Special features of the microstructure of the surface of 'fir tree' locks of the single crystal blades of ZhS36VI alloy.
The chemical and phase composition of the alloys have a strong effect on the nucleation of fatigue cracks and their propagation resistance, and also the absolute values of the endurance limit [[sigma] sub.-1] (1). The movement of the compositions and structure of the creep-resisting alloys, improvement of the level of the static characteristics of the long-term and short-term strength help in increasing the resistance to high cycle fatigue of the material. As reported in [2, 3], the behaviour of metals under cycling loading is determined by their microstructural and macrostructural special features.
The main areas of the initiation of fatigue microcracks in the alloys are the particles of carbides and borides and micropores. In carbon-free single crystal palace, for example, ZhS 36 and ZhS40, the amount of the carbide phase is extremely small. Therefore, the fatigue limit of these alloys at room temperature is higher than in the carbon alloys with the equiaxed structure (4). However, in comparison with these alloys, the single crystal alloys are characterised by the high sensitivity to stress concentration, especially the single crystal carbon-containing alloys of the type ZhS6F and ZhS32, whose stress concentration coefficient 11smooth1notch/K[sub.[sigma]-1]=[sigma][ sub.-1smooth] / [[sigma] sub.-1notch] ([sub.[sigma]-1]) [sub [sigma]-1]- the endurance limit of the respectively smooth (no notch) and notched specimens) is 2.5-3.0 at 20[degrees]C, whereas for the carbon-free single crystal alloys (ZhS36 and ZhS40) K[ sub.[sigma]-1] = 1.5-1.6 (1). The common feature of all the alloys is the reduction of the sensitivity to stress concentration at elevated temperatures. [sigma]
The fatigue resistance of the single crystal creep-resisting alloys is determined by their crystallographic orientation. For example, at 20[degrees]C, the  orientation for all the alloys is stronger than the  orientation. However, the fatigue limit of ZhS36 alloy at a temperature of 1100[degrees]C is the same for both orientations. Consequently, the anisotropy of the values of [sigma]-1 for the single crystal creep resisting alloys at high temperatures is eliminated as in the case of the short-term and long term static tests. The single crystals, cast in the furnaces with high thermal gradient, are characterised by the higher values of [sigma]-1 because they have a dispersed dendritic structure contains a smaller number of micropores .
The high sensitivity of the single crystals to stress concentration is attributed to the absence of the grain boundaries which is to some extent the obstacle in the path of propagation of the main fatigue crack. In the polycrystalline alloys, the crack changes the direction of propagation every time when a grain boundary is intersected. This determines the increase of the length of the trajectory of the fatigue cracks.
In the notched single crystals, the fatigue crack propagates almost without obstacles throughout the entire specimen. Therefore, a single crystal blade is characterised by high sensitivity to stress concentration, for example, casting defects.
In the working blades of the gas turbine engines, the 'fir tree' tailpiece is also an important part in comparison with the blade. In particular, the lock sections, operating at temperatures not higher than 600-750[degrees]C, carried the highest load and are subjected mainly to fatigue failure. For example, in stand tests of the working blades of ZhS 36VI alloy the regions of nucleation of the fatigue cracks were found in the area of exit of the depression of the first tooth to the end part of the tailpiece.
The surface layer of the tailpiece of the blade plays a significant role in the formation of the load carrying capacity of the components . Fatigue failure of the turbine blades amounts to 33% of the total range of failures of this type in the components of the aviation gas turbine engines. One of the main reasons for failure is the reduction of the endurance limit as a result of the formation of microcracks in the surface layer, damaged in the non-optimum conditions of machining of the component which include the conditions of cutting the profile of the tailpiece, the nonuniform distribution of the load in the teeth of the tailpiece, metallurgical defects and residual tensile stresses.
The aim of the present work is the determination of microstructural reasons for the failure of single crystal blades in fatigue testing; it was required to determine metallurgical, structural or procedural factors leading to a decrease of the endurance of the metal under alternating loading.
The working blades were produced by investment casting of ZhS36VI single crystal alloy (wt.%: Ni-9Co-4Cr-6 Al-1Ti-1Mo-12W-2Re). The cast blanks of the blades were subjected to the total cycle of thermomechanical treatment after manufacture. Finishing treatment of the lock, i.e., surface plastic deformation (SPD) consisted of shot blast hardening with steel microspheres (SBH) of the 'fir tree' tailpiece.
The experiments were carried out on fragments of the lock section after different stages of thermomechanical treatment of the blades. The microsections were used to examine the structure of the surface of the metal in the main part in the vicinity of the first tooth and the depression of the tale place. Examination was carried out in a CamScan-4 electron microscope, chemical composition was determined using x-ray spectrum attachment Energy-200. The microhardness of the metal of the subsurface layers was determined in equipment Micro-Duromat 4000E, load 0.1-0.2 N.
The experimental results
The main structure investigations were carried out on the lock part of the working blades are subjected to all stages of thermomechanical treatment prior to the stand tests. The technological cycle of production of the blades included homogenising of the blanks at 13 20[degrees]C, 4 h; cooling in argon; cutting the lock by deep grinding (5); high-temperature intermediate ageing at 1030[degrees]C, 4 h; final ageing at 870[degrees]C, 24-32h and hardening by shot blasting.
Figure 1 shows the developed rough surface of the first tooth of the tailpiece with the depths of microirregularities of 10-20 [micro]m. Irregular 'notches' and depressions on the surface of the tooth indicated the relatively hard conditions of cutting the lock and surface plastic deformation in the process of shot blast hardening of the blade. The graph of the variation of microhardness of the surface of the metal of the first tooth at a depth of 150-200 [micro]m confirms the presence of the layer subjected to surface plastic deforma-tion in the tailpiece of the blade (Fig. 2, 1).
[FIGURE 1 OMITTED]
[FIGURE 2 OMITTED]
The microhardness of the alloy of the ZhS36VI blades in the volume of the tail-piece was approximately 3400 MPa. The maximum surface hardening after all the stages of thermomechanical treatment of the blades was recorded at a depth of up to 30 [micro]m. The maximum microhardness in this case was 5700-5800 MPa.
The nature of distribution of microhard-ness in the thickness of the surface layers indicates that the effect of the surface plastic deformation in the treatment of the lock is evident at a depth of up to 80 [micro]m. This was not always the reflected in the microstructure of the surface layers of the first tooth. The resultant range of the microhardness values (Fig. 2, 1) indicated the small scatter of the measured values, obtained under surface in the region of the tip and the depression of the first tooth.
Examination of the structure of the surface of the tooth in back-the reflected electrons showed the presence of an external layer containing light inclusions. The thickness of the layer of the unetched section was in the range 5-25 [micro]m (Fig. 1a, d). The largest thickness of the characteristic layer was recorded in the areas of transition from the surface of the wall of the tooth to the tip; microcracks were periodically detected. The light inclusions were in the form of carbide phases based on tungsten (18-30 wt.%) of the type ([Ni sub.3], [W sub.3]) C (Table 1, Fig. 3). In addition to tungsten, the composition of the carbide phases contained nickel and rhenium (4-11 wt.%). The size of the carbide phases was 1-2 [micro]m.
[FIGURE 3 OMITTED]
Table 1. Chemical composition of the metal of the surface layer of the tailpiece of the blade made of ZhS36VI alloy after the complete cycle of thermomechanical treatment, including shot blasting surface hardening of the lock Analysis Mass fraction of components, % spectrum C* O Al Ti Cr Co Ni Fe Nb Mo Tip of first tooth 1 6.20 2.51 0.78 6.47 7.99 41.02 0.67 0.81 2.28 2 8.60 2.92 0.76 6.47 7.52 46.31 0.68 0.57 2.23 3 8.94 -- 1.73 0.61 8.38 5.90 32.07 1.10 1.60 3.78 4 17.93 0.35 1.39 0.42 8.33 7.19 32.41 0.80 1.02 2.99 5 11.62 2.50 2.21 0.93 8.64 7.76 43.59 0.34 1.02 1.65 6 5.82 -- 2.71 0.52 12.12 10.73 52.91 -- 0.70 1.52 7 6.15 -- 5.46 1.41 4.32 6.16 64.21 -- 0.77 0.29 Transition from first tooth to depression 1 33.30 4.09 2.25 0.39 7.53 6.95 33.15 0.49 0.11 1.15 2 25.93 8.89 2.41 0.39 7.59 7.23 32.62 1.77 -- 1.02 3 6.23 -- 1.20 0.26 7.23 8.07 27.51 -- 0.24 2.56 Analysis spectrum W Re 1 20.88 10.38 2 16.49 7.45 3 30.48 5.41 4 23.57 3.61 5 17.02 2.69 6 10.09 2.88 7 9.31 1.92 1 8.68 1.91 2 10.36 1.80 3 33.74 12.94 *Carbon content values are qualitative
The presence of the carbide particles on the working surfaces of the tooth made of the single crystal alloy may be regarded as a possible area of initiation of the fatigue microcracks in tests or service of components in a turbine. The carbides are sources of stress concentration in the matrix under loading as a result of the difference in the elasticity moduli of the carbides and the matrix solution (2).
The surface of the tooth of the blade contained remnants of the carbide film whose formation was associated in all likelihood with the process of deep drawing (Table 1, Fig. 3). Evidently, the presence of excess carbon on the surface, and a relatively high temperatures of the first ageing treatment (10 30 [degrees]C). Maybe accompanied by the reaction of interaction of the active carbide-forming (tungsten, rhenium) elements with three carbon and the formation of carbide phases. In subsequent stages.
Since the metallic surface of the tailpiece of the tooth and of the depression is activated by the given conditions of the drawing, the diffusion interaction of the components of the cooling medium with the surface layers of the tailpiece in grinding was relatively fast.
A special feature of the structure of the subsurface layer of the truth of the tailpiece was that as a result of diffusion interaction of the components of the alloy with carbon of the cooling medium the matrix solution (at a depth of up to 30 [micro]m) was depleted in alloying elements tungsten (from 12.7 to 9.67%) and rhenium (from 3.67 to 1.84%), which diffused to the surface and took part in the formation of carbides.
The stabilisers of the [UPSILON]'-phase (aluminium, titanium) with the distribution coefficient [k sup.i[UPSILON]'/[UPSILON]] = [ci sup.i[UPSILON]'/[UPSILON]] , higher than unity, penetrated into the depths of the layer.
Thus, in chemical etching. The microstructure of the upper layer of the metal of the tooth with a thickness of 10-25 [micro]m consisted of a two-layer composition, formed by the external layer (thickness 3-10 [micro]m) of the solid solution, hardened with the secondary carbide phases (Fig. 4a) and the internal layer with the developed coarse-grained [UPSILON]'-phase of the sharp angle form which, playing the role of notches, may reduce the strength of the metal of the lock in high-cycle fatigue testing (Fig. 4b).
[FIGURE 4 OMITTED]
Below the diffusion zone and the external carbide layer there was an internal zone, which received mostly plastic deformation during shot blasting. The depths of the zone from the surface was 10-60 [micro]m. The zone is characterised by the deformed ([UPSILON]-[UPSILON]') structure with blurred distorting boundaries of the hardening phase (Fig. 4c). Plastic deformation, determined by treatment (blowing) of microspheres, resulted in yielding and the matrix [UPSILON]'-phase boundary and, consequently, resulted in the distortion of the contact boundaries of the particles of the [gamma]'-phase with the matrix solution.
The structure of the single crystal alloy in the body of the log below the zone of high-intensity deformation was characterised by a distinctive cubic structure with the particle size of the [gamma]'-phase of 0.3-0.6 [micro]m which formed as a result of standard three-stage heat treatment of the ZhS36VI alloy (Fig. 4d).
After chemical etching, the metal of the local the blade contained a distinctive struc-ture with the subgrain size of 15-40 [micro]m. This microstructure corresponding to the substructure of level III and was the block structure of the general type with misorientation of blocks of tens of minutes . The boundaries of the subgrains were decorated by a chain of non-ordered precipitates of the hardening phase (Fig. 5b). The dimensions and scatter of the subgrains decreased closer to the surface.
[FIGURE 5 OMITTED]
To investigate which stage consisted of the formation of changes in the surface layers of the tailpiece, the structure of the 'fir tree' locks of a number of blades was investigated after individual stages of mechanical-thermal treatment.
No differences were found in the structure of the metal of the lock of the blade examined as a blank after homogenising (1320[degress]C, 4-7 h) with cooling in our gun on the surface, and in the body of the blade (Fig. 6). As a result of rapid cooling (100[degress]C/min) the structure of the alloy consisted of the matrix solution with the precipitates of the hardening [gamma]'-phase in the form of sub-cubes. The size of the phase was 0.3-0.5 [micro]m. The external carbon-enriched layer with a thickness of approximately 3-5 [micro]m is evidently, the result of ??? of oil vapours of the diffusion pumps into the vacuum space. This was not critical for further mechanical treatment and manufacture of the lock blade.
[FIGURE 6 OMITTED]
The microhardness, measured to a depth of 200 [micro]m from the surface, was in the range 3370-3500 MPa (Fig. 2, 2).
In investigation (after etching) of the cross-section of the lock part of the blade of the first intermediate annealing examination showed the structure of the zone of plastic deformation which was 10-15 [micro]m thick (Fig. 7). The typical pattern of the stress layer was detected both on the surface of the tooth and in the depression.
[FIGURE 7 OMITTED]
This subsurface zone was characterised by distortion of the regular cubic structure of the [gamma]'-phase.
The size of the hardening phase in the metal of the lock after the first ageing treatment was 0.2-0.9 [micro]m. This scatter can be explained by the selective dissolution of the particles at the ageing temperature, by the growth of the particles present in the matrix, and by the precipitation of the new sub-dispersed particles of the [gamma]'-phases. The results of annealing at 1030[degress]C.
The microhardness of the surface layers of the tooth and of the depression different only slightly. In most cases, the distribution of microhardness was linear throughout the entire measured debts (200 [micro]m) of the material (Fig. 8, 1). There was no large increase of the values of microhardness on the investigated surface as a result of the compressive stresses, determined by the coating of the tooth with the formation of the deformation zone, because of the small thickness (5-10 [micro]m) of the zone.
[FIGURE 8 OMITTED]
After second ageing at a temperature of 870[degress]C for 30 2h, the surface of the tooth of the tailpiece of the blade retained the region subjected to deformation (as a result of the formation of the 'fir tree' low), which greatly differs in the morphology from the hardening phase (Fig. 9).
[FIGURE 9 OMITTED]
In the vicinity of the surface there were carbide phases of the type [Me.sub.6]C based on tungsten and rhenium with a particle size of up to 1 [micro]m (Fig. 9a, b).
Partial removal of the refractory components and the matrix solution (tungsten, rhenium), with the minimum coefficients of distribution, resulted in the changes in the morphology of the [gamma]'-phase, i.e., increase of the grain size (Fig. 9).
After low-temperature ageing the subsurface zone of the effect of the formation became wider, reaching 15-30 [micro]m (in comparison with 10-15 [micro]m after first ageing). This was quite evident on the level of the microhardness measured in the cross-section of the tooth to a depth of up to 200 [micro]m (Fig. 8, 2).
Effect of heat treatment on the fine structure of the ZhS36VI single crystal alloy.
The process of homogenising at high temperatures ([less than or equal to]1310[degress]C) the rapid cooling in argon of the blanks for the blade produce produce the dispersion-hardened ([gamma]-[gamma]') structure of the single crystal. The hardening phase has the formal subcubes with the size of 0.2-1.0 [micro]m (Fig. 10a). After high-temperature ageing (1030[degress]C, 4 h) the structure of the single crystal consists of less ordered cuboidal particles, with the size of 0.2-0.6 [micro]m (Fig. 10b).
[FIGURE 10 OMITTED]
After heat treatment (high-temperature quenching, ageing at 1030[degress]C, 32 h, final treatment), and at 870[degress]C the structure of the single crystal of the ZhS36VI creep-resisting alloys has the form of a matrix solution with the uniformly ordered cubic precipitates of the hardening [gamma]'-phase (Fig. 10c). In particular, this structure with the volume fraction of the [gamma]'-phase of approximately 65% in the particle size of the [gamma]'-phase of 0.3-0.5 [micro]m ensures the maximum creep resistance and long-term strength of the single crystal alloy at high temperatures.
On the basis of the experimental results. It may be concluded that shot blasting of the surface of the 'fir tree' lock of the blade with microspheres leads to the formation of a contact layer on the external surface of the tooth above the main deformed layer. The thickness of this contact layer reaches a maximum of 30 [micro]m, with the changed chemical composition and structure. The presence of such a contact layer, containing the carbide phases and the coarse-grained [gamma]'-phase of the acicular, has an unambiguously detrimental effect on the fatigue strength of the blades in high cycle fatigue and results in early failure of the components.
The 'fir tree' lock of the blade in the production cycle is produced by deep grinding which produces the required profile of the lock part in three passes. The process is relatively thermally stressed. The surface of the tailpiece is characterised by the formation of the maximum values of the pulses of instantaneous contact temperature acting on the upper layers of the processed metal.
The lower zone of the effect as a result of the heat exchange between the components are calling liquid is characterised by the formation of conditions leading, together with the force effect, as to the formation of residual compressive stresses on the surface layer. These stresses, it is 150-250 MPa and the depth of propagation of the compressive stresses and the profile of the lock is 40-80 [micro]m (5).
In the formation of the 'fir tree' tailpiece. Examination showed of formation of a surface layer with unstable structural parameters in the thickness which define the degree of cold working, surface roughness, the presence of depressions and the level of residual stresses.
A water emulsion of emulsifying oil was used for cooling. At the carbon content of the emulsion and the absence of the washing operation, i.e., the the removal of the newlyformed carbon film from the surface of the blade tools, resulted in the formation of carbides of the type Me6C (in the activated surface layer) based on tungsten and the rhenium in the process of high-temperature ageing at 10 30[degress]C, 4 h. The carbide particles, with the size of 0.5-3.0 [micro]m are distributed discreetly and form under the residual carbon film. Subsequent annealing and shot blasting was accompanied by the enlargement of the carbides and the increase of the width of the zone subjected to surface plastic deformation. Deformation hardening resulted in the intensification of the activation of the processes in the subsurface volumes of the tailpiece, and the particles of the carbide phases were ??? into the metal to a large depth.
The final stage of treatment of the tailpiece of the blade is shot blasting hardening of the surface. The residual surface compressive stresses, formed as a result of this treatment, slowdown the propagation of the micro-crack in the body of the component, but do not prevent the formation of these microcracks. The results of the modern experiments with the hardening of the 'fir tree' of the lock were used to select the special conditions of treatment of the tailpiece is of the blades in two positions. The minimum diameter of the nozzle resulted in the maximum speed of discharge of the flow of the spheres (kinetic energy) and, consequently, high-intensity of hardening was achieved to a depth of up to 100 [micro]m. The propagation of the induced stresses to the depth of the metal was 35-40 [micro]m greater than in other treatment methods. This was confirmed by the microhardness measurements results (Fig. 11).
[FIGURE 11 OMITTED]
The mean curve (Fig. 11) describes the nature of strain hardening of the tooth of the tailpiece after treatment of the component with the flow of microspheres in the selected conditions. In this case, hardening (increase of hardness) of the parent metal in the cross-section at a depth of 200 [micro]m from the surface along the circumference of the tooth, the walls and the depression was approximately the same.
Figure 12 shows the fine structure of the subsurface layers of the metal of the fragment of the tailpiece of the blank after homogenising and shot blasting by microspheres. The results showed slight plastic deformation of the particles of the [gamma]'-phase in the matrix immediately and the surface of the metal of the lock. No carbide phases were found in the external layer.
[FIGURE 12 OMITTED]
However, as a result of the formation of the lock in the subsurface layer of the blade with the cut teeth there was (and also in other blades) a region with white globular particles of the [Me.sub.6]C carbides based on tungsten, rhenium and molybdenum. The formation of these particles was caused by the interaction of ZhS36VI alloy with the material of the carbon-containing medium in deep grinding of the 'fir tree' lock resulting in overheating (burning) of the surface of the tailpiece.
Shot blasting with microspheres may have a negative effect for casting alloys, containing large carbide phases or other defects of the notch type (2). As a result of the formation of the subsurface zone of the tensile residual stresses, the region with the defects may be characterised by the formation of cracks causing in later stages premature failure of the components during testing.
The working single crystal blades of ZhS-36VI alloyed after the complete cycle of heat treatment without shot blasting of the 'fir tree'lock, tested in the testing stand of the Institute of strength of materials, National Academy of Sciences of the Ukraine, shows the high cycle fatigue limit at N = 20X[10.sup.6] cycles of 50--70 MPa.
The blade, subjected to 2-position treatment of the tailpiece with the microspheres has a higher endurance limit [[sigma].sub.-1], equal to 95--10 MPa. The depression of the first tooth was not a critical area of the initiation of fatigue failure (fatigue failure was transferred to the near-root part above the lock). The value [[sigma].sub.-1] was 15 MPa higher than the fatigue limit, obtained by the customer as a result of shot blasting with microspheres of the tailpiece of the blades for the same test time.
1. Investigations were carried out into the structure and chemical composition of the subsurface layer of the metal of the lock part of the non-conditioned single crystal working blades made of ZhS36VI alloyed after different stages of thermomechanical treatment of the blanks and directly of the blade. The experimental results showed the presence on the surface of the 'fir tree' lock of a layer containing carbide phases of the type [Me.sub.6]C based on tungsten and rhenium and the enlarged [gamma]'-phase of the irregular shape. The subsurface layer consisted of the external zone from which alloying elements were partly removed with a large number of dispersed carbide phases and the internal zone with the coarsened hardening phase, also containing the carbide phases.
2. It was found that the manufacture of the lock by deep grinding and the formation of compressive stresses leads to the formation in the subsurface layer of the 'fir tree' tailpiece of the zone of induced plastic deformation and this is the reflected in the change of the shape, ordering and regularity of the [gamma]'-phase, and also in the higher microhardness of the parent metal. The application of the water emulsion of the oil as a result of the cooling liquid in deep grinding results in the formation of a thin passivating carbon film on the surface of the completed lock. In high-temperature ageing at 1030[degress]C (4 h) the residual carbon in the activated surface layer of the lock interacts with the refractory alloying elements of the alloy with the formation of the carbide phases which are unacceptable in the structure of the single crystal because, being on the surface, they are the regions of nucleation of microcracks, causing premature failure of the blades in testing.
3. It has been shown that the deformation hardening of the surface in shot blasting of the 'fir tree' lock produces in the metal a field of compressive stresses but also leads to the nucleation of microcracks have the matrix-carbide particle interface. These microcracks may lead to the initiation of the main high cycle fatigue cracks in the region of the depression of the first tooth in vibration testing of the blades and, therefore, the hardening of the tailpiece of the blade with the microspheres should be carried out in the controlled conditions.
(1.) Shalin R.E., et al., Single crystals of nickel creep-resisting alloys, Mashinostroenie, Moscow, 1997.
(2.) Verin D.G., Microstructure and properties of creep-resisting alloys: in: Creep-resisting steels, Metallurgiya, Moscow, Moscow, 1976, 217-241.
(3.) Petukhov A.N., Fatigue resistance of the components of gas turbine engines, Mashinostroenie, Moscow, 1993.
(4.) Kablov E.N. and Alekseev A.A., Physics of the creep strength of heterophase alloys, Nauka, Moscow, 2006, 44-45.
(5.) Boguslaev V.A., et al., Technological parameters ensuring the required service characteristics of the components of gas turbine engines, Motor Sich, Zaporozh'e, 2003.
V.V. Kurenkova E.O. Paton Electric Welding Institute Pratt and Whitney Paton Research Centre, Kiev
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|Title Annotation:||GENERAL PROBLEMS OF METALLURGY|
|Publication:||Advances in Electrometallurgy|
|Date:||Jul 1, 2010|
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