Crack nucleation by oxide ridging at surface grain boundaries has been 

- 22 - 
0 0 [ observed to be more characteristic of cast Ni-base supetalloys at 1600 F (871 C). 

1341 In all cases, crack nucleation is intergranular, with oxide ridging being 

the principal mechanism. Propagation, on the other hand, is transgranular at all 

frequencies (0.03 - 36.2 cpm) tested. This transgranular crack propagation is 

believed to account for the frequency insensitivity found at high strain and 

short fatigue lives, through crack tip blunting due to oxidation with decreasing 

frequency. Transgranular crack growth acts to enhace the resistance of this alloy 

11341 to fatigue at high-strain low-cycle frequency and hold time. P 

Among the nickel-base superalloys (IW-100, IN-738, Mar-M 200, ~ene' 77, Rene 

80, en: 95, Inconel alloy 600, Inconel alloy 718, Udimet 500, and Udimet 700) 

tested by several researchers cited in this chapter, Inconel alloy 718 and Udimet 

500 contain 18 - 19%Cr (see Table 7). This amount of chromium is nearly enough 

to provide a protective surface oxide film, when a specimen is stress free, as 

was discussed in chapter 11. However, under a cyclic load, cracks are initiated 

o [114, at grain boundaries and propagated intergranularly in Udimet 500 above 815 C 

'lg1 and crack initiation is observed at oxidized carbides in Inconel alloy 718 

o 11241 

at 850 C. These observations suggest that on material containing about sufficient chror~iux c3::tent fcr the £0-rnztiun sf prctecti-re oxide film, suc5 files 

are not protective against fatigue cycling. . 

In general, the effects of temperature is to cause an increase in the fatigue 

crack growth rate, a slight decrease in the threshold AK, and a greater sensitivity to frequency effects. At elevated temperatures in air, the cylcic crack 

growth rates da/dN increase with decreasing frequency. In several nickel-base 

-5 superalloys, the fatigue crack growth rates range from 2 x 10 to 2 x in./ 

1135 I cycle. 

Gel1 and Duquette interpreted the variations of high-temperature fatigue 

life in terms of creep-oxidation interaction. They proposed that in the regime 

of high stresses and short fatigue lives, crack initiation and crack propagation 

are intergranular because of the high creep component during a fatigue cycle. 

For tests run in air, oxidation of grain boundary phases accelerates the rates 

of intergranular crack initiation and propagation and reduces the fatigue life. 

But, as the stress is lowered, the creep component in a fatigue cycle is reduced, 

' El371 which favors a transgranular mode of cracking and longer fatigue life. 

~carlin"~~~ has shown that the crack propagation rates for both Nimonic 

105 and IN 738LC increased as the cyclic frequency decreased for high-temperature 

tests in air. The data also showed a higher rate of crack growth in air than in 

- 23 - 
a vacuum. Scarlin concluded that oxidation was the controlling factor in increasing the propagation rates in air. J,ames has also studied the effect of oxidation on Ni-base superalloys (Inconel 600 ans Inconel 718). Crack growth tests 

[I391 The at elevated temperatures were performed in air and in liquid sodium. 

growth results showed that crack propagation in liquid sodium was slower than in 

air, especially in the lower AK regime. The decrease in crack growth rate in 

liquid sodium was associated with the presence of an oxygen free environment. 

Recently, ~eerasoori~a~~~~~stud~in~ fatigue crack growth rate in IN 718 and 

IN 100 as a function of temperature (21 - 648'~) and frequency (0.001 - 10 Hz), 

reported that the fracture surface was entirely covered with oxide and that the 

color (which is related to the thickness of oxide film) changed continuously 

depending on temperature and frequency (i.e:, oxidation time). 

Oxide films are protective if they adhere to the substrate. The film may 

crack or spall, however, if the sum of the internal and applied stresses exceed 

the fracture strength of the oxide or the adhesion strength. Both fracture and 

adhesion strengths can be affected by the temperature, the time at the temperature, 

and possibly the total accumulated cyclic strain. 

Gr*in tobndaries are sites of higlier energy, mors com2lex composition, iind 

higher chemical reactivity than the grain matrix. Therefore, the rate of oxidation is faster at the grain boundary than within a grain. Accordingly, the 

oxide is much thicker in the grain boundary region, and the grain boundaries are 

the location of potential layer oxide cracks. The nonuniform oxidation rate on 

the surface of an alloy causes complex internal stresses. Spalling will produce 

geometric notches in the oxide layer. Both cracks and notches will expose "fresh" 

substrate material for further.oxidation, and the strain localization at the 

notches and the crack tips will accelerate the crack nucleation process within 

the matrix material. 

Oxidation depletes A1 and Cr, which are the major constituents for the y 

and y' strengthening phases. The depletion of A1 and Cr weakens the resistance 

to plastic deformation, including resistance to cyclic plastic deformation, which 

is one of the principal mechanisms of fatigue crack initiation. 

When crack growth during a fatigue cycle (da/dt)At by the mechanism of oxidation is slower than crack growth by the "regular" cyclic deformation mechanism, 

then the cyclic crack growth.wil1 take over as the major damaging mechanism. Once 

a crack starts to grow by the cyclic mechanism, the growth rate at high temperature can be accelerated by oxidation and/or creep. 
Let us define the cyclic crack growth rate without environmental and creep 

effects as Aa To affect oxidation, oxygen has to be transported to the crack f ' 

tip. Let Aa- be the depth of penetration by the detrimental chemical species, i. 

Y 

e. , oxygen. According to Gel1 et al! the rare-controlling mechanism is either 

the surface diffusion of oxygen to the crack tip or the diffusion of oxygen to 

the region ahead of the crack tip, either through the lattice or along grain 

boundaries. In addition, the mechanism of direct adsorption and diffusion along 

slip bands could also be the rate-controlling mechanism. 

When Aaf >> Aa the chemical environment will have a negligible effect on 

P ' 

the overall rate of crack growth. This could be the case of low temperature and 

high cyclic frequency. This is the region of nil-environmental effect, as sugget- 

[141,142] ed by Liu and McGowan. In this region, the cyclic crack growth rate is 

not temperature- and frequency-dependent. 

When Aa >> A the detrimental chemical species reaches far ahead of the 

P f' 

crack tip. In this case, a crack is propagating in a fully embrittled solid. 

Again da/dN is not temperature- and frequency-dependent. This is called the 

fully-environmental effect region. Between these limits, Aa is comparable f 

+ ,, . ,,_, n ., and dz/dN is both temperat1:re- and frequency-dependent, Fignre 9 ill:^-. r 

strates schematically the fatigue crack growth rate as a function of AK, tem- 

[I421 perature (T), and cyclic frequency (v). 

Cyclic plastic deformation at the crack tip may fracture the oxide film 

and expose nascent surface for oxidation attack. 

When creep accelerates fatigue crack growth, there seems no valid reason 

for a limit line to which all the crack growth rate data at various temperatures 

and frequencies converge. 

When the cyclic plastic zone size is small relative to the grain size, the 

crack plane follows specific metallographic planes, crack growth is faceted, and 

the rate of fatigue crack growth is dependent on crystal orientation and may 

vary from grain to grain. 

Consequently, there are two main competing mechanisms involved in high temperature low-cycle fatigue. To determine which is dominant, a comparison should 

be made between the crack growth rate (daldt) during a fatigue cycle by the mechanism of the grain boundary oxidation and the crack growth rate v(da/dN) by the 

I I regular" cyclic deformation mechanism. 

So far, this review has concentrated on general oxidation behavior and localized oxidation (e.g., selective oxidation, internal oxidation, and grain 

- 25 - 
boundary oxidation) of nickel-base superalloys as well as stress generation and 

microstructural changes during the oxidation process. It has also been shown that 

grain boundaries as well as carbides will be sites for fatigue crack initiation, 

and that grain boundary oxidation has strong effects on not only fatigue crack 

initiation but also its propagation stage. Figure 10 summarizes the information 

presented in this review and illustrates the interrelationships among the factors 

that are important for high-temperature fatigue phenomena. 
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Table 1 Heats of decomposition of individual oxide 

of alloying elements in Nickel-base alloys [8] 

metal 

Ni 

Cr 

Co 

Mo 

W 

Ta 

bib (Cb) 

A1 

Ti 

Fe 

oxide 

NiO 

Cr203 

coo 

Moo2 

Moo3 

W02 

W03 

Ta205 

Nb 2'5 

*'2'3 

Ti02 

FeO 

minor alloying element 
Table 2 

activation energy for self-diffusion, -Q [~cal/mole] 

system calculated from fig .3 ref. data 

Co / coo 

Ni /NiO 

Table 3 

activation energy for high temperature oxidation, -Q [Kcal/mole] 

weight % 

material calculated Q Cr Co Mo A1 Ti Nb 

from fig. 4 

INCO 713 87 

Table 5 Volume ratio of oxides 

molecular volume of oxide 

metal oxide volume ratio ($) = atomic volume of metal 

K K2° 

Ba BaO 

Cr Cr203 

M MO, - 

W w02 

Fe Fe203 

Fe304 

Co coo 

Ni NiO 1.65 
Table 6 Effects of cyclic air oxidation on spalling 

(400 hr at 1100~~) [421 

Static Exposure Cyclic Exposure ' 

Total ~ffccted~ Total hffcstcdh 

Spallcd Oxide irlctai Spaled Ox~de Metal 

(mg/cm2 ) (niil/side) (mg/cm2) (mil/side) 

RenC 80 15.2 7.3 101.0 12.9 

IN-I00 3 .O 4.6 8.7 12.5 

713C 2.7 1.8 36.8 9.5 

Hastclloy X 0.5 2.5 3.4 4 .O 

Ren6 Y 0 1.8 0.07 2.0 

TD Ni-Cr 0 0.8 0.3 1 .O 

16-25 hr cyclcs to room temperature. 

lncludcs metal that converted to oxidc and 10 (infernal oxides) 
Table 7 Chemical compositions of Nickel-base superalloys (~10) 

Alloy Designation 

Nickel dase 

~lloy 713C 

Alloy 713LC 

B-1900 

FORD 406 

1~-1001~en~ 100 

:IN-162 

IN-731 

11-738 

IN-792 

HZ2 

MAR-H 200 

n~r-H 200(DS) 

MAR4 246 

UAR-H 421 

MAR-H 432 

NX188(DS) 

REX' 77 

RENE' 80 

S EL 

SEL-15 . . TAE-8A 

TRW-NASA VIA 

UDIHET 500 

WAZ-2O(DS) 

Aatroloy 

D-979 

Ilastelloy alloy X 

tnconel alloy 600 

Inconel alloy 601 

- Inconel alloy 625 

lnconel alloy 706 

tnconel alloy 718 

Inconel alloy X-750 

IN-102 

IN-587 

IN-597 

IN-853 

H-252 

Nimonlc alloy 75 

Nimontt alloy BOA 

Nimonie alloy 90 

Nlmonic alloy 105 

Nimonlc alloy 115 

Nimonic alloy PE.11 

Nimonlc alloy PE.16 

Nimonlc alloy PK.33 

Nimonlc alloy 120 

Nimonlc alloy 942 

Pyro~rr 860 

M-333 

 en: 41 

 en; 95 

TD Nichal 

TD NtC 

Udlmet 500 

L'dimrt 520 

Udimct 700 

Udloec 710 

Unit*ap AI2-1DA 

Ya*paloy 

- Nominal 

-- Cb A1 

2.0 6.1 

2.0 5.9 

- 6.0 

2.0 4.5 

- 5.5 

1.0 6.5 

- 5.5 

0.9 3.4 

- 3.1 

- 6.3 

1.0 5.0 

1.0 5.0 

- 5.5 

2.0 4.3 

2.0 2.8 

- 8.0 

- 4.3 

- 3.0 

- 4.4 

0.5 5.4 

2.5 6.0 

0.5 5.6 

. - 3.0 

- 6.5 

Compoaition. Weight X 

-- Zr Other 
I I 1 I I . 10 100 loo0 10,000 

Time. min 

(a) ~en6-41 

19%Cr, 11%Co, 10%Mo, 3.1%Ti, 1.5%A1, 

so01 I 1 I 1 1 10 100 1000 10,000 

Time, min 

(b) Udimet-700 

15%Cr, 19%Co, 5%Mo, 3.5%Ti, 4.5%A1 

Note: i) phases in the numerators represent the external scale, 

ii) phases in the denominator represent the internal scale, 

@ represents the linear oxidation region, 

@and@ represent the parabolic oxidat ion regions, and 

@represents a constant weight region after 1000 min. 

Fig. 1 schematic representation of reaction products formed 

as a function of time and temperature in air oxidation [lo] 
compact NiO layer 

[NiCrCoMoTiAl ] alloy 

(a> 

porous NiO layer 

1 spinel type 

! NiO -Cr203 

oxide y NiO =(Cr ,Al) 2 0 3 

sesquioxide \ 

d L 

internal oxide layer external oxide layer 

Fig.2 Schematic illustrations for complex oxides formation 

on Nickel-base superalloys 

-41- 
Fig. 3 Self-dif fusion coefficients of the cations in oxides [411 
Fig.4 Parabolic rate constants for oxidation on several Ni-base superalloys 

[411 
Temperature (-C) 

Fig. 5 

Co Uo W TA Cb A1 Ti (vtX) 

Pane 41 - - - t.5 3.1 

Uar M2OO 12- 15 2 

IN - 100 60 10 15 - - - 5.5 4.7 . .. - - - - 0.4 

Parabolic rate constants for oxidation.on several Ni-base superalloys 

-44- 
Fig.6 The critical applied stress required to produce 

acceleration oxidation as a function of 

exposure temperature for Ni-base superalloys 
oxide thickness (SO 

penetration depth 

of oxide (PDO

internal oxide particles 

(b 

grain boundary oxide thickness (=2 x GBO) 

Fig.7 Schematic illustrations for oxide growth in cross-sectional view 
- Fat cdhg -- Sow sodlng 

Fig.8 The effect of cooling rate on the cyclic oxidation of 

some high-temperature alloys 

Note:1hratTmaxand0.5hratT percycle min [931 
Fig. 9 Cyclic crack growth model schematic. 

[I431 
t (STRESS CYCI.INC;) 

Fig.10 Interrelationships among various important factors 

for high temperature low cycle fatigue 
1. Report No. 

NASA CR-174639 

,, t ., + I March 1984 

I I 

Literature Survey on Oxidations and Fatighc LIVc3 aL 

Elevated Temperatures 6. Perform~ng Organ~zat~on Code 

2. Government Accession No. 

4. T~tle and Subtitle 

3. Rec~pient's Catalog No. 

5. Report Date 

9. Perforrnlng Organization Name and Address 

Syracuse University 

Dept. of Chemical Engineering and Materials Science NAG 3-348 

I 

7 Author@) 

H. W. Liu and Y. Oshida 

8. Performing Organization Report No 

None 

10. Work Unit No. 

National Aeronautics and Space Administration 

Washington, D .C . 20546 

Syracuse, New York 13210 

12. Sponsoring Agency Name and Address 

I 

15. Supplementary Notes 

13. Type of Report and Period Covered 

Contractor Report 

Final report. Project Manager, John L. Shannon, Jr., Structures and Mechanical 

Technologies Division, NASA Lewis Research Center, Cleveland, Ohio 44135. 

16. Abstract 

Nickel-base superalloys are the most complex and the most widely used for 

high-temperature applications such as aircraft engine components. The desirable 

properties of nickel-base superalloys at high temperatures are tensile strength, 

thermomechanical fatigue resistance, low thermal expansion, as well as oxidation 

resistance. At elevated temperature, fatigue cracks are often initiated by grain 

boundary oxidation, and fatigue cracks often propagate along grain boundaries, 

where the oxidation rate is higher. Oxidation takes place at the interface 

between metal and gas. Properties of the metal substrate, the gaseous 

environment, as well as the oxides formed all interact to make the oxidation 

behavior of nickel-base superalloys extremely complicated. The important topics 

include general oxidation, selective oxidation, internal oxidation, grain 

boundary oxidation, multi-layer oxide structure, accelerated oxidation under 

stress, stress-generation during oxidation, composition and substrate 

microstructural changes due to prolonged oxidation, fatigue crack initiation at 

oxidized grain boundaries and the oxidation accelerated fatigue crack propagation 

along grain boundaries. These factors are reviewed in this report. 

17. KW Wwdr (Su~Qmtrd by ~~thOf(8)) 18. Dlstributlon Statement 

I Superal loys; Oxidation; Fatigue; Crack 

initiation; Crack propagation 

Unclassified - unlimited 

STAR Category 26 

I 

19. S.curfty Clurif. (of thle nport) 

unclassified 

 or sale by the National Technical Information Service. Springfield. Virginia 22161 

20. Sacurlty Ckrrlf. (of thlr page) 

Unclassified 

21. No. of pages 

5 1 

22. Prlc.. 

A04