High Temperature Corrosion Failure of Super Alloy Turbine Blades

 V. Vijay Raghavan

Qatar PetroleumP.O.Box 100001Dukhan, Qatar

Fax No: (0974) 4717488E-mail: v_vijayaraghavan@qp.com.qa

ABSTRACT:Two first-stage turbine blades catastrophically failed in operation. The Fractured surface indicated that the crack was propagated by high cycle fatigue (HCF). The presence of corrosion products of sulfur on the blades pointed out a possibility of Type 2 high temperature corrosion. The Manufacturers recommendation was to using a heavy-duty filtration to remove contaminants to eliminate the High temperature corrosion. The alloy used in the blades was a second generation CMSX-4, a product especially developed for high temperature operations for aircraft turbines. This means reason for the failure needed further investigation. This paper is a brief presentation of the corrosion failure study conducted by the Qatar Petroleum’s Corrosion Section.

Contents:1. Super alloys -History.

2. Nickel Based Super alloys

3. Application of Nickel based super alloys

4. Super Alloy Casting

5. 1st , 2nd and 3rd Generation single crystal Nickel based Super alloys.

6. Microstructure of Single crystal super alloy casting.

7. What really happened? The failure

8. Investigation and Test Methodology

9.

1. Super alloys -History.

The term "superalloy" was first used shortly after World War II to describe a group of alloys developed for use in turbo-superchargers and aircraft turbine engines that required high performance at elevated temperatures. The range of applications for which superalloys are used has expanded into many other areas and now includes aircraft turbines, land-based gas turbines, rocket engines, chemical, and petroleum plants. They are particularly well suited for these demanding applications because of their ability to retain most of their strength even after long exposure times above 650°C (1,200°F). Their versatility stems from the fact that they combine this high strength with good low-temperature ductility and excellent surface stability.

Super-alloys are based on Group VIIIB elements and usually consist of various combinations of Fe, Ni, Co, and Cr, as well as lesser amounts of W, Mo, Ta, Nb, Ti, and Al. The three major classes of super-alloys are nickel, iron, and cobalt-based alloys. As CMSX 4 is a Nickel alloy we limit our discussion to Nickel based super alloys.

2. Nickel based Super alloys

Nickel-based alloys can be either solid solution or precipitation strengthened/Hardened. Solid solution strengthened alloys, such as Hastelloy X, are used in applications requiring only modest strength. In the most demanding applications, such as hot sections of gas turbine engines, a precipitation strengthened alloy is required. Most nickel-based alloys contain 10-20% Cr, up to 8% Al and Ti, 5-10% Co, and small amounts of B, Zr, and C. Other common additions are Mo, W, Ta, Hf, and Nb (often still referred to as "columbium" although the name "niobium" was adopted by the International Union of Pure and Applied Chemistry in 1950 after more than 100 years of controversy). In broad terms, the elemental additions in Ni-base superalloys can be categorized as being i) formers (elements that tend to partition to the matrix, ii) ' formers (elements that partition to the ' precipitate, iii) carbide formers, and iv) elements that segregate to the grain boundaries. Elements which are considered formers are Group V, VI, and VII elements such as Co, Cr, Mo,W, Fe. The atomic diameters of these alloys are only 3-13% different than Ni (the primary matrix element). ' formers come from group III, IV, and V elements and include Al, Ti, Nb, Ta, Hf. The atomic diameters of these elements differ from Ni by 6-18%. The main carbide formers are Cr, Mo, W, Nb, Ta, Ti. The primary grain boundary elements are B, C, and Zr. Their atomic diameters are 21-27% different than Ni.

The major phases present in most nickel superalloys are as follows:

Gamma ( ): The continuous matrix (called gamma) is an face-centered-cubic (fcc) nickel-based austenitic phase that usually contains a high percentage of solid-solution elements such as Co, Cr, Mo, and W.

Gamma Prime ( '): The primary strengthening phase in nickel-based superalloys is Ni3(Al,Ti), and is called gamma prime ( '). It is a coherently precipitating phase (i.e., the crystal planes of the precipitate are in registry with the gamma matrix) with an ordered L12 (FCC) crystal structure. The close match in matrix/precipitate lattice parameter (~0-1%) combined with the chemical compatibility allows the '

to precipitate homogeneously throughout the matrix and have long-time stability. Interestingly, the flow stress of the ' increases with increasing temperature up to about 650oC (1200oF). In addition, ' is quite ductile and thus imparts strength to the matrix without lowering the fracture toughness of the alloy. Aluminum and titanium are the major constituents and are added in amounts and mutual proportions to precipitate a high volume fraction in the matrix. In some modern alloys the volume fraction of the ' precipitate is around 70%. There are many factors that contribute to the hardening imparted by the ' and include ' fault energy, ' strength, coherency strains, volume fraction of ', and ' particle size.

Extremely small ' precipitates always occur as spheres. In fact, for a given volume of precipitate, a sphere has 1.24 less surface area than a cube, and thus is the preferred shape to minimize surface energy. With a coherent particle, however, the interfacial energy can be minimized by forming cubes and allowing the crystalographic planes of the cubic matrix and precipitate to remain continuous. Thus as the ' grows, the morphology can change from spheres to cubes (as shown in this figure) or plates depending on the value of the matrix/precipitate lattice mismatch. For larger mismatch values the critical particle size where the change from spheres to cubes (or plates) occurs is reduced. Coherency can be lost by overaging. One sign of a loss of coherency is directional coarsening (aspect ratio) and rounding of the cube edges. Increasing directional coarsening for increasing (positive or negative) mismatch is also expected.

Carbides: Carbon, added at levels of 0.05-0.2%, combines with reactive and refractory elements such as titanium, tantalum, and hafnium to form carbides (e.g., TiC, TaC, or HfC). During heat treatment and service, these begin to decompose and form lower carbides such as M23C6 and M6C, which tend to form on the grain boundaries. These common carbides all have an fcc crystal structure. Results vary on whether carbides are detrimental or advantageous to superalloy properties. The general opinion is that in superalloys with grain boundaries, carbides are beneficial by increasing rupture strength at high tempeature.

Topologically Close-Packed Phases: These are generally undesirable, brittle phases that can form during heat treatment or service. The cell structure of these phases have close-packed atoms in layers separated by relatively large interatomic distances. The layers of close packed atoms are displaced from one another by sandwiched larger atoms, developing a characteristic "topology." These compounds have been characterized as possessing a topologically close-packed (TCP) structure. Conversely, Ni3Al (gamma prime) is close-packed in all directions and is called geometrically close-packed (GCP).

TCPs ( , µ, Laves, etc.) usually form as plates (which appear as needles on a single-plane microstructure.) The plate-like structure negatively affects mechanical properties (ductility and creep-rupture.) Sigma appears to be the most deleterious while strength retention has been observed in some alloys containing mu and Laves. TCPs are potentially damaging for two reasons: they tie up and ' strengthening elements in a non-useful form, thus reducing creep strength, and they can act as crack initiators because of their brittle nature.

3. Strength versus Temperature

Gas turbine thermal efficiency increases with greater temperature of the gas flow exiting the combustor and entering the work-producing component—the turbine. Turbine inlet temperatures in the gas path of modern high-performance jet engines can exceed 3,000°F, while nonaviation gas turbines operate at 2,700°F or lower. In high-temperature regions of the turbine, special high-melting-point nickel-base superalloy blades and vanes are used, which retain strength and resist hot corrosion at extreme temperatures. These superalloys, when conventionally vacuum

cast, soften and melt at temperatures between 2,200 and 2,500°F. That means blades and vanes closest to the combustor may be operating in gas path temperatures far exceeding their melting point and must be cooled to acceptable service temperatures (typically eight- to nine-tenths of the melting temperature) to maintain integrity.

Thus, turbine airfoils subjected to the hottest gas flows take the form of elaborate superalloy investment castings to accommodate the intricate internal passages and surface hole patterns necessary to channel and direct cooling air (bled from the compressor) within and over exterior surfaces of the superalloy airfoil structure. To eliminate the deleterious effects of impurities, investment casting is carried out in vacuum chambers. After casting, the working surfaces of high-temperature cooled turbine airfoils are coated with ceramic thermal barrier coatings to increase life and act as a thermal insulator (allowing inlet temperatures 100 to 300 degrees higher).

The strength of most metals decreases as the temperature is increased, simply because assistance from thermal activation makes it easier for dislocations to surmount obstacles. However, nickel based superalloys containing γ', which essentially is an intermetallic compound based on the formula Ni3(Al,Ti), are particularly resistant to temperature.

Ordinary slip in both γ and γ' occurs on the {111}<110>. If slip was confined to these planes at all temperatures then the strength would decrease as the temperature is raised. However, there is a tendency for dislocations in γ' to cross-slip on to the {100} planes where they have a lower anti-phase domain boundary energy. This is because the energy decreases with temperature. Situations arise where the extended dislocation is then partly on the close-packed plane and partly on the cube plane. Such a dislocation becomes locked, leading to an increase in strength. The strength only decreases beyond about 600oC whence the thermal activation is sufficiently violent to allow the dislocations to overcome the obstacles.

To summarise, it is the presence of γ' which is responsible for the fact that the strength of nickel based superalloys is relatively insensitive to temperature. The yield strength of a particular superalloy containing only about 20% of γ'. The points are measured and the curve is a theoretical prediction. Notice how the strength is at first insensitive to temperature.

When greater strength is required at lower temperatures (e.g. turbine discs), alloys can be strengthened using another phase known as γ''. This phase occurs in nickel superalloys with significant additions of niobium (Inconel 718) or vanadium; the composition of the γ'' is then Ni3Nb or Ni3V. The particles of γ'' are in the form of discs with (001)γ''||{001}γ and [100]γ''||<100>γ

The crystal structure of γ'' is based on a body-centred tetragonal lattice with an ordered arrangement of nickel and niobium atoms. Strengthening occurs therefore by both a coherency hardening and order hardening mechanism. The lattice parameters of γ'' are approximately a=0.362 nm and c=0.741 nm

4. Application of Nickel based Super alloys

Nickel-based superalloys are used in load-bearing structures to the highest homologous temperature of any common alloy system (Tm = 0.9, or 90% of their melting point). Among the most demanding applications for a structural material are those in the hot sections of turbine engines. The preeminence of superalloys is reflected in the fact that they currently comprise over 50% of the weight of advanced aircraft engines. The widespread use of superalloys in turbine engines coupled with the fact that the thermodynamic efficiency of turbine engines is increased with increasing turbine inlet temperatures has, in part, provided the motivation for increasing the maximum-use temperature of superalloys. In fact, during the past 30 years turbine airfoil temperature capability has increased on average by about 4°F per year. Two major factors which have made this increase possible are

1. Advanced processing techniques, which improved alloy cleanliness (thus improving reliability) and/or enabled the production of tailored microstructures such as directionally solidified or single-crystal material.

2. Alloy development resulting in higher-use-temperature materials primarily through the additions of refractory elements such as Re, W, Ta, and Mo.

About 60% of the use-temperature increases have occurred due to advanced cooling concepts; 40% have resulted from material improvements. State-of-the-art turbine blade surface temperatures are near 2,100°F (1,150°C); the most severe combinations of stress

and temperature corresponds to an average bulk metal temperature approaching 1,830°F (1,000°C).

Although superalloys retain significant strength to temperatures near 1800°F, they tend to be susceptible to environmental attack because of the presence of reactive alloying elements (which provide their high-temperature strength). Surface attack includes oxidation, hot corrosion, and thermal fatigue. In the most demanding applications, such as turbine blade and vanes, superalloys are often coated to improve environmental resistance

5. Super alloy Casting The material and casting technique improvements that have taken place during the last 50 years have enabled superalloys to be used first as equiaxed castings in the 1940s, then as directionally solidified (DS) materials during the 1960s, and finally as single crystals (SC) in the 1970s. Each casting technique advancement has resulted in higher use temperatures.

In DS processing, columnar grains are formed parallel to the growth axis. In nickel-based alloys, the natural growth direction is along the <100> crystallographic direction. This morphology is accomplished by pouring liquid metal into a mold that contains a water-cooled bottom plate. Solidification first occurs at the bottom plate, after which the mold is slowly withdrawn from the furnace, allowing the metal inside to directionally solidify from bottom to top. The exceptional properties of DS and SC alloys is due to

1. The alignment or elimination of any weak grain boundaries oriented transverse to the eventual loading direction.

2. The low modulus associated with the <100> directions enhances thermal mechanical fatigue resistance in areas of constrained thermal expansion—particularly turbine vanes. In general, the lack of transverse grain boundaries coupled with the lower modulus can result in 3-5 times improvement in rupture life.

SC casting were developed during the 1970s and were a spin-off from the technological advances made in the DS casting processes. SC casting are produced in a similar fashion to DS by selecting a single grain, via a grain selector. During solidification, this single grain grows to encompass the entire part. Single crystals obtain their outstanding strength through the elimination of grain boundaries that are present in both equiaxed and directionally solidified materials. In addition, the elimination of grain boundary strengtheners such as C, B, Si, and Zr raises the single crystal's melting point. By increasing the alloy's melting point, the homogenization heat-treat temperature can be increased without fear of incipient melting, thus allowing for more complete solutioning of the ' and thereby increasing alloy strength and maximum use temperature.

6. 1 st , 2 nd and 3 rd Generation single crystal Nickel based Super alloys

The single-crystal superalloys are often classified into first, second and third generation alloys. The second and third generations contain about 3 wt% and 6wt% of rhenium respectively. Rhenium is a very expensive addition but leads to an improvement in the creep strength and fatigue resistance. It is argued that some of the enhanced resistance to creep comes from the promotion of rafting by rhenium, which partitions into the γ and makes the lattice misfit more negative. Atomic resolution experiments have shown that the Re occurs as clusters in the γ phase. It is also claimed that rhenium reduces the overall diffusion rate in nickel based superalloys.

The properties of superalloys deteriorate if certain phases known as the topologically close-packed (TCP) phases precipitate. In these phases, some of the atoms are arranged as in nickel, where the close-packed planes are stacked in the sequence ...ABCABC.. However, although this sequence is maintained in the TCP phases, the atoms are not close-packed, hence the adjective 'topologically'. TCP phases include σ μ. Such phases are not only intrinsically brittle but their precipitation also depletes the matrix from valuable elements which are added for different purposes. The addition of rhenium promotes TCP formation, so alloys containing these solutes must have their Cr, Co, W or Mo. concentrations reduced to compensate. It is generally not practical to remove all these elements, but the chromium concentration in the new generation superalloys is much reduced. Chromium does protect against oxidation, but oxidation can also be prevented by coating the blades.

Single-Crystal Superalloys, nominal compositions (wt%).

A nickel-base single crystal superalloy having high creep rupture strength at high temperatures, said superalloy being obtained by subjecting a single crystal alloy having a composition consisting essentially of, by weight,

Cr 4.5-10%,

W 7.5-20%,

Al 4.5-6%,

Ta 2-12%,

Co 5-10%,

Ni substantially being the balance,

W+Ta=17-24%,

to a solutionizing treatment, an air cooling treatment and an aging treatment; and a process for producing the same.

7. Microstructure of Super alloy Casting

CMSX-4 is an ultra high strength, single crystal alloy development of the Cannon Muskegon Corporation. This second generation rhenium-containing, nickel-base single crystal alloy is capable of higher peak temperature/stress operation of at least 2125°F (1163°C). Large tonnages of CMSX-4 have been produced. Solar Turbines report blade lives to overhaul of 25,000 - 30,000 hrs in their 15,000 hp Mars 100 industrial gas turbine.

On the effect of orientation on the primary-creep mechanism of the Ni-based single-crystal super-alloy CMSX-4, both the elastic and the mechanical properties are anisotropic. Anisotropy is the property of being directionally dependent, as opposed to isotropy, which means homogeneity in all directions. The young modules off these items are anisotropic as well it varies from 130 Gpa (<001> orientation) to 290 Gpa (<111>orientation). The effect of crystalline anisotropy is dependent on the chemical composition, heat treatment, testing mode and temperatures.

Turbine blades in gas turbine engines are subjected during operation to tri-axial stress fields. For the description of the deformation behavior of anisotropic single-crystal blades, constitutive equations are required which take account of modifications to the deformation processes caused by evolution of the γ/γ' microstructure during service (γ' rafting). A microstructure-dependent, orthotropic Hills potential, whose an isotropy coefficients are connected to the edge length of γ' particles, has been applied. The shape of γ' particles remains cubic below exposures at 700 °C. At high temperatures (above 850 °C) the γ' particles coalesce to rafts, and the visco-plastic response of the superalloy is continuously modified. This reduces the creep resistance of 001 orientated specimen. After tensile loading of the 001 -orientated specimens at 1000 °C, the rafting of γ' in the (100) plane was observed as expected, whereas the 111 specimens did not reveal γ' rafting. Torsionally loaded specimens exhibited rafting only in the near 100 -orientated surface regions of the specimen. The deformation in the 111 tensile and 001 torsion specimens occurred by octahedral slip of dislocations and not by cubic slip, as expected from theoretical considerations. Rafting did not occur in the 111 -orientated specimens.Tensile behavior of a new single-crystal nickel-based super-alloy with rhenium (CMSX-4) was studied at both room and elevated temperatures. The investigation also examined the influence of γ' precipitates (size and distribution) on the tensile behavior of the material. Tensile specimens were prepared from single-crystal CMSX-4 in [001] orientation. The test specimens had the [001] growth direction parallel to the loading axis in tension. These specimens were given three different heat treatments to produce three different γ' precipitate sizes and distributions. Tensile testing was carried out at both room and elevated temperatures. The results of the present investigation indicate that yield strength and ultimate tensile strength of this material initially increases with temperature, reaches a peak at around 800 °C, and then starts rapidly decreasing with rise in temperature. Both yield and tensile strength increased with increase in average γ' precipitate size. Yield strength and temperature correlated very well by an Arrhenius type of relationship. Rate-controlling process for yielding at very high temperature ( T ≥ 800 °C) was found to be the dislocation climb for all three differently heat-treated materials. Thermally activated hardening occurs below 800 °C whereas above 800 °C thermally activated softening occurs in this material.

8. What really happened? The failure :

Two first-stage turbine blades catastrophically failed in operation after around 9500 h service (approximately 12 months). The expected service life was 40 000 h the failure was visually analyzed by optical microscopy, scanning electron microscopy (SEM)and X-ray diffraction (XRD) and dimensional metrology. The Fractured surface indicated that the crack was propagated by high cycle fatigue (HCF). The presence of corrosion products of sulfur on the blades pointed out a possibility of Type 2 high temperature corrosion. The blades, manufactured in the nickel super-alloy CMSX-4, did not have any protective coating the un-protected surfaces was susceptible for mechanical damage due to particle impact and high-temperature hot corrosion (Type-II corrosion).

Hot corrosion may be defined as an accelerated corrosion, resulting from the presence of salt contaminants such as Na2SO4, NaCl, and V2O5 that combine to form molten deposits, which damage the surface oxides. Hot corrosion occurs when metals are heated in the temperature range 700–900C in the presence of sulphate deposits formed as a result of the reaction between sodium chloride and sulphur compounds in the gas phase surrounding the metals. At higher temperatures, deposits of Na2SO4 are molten (m.p. 884C) and can cause accelerated attack on Ni- and Co-based super-alloys. This type of attack is commonly called ‘hot corrosion’. Accelerated corrosion can also be caused by other salts, viz. vanadates or sulphates– vanadate mixtures and in the presence of solid or gaseous salts such as chlorides. Contaminants such as Vanadium and sulfur from the fuel gas/oil react to form Na2SO4 in the combustion system. During combustion of the fuel, vanadium reacts with oxygen to form an oxide V2O5 (m.p. 670°C). Thus V2O5 is a liquid at gas turbine operating temperature. These compounds, known as ash, deposit on the surface of materials and induce accelerated oxidation (hot corrosion) in energy generation systems. Corrosion occurs when these molten compounds dissolve the protective oxide layers, which naturally form on materials during gas turbine/boiler operation. Further, Vanadium compounds are good oxidation catalysts and allow oxygen and other gases in the combustion atmosphere to diffuse rapidly to the metal surface and cause further oxidation. As soon as the metal is oxidized, the cycle starts over again and high corrosion rates occur.

Two types of sulfate-induced hot corrosion are generally identified. Type I takes place above the melting point of Na2SO4 and Type II occurs below the melting point of Na2SO4

but in the presence of small amounts of SO3. In Type I the protective oxide scale is dissolved by the molten salt. Sulfur is released from the salt and diffuses into the metal substrate forming discrete grey/blue colored aluminum or chromium sulfides so that, after the salt layer has been removed, the metal cannot rebuilt a new protective oxide layer.

How does a well-designed product like this can fail?We asked ourselves a few questions:Will high-temperature oxidation will take place in this material and will it damage the surface? YesWill mechanical damage will occur on the surface of the ballade due to particle impinch ment? YesWill these lead to catastrophic failure? No - it shouldn'tIf the operation is under the threshold limit will any damage on the surface will lead to catastrophic failure? If NO then what could caused it.

An investigation has been undertaken into the creep behavior of the single-crystal super-alloy CMSX-4. Creep deformation in the alloy occurs largely through dislocation activity in the γ channels. Shearing of the γ′ dislocations is observed, but, at higher temperatures, this does not occur until late in life via the passage of super-partial dislocation pairs. At lower temperatures (1023 K) and high stress levels, shearing of the γ′ precipitates is

observed relatively early in the creep curve through the passage of {111}⟨112⟩ dislocations, which leave super-lattice stacking faults (SSFs) in the precipitates.

The stress-rupture behavior of CMSX-4 has been modeled using a damage-mechanics technique, where the level of damage required to cause failure is defined by the effective stress reaching the material’s ultimate tensile strength (UTS). This technique ensures that short-term rupture data extrapolate back to the UTS. High-temperature steady-state and tertiary creep are modeled using modified damage-mechanics equations, where the strain and damage rates are similar functions of stress. At intermediate operating temperatures of 1023 to 1123 K, the material exhibits pronounced sigmoidal primary creep of up to 4 pct strain, which cannot be modeled using a conventional approach. This transient behavior has been explained by the effect of internal stresses acting on dislocations in the gamma matrix; such an internal stress has been included in the creep law and evolves as a function of the damage-state variable.

9. Investigation and Test Methodology Specimens 50mm by 8mm by 8mm were cut using electrical discharge machining. A notch diameter of 4mm was chosen in order to make it possible to polish the notch. Finite element analysis was carried out to determine the required depth of a 4mm diameter notch to achieve a stress concentration of 2 (typical of stress concentrations at the notch root fixing in service). Notches were then polished to a 1 _m finish using a specially designed rig in order to remove oxide scale left over from machining process. Specimens were tested in 3 point bend using an Instron 8501 servo-hydraulic machine fitted with a high temperature chamber. Low frequency tests (0.25Hz) were conducted at an R ratio of 0.1 and a test temperature of 650°C using a 1–1–1–1 trapezoidal waveform (where ramps up and down and dwells at maximum and minimum load were all of 1 second). Temperature was controlled to within 1°C. A finite element model based on S-N data supplied by ALSTOM was used to identify a strain range in the notch root that would give around 10,000 cycles to failure (a typical service lifetime, and a test-time that would allow replication within reasonable testing timescales). The model used was an elasto-plastic 2D monotonic model using ANSYS finite element software and monotonic CMSX4 material properties supplied by ALSTOM. Results for air tests at 650°C and 725°C are discussed within the results and analysis section of this report. Scanning Electron Microscopy (SEM) of the fractured surfaces was used to identify crack initiation points and determine fracture modes. The SEM was also used to give topographical and compositional scans of the fracture surface. Energy dispersive x-ray (EDX) compositional mapping was conducted on sites of particular interest on the fracture surface using a Jeol JSM-6500F FEG SEM in conjunction with Oxford Inca 300 software. Material left over from fatigue specimen machining was cut into ~8mm square samples. Plain polished samples and polished and etched samples were prepared and exposed at 650°C for 1, 2, 4, 8, 16, 32, 64, 128 and 256 hours. These were then available for examination by SEM after exposure in the furnace at 650°C.

10. Findings

Lifetimes for each test are shown in Table 1. There is no clear difference between orientation A and B over the range of temperatures tested within this limited test matrix.

Temp (°C) Orientation

Pmax

(kN)

(%) Cycles to

Failure650 B 6.

21.38 6500

650 A 6.2

1.38 25,500

725 A 6.2

1.38 5,271

725 B 6.2

1.38 13,717

Crack initiation at high temperatures occurred at sub surface pores in all cases. Cracks in the surface oxide do penetrate the substrate but do not initiate the critical crack. All initiating sub surface pores were encircled by a halo (Figure 2). Subsurface pores were predominantly irregular shapes consistent with interdendritic spacing both in size and shape. The texture of the fracture surface within the halo differs from that seen in the surrounding area. This is better observed using backscattered electron imaging (Figure 3) to look at topographical features on the fracture surface. Using this method, several new crack initiation points were identified. A compositional scan also picked up differences within the halo region compared with the surrounding area.

11. Conclusions and recommendations

Crack initiation in all but one high temperature test occurred at subsurface pores, this agrees with observations in the literature [i, ii, iii]. Such initiation is characterised by a halo around the pore where initial crack growth has occurred in vacuum. The proposed mechanism that causes this halo effect is due to the crack initiating and propagating in vacuum until it breaks the surface of the sample and is exposed to air. Initial crack propagation conditions are that of fatigue crack propagation in vacuum. Once air can enter the crack, the initial fatigue area within the halo undergoes oxidation, after failure has occurred. The boundary of the halo marks the point at which the crack continues to propagate, but now under combined fatigue and oxidation conditions. The two mechanisms described give rise to the change in texture and composition of the oxidised fatigue crack (Figure 2 & Figure 3). Similar halo effects have been seen in polycrystalline disk alloys [iv]. It is not yet fully understood why cracks do not initiate at the notch surface at higher temperatures. There may be sufficient

oxidation of the surface pores to effectively plug them up therefore reducing the stress concentration induced by the pore - similar to oxide induced closure effects. The stress/strain fields below the notch surface may change/redistribute at higher temperatures thus making sub surface initiation more likely. Some initial analytical work has been done to understand the effects of sub surface pore geometry. A simple analysis of pore aspect ratio has been conducted using Scott and Thorpe’s approximation for a semi-elliptical surface crack [v]. A pore area was selected that was representative of those observed to initiate cracks. The ratio of a/c was varied whilst keeping the area of the pore constant. A pore with high aspect ratio with major axis parallel to the notch surface gives the highest value of Kd (stress intensity at the maximum depth position). The secondary orientation of the single crystal will affect the direction in which interdendritic pores form and therefore could affect initiation behaviour if further analysis shows that the angle/orientation of the pore is important.

CONCLUSIONS

In conclusion it was derived that the turbine blades failure was due combination of factors given below they are: Time to time turbine blades were operating outside their thresh hold limits.Uncoated surfaces and poor filtration of airflow was contributing for surface damage due to high temperature corrosion and abrasion. Continues operation of the turbines was contributing cumulative strain holding on the blades while freezing stresses in the matrix.

With these findings in the hand the manufacturer was directed to establish the best-fit solution for the present problems.

Initiation of fatigue cracks is affected by critical pore size, angle, position and aspect ratio. The angle of the pores may be a direct effect of the secondary orientation of the specimen as pores are normally interdendritic.

Coalescence of the fatigue cracks depends on initial distribution of initiating features (porosity). Pores have also been seen to cause temporary crack arrest which can add to the specimen lifetime.

FURTHER WORK What was our recommendations :

Limit the operational loading limit 15% under the fatigue thresh hold.  Selecting the right coating for application

 Introduce controlled intervals between maintenance.  Regular Condition assessment.  Dynamic Testing and Monitoring  Corrosion Control and Environmental Chemistry Monitoring  Remaining life prediction – Routine lifetime evaluations based upon condition.

Further work relevant to this report consists of: sectioning and observation in SEM of some current test specimens and oxidation study specimens, a literature search for Oxidation of CMSX4. Also a further set of run out tests (650°C orientation A and B) on sub size CMSX-4 test specimens has been planned in order to generate S-N data and more porosity data from fracture surfaces.