Grain boundary microstructure and fatigue crack growth in Allvac 718Plus superalloy

Materials Science and Engineering A 528 (2011) 2570–2580

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Grain boundary microstructure and fatigue crack growth in Allvac 718Plus superalloy L. Viskari a,∗ , Y. Cao a , M. Norell a , G. Sjöberg b , K. Stiller a a b

Chalmers University of Technology, SE-412 96, Göteborg, Sweden Volvo Aero Corporation, SE-461 81, Trollhättan, Sweden

a r t i c l e

i n f o

Article history: Received 4 August 2010 Received in revised form 25 November 2010 Accepted 25 November 2010 Available online 3 December 2010 Keywords: Delta-phase Fractography Oxidation Hold-time fatigue Electron microscopy X-ray photoelectron spectroscopy Auger electron spectroscopy

a b s t r a c t The correlation between grain boundary microstructure and fatigue crack growth with hold-times was investigated for two conditions of the superalloy Allvac 718Plus; a Standard condition with the recommended distribution of grain boundary phases and a Clean condition with virtually no grain boundary phases. Fatigue testing was performed at 704 ◦ C using 10 Hz cyclic load with intermittent hold-times of 100 s at maximum tensile load. Microstructural characterization and fractography were conducted using scanning- and transmission electron microscopy techniques. Auger electron- and X-ray photoelectron spectroscopy techniques were used for oxide analyses on fracture surfaces. It was found that in the Standard condition crack growth is mostly transgranular for 10 Hz loading and intergranular for hold-times, while for the Clean condition crack growth is intergranular in both load modes. The lower hold-time crack growth rates in the Standard condition are attributed to grain boundary ␦-phase precipitates. No effect of ␦-phase was observed for 10 Hz cyclic loading crack growth rates. Two different types of oxides and oxide colours were found on the fracture surfaces in the Standard condition and could be correlated to the different loading modes. For cyclic loading a bright thin Cr-enriched oxide was dominate and for hold-times a dark and slightly thicker Nb-enriched oxide was dominant These oxide types could be related to the oxidation of ␦-phase and the matrix respectively. The influence of ␦-phase precipitates on crack propagation is discussed. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Allvac 718Plus is a recently developed precipitation hardened Ni-base superalloy which has received some attention by being what may be described as a cross-over between the two well established superalloys Alloy 718 and Waspaloy. Superalloys are commonly employed in gas turbine engine applications where they are exposed to high temperatures, high loads and chemically harsh operating environments. The loads experienced by these alloys are in most operating conditions non-uniform such that the load amplitudes and frequencies vary during operation. The effect of such load cycles deems uniform high temperature low cycle fatigue (LCF) testing to be insufficient for predicting component lifetimes, as this does not account for the detrimental effect of low frequencies or hold-times. For Allvac 718Plus, isothermal crack propagation at temperatures exceeding 700 ◦ C in atmospheric conditions has been shown to be predominantly intergranular when subjected to mechanical testing by, e.g. stress-rupture testing as well as fatigue testing with and without hold-times [1–3]. Correlating results from

∗ Corresponding author. E-mail address: [email protected] (L. Viskari). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.11.080

high temperature hold-time fatigue testing with analysis of grain boundary microstructure contributes significantly to the general understanding of hold-time crack growth. Fatigue crack growth modes and rates at high temperatures have been shown to highly depend on the loading conditions used. For a large number of superalloys the inverse proportionality between the LCF frequency used and crack growth rates (CGRs) per cycle (da/dN) at high temperatures has been shown, well summarized in [4], as an effect of crack growth being time dependent at these temperatures. Low frequencies also include hold-times, i.e. trapezoidal periods of static tensile loading introduced into cyclic loading sequences. Two main damage mechanisms are often considered to occur during holdtimes; Dynamic Embrittlement (DE) [5,6] and Strain Assisted Grain Boundary Oxidation (SAGBO) [4]. Both have similar precursors for occurring, however the principal difference is that while fracture by DE occurs through decohesion of grain boundaries, fracture by SAGBO occurs in oxidized and thus embrittled grain boundaries. A general term for this kind of embrittling damage has also been coined, Gas Phase Embrittlement (GPE) [4]. In this work two conditions of Allvac 718Plus with nominally same chemistry are investigated, namely a Standard heat treatment with grain boundary phases and also a Clean heat treatment with virtually no grain boundary phases. The influence of the grain boundary precipitates

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on fatigue crack propagation mechanisms, rates, and oxidation phenomena in 718Plus are presented in this paper. Specific interest is directed towards the in-detail role of grain boundary ␦-phase precipitates on crack propagation and crack surface oxidation. 2. Experimental 2.1. Materials Allvac 718Plus in the form of a 2.9 mm thick hot rolled sheet was used for all specimens. The composition of this sheet is shown in Table 1. The main hardening phase in 718Plus is the ␥ -phase, Ni3 (Al,Ti,Nb) having the L12 structure. The ␥ -phase, Ni3 (Nb,Al,Ti) with D022 -structure is not generally expected based on experimental observations, however, simulations have predicted that very limited quantities of this phase may precipitate. The orthorhombic Ni3 (Nb,Al,Ti) grain boundary ␦-phase is precipitated for grain size control and increased stress-rupture notch ductility. Centre-notched fatigue specimens were obtained from the sheet using electric discharge machining (EDM), the specimen geometry is specified in Fig. 1. The specimens were heat treated to produce a Standard condition containing ␦-phase and a Clean condition without ␦-phase. The Standard condition was heat treated according to recommendations by Allvac [7] using water quenching (WQ) and furnace cooling (FC) between the respective steps; 954 ◦ C/1 h (WQ) + 788 ◦ C/8 h (FC) + 704 ◦ C/8 h (FC). The purpose of the first stage of this heat treatment is to solve hardening phases and concomitantly produce proper distribution of ␦-phase at grain boundaries. The 788 ◦ C/8 h (FC) + 704 ◦ C/8 h (FC) aging treatments precipitate the ␥ hardening phase. The Clean condition was first solution heat treated at 1040 ◦ C/0.5 h (WQ) followed by 954 ◦ C/4 h and aging as above. The 1040 ◦ C/0.5 h-treatment dissolves the ␦phase whilst the subsequent 954 ◦ C/4 h-treatment was initially expected to re-precipitate excessive amounts of ı-phase [7]. However, an affiliated study [8] on this specific heat and condition of 718Plus found this heat treatment to produce a very clean microstructure in terms of grain boundary precipitates. The solution of ␦-phase in the Clean condition is accompanied by grain growth as the pinning of grain boundaries by ␦-phase is removed. Thus, the average grain size of the Clean condition is slightly larger (55 ␮m or ASTM 5.5) than that of the Standard condition (32 ␮m or ASTM 7). The hardness of the Standard and Clean conditions are highly similar at 294 and 292 HV10 kg, respectively.

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2.3. Electron microscopy and X-ray diffraction Several characterization methods were used in this work. X-ray diffraction (XRD) was used for bulk identification of main phases. Scanning electron microscopy (SEM) was applied for identification and localization of main phases and for fractography. Transmission electron microscopy (TEM) was employed for more detailed investigation of grain boundary precipitates. Energy dispersive Xray (EDX) analyses were performed using both SEM and TEM. For selected samples and regions of interest, thin foil TEM samples were analysed in SEM using in-lens and scanning transmission electron microscopy (STEM) imaging detectors and also EDX. Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS) were used for studies of the oxides formed on the crack surfaces. Cross-sections of grain- and twin boundaries and ␦-phase were milled using Focussed Ion Beam (FIB). XRD analyses were performed using a Philips D5000 diffractometer in Bragg-Brentano configuration, Cu X-ray source and 40 kV acceleration voltage, with 2 scanning angle spanning 20–100◦ in 0.05◦ increments. For SEM analyses a FEI Quanta 200 field-emission gun (FEG) SEM equipped with Oxford Inca EDX-system was used. All SEM-EDX analyses were performed using 20 kV acceleration voltage. SEM on thin foils was performed using a Leo Ultra55 SEM at 30 kV acceleration voltage. FIB crosssections were milled and imaged using a FEI Strata 230 Dual Beam using ion mode at 30 kV. TEM was performed using a Philips CM200 FEG-TEM operated at 200 kV, equipped with a Oxford Inca EDX-system. The samples for characterization of microstructure were excised from the grips of the tested specimens (i.e. furthest away from the crack with very limited effect of heating and strain) and prepared using standard metallographic sample preparation. Samples for crack surface analyses by SEM and AES and secondary crack analyses by SEM were sectioned 10 mm below the crack surface. No etching was performed on the samples used for EDX analyses as etching may significantly alter the sample chemistry. These samples were also used for XRD analyses. For imaging purposes, additional samples of both conditions were etched using Kallings no 2 etchant to relieve precipitates. These etched samples were also used for FIB cross-sections as they allowed for more precise targeting of regions of interest. TEM and SEM-STEM samples were produced using twin-jet electropolishing of punched metallographically polished 3 mm discs. 2.4. Auger electron- and X-ray photoelectron spectroscopy

2.2. Hold-time fatigue testing The centre-notched samples were pre-cracked a total of 0.4 mm (0.2 mm on each side) using cyclic loading at room temperature, following which the temperature was increased to 704 ◦ C. Maximum and minimum loads of Pmax = 10 kN and Pmin = 0.5 kN were used (giving  0 = 167 MPa, R = 0.05). The frequency used for cyclic loading was 10 Hz and several hold-times of 100 s at Pmax were introduced intermittently in blocks into the cyclic loading sequences. Final fracture was performed at room temperature. Direct Current Potential Drop (DCPD) technique was used to monitor crack growth. CGR as per da/dN was calculated based on the PD measurements. da/dt was calculated by multiplying da/dN with the frequency used. K was calculated according to Eqs. (1) and (2) based on reference [9]. In these equations  0 is the initial stress, a is crack radius and b is specimen width. √ K = 0 af (a/b) (1) where 2

f (a/b) = 1 + 0.128(a/b) − 0.288(a/b) + 1.525(a/b)

3

(2)

The oxide thickness and compositions on the fracture surfaces were determined using both Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS). For the Standard condition both methods were used in areas formed during hold-time and cyclic testing. For the Clean condition only XPS was applied. Elementary AES depth profiles were obtained from the two loading areas (from micron-sized areas in several facets) by ion etching using a ThermoFisher Microlab 350 at 10 kV acceleration voltage. When depth profiles on irregular fracture surfaces were generated, the topography had to be considered. Here the sample was tilted in 10◦ steps from 0◦ to 40◦ and the projected dimensions of the selected facets were checked to make sure that they were parallel to the top surface. A 40◦ tilt was used both for the analyses and the calibration of the etch rate on a standard Ta2 O5 sample. Further information on AES depth profiling of irregular samples is given in ref. [10]. Sensitivity factors were deduced from the elementary peak areas as measured in the bulk and the nominal alloy composition. The intensity determination for Co was not straightforward as all its main peaks overlap with those of other alloying elements. Thus, intensity of Co was estimated by mea-

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Table 1 Studied composition of Allvac 718Plus in wt%. Ni

Cr

Fe

Co

Nb

Mo

Al

W

Ti

C

P

B

51.9

17.86

9.59

8.97

5.49

2.7

1.49

0.99

0.76

0.024

0.006

0.004

Fig. 1. Geometry of specimens used for mechanical testing. The thickness of the specimen is 2.9 mm.

suring the total peak area at 649 eV and subtracting the intensity contribution from Fe in proportion to the Fe intensity at 598 eV, using Eq. (3)

ICo = I649eV − IFe649eV = I649eV − 1.24IFe598.5eV

(3)

The constant 1.24 is the ratio of sensitivity factors for Fe at 649 eV and 598.5 eV provided by the instrument manufacture. The XPS depth profiles were obtained from areas (diameter 0.8 mm) by means of successive Ar ion etching using a PHI5500 instrument with monochromatized AlK˛ source. The Ni 2p3/2 , Cr 2p3/2 , Fe 2p1/2 , Nb 3d, O 1 s and Co 2p1/2 photoelectron peaks were analyzed with 23.5 eV pass energy and 45◦ take off angle. Sensitivity factors provided by the instrument manufacturer were used for quantification and the etch rate was calibrated as for AES. For both methods the oxide thickness was estimated from to the depth where half the O intensity decrease occurred.

3. Results 3.1. Characterization of microstructure Fig. 2 shows the obtained XRD spectra in the 2 range of 40–55◦ for the two material conditions. For comparison, the spectrum of the non-heat treated As-Received material is also included. The spectra are proportionally scaled but offset in terms of intensity for clarity. The largest peaks at 2 = 43.5 and 2 = 50.7 both correspond to the ␥ and ␥ phases and are present for all conditions. There are minor (and in these spectra indistinguishable) differences between the d-spacings of the highly coherent ␥ and ␥ phases. ␦-phase peaks are clearly visible for the Standard condition in which also minor peaks corresponding to HCP Ni3 Al0.5 Nb0.5 -phase are present. Elemental mapping performed by SEM–EDX on the Standard condition shows the presence of grain boundary precipitates rich in Nb- and Ni- and poor in Cr- and Fe as seen in Fig. 3. In addition to these elements Al, Co and Ti contents were also found to be increased in relation to the surrounding matrix. Corresponding analyses performed on the Clean condition (Fig. 4) show a very homogeneous elemental distribution lacking grain boundary pre-

Fig. 2. XRD spectra for the As-Received, Standard and Clean conditions in the 2 range 40–55◦ . The spectra are offset in terms of intensity. The vertical 2-indicators correspond to ␦ (square), HCP Ni3 (Al0.5 Nb0.5 ) (circle) and ␥  (triangle) phases.

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Fig. 3. SEM image (left) and SEM-EDX elemental maps for Cr, Fe, Nb and Ti for the corresponding area in the Standard condition.

Fig. 4. SEM image (left) and SEM-EDX elemental maps for Cr, Fe, Nb and Ti for the corresponding area in the Clean condition. Dashed lines indicate the positions of the here indistinguishable grain boundaries.

cipitates. TEM and SEM–STEM EDX line-scans do not reveal any apparent segregation to grain boundaries in either condition. Fig. 5 additionally shows non-etched As-received microstructure as a reference to the material’s response to the heat treatments in the aspect of ␦-phase distribution. Fig. 6 a shows a typical TEM image of ␦-phase precipitates residing at grain boundaries in the Standard condition. The composition and structure of this phase were identified using TEM–EDX and SAED. The EDX results in Table 2 represent the average composition of several precipitates, the major phase constituents being Ni, Nb, Al, Co and Ti. The plate-like morphology of ␦-phase is apparent in the FIB cross-section in Fig. 7 which allows a three dimensional perception of its morphology. In this figure it is also apparent that the actual grain boundary has accommodated to the presence of ␦phase by having a serrated appearance, also to be observed in Fig. 6. Furthermore the volume surrounding the ␦-phase is denuded of large ␥ hardening precipitates (see Fig. 6). A SEM in-lens image of such a zone in a thin foil specimen is shown in Fig. 8. No denuded zones were observed in the Clean condition. Fig. 9 shows a SEM image of the Standard condition following etching, revealing that

Table 2 Major constituents in ␦-phase precipitate, obtained by TEM–EDX (at%). Ni

Nb

Al

Co

Ti

Cr

Fe

W

Mo

66.3

11.5

7.2

5.9

4.9

1.8

1.6

0.5

0.4

Fig. 5. SEM image of the cipitates, identified by EDX visible.

As-received and XRD

material. Plate-like preas ␦-phase, are clearly

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Fig. 6. TEM bright field image showing plate-like ␦-phase precipitates at grain boundary.

Fig. 8. SEM in-lens image of thin foil sample showing a ␥ -denuded zone in the Standard condition. ␥ precipitates appear as bright spots (approximately 80–100 nm in size) and the absence of these near the ␦-phase precipitates is apparent. Black dotted lines mark the perimeter of this zone.

3.2. Hold-time fatigue testing

while the major precipitation of ␦-phase occurs in grain boundaries some precipitation also takes place in twin boundaries and also what appears as intragranular (see further discussion below). The angle of the precipitates in relation to the boundary or plane in which they reside varies greatly. When precipitated in grain boundaries the angle of the plates, as viewed in the plane of the image, ranges from parallel to orthogonal while for twin boundaries the precipitates are most frequently found fully or near-parallel with the twin boundary, also confirmed by FIB cross-sections for three dimensional orientation perception.

Fig. 7. FIB cross-section of ␦-phase plates at a grain boundary. The angle between the top surface and the plane of cut is 90◦ and the viewing angle in respect to the top plane is 38◦ . ␥ precipitates appear as dark spots in the cross-section. A zone denuded of these precipitates (indicated by arrows) surrounding the ␦-phase precipiates is clearly visible.

It is shown in Fig. 10 a that da/dN increases severely during hold-times for both for Standard and Clean conditions. Upon shifting from hold-time to cyclic loading transitional effects can be observed, where the CGRs are initially at the same elevated level as during hold-times but then decline to a steady growth rate convergent with the 10 Hz loading CGR. This is valid for all presented circumstances with the exception of K < 15 for the Standard Condition. The distances over which the transitions occur span up to approximately 1 mm in the Clean condition and 0.5 mm in the Standard condition and are throughout the range longer for the Clean condition. This is shown in Fig. 10(c) and (d), showing plots of da/dt vs. the total crack length (TCL) of one half of the centre-notched

Fig. 9. SEM image of etched Standard condition, relieved ␦-phase precipitates appear bright.

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Fig. 10. Plots of (a) da/dN vs. K (b) da/dt vs. K (c) da/dt vs. total crack length (TCL) for the Standard condition, and (d) da/dt vs. TCL for the Clean condition. Regression lines (without consideration for transients) are shown for Figures (a) and (b).

samples. In terms of da/dN the convergent levels of cyclic CGRs are very similar for both conditions (marked by the superimposed regression line in Fig. 10(a)), while the hold-time CGRs differ by those for the Clean condition being slightly higher than the Standard condition. Regarding da/dt as shown in Fig. 10(b), the intergranular crack growth in the Clean condition occurs at the same rate regardless of whether the loading is cyclic at 10 Hz or 100 s hold-time, when excluding any transitional effects. Also, the intergranular CGR during hold-times for the Standard condition is lower than that for the Clean condition, while the transgranular crack propagation in the Standard condition during cyclic loading occurs at the same rate as the intergranular CGR for the Clean condition. 3.3. Fractography Fig. 11 shows the crack surfaces of the Standard and Clean conditions. For the Standard condition there are clearly distinguishable alternating dark and bright sections along the crack path. These zones are also identifiable by contrast in SEM using BSE mode. Correlation of the crack lengths in the different loading modes shows the dark sections to correspond to hold-time loading and the bright sections to cyclic loading. SEM images of these sections in the Standard condition are seen in Fig. 12. The sections corresponding to cyclic loading show mixed inter- and transgranular crack propagation (mostly transgranular), while the sections corresponding to hold-time loading show predominately intergranular crack propagation. Also the hold-time sections are far more rough (varying topography in respect to the general crack growth plane), than the relatively flat cyclic regions. No such sections were distinguished between the cyclic- and hold-time loadings in the Clean condi-

tion where crack propagation was intergranular for both modes, as evident in Fig. 13. By comparing the intergranular crack surfaces shown in Fig. 12(a) and Fig. 13(a) and (b), it is seen that these are very dissimilar for the two conditions. The crack surface of the Standard condition has heavily serrated facets while the Clean condition has flat facets. Both conditions contain significant quantities of secondary cracks and by investigating these secondary cracks in cross-sections the details of crack paths were obtained. As seen in Fig. 14, secondary crack propagation in the Standard condition is serrated by grain boundary ␦-phase in contrast to the non-serrated propagation in the Clean condition. Secondary cracks in the Standard condition were furthermore found to preferentially initiate at and propagate along the interfaces between ␦-phase and the matrix. Especially grain boundaries with the ␦-phase precipitates seemingly parallel or near-parallel to grain boundary planes were found to exhibit this behavior. Additionally, several observations were made of fractured ␦-phase precipitates in intergranular secondary cracks. Some oxidized and fractured Ti- and Nb-carbides were further observed, primarily on the facets of the Clean condition. 3.4. Oxide analyses In the Standard condition darker oxides formed on the fracture surface in the hold-time sections than in the sections formed during cyclic testing. The AES profiles showed that the same two types of oxides formed in different facets on the fracture surface in both the cyclic and the hold-time sections. Fig. 15 shows typical AES depth profiles for the two oxides. Type 1 oxide (Fig. 15(a)) was a thicker oxide with Ni-rich outer layer and Nb, Al-rich inner layer. No obvious Cr-oxide was found. Type 2 oxide (Fig. 15(b)) was a thinner

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Fig. 11. Fracture surfaces of the Standard condition (above) and the Clean condition (below). The alphabetical identifications of areas shown are referred to in oxide analyses.

Fig. 12. SEM images of fracture surfaces in the Standard condition for Hold-Time (left) and cyclic (right) loading.

oxide and composed of multiple layers. An outer thin Ni-rich layer was formed on top of a NiFe-oxide and a Cr-oxide at the interface to the metal. A close examination of the cation profile showed a Co enrichment between Fe-rich and Cr-rich oxides (lower part of Fig. 15(b)). Slight enrichment of Nb and Al below the Cr-oxide was also observed in type 2 oxide. The question arises whether the area proportions of these two oxides were identical in different bands with altered loading con-

ditions. The selected facets (suitable for the AES profiling) will not necessarily give representative proportions. Therefore XPS depth profiles over a larger analysis area were determined, as shown in Fig. 16. Profiles from both the cyclic and hold-time (Fig. 16(a)) bands were alike at the first glance and consistent with the presence of two types of oxide observed in the AES study. It was found that a multi-layer structure similar to type 2 oxide existed together with a clear Nb-rich layer which was considered to be representative of

Fig. 13. SEM images of fracture surfaces in the Clean condition for Hold-Time (left) and cyclic (right) loading.

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Fig. 14. SEM BSE images of secondary cracks branching from the main cracks in (a) the Standard condition and (b) the Clean condition. While the secondary crack in the Standard condition propagates in serrated manner, caused by the ␦-phase precipitates, the secondary crack in sample the Clean condition propagates in a straight manner.

Fig. 15. Typical AES depth profiles from the standard condition showing (a) oxide type 1 (b) oxide type 2. The lower part in (b) is from a cation profile disregarding O.

type 1 oxide. However, a close comparison of the cation profiles for Nb and Cr (Fig. 16(b)) showed that the hold-time area had less Cr and more Nb over a thicker layer. From these observations it is concluded that the proportions of two types of oxides differed between hold-time and cyclic areas. In the hold-time area, more Nb in the inner part was consistent with a larger area fraction of type 1 oxide. Similarly, the cyclic area had more type 2 oxide with an inner Cr-rich oxide. Table 3 gives the average oxide thickness as estimated from AES and XPS depth profiles. The band designations are shown in Fig. 11. It can be seen that the type 1 oxide was thicker than type 2 oxide in all bands. It was also found that the cyclic loading band (C) had a slightly thinner oxide compared to that of the adjacent holdtime area (B), despite a longer exposure time following fracture. The final cyclic band (A) had the thinnest oxide and the shortest exposure time at high temperature. The XPS profiles were carried out over larger non-planar analysis areas with partly less favorable

orientations relative to the argon ion incidence and thus lower etch rate, resulting in overestimated thickness values. The C-band was too narrow to be analyzed by XPS without interference from the adjacent bands. As the fracture for the Clean condition was intergranular, irrespective of the loading condition with no obvious bands of different oxide colour, a rather uniform oxide can be expected, further affirmed by the absence of grain boundary phases. An XPS

Table 3 Estimated oxide thickness of different bands in the Standard condition. Band AES results XPS results a b

Oxide type 1. Oxide type 2.

A (nm) 107a , 75b 120

B (nm) 230a , 165b 180

C (nm) 210a , 120b N/A

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Fig. 16. XPS area profiles for the standard conditions (a) hold-time area (band B) and (b) comparison of Cr and Nb cation profiles under cyclic (band A) and hold-time (band B) loading.

depth profile is shown in Fig. 17. From the surface and inwards metal enrichments in the order Ni–Fe–Co–Cr–Nb can be observed. Despite the different analysed areas this has a striking resemblance with the type 2 oxide as shown in Fig. 15. 4. Discussion 4.1. General fatigue behavior The 100 s hold-times show severely increased da/dN in comparison to 10 Hz cyclic loading for both conditions, clearly demonstrating the time dependent crack growth behavior for this type of loading. Moreover it is interesting to note that for CGRs expressed as da/dt, the Clean condition exhibits nearly the same CGRs for both hold-time and cyclic loading while for the Standard condition there is a clear reduction of da/dt during hold-times (Fig. 10(a) and (b)). The difference in hold-time CGRs for the two conditions is attributed to the presence of ␦-phase in the Standard condition grain boundaries. There is no apparent effect of ␦-phase on cyclic loading CGRs, which are believed to be mostly

Fig. 17. XPS area profiles from the Clean condition.

cycle dependent and also less influenced by microstructure. The actual crack propagation mode at 10 Hz is however clearly affected by the microstructure, being mostly transgranular for the Standard condition and intergranular for the Clean condition. Furthermore, even if the calculated da/dt values may be influenced by the frequencies used they indicate the following important trend: The intergranular hold-time crack propagation in the grain boundaries of the Clean condition, i.e. in ␥/␥ boundaries, may be stated to occur faster than in what are primarily ␥/␥ grain boundaries with ␥/␦ phase boundary segments in the Standard condition. Some of the possible mechanisms behind the importance of ␥-phase on crack propagation are discussed below.

4.2. The influence of ı-phase on the local chemistry and morphology of the grain boundaries It is believed that in the Standard condition the majority of ␦-phase resides in grain- and twin boundaries as the ␦-phase distribution shown in Fig. 9 needs to be interpreted with some caution. What is shown in this figure is an arbitrary two-dimensional plane of cut from a three dimensional body. Thus, it may well be that what appear as intragranular precipitates are in fact grain boundary precipitates protruding to the grains interiors and being intersected by the cut. The volumes denuded of large ␥ in this condition, addressed in Fig. 7, are without exception related to the presence of ␦-phase and were observed using FIB (Fig. 7), TEM (Fig. 6) and also SEM on thin foil samples (Fig. 8). At this stage analyses do not conclusively allow distinction of whether all ␥ precipitates have been denuded or if this is applicable solely for the larger ␥ precipitates clearly visible here. The denuded zones, some micron thick, are most likely more ductile than the surrounding matrix, which would undoubtedly have an influence on the micromechanisms of intergranular crack propagation. Whilst attempts were made at measuring the hardness locally by Vickers hardness testing (5- and 10 g loads), the results proved inconclusive and this was ascribed to the small volumes of the targeted zones. Further work is to be performed on these denuded zones. Moreover, the serrated shape of the grain boundaries found in the Standard condition as the result of ␦-phase (Figs. 6 and 7) is also likely to affect crack

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growth rates. Whilst the serration is only at a very small scale, it is a phenomenon comparable to the sought effects of serrated grain boundaries by grain boundary engineering [11,12]. Thus, ␦-phase has an important impact on crack propagation in Allvac 718Plus as it locally alternates the chemistry, microstructure and morphology near and at grain boundaries. 4.3. The influence of ı-phase on oxidation Similar frequency dependent oxide colour bands as the ones observed in this study for the Standard condition have been reported by, e.g. Ohmura et al. [13] for Co-based superalloy HS188. In the referred study the colour bands were attributed to the oxide thickness. However, this statement was not supported by any experimental evidence. Our results further clarify thicknessand compositional attributes of the oxides for Allvac 718Plus. The type 1 and 2 oxides found can be related to the oxidation of ␦-phase and ␥ matrix respectively. The Nb-rich ␦-phase phase is shown to reside primarily at grain boundaries in the Standard condition. In the mostly intergranular hold-time fracture surfaces of this condition a higher area fraction of the thicker type 1 oxide was found, characterized by a Nb-rich inner part. An Al-rich oxide layer was also always found for this oxide type, explainable by the elevated levels of Al found for the ␦-phase and the high O affinity of Al. Furthermore, little or no Cr is present in neither the ␦-phase nor in the type 1 oxide. In the mostly transgranular 10 Hz cyclic loading areas in the standard condition the average oxide was found to be thinner with a Cr-rich inner layer, which correlates to a larger area fraction of type 2 oxide. The enrichment of Cr is ascribed to the matrix being capable of supplying Cr more readily in the event of transgranular fracture. The AES depth profile for type 2 oxide in Fig. 15(b) confirms a slight enrichment of Nb below the Cr-oxide also for this type of oxide. This can be explained by the ␥ matrix not being fully depleted of Nb allowing for a thin Nb-rich oxide layer. Also for the laterally less confined XPS analysis, as the crack propagation is only mostly transgranular for 10 Hz loading in the standard condition there are quite possibly also some areas of intergranular fracture surfaces with oxidized ␦-phase included in the analysis. The intergranular crack surfaces of the Clean condition have similarities to that of the type 2 oxide in the Standard condition. In the Clean condition Nb-levels in the matrix are higher than in the Standard condition matrix due to the solution of the Nb-rich ␦-phase by heat treatment. This means more Nb is readily available during oxidation of the matrix in the Clean condition. The observed difference in oxidation of ␦-containing and ␦-free grain boundaries in the Standard and Clean conditions respectively indicate that if oxygen diffusion or oxidation ahead of the propagating crack does indeed occur, it is affected by the presence of this phase. In other words, the course of events for DE or SAGBO is different in the Standard and Clean conditions. 4.3.1. Transition zones The transition zones observed in the results from mechanical testing, up to 1 mm in length, are highly interesting. It appears as if a gradient damage has occurred ahead of the main crack, possibly by environmental effect. Once the main crack propagates through this damaged region by 10 Hz cyclic loading, the CGRs successfully decline (reflecting the gradient damage) until reaching a steady rate of growth. The zones increase in size with increasing K and were at a given K consistently larger for the Clean condition. No connections were found between the size of the transition zones and microstructural features (e.g. grain size) of either condition or inter-precipitate distance of grain boundary ␦-phase in the Standard condition. Ma and Cheng [14] have reported similar effects for Alloy 783, as have Pfaendtner and McMahon [5] for Alloy 718. However, these

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reports have contradicting interpretations concerning the mechanism for the occurrence of these zones. While Ma and Cheng [14] ascribe this to SAGBO, Pfaendtner and McMahon [5] point out DE as the responsible mechanim. In the model presented for DE the damaged zone is said to consist of some embrittled grain boundaries ahead of the main crack, mixed with uncracked ligaments. This model, suggesting partial cracking along preferential grain boundaries (explained in more detail in McMahons comments [15] on the work by Ma and Cheng), would allow an explanation of the mixed mode crack propagation observed for the cyclic regions in the Standard condition. Whether these mixed zones were created ahead of the main crack tip could not be discerned in the present study as the samples were fully fractured (not allowing for the same fractographic methods as in [5]). In the Clean condition no correlations could be made to uncracked ligaments as here also 10 Hz cyclic loading exhibits fully intergranular fracture. It is worth considering that the temperatures used in the referred work [5,14] were up to 650 ◦ C while in the current work 704 ◦ C was used. For both SAGBO and DE thermal activation has been shown to be of great importance and also creep may play a greater role at this higher temperature. The higher temperature used here may alter the extent of damage and/or the damaging mechanisms themselves. Furthermore, the oxidation phenomena investigated here allow no basis for distinction between the proposed mechanisms. Further work is being performed on this subject. 4.4. Additional aspects There are other factors that may have an influence on crack propagation, such as grain boundary carbides, grain size and grain boundary segregation. In this study the oxidized Nb-carbides found on fracture surfaces are relatively small and rare, thus not expected to play a major role. Moreover, the average grain sizes of the conditions used in this work differ by 18 ␮m which unlikely is the dominant factor for the differences observed. Several observations were made of grain boundary ␦-phase fractured by secondary cracks in the Standard condition. These observations show that deflection of the crack path by ␦-phase is not fully consistent. No segregation of elements to the grain boundaries nor the absence of elements in these could be observed by SEM– and TEM–EDX techniques used. However, these techniques are limited in terms of light element analysis and thus full analysis would necessitate the use of other techniques capable of this, e.g. secondary ion mass spectroscopy (SIMS) or atom probe tomography (APT). For instance, grain boundary segregation of the light element B has been observed by SIMS for solution heat treated Allvac 718Plus [16] and B is known to affect the mechanical properties of superalloys. Consequently, further analyses by APT are being performed on Allvac 718Plus and Alloy 718. 5. Conclusions • The CGRs are time dependent for 100 s hold-times and mostly cycle dependent for 10 Hz loading. Fracture occurred intergranularly for all loading in the Clean condition and for hold-times in the Standard condition. For 10 Hz cyclic loading of the Standard condition fracture was mostly transgranular. • The presence of grain boundary ␦-phase in Allvac 718Plus is beneficial for arresting hold-time crack propagation whilst cyclic loading at 10 Hz is nearly unaffected. • In the Standard condition, the presence of ␦-phase causes the volumes surrounding the phase to be denuded of ␥ hardening precipitates, creating a local microstructure more ductile than the general bulk of ␥ hardened ␥ matrix. This is believed to have an impact on crack growth micromechanisms.

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• Two types of oxides were found on the Standard condition crack surfaces, correlated to loading and fracture modes; a Nb-enriched oxide was dominate for hold-times and a Cr-enriched oxide was dominate for cyclic loading. The enrichments are attributed to the oxidation of the Nb-rich ␦-phase for intergranular fracture and the Cr-rich ␥ matrix for transgranular fracture, respectively. For Clean condition only the Cr-enriched oxide type was found. This indicates that if occurring, the oxidation and O diffusion ahead of the propagating crack is different for the Standard and Clean conditions. • Up to 1 mm long transition zones of elevated yet declining CGRs were found when changing from hold-time to 10 Hz cyclic loading. It is believed that the hold-times cause damaged zones ahead of the main crack tip and the main crack propagates through these zones during 10 Hz cyclic loading.

Acknowledgments The authors wish to acknowledge; Mr Joel Andersson at Volvo Aero Corporation for his assistance with heat treatments and mechanical testing layout; Dr Thomas Hansson at Volvo Aero Corporation for his assistance with analyses of mechanical testing data; Professor Vratislav Langer at Chalmers University of Technology,

the Department of Chemical and Biological Engineering, for his assistance with XRD analyses. References [1] J. Andersson, G. Sjöberg, S. Hatami, ISABE 2007, Beijing, 2007. [2] X. Liu, S. Rangararan, E. Barbero, K.M. Chang, W.D. Cao, R. Kennedy, T. Carneiro, Proceedings of the International Symposium on Superalloys, Champion, PA, 2004, pp. 283–290. [3] X. Liu, J. Xu, E. Barbero, W.D. Cao, R.L. Kennedy, Materials Science and Engineering A 474 (2008) 30–38. [4] D.A. Woodford, Energy & Materials 1 (2006) 59–79. [5] J.A. Pfaendtner, C.J. McMahon Jr., Materials Science Forum 294–296 (1999) 743–746. [6] U. Krupp, C.J. McMahon Jr., Journal of Alloys and Compounds 378 (2004) 79–84. [7] W.-D. Cao, R.L. Kennedy, ATI Allvac, Monroe, NC, 2006. [8] J. Andersson, G. Sjöberg, L. Viskari, A. Brederholm, H. Hänninen, C.S. Knee, Proceedings of the 47th Conference of Metallurgists, Winnipeg, Canada, 2008, pp. 401–413. [9] Tada H. Irwin, Paul C. Irwin, R. George, The Stress Analysis of Cracks Handbook, third ed., 2000. [10] M. Norell, L. Nyborg, T. Tunberg, I. Olefjord, Surface and Interface Analysis 19 (1992) 71–76. [11] R.C. Reed, The superalloys: Fundamentals and Applications, Cambridge University Press, Cambridge, UK, New York, 2006. [12] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, Wiley, New York, 1987. [13] T. Ohmura, R.M. Pelloux, N.J. Grant, Engineering Fracture Mechanics 5 (1973). [14] L. Ma, K.M. Chang, Scripta Materialia 48 (2003) 1271–1276. [15] C.J. McMahon Jr., L. Ma, X. Liu, K.M. Chang, Scripta Materialia 54 (2006) 305–311. [16] K. Vishwakarma, M. Chaturvedi, Proceedings of the International Symposium on Superalloys, Champion, PA, 2008, pp. 241–250.