failure analysis of a first stage high pressure turbine ... · failure analysis of a first stage...

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Failure Analysis of a First Stage High Pressure Turbine Blade in an Aero Engine Turbine on PK-GSG Boeing B747-400

Arif Sugianto1, Reza Jaya Wardhana2, Nanang Yulian1, I Gede Kusuma Jaya Wardana3, Muchtar Karokaro4, Hariyati Purwaningsih4

1 Dept of Aircraft Engineering, Research and Development Material Process - Engineering Services GMF Aero Asia, Soekarno-Hatta Intl Airport, Tangerang 19103, Banten, Indonesia

Phone: +62-21-550-8008 & 8188, Fax: +62-21-550-1282, E-mail: [email protected] & [email protected] 2 Dept. of Materials and Metallurgical Engineering, Faculty of Industrial Technology

Institut Teknologi Sepuluh Nopember (ITS), Surabaya 60111, East Java, Indonesia Phone:+62-81-809808908, E-mail: [email protected]

3Dept. of Powerplant Engineering Program Garuda Indonesia, Soekarno-Hatta Intl Airport, Tangerang 19103, Banten, Indonesia

Phone: +62-21-550-8172, E-mail:[email protected] 4 Dept. of Materials and Metallurgical Engineering, Faculty of Industrial Technology

Institut Teknologi Sepuluh Nopember (ITS), Surabaya 60111, East Java, Indonesia E-mail:[email protected]

Abstract

The failure of a First Stage High Pressure Turbine (HPT) Blade in an Aero Engine Turbine was investigated by metallurgical investigation and stress analysis of the failed blade. The blade was made of a nickel-based superalloy, Columnar Grain Directionally Solidified Rene 142 (DSR 142). The turbine engine has been in service for about 60382 hours and the failure was known after maintenance on 6 January 2009. Due to the blade failure, the turbine engine was damaged severely. The investigation was started with a thorough visual inspection of the blade surfaces followed by the fractography of the fracture surfaces, microstructural investigations, and chemical analysis. The blades experience internal cooling hole cracks in different airfoil sections assisted by a coating and base alloy degradation due to operation at high temperature. A detailed analysis of all elements which had an influence on the failure initiation was carried out, namely: loss of aluminum from coating due to oxidation and coating phases changing; decreasing of alloy ductility and toughness due to carbides precipitation in grain boundaries; degradation of the alloy gamma prime (γ’) phase (aging and coarsening); blade airfoil stress level; evidence of intergranular creep crack propagation.

It was found that the coating/substrate crack initiation and propagation was driven by a mixed fatigue/creep mechanism. The coating degradation facilitates the crack initiation due to thermo-mechanical fatigue. The intergranular crack initiation and propagation in the substrate were due to a creep mechanism which was facilitated by grain boundary brittleness caused by formation of a continuous film of carbides on grain boundaries, the degradation of γ’ due to elongation (rafting) and coalescence, and high thermo-mechanical stress level.

1. Introduction Aero engine turbine blades are made of nickel-base and cobalt base superalloys principally. During the operation of aero engine turbine, the blades and other elements of hot gas path suffer service induced degradation which may be natural or accelerated due to different causes. The degradation or damage may have a metallurgical or mechanical origin and results in reduction of equipment reliability and availability.

It also increases risk of failure occurring. Also, due to metallurgical deterioration of the blade material, the material creep, fatigue, impact and corrosion properties decrease. There are different factors

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which influence blade lifetime as design and operation conditions but the latter are more critical. Generally speaking, most blades have severe operation conditions characterized by the following factors: Operation environment (high temperature, fuel and air contamination, solid particles, etc.). High mechanical stresses (due to centrifugal force, vibratory and flexural stresses, etc.). High thermal stresses (due to thermal gradients).

The type of damage which occurs in gas turbine blades and nozzles after a service period can be divided into: External and internal surfaces damage (corrosion, oxidation, crack formation, erosion, and

foreign object damage). Internal damage of microstructure as γ‟ [Ni3(Al,Ti)] phase aging (rafting), grain growth, grain

boundary creep voiding, carbides precipitation and brittle phases formation[1]. Surface damage produces blades/nozzles dimensional changes which result in operational stress

increase and turbine efficiency deterioration. In service, blade material deterioration is related to the high gas temperature, high steady state load levels (centrifugal load) and high thermal transients loads (trips, start-ups and slowing downs). However, the degree of deterioration in individual blades differs due to several factors such as: Total service time and operation history (number of start-ups, shut-downs and trips). Engine operational conditions (temperature, rotational speed, mode of operation (base load, cyclic

duty). Manufacturing differences (grain size, porosity, alloy composition, heat treatment). The DSR 142 alloy commonly used for gas turbine blades is strengthened by precipitation of γ‟

phase. The microstructural changes due to blade operation at high temperature include irregular growing of γ‟ particles (rafting) and formation of carbides in grain boundaries and matrix[2]. This leads to alloy creep properties reduction[3]. In order to have an instrument for the deterioration evaluation of gas turbine blade alloy, it is necessary to correlate the influence of service induced microstructural degradation to the change in mechanical properties.

2. Background

Report of „CF6-80C2B1F HPT Stage 1 Blade Project’[5] by Kusuma Jaya Wardana (GA) and Dedek Zuldin (GMF), has known that between 1995 and 2009, Garuda has had 9 UER‟s due to HPT Stage 1 Blade defects. HPT Module has contributed approx. 32.4% total engine removals (Fig.1) and approx. 43% of HPT Module caused engine removals are driven by HPT Stage 1 Blade liberation. Preliminary on-site investigation concluded that first stage of HPT Blade, were the physical causes of those engine failures. It was reported that visual and Borescope Inspection performed during the last overhaul did not detect any defects on the blade surfaces. This present investigation is aiming to determine the failure cause and to explain why fatigue fracture occurred without any defects on the turbine blade surfaces in order to reduce the number of UER‟s caused by HPT Stage 1 Blade defects and evaluation of extent and degree of material damage, lifetime consumed to obtain recommendations for blade rejuvenation treatments, operation and reposition. Application of effective methods of material deterioration evaluation can be used for practical lifetime prediction, just in-time blade rehabilitation (rejuvenation), safe and cost-effective lifetime extension and to avoid blade catastrophic failure. The blade under evaluation was the first stage blade of a High Pressure Turbine (HPT) blade from ESN 704136 with P/N 1538M90P12, S/N PCMILU4S for CF6-80C2 #1, PK – GSG Boeing 747 - 400

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with gas inlet temperature of 893oC. The evaluation was carried out after 60382 hour and 10536 cyclic period. The Blade is made of nickel-base “Rene 142” superalloy by means of columnar grain directionally solidified (CGDS) polycrystalline investment casting, coated by Thermal Barrier Coating (TBC) over Pt-Al coating[6]. The location of HPT Blades stage 1 at turbine engine is shown in Fig.2.

Figure 1 a) Percentage Garuda Airways Engine Removal b) Percentage Garuda Airways HPT Module Engine defect

Fig 2 Location of HPT Blades Stage 1 at Turbine Engine[7]

HPT Rotor Disk

Diffuser Assy Air Seal

Ring Piston

Turbine Engine

Seal Ring

Seal Ring

Disk Stage - 1

Ring-Rotor Stage 1 Disk

HPT Blade Stage 1 HPT Blade Stage 2

Nut Self Locking

Cup Windage

Disk Stage 1 Blade Assy-Stage 1

Damper Assy-Stage 1

Seal Body Stage - 1

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3. Experimental Procedure

In addition to the failed blade, two adjacent blades were subjected to the laboratory for extensive failure analysis. The following sequences of examination were performed on the blades:

1. Visual examination and photographic the documentation 2. Examination of the fracture surface by means of optical and electron microscopes 3. Chemical analysis of the blade material 4. Metallographic examination of the cross-section airfoil beneath the fracture surface area 5. Profile measuring for load identification

4. Experimental Results

4.1 Visual Examination

Figure 3a. shows the comparison between complete blade and failure blade. As a result of failure blade, nearly 1/3 airfoil area has been damaged from leading edge until trailing edge region. Some coating has peeled from surface area. On these fracture surfaces (Fig. 3b), it was readily distinguished that there are two type surface areas i.e. flat shiny and dark fibrous blur areas and those can be interpreted as fatigue fracture and static final fracture areas, respectively. For the static final fractured, was seen in cavity 2, 7, and 10 that looked dark (Fig. 3b). This can be interpreted that at a time just before final fracture, the blade were operated at relatively high stress since the remaining area of static final fracture is comparably large [8].

Figure 3. a) Complete Blade and Failed Blade b) Fractured area of failed blade

4.2 Analysis of Fracture Surface

Fracture surfaces of the failed blade was examined by stereomicroscope using Wild Heerburg and SEM using JEOL JSM-6510LA. General view of the fracture surface is shown in figure 3b. The fracture surfaces have a dendritic appearance without evidence of fatigue features. It is suggested that these fractures were formed by fatigue that washed away from thermo-mechanical action and then finally ruptured. The crack initiation is indicated on concave area at 6th cavity (Fig.4a and 4b). From SEM

a) b) Flat Shiny

Peeled Coating

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photograph (Fig.5), crack started from inside cavity wall. It indicates that after initiation of the crack, propagation has been conducted by high mechanical stresses due to the vibration of the blade. It was shown that crack propagation appeared after coating was peeled. Under the action of bending stress, spin centrifugal tensile stress, and the overheat damage, the microcracks propagated and converged so that the transverse fracture occurred at the mid-airfoil location where the maximum stress acted [9]. When the crack has propagated so far, so that remaining cross-sectional area could no longer withstand the loads, then the blade will have a static fracture (final fracture) in the leading edge and trailing edge [10].

Figure 4. Optical Stereomicroscope showing crack at concave area a) 6x magnification b) 25x magnification c) SEM Morphology of turbine blade top surface, the crack initiation from inner cavity and propagate to the outside

A similar incident occurred on the convex area (Fig 6a). Initial crack with direction of propagation look like axially, indicates that crack grow transgranularly (Fig 6b). From SEM photograph (Fig.7), it was indicated that crack occurred after catastrophic fracture. From Overhaul Manual Data, it is informed that this blade has found in a state of missing material when doing maintenance on the aircraft, so that it can be concluded that the turbine blade is still operating despite being in a damaged condition. The loads is still working despite a turbine blade has been damaged due to airfoil area and is not distributed evenly on the airfoil area but directly on cavity cooling holes.

Crack Initiation

a) b)

c)

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Figure 6. Optical Stereo Microscope showing crack at convex area. Propagation look like axially, indicates that crack grow transgranularly a) 6x magnifier b) 25x magnification c) SEM micrograph of fracture surface at convex area, showing the location of crack initiation

4.3 Chemical Analysis

To determine the chemical composition of the blade material a bulk analysis was performed by an Optical Emission Spectroscopy using Foundry Master. The result of chemical analysis showing that material of blade is accordance with Rene 142. Table 1 Chemical composition of Nickel-Based Superalloy Rene 142 (%wt)

Ni Based (Rene 142) Ni Mn Cr Mo Cu Fe Co Ti Al Nb W V Si Rc Ta Zr Hf

OES Test 60 0.04 7.68 1.4 0.07 0.117 11.1 0.05 6.19 0.01 5.88 0.03 0.16 2.6 3.93 0.01 0.7 ASM International Bal 0.03 7.8 1.5 … … 12 0.04 6.15 0.01 4.9 0.02 0.15 1.8 4.1 0.02 1.5

Crack Propagation

Crack Initiation

a) b)

c)

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4.4 Metallographic Examination

Sample for microstructural evaluation were taken from the cross-section airfoil beneath the fracture surface area. Sample were observed by using Nikon Clemex Scan Micrograph. The microstructure photograph was taken under unetched and electro-etched with Buehler Electromet Etcher per ASTM E 407 by a solution of 5 ml H2SO4, 8 gr Cr2O3, 85 ml H2SO4 and electrolytic at 10 V for 5-30s for reveals γ‟. Fig. 7 illustrates the micrograph of the material. As to be observed the microstructure has composed of dendritic grains, which is a characteristic of cast structures.

Figure 7. Optical micrograph of blade with a) unetched and b) electroetched condition, 50x magnifition, show dendritic microstructure. 4.4.1 Coating Evaluation

From SEM-EDS and XRD results, it can be obsreved that this system consist of an insulating ceramic outer layer (top coat) and metallic inner layer (bond coat) between the ceramic and the substrate. The blade airfoil has Thermal Barrier Coating (TBC) as top coat and oxidation resistant as bond coat for hot corrosion and oxidation resistance. This TBC have characteristic sufficiently thick, and have a low thermal conductivity, high thermal shock resistance, and high concentration of thermal expansion[11]. . Current TBCs are yttria-stabilized zirconia (YSZ), that is, zirconium oxide (ZrO2-zirconia) with 6-8% wt of yttrium oxide (Y2O3-yttria) to partially stabilize the ZrO2 for good strength, fracture toughness, and resistance to thermal cycling and Pt – Al coating as bond coat. The metallic inner layer (Pt-Al bond coat) aids in the adhesion of the ceramic top coat, protects the substrate from hot corrosion and oxidation, and helps in handling expansion mismatch between the ceramic and superalloy[12]. This coating is formed by the diffusion interaction of aluminium with surfaces of the nickel base superalloy. Bond coat region consists of three zone (Fig.8): External zone of PtAl2 Intermediate zone of NiAl or β Internal or Interdiffusion Zone

A TBC typified by the structure and composition is illustrated in figure 8. The microstructure of top coat consists of columnar grains that provide for strains tolerance, can be concluded that the technology of TBC‟s, which used for this blade is using Physical Vapour Deposition (PVD)[11], with 129 μm thick which provides oxidation protection for the substrate and forms an adherent Al2O3 scale to anchor the zirconia top coat.

a) b)

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Fig.8. Microstructure of typical TBC with a 129 μm top coat thick PVD applied ZrO2-8% Y2O3 and Pt-Al coating with a 96 μm bond coat thick by diffusion processing From the analysis results of EDS in figure 8 and 9a, it is observed that in intermediate zone of NiAl or β mostly in the form of oxide and Al, then there is Ni, and some other small form of Co, Cr, and Zr. There is also element of S, which is a sign of hot corrosion. Meanwhile, for the interdiffusion zone, composed mainly of Ni and Al, and some other small form of Co, Cr, Pt, and Zr. For the substrate (base metal), most of the compositions are Ni, Al, Cr, and Co. There was also C element that may be due to the influence of the diamond compound for polishing process. For Ni-Based Superalloy, the level of oxidation resistance at temperatures below about 1600 to 1800o F (871 to 982o C) is a function of chromium content (forms as a protective oxide); at temperatures above about 1800oF (982oC)[13], aluminum content becomes more important in oxidation resistance (Al2O3 forms as protective oxide). EDS observations in figure 9, shows that the content of Al will tend to increase in dense inner layer zone , while for the content of Cr seems likely to decline. This shows that element of Al diffuses outward to form the oxide layer and distributed to the base metal. While elements of Cr and Co in the base metal, distributed toward the surface, it can be known observed from the elements Cr and Co on the external and internal intermediate zone of NiAl or β. Interdiffusion zone (fig. 9e), consist of refractory metal (iron, molybdenum, tantalum,etc) carbides and/or complex intermetallic phase in a NiAl and/or Ni3Al matrix (Fig. 9e), formed by the removal of nickel form the underlaying alloy, thereby converting its Ni-Ni3Al structure to those phase[14]. The deposition of platinum by electroplating is a simple and well established process but deposits are usually hard, may be highly stressed and are particularly prone to porosity (Fig. 8 & 9d). Additionally, poor cleaning procedures prior to electroplating can cause entrapment of nonmetallic deposites such as sulfur and zircon, leading to poor adhesion and incomplete diffusion of aluminum and platinum during the aluminizing process[15].

Insulative Ceramic

Top coat

(Yttria-Stabilized

Zirconia)

Oxidation Resistant Bond

Coat (Pt –Al Coating)

Interdiffusion Zone

Dense Inner Layer (ZrO2)

129 μm

96 μm

Base Metal Porosity

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Element Mass% O (K) 42.11 Al (K) 5.13 Cr (K) 1.027 Co (K) 1.68 Ni (K) 8.83 Zr (L) 41.28 Total 100

Element Mass% O (K) 43.17 Al (K) 26.96 Cr (K) 1.1 Co (K) 1.8 Ni (K) 10.14 Zr (L) 16.68 Total 100

a)

b) c)

d) e)

a)

b)

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Fig 9. Schematic of EDS coating and ZAF method standardless quantitative analysis a) TBC coating b) Dense Inner Layer c) Pt-Al coating d) Porosity in the Pt-Al Coating e) Interdiffusion Zone On the top of interdiffusion layer appears a topological closed packed (TCP) (Fig. 10), which can reduce the mechanical properties of the substrate. The emergence of TCP is likely caused by oxidation processes that cause changes in chemical composition to the substrate, which raised a new phase, and

Element Mass% O (K) 15.41 Al (K) 8 S (K) 0.72 Cr (K) 2.42 Co (K) 5.87 Ni (K) 33.25 Zr (L) 9.98 Pt (M) 24.34 Total 100

Element Mass% O (K) 30.13 Al (K) 16.53 S (K) 0.7 Cr (K) 1.98 Fe (K) 0.34 Co (K) 3.164 Ni (K) 18.34 Zr (L) 12.93 Pt (M) 13.74 Tl (M) 2.15 Total 100

Element Mass% O (K) 24 Al (K) 9.65 Cr (K) 7.18 Co (K) 9.1 Ni (K) 31.4 Zr (L) 4.55

Mo (M) 2.48 Pt (M) 11.65 Total 100

c)

d)

e)

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indeed in the super alloys of Ni and Co, both of which are meta-stable alloys, more easily arise TCP phase and this is one of the weaknesses of the two alloys[16]. On the internal part of intermediate zone of NiAl, phase changes from β NiAl to β + γ‟(Ni,Pt)3Al.

Fig 10. TCP phase at the boundary between coating and base metal

4.4.2 Crack Evaluation

Coatings must be capable of tolerating strain due to thermal expansion mismatch and mechanical loads, in order to retain coating-substrate integrity and as low as possible so that cracking in the coating does not occur in service, because the cracks may then propagate into the substrate. Fig. 11a&b, show coating deterioration in internal cooling holes of the airfoil. The occurrence of cracks on the coating, constituent due to oxidation can be seen. Oxidation that occurs is a kind of internal oxidation in which taking place below the external surface but not just at to substrate area[17]. In this particular case of the Pt-Al coating, it occurs due to loss of aluminium from the coating to form protective Al2O3 on the surface. In turn the coating provides additional aluminium, to form new aluminium oxide. According to the aluminum content (also Pt) reduction, the coating is consumed and the phase change as follows (fig 11): PtAl2 particles present in external surface layer dissolve to NiAl to form singular phase (Ni,Pt) Al

or β phase This phase decomposes to biphase structure γ‟ (Ni,Pt)3Al+β (TCP phase) Because aluminum is consumed, the β phase cannot survive and only γ‟ exists This decomposes to γ and γ‟ (Ni and other elements into solid solution) Finally only γ exists and coating is consumed completely

Practically, coating protects base metal when it is formed by only singular phase β and loss its protective characteristics when it transforms to a biphase structure γ‟+β. This stage is optimum for coating restitution (stripping deteriorated coating and recoating). In the case of a gas turbine fuelled by natural gas, the coating deterioration occurs due to oxidation in the blade airfoil hottest zones, which typically may be leading edge or other zones depending on blade design. There were found many different airfoil zones with different degrees of coating deterioration[18].

TCP Phase

Porosity

Base Metal

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Fig 11. Oxidation cracks on the coating that occurs in internal cooling holes area; a) Optical microscope with 100x

magnification on the upper side wall at 6th cavity (unetched); b) Using SEM in the same location; c) Optical microscope with 200x magnification on the bottom side wall at 4th cavity (unetched); d) Using SEM in the same location

This event led to the emergence of grain boundary crack initiation and propagation in the substrate. The crack in the internal cooling holes of the blade was detected (Fig.12). Cracks initiate in the cooling hole coating and propagate into the substrate following grain boundary trajectories. The crack size reaches 0.2 mm. On the basis of these evidences, it was evaluated that the crack initiation propagation was derived by a mixed fatigue/creep mechanism. The coating crack initiation was probably due to thermal fatigue mechanism as a result of high thermal transient loads (trips, start-ups, and slow-downs), and crack grain boundary initiation and propagation in the substrate by a creep mechanism due to high temperature gradients across the airfoil wall (between vane external surface and surface of internal cooling holes) Crack initiation in internal cooling holes has the following sequence:

Coating degradation and cracking due probably to a thermal fatigue mechanism and environmental attack

Crack initiation in substrate Generally speaking is that the degradation of coatings resistant to high temperature proceeds in two ways: Loss of coating constituents due to oxidation and corrosion, which results in loss of Al from

coating to create a thin protective layer of Al2O3 on the surface Interdiffusion of coating-substrate, which results in diffusion of alloy elements into the coating

[18]

a) b)

c) d) PtAl2

NiAl

Interdiffusion

Oxidation Crack

TCP Phase

Internal Cooling Hole

Internal Cooling Hole

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Fig 12. Thermomechanical fatigue crack in the cooling holes area a) Unetched b) Etched c) With SEM

In Fig.13, the microstructure can be revealed after electro etching. The microstructure consists of grains of γ and carbide particles precipitated in the matrix and grain boundaries. In grain boundaries was found a continuous film of carbides of 1.5-3 μm thickness. The presence of a continuous film of carbides, is a result of transformation of carbides of MC type to carbides of M23C6 type due to high temperature operation of the blade. This dense and continuous network of carbides reduces ductility and toughness of the alloy to 30% of initial value and facilitates crack initiation and propagation which leads to reduced lifetime [18]. The degradation of γ‟ due to elongation (rafting) and coalescence (growing) originates a reduced alloy creep lifetime. The predicted airfoil maximum stress location on the internal cooling hole surface is consistent with the crack location. Substrate crack initiation and propagation is facilitated also due to grain boundary brittleness caused by the formation of a grain boundary continuous film of carbides as mentioned before. From figure 13c & d, it can be seen that an undamaged structure of γ‟ quasi cuboids has grown into a little rafted γ‟ resulting from exposure to stress at high temperature [18]. It should be that noted that voids on the grain boundaries were observed in figure 13, and grain boundaries in some locations were found open.

a) b)

c)

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Fig.13 Microstructure nearby the crack area; a) Show formation of grain boundary continous film; b) Void in grain boundary area

c) Solutioned and re-precipitated γ‟, degradated carbide and melting boundary; d) Gamma prime (γ‟) phase morphology The magnified morphology reveals localized melting of the boundary and degradation of interdendritic carbides. The microstructural degradation served as subsurface crack in blade is kind of

Melting Boundary

Void

M23C6

MC

γ+ γ’ eutectic

γ+ γ’ eutectic

a) b)

c)

d)

Grain boundary pattern

γ’

Solutioned and re-precipitated γ‟

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decohesion between carbide and its surrounding matrix. The interface decohesion was formerly caused by precipitation, lengthening and coarsening of carbides and this can be interpreted that blade was often operated at high temperature for a relatively long time under high stress[8].

The microstructural degradation phenomena during prolonged time exposures to high temperature were similar to those observed by Floyd et[19]. It is suggested that the operating temperature on the located regions of the failed blade had exceeded the specified maximum temperatures of operation. The instantaneous operating temperature in the vicinity of the solvus temperature of γ‟ phase or even over the solidus temperature will cause overheat damage. 4.5 Profile Measuring

The HPT Blade, which is in good condition, is taken for profile measurement using Hexagon Metrology Bridge CMM Machine Series 12-15-10. Measurement for this profile was conducted by defining point by point along the cross-section area of airfoil and used movement-gap of the probe tip 1 inch along the lengthwise area of airfoil (fig 14 a&b). The results for this measurement are coordinate points for 3D axes (x, y, z axes). The airfoil cross-sectional profile can be obtained by converting coordinate points into 2D curve. Integral method is used to obtain the equation of cross sectional area. By means of Microsoft excel, 2D-curve profile approach of airfoil cross-sectional area was obtained, as shown in Fig.15.

Figure 14. a) CMM Machine b) Location of CMM measurement in radii-dependent HPT blade on un-ruptured blade

Scan Line #1

Scan Line #2

Scan Line #3

Scan Line #4

Scan Line #5

1 inch

a)

b)

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Fig. 15. Profile Approach of Airfoil Cross-Sectional Area

Utilizing fourth exponential equation, the function of airfoil cross-section area is:

4.6 Analytical Calculations of Centrifugal Forces

One of the most critical loads in rotating component of gas turbine is the centrifugal force, which generates tensile stress along the blade length in the radial direction. It depends on two variable parameters; the whirling speed of rotor and the distance of each position from the rotating axis.

Basic Equation :

....... Eq. (1)

Where: σ(r) : Tensile stress (N / m2) F(r) : Centrifugal Force (N) S(r) : Airfoil cross-section area (m2)

S(ξ) = - 459.7x3 + 941x2- 643.3x + 146.12

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Determine the centrifugal force (F(r) ):

Fig. 16 A moving blade that is detailed for selecting a radial element[20]

As shown in fig.16, the applied centrifugal force in blade of turbine effected by centrifugal force due to beanding ( Fb ), and centrifugal force due to portion of the blade confined between the section r-Rh and Rt ( Fbl(r) ). So the basic equation for determining the centrifugal force is:

...... Eq.(2)

Centrifugal Force due to bending :

Where : m : blade mass (kg) = 0.2765 kg n : rotation of axial impeller = 10540 rpm ω : angular velocity of rotor (rad/s) = (2π x 10540)/60 = 1103.1 rad/sec Rtip : tip radius (m) = 0.106 m δ : bending thickness (mm) = 0.0127 m

Centrifugal Force due to portion of the blade confined between r-Rh and Rt : By taking an infinitesimal, element dξ is separated into section ξ, the force Fbl(r) can be found as:

...... Eq. (4)

Where: l = Rtip - Rhub = (0.106– 0.044) m = 0.062 m S(ξ) = Function of airfoil cross-section area

ρ = 13825 kg/m3 Rhub = 0.044 m

Fb= mb ω2(RTip+δ/2)

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As follows equation [2], the total centrifugal force is :

.... Eq.(5)

Finally by substituting the calculated functions in equation [1], the functions of tensile stress will be resulted as follows:

σr = 37806.708 + (1.55 1012r5 – 4.2 1012r4 + 4.08 1012r3 – 1.47 1012r2 – 11035450r + 1.23 1010)

- 459.7(r-0.044)3 + 941(r-0.044)2 – 643.3(r-0.044) + 146.12

The plot of centrifugal stress versus rotating radius is shown in Fig.17. By attention the position of fracture on the broken blade, the stress value is about 36.98 MPa

Fig 17. Graph of tensile stress distributed along the blade

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4.7 Cycle Analysis

The turbine blade is a component that works on a fairly extreme conditions, especially for first stage of HPT Turbine Blade. This component is located after the combustion chamber, resulting in pressure and temperature that high enough due to process of gas combustion products that would affect the work of this component. To determine the size of pressure and temperature, it is necessary to perform calculations with the cycle analysis. Cycle analysis studies the thermodynamic behavior of air as it flows through an engine, without regard for the mechanical means used to affect its motion. From Fig. 28, it is known that the CF6-80C2 engine is a type of Turbo Fan engine, with its main components consist of:

Fan Compressor Combustor Turbine Nozzle

Figure 18. Schematic of CF6-80C2 Turbofan Engine[21]

Compressor Combustor

Fan Inlet

HPT Turbine

Blade

Compressor Outlet

Compressor Inlet

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By attention to fig 18, the value of Pressure and temperature at HPT Turbine Blades, can be estimated from the following equation[22]:

Fan

…Equation 6

To get the value of V, then calculated with equation:

M = Mach

a = Speed of Sound (m/s)

Inlet Formula Outlet Ta, Pa

T0if = Ta+ V2

1111

ff

f

T0ef = 𝛕fToif

M / V

P0if = Pa

P0ef = πfP0if

General Electric CF6-80C2 Turbofan Engine, with:

Overall compressor pressure ratio = 30:1 Fan pressure ratio (πf) = 1.5 Compressor pressure ratio (πc) = 19 Combustion Chamber pressure ratio (πb) = 1 Ambient Temperature (Ta) = 298 K (Room

Temperature) γ = 1.4 R = 287 m2/s2K Pa = 101325 Pa (ISA Standard) cp = 1005 J/kgK (ISA Standard) M = 0.279 ηf = 100% ηc = 100% T0b = 1167 K

Assumption:

Engine work under ideal conditions, f <<1, so that losses (Δp loss) due to the combustion process is almost close to 0

By attention to fig x, the value of Pressure and temperature

at HPT Turbine Blades, can be estimated from the following

equation:

2 cp

T0if

Ta

γ

γ - 1

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With a:

As follows equation [6], the value of T0if, T0ef, P0if, and P0ef are:

T0if (K)

T0ef (K)

P0if (Pa)

P0ef (Pa)

302.63 371.63 106943 160414.2

Compressor

Inlet Formula Outlet

As follows equation [7], the value of T0ic, T0ec, P0ic, and P0ec are:

T0ic

(K) T0ec

(K)

P0ic

(Pa) P0ec

(Pa) 371.63 861.89 160414.2 3047869.8

Combustion Chamber

Inlet Formula Outlet

As follows equation [8], the value of T0ib, T0eb, P0ib, and P0eb are:

T0ib

(K) T0eb

(K) P0ib

(Pa) P0eb

(Pa) 861.89 1167 3047869.8 3047869.8 Therefore Toeb=Tit and Poeb=Poit, can be concluded that temperature was suffered at HPT Blade is 1167 K (894oC) with pressure 3047869.8 Pa (442 psi).

efic TT 00

1111

cc

ciccec TT 00

efic pp 00 iccec pp 00 … Equation 7

ecib TT 00 beb TT 00

ecib pp 00 lossb p 1 ibbeb pp 00 … Equation 8

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5. Conclusions

This failure analysis study is based on evidence of metallurgical investigation and stress analysis. The following is concluding remarks:

The Blade is made of nickel-base “Rene 142” superalloy by means of columnar grain directionally solidified (CGDS) polycrystalline investment casting, coated by Thermal Barrier Coating (TBC) over Pt-Al coating.

For visual examination on fracture surfaces, it was readily distinguished that there are two type surface areas i.e. flat shiny and dark fibrous blur areas and those can be interpreted as fatigue fracture and static final fracture areas[8].

Analysis of the fracture surface from SEM photograph, crack started from inside cavity wall. It indicates that after initiation of the crack, propagation has been conducted by high mechanical stresses due to the vibration of the blade. It was shown that crack propagation appeared after coating was peeled. Under the action of bending stress, spin centrifugal tensile stress, and the overheat damage, the microcracks propagated and converged so that the transverse fracture occurred at the mid-airfoil location where the maximum stress acted[9].

Crack was found in the cooling hole to 0,2 mm deep. On the basis of these evidences, it was evaluated that the crack initiation /propagation was derived by a mixed fatigue/creep mechanism. The coating crack initiation was probably due to thermal fatigue mechanism as a result of high thermal transient loads (trips, start-ups, and slow-downs), and crack grain boundary initiation and propagation in the substrate by a creep mechanism due to high temperature gradients across the airfoil wall (between vane external surface and surface of internal cooling holes).

Crack initiation in internal cooling holes has the following sequence: Coating degradation and cracking due probably to a thermal fatigue mechanism and

environmental attack Crack initiation in substrate

The presence of a continuous film of carbide, is a result of transformation of carbides of MC type to carbides of M23C6 type due to high temperature operation of the blade. This dense and continuous network of carbides reduces ductility and toughness of the alloy to 30% of initial value and facilitates crack initiation and propagation which leads to reduced lifetime[18].

The magnified morphology reveals localized melting of the boundary and degradation of interdendritic carbides. The microstructural degradation served as subsurface crack in blade is kind of decohesion between carbide and its surrounding matrix. The interface decohesion was formerly caused by precipitation, lengthening, appearance of the located melt boundary, remelt eutectic phase (γ+γ‟) in the mid-airfoil regions of the failed blades and coarsening of carbides can be interpreted that the blade was often operated at high temperature for a relatively long time under high stress[8].

Overheating reduced the strength of the interdendritic region so that microcracking occurred in that area. Under the action of bending stress and the centrifugal tensile stress, the microcracks propagated within the interdendritic regions and formed tortuous short crack on the airfoil surface. The neighbored short crack with identical orientation converged to form a long crack, and the transverse fracture occurred at the mid – airfoil location. The fragments from the previously fractured blade would hit the neighbored blades to lead to the damage of the turbodisk.

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6. Recommendations Recommendation for preventing failure of the High Pressure Turbine (HPT) Blade suggests: The cracks penetrate the coating and substrate significantly in highly stressed areas (airfoil), it

can be concluded that the blade lifetime was consumed and it is not possible to apply a repair process (recoating, rejuvenation heat treatment, etc.) to restore blade original characteristics and extend lifetime[18], so it must be replaced with new ones.

Make possible refurbishing and lifetime extending blades should be retired from service before cracks initiate in the substrate.

To accomplish blade lifetime extension, the blade cooling system should be improved to prevent failures by reducing the airfoil thermal gradients, minimizing the airfoil thermal stresses.

Report of „CF6-80C2B1F HPT Stage 1 Blade Project’ by Kusuma Jaya Wardana (GA) and Dedek Zuldin (GMF), has known that the mean life of this blade is range about 11500-19400 h, It is therefore after 15000 of service life of TSN, the blade is better to be discarded. There were same existed mechanism in failure HPT Blade due to operation condition, which occurred coating peeled and continued with crack initiation in concave area especially in mid airfoil region, so that it is required special attention and maintenance in those area.

Before coating process is carried out, as good as possible carried out by cleaning process at substrate area, because poor cleaning procedures prior to electroplating Pt-Al can cause entrapment of nonmetallic deposites that will cause leading to poor adhesion and incomplete diffusion of aluminum and platinum during the aluminizing process [15].

Acknowledgement

The authors acknowledge Mrs. Dede and Mr. Bayu Aji for their support and laboratory practice in GMF AeroAsia. Mr. Arby for consultation about SEM HPT Blade. Mr. Anni Rahmat for modeling support.

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