Acta Materialia 155 (2018) 362e371
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On the diffusive phase transformation mechanism assisted by extended dislocations during creep of a single crystal CoNi-based superalloy Surendra Kumar Makineni a, *, Ankit Kumar a, Malte Lenz b, Paraskevas Kontis a, Thorsten Meiners a, Christopher Zenk c, Stefan Zaefferer a, Gunther Eggeler d, Steffen Neumeier c, Erdmann Spiecker b, Dierk Raabe a, ***, Baptiste Gault a, ** a
Max-Planck-Institut für Eisenforschung GmbH, 40237, Düsseldorf, Germany €t ErlangenInstitute of Micro- and Nanostructure Research & Center for Nanoanalysis and Electron Microscopy (CENEM), Friedrich-Alexander-Universita Nürnberg, Cauerstraße 6, 91058, Erlangen, Germany c €t Erlangen-Nürnberg, Martensstr. 1, 91058, Erlangen, Germany Institute of General Materials Properties, Friedrich-Alexander-Universita d €t Bochum, Universita €tsstrasse 150, D-44 780, Bochum, Germany Institut für Werkstoffe, Ruhr-Universita b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 12 March 2018 Received in revised form 28 May 2018 Accepted 31 May 2018 Available online 14 June 2018
We propose here a deformation-induced diffusive phase transformation mechanism occurring during shearing of g0 ordered phase in a g/g0 single crystalline CoNi-based superalloy. Shearing involved the creation and motion of a high density of planar imperfections. Through correlative electron microscopy and atom probe tomography, we captured a superlattice intrinsic stacking fault (SISF) and its associated moving leading partial dislocation (LPD). The structure and composition of these imperfections reveal characteristic chemical e structural contrast. The SISF locally exhibits a D019 ordered structure coherently embedded in the L12 g0 and enriched in W and Co. Interestingly, the LPD is enriched with Cr and Co, while the adjoining planes ahead of the LPD are enriched with Al. Quantitative analysis of the threedimensional compositional field in the vicinity of imperfections sheds light onto a new in-plane diffusion mechanism as the LPD moves on specific {111} planes upon application of stress at high temperature. © 2018 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Superalloys Creep deformation and stacking faults Electron channeling contrast imaging (ECCI) Transmission electron microscopy (TEM) Atom probe tomography (APT)
1. Introduction The g/g0 microstructure in a Co-Al based system was first reported in 1971 [1], with a surge of interest in the past decade [2,3]. The L12 ordered g0 phase is stabilized by addition of W [1,4,5], Mo/ Nb [6,7] or Mo/Ta [8] into the binary Co-Al system. The microstructure resembles the well-established high-temperature Ni-Al based superalloys that are commercially used for e.g. hot gas turbine engines. The g matrix is a face-centered-cubic (fcc) solid solution while the g0 phase is ordered with a L12 structure that shares a coherent interface with the g matrix. The g/g0 lattice misfit in Co-
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses:
[email protected] (S.K. Makineni),
[email protected] (D. Raabe),
[email protected] (B. Gault). https://doi.org/10.1016/j.actamat.2018.05.074 1359-6454/© 2018 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Al based superalloys is positive [9e11] i.e. the g0 phase has a larger lattice parameter than the g matrix while most of the Ni-Al based superalloys have negative lattice misfit [12,13]. Because of this, the rafting of g0 during tensile creep deformation occurs parallel to the stress axis for Co-Al based superalloys (P-type) while it is perpendicular to the stress axis for Ni-Al based superalloys [14]. P-type g0 rafts are expected to reduce the creep rate because of a slower movement of the rate-determining glide-climb interfacial dislocations along vertical g/g0 interfaces [9,14,15]. They also act as strong obstacles for fatigue crack growth, thereby enhancing fatigue resistance [14,16]. Several new compositions of Co-Al based superalloys and their mechanical/creep properties [17e21], microstructural stability, mass density, solvus temperature [22e26] and oxidation resistance [27e29] were investigated. The addition of Ni increases the solvus temperature and reduces the g/g0 lattice misfit to a smaller positive value [30,31] or even leads to negative lattice misfits [32]. It also
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allowed further transition metal additions to this alloy class by increasing the stability of the metastable g0 phase-field that is usually narrow in Co-Al-W [33,34] or Co-Al-Mo-Nb/Ta alloys [35]. The new CoNi-based superalloys showed significant improvement in creep resistance at 900 C with creep properties comparable to 1st generation Ni-Al based superalloys [10,36,37]. Creep deformation at 900 C of CoNi-based alloys causes shearing of the g0 phase by a/2 matrix dislocations [38,39], creating antiphase boundaries (APBs) in their wake. Later, it was also shown that a/3 partial dislocations create a superlattice intrinsic stacking fault (SISF) during the shearing of the g0 phase, followed by glide of an ensuing a/6 trailing partial dislocation transforming the SISF into an APB [40]. Additionally, chemical enrichment to these planar faults was observed by energy dispersive spectroscopy (EDS) in transmission electron microscopy (TEM), with evidence of local reordering and solute diffusion occurring during g0 shearing [40,41]. Similar evidence of solute segregation to creep-induced stacking faults were reported recently in Ni-Al based [42e44] and Co-Al based superalloys [41,45]. This segregation led us to the proposition that the rate limiting step for creep deformation is directly linked to the diffusivity of solute species to these lattice defects. Here, we reveal the atomic structure and local composition of a superlattice intrinsic stacking fault (SISF) and its associated leading partial dislocation (LPD) in three dimensions. The partitioning of the solute and the compositional gradients in their vicinity led us to propose an in-plane diffusive phase transformation mechanism operating under stress at high temperature. This phenomenon is characterized by the confined transformation of L12 into D019 inside of stacking faults. The mechanism is assisted by the displacement of the LPD on specific {111} planes that drives solute partitioning along the fault plane. The effect is evidenced in a Co-Ni-Al-W-based single crystal g/g0 superalloy. The material also contains Ti and Ta for improved high temperature stability and Cr for better high temperature oxidation resistance. A correlative approach combining controlled electron channeling contrast imaging (cECCI) [46], conventional and high-resolution scanning transmission electron microscopy (HR-(S)TEM) and atom probe tomography (APT) was used to explore the full chemical-structural nature of the individual features associated with this extended defect. 2. Experimental 2.1. Materials and alloy preparation A CoNi-based single crystal alloy, called ERBOCo-1, with composition Co-32Ni-8Al-5W-6Cr-2.5Ti-1.5Ta (at.%) with small additions of 0.1Hf and 0.4Si was prepared by the Bridgman process. The heat treatment steps were 1280 C/8 h þ 1050 C/5 h þ 900 C/ 16 h þ water quenched to obtain a uniform g/g0 microstructure. Tensile specimens were cut and crept at 850 C for 380 h at an applied stress of 400 MPa along the [001]-direction up to 4.6% plastic strain. Samples oriented close to the [110] direction were cut and mechanically/electrochemically polished for further microstructural characterization. 2.2. Electron channeling contrast imaging (ECCI) A Zeiss Merlin scanning electron microscope (Carl Zeiss SMT, AG, Germany) equipped with a Gemini-type field emission gun electron column was used to perform electron backscattered diffraction (EBSD) and ECCI measurements. The operating accelerated voltage of was 30 kV with a probe current of 4 nA. The working distance was kept at 6 mm for imaging. Controlled ECCI is a combination of
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EBSD orientation mapping and simulation of electron channeling pattern (ECP) for different sample tilts using TOCA (Tools for Orientation Determination and Crystallographic Analysis) program [47,48]. First the region of interest (ROI) on the bulk sample surface is located and the orientation measured by EBSD. The crystal orientation information is fed to TOCA that calculates the accurate settings for the sample stage of the microscope that allows to orient the crystal in a specific two-beam conditions. Using this orientation, diffraction patterns were calculated using TOCA that in turn provides the tilt and rotation angles needed to orient the crystal into optimal diffraction conditions. 2.3. Correlative transmission electron microscopy (TEM) and atom probe tomography (APT) Specimens for correlative investigation by TEM and APT were fabricated using a dual beam SEM/focused-ion-beam (FIB) instrument (FEI Helios Nanolab 600) via an in-situ lift-out protocol. The adopted detail procedure can be found in Ref. [49]. Regions containing imperfections were extracted from the bulk and subsequently attached onto the electropolished tips of a halved TEM Mogrid that was to mounted in a special correlative holder designed in-house [50]. These regions were sharpened by FIB milling at 30 kV followed by a final cleaning procedure at 2 kV and 16 pA current to remove severely damaged regions induced previously by a highenergy Ga ion beam. TEM on the APT needle-shaped specimens was carried out using a Philips CM-20 instrument operated at 200 kV. High resolution imaging was carried out in STEM mode using an aberration-corrected FEI TITAN microscope operated at 300 kV, similarly to the protocols described in Ref. [51]. Nearatomic-scale compositional analysis was done using atom probe tomography (APT). APT measurements were conducted using a LEAP™ 5000X HR (Cameca Instruments), equipped with a reflectron lens. The instrument was operated in laser pulsing mode at a pulse repetition rate of 200 kHz and a pulse energy of 40 pJ. The specimen's base temperature was kept at 40 K and the target detection rate was set to be 5 ions detected every 1000 pulses. Data analysis was performed using the software package IVASTM 3.8.0. 3. Results 3.1. Microstructure after creep deformation ehigh density of stacking faults Fig. 1(a) shows an overview cECC image of a metallographic cross section from the bulk crept sample surface with its crystal oriented near to the [110] direction of the cubic lattice. A strong two-beam diffraction condition with g ¼ 002 (where g is the diffraction vector) was established by controlled rotation/tilt of the crystal. The corresponding stereogram is also shown in Fig. 1(b). The image clearly shows atomic density contrast between the fcc g (dark) and the rafted L12 g' (grey) phases parallel to the tensile loading axis. A higher magnification image in Fig. 1(c) shows significant channeling contrast from defects that appear as bright intensity-oscillating areas in the rafted g0 phase. These bright areas were marked as A, B, C and D, respectively, according to their crystallographic orientation relative to the load axis. These are typical characteristics of stacking faults (SFs) where there is a relative shift of atomic columns above and below the fault plane [46,52]. The bright line at the intersection of the SF and the sample surface occurs due to the rapid drop in backscattering signal obtained from a SF with increasing depth of the SF below the surface. For types A and D, the contrast fades with distinct intensity oscillations towards the left and the right, respectively. The wavelength
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Fig. 1. (a) An overview of the crept microstructure: cECC image of the crept CoNi-based superalloy taken under two-beam diffraction condition near the [110] pole of the cubic lattice showing rafted g0 phase parallel to the tensile loading direction and (b) Corresponding stereogram of plane normals viewed in the same orientation as the specimen's surface. (c) High magnification cECC images taken two-beam conditions (c) with g ¼ 002 e All SFs are visible and have been marked A, B, C, D according to their plane normals in stereogram highlighted by thin blue, red, green and golden lines (d) with g ¼ 111 e SF B is invisible (e) g ¼ 111 e SF C is invisible and (f) g ¼ 220 e SFs A and D are invisible. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
of the intensity oscillations corresponds the extinction distance,
xeff g ; similar to the SF observations in transmission electron microscopy [46]. From the stereogram, Fig. 1(b), their traces were identified to pertain to {111} cubic planes with displacement vectors of the type R ¼ ±13 112. This magnitude of the stacking fault vector is assumed based on the detailed analysis carried out by Eggeler et al. [40]. Using large angle convergent beam electron diffraction (LACBED) in a similar crept CoNi-based superalloy. The faults A, B, C and D are associated with the planes ð111Þ, ð111Þ, ð111Þ, and ð111Þ respectively and have been confirmed through g.R invisibility criteria. Fig. 1(def) show ECCI micrographs taken using two beam 111, 111 and 220 conditions near the [110] zone direction respectively. In Fig. 1(d), the B type SF fulfills g.R ¼ 0 and becomes invisible. Similarly, the C and A/D type becomes invisible for the 111 and 220 two-beam conditions. From the stereogram, Fig. 1(b), aligned according to the sample orientation, the
stacking faults (SFs) A and D belonging to ð111Þ and ð111Þ respectively, forms angles of 38 and 36 with the surface plane. In contrast, the SFs B and C appear as sharp lines (with no fading contrast) and belong to ð111Þ and ð111Þ with the angles 80 and 81 from the sample surface plane, respectively. Thus, the B and C SFs appear as sharp white straight lines since the view is nearly edge-on. 3.2. Crystallography and atomic structure of a stacking fault An APT specimen was extracted from a region like the one indicated by the thin dashed yellow line in Fig. 1(c), and in this specific orientation. Fig. 2(a) shows a bright-field (BF) TEM image of the as-prepared APT specimen taken in [110] zone axis. The two stacking faults, SF1 and SF2, are visible and are oriented in the same direction, each exhibiting two darker contrast lines. The line contrast originates from the intersection of the SF plane with the curved surface of the APT specimen. To identify the type of stacking
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Fig. 2. (a) A TEM BF image of an APT tip taken along [110] zone direction showing two dark contrast lines corresponding to planar stacking faults SF1 and SF2 (b) BF images taken in two-beam conditions with g ¼ 220 (c) g ¼ 002 (d) g ¼ 111 (e) g ¼ 111 where the SFs are invisible or exhibit very weak contrast.
fault (A, B, C or D type), similar g:R ¼ 0 or ±n invisibility criterion is used. Fig. 2(bee) show bright field micrographs of the same tip imaged under different two-beam diffraction conditions. Both the faults become effectively invisible when imaged with g ¼ 111 and exhibit very weak (only show residual) contrast in g ¼ 220 condition. Hence, the visible SFs correspond to those labelled C in Fig. 1(c).
To obtain atomically-resolved structural information, aberration-corrected HR-STEM was performed directly on the APT specimen. Fig. 3(a) and (b) shows the BF TEM and a lowmagnification HAADF-STEM image of the specimen taken close to the [110] zone axis showing the stacking fault edge-on within the thin dashed yellow square. The stacking fault is visible as a line with a bright contrast. Fig. 3(c) shows an aberration corrected HR-STEM
Fig. 3. (a) A TEM BF and (b) STEM HAADF image of the APT tip taken along [110] zone direction showing bright contrast line corresponding to planar stacking fault SF2 (c) High resolution STEM HAADF image showing different atomic structure at the SF from the surrounding lattice (d) Simulated atomic structure of SISF with D019 structure coherently embedded in L12 ordered lattice (e) The HR-STEM image showing slightly increased intensity (Z-contrast) in SISF indicating segregation of heavy elements to the SF (f) The sequence of atomic planes across the SF showing intrinsic nature of the fault plane (missing A plane).
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image of the fault with atomic-column resolution. The atomic Zcontrast is clearly visible along the fault with respect to the surrounding lattice. The HR-STEM image in Fig. 3(e) show slightly increased intensity (Z contrast) in the SF region indicative of segregation of heavy alloying elements. Notation A, B and C in Fig. 3(f) corresponds to atomic planes. Outside of the fault, an FCC ABCABC stacking sequence is present while at the fault, HCP BCBC stacking was observed. This sequence confirms the presence of a local D019 structure corresponding to a superlattice intrinsic stacking fault (SISF) coherently embedded in an L12 ordered lattice. The intrinsic nature of the fault is confirmed from the missing A plane in the SF region. A schematic view of the atomic structure of the SISF with DO19 ordering [40,41] embedded in the L12 ordered g0 phase is shown in the form of a [110] projection in Fig. 3(d). 3.3. Compositional gradients in the vicinity of a stacking fault Subsequent to TEM, APT was carried out on the specimen to obtain compositional information at the near-atomic scale. Fig. 4(a) shows a bright field (BF) TEM image of the as-prepared APT specimen taken in a two-beam diffraction condition with g ¼ ð111Þ near to the [110] zone axis. Fig. 4(b) shows the reconstructed APT point-cloud with only a fraction of the Co atoms displayed. The overall composition is that of a g0 precipitate. We observed a highly confined linear Cr decoration at the bottom of the tomogram, highlighted by an iso-composition surface, shown in pink, with a threshold of 5.6 at.% Cr. The overlay of the 3D atom map and the BF image, as displayed in Fig. 4(c), points out that it belongs to SF2 and illustrates that the confined linear Cr decoration occurs along a partial dislocation. Further validation came from plotting 2D elemental composition maps projected onto the faces (xy, xz, yz) of a cuboidal region-of-interest centered on the feature. Fig. 4(d) shows an example for the solute Al with the corresponding composition color-scale. Additional projections perpendicular to the yz of the 3D point cloud highlights a local depletion of Al along a straight plane terminated by a chemically distinct linear dislocation distinct from the ordered L12, where an enrichment of Cr atoms is readily visible. We successfully captured in a single APT experiment
the SISF, its leading partial dislocation (LPD) and part of the parent L12 lattice ahead of the LPD. However, due to the limited field-ofview of APT, the trailing partial dislocation (TPD) was not observed. Inside the SISF the compositions of both W and Co increase by approx. þ1.1 at.% and þ1.6 at.%, respectively, while Al and Ni are depleted by approx. 1.2 at.% and 1.9 at.% relative to the composition of the surrounding precipitate (Fig. 5(b)). There is no significant variation in the compositions of Cr, Ti or Ta. In contrast, at the bounding LPD (Fig. 5(c)), the Cr and Co compositions increase by ~ þ4.1 at.% and þ4.8 at.%, respectively, while Al and Ni are depleted by 2.4 at.% and 4.2 at.%, respectively. Additionally, W and Ti were also found to be depleted by approx. 1.4 at.% and 0.5 at.% relative to the surrounding lattice. The composition profile in Fig. 5(d), measured along the SISF plane (marked as CC0 in Fig. 5(a)), shows that in the adjoining g0 planes, the composition in Al increases ahead of the leading partial over approx. 4 nm and drops thereafter, in contrast with the W content that first drops and then increases. The compositions of the surrounding L12 lattice, SISF and LPD are summarized in Table 1. 4. Discussion The cECC imaging shows the plastic deformation for the present CoNi-based superalloy ERBOCo-1 during tensile creep at 850 C occurs through formation of high density of stacking faults on {111} planes by shearing of g0 phase. Quantitatively, the number density is estimated to be ~2.6 (±0.6) x 1013 m2 measured with the method introduced in Ref. [47] by using multiple ECC images. We have shown a successful implementation of correlative approach, from selecting a single imperfection in a bulk sample, determining its full crystallographic nature by electron microscopy techniques, as well as the three-dimensional compositional field in its vicinity by APT. We captured and measured the near-atomic-scale composition of the SISF, its associated LPD and the adjacent L12 lattice ahead of the partial. The details pertaining to the atomic structure of the LPD by HRSTEM could not be accessed in this investigation but will be studied in the future. The compositional variations allow two conclusions to be drawn. First, the solutes
Fig. 4. (a) BF TEM image of an APT specimen taken in a two-beam condition with the diffraction vector g ð111Þ near to the [110] zone direction showing SF planes. (b) APT reconstruction point-cloud for the same APT specimen showing confined linear Cr decoration at the bottom of the tomogram. (c) Overlay of a region-of-interest (ROI) within the APT data and BF TEM image. (d) 2D elemental compositional maps projected onto the xy, xz and yz planes as faces of a cuboidal ROI centered on the SF2 for the solute Al with the corresponding composition color-scale. The distribution of Al (brown) and Cr (pink) atoms viewed perpendicular to the yz plane are also shown. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 5. (a) Schematic illustration of the planar imperfection consisting an SISF (dark orange), leading partial dislocation (LPD, pink), Al solute pileup region ahead of the LPD (dark green) and the L12 lattice (green). Solute composition profiles across (b) the SISF (AA0 ), (c) leading partial dislocation (LPD, BB0 ) and (d) along the fault plane (CC0 ). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
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Table 1 Compositions measured by APT for the surrounding L12 lattice, leading partial dislocation (LPD) and the stacking fault (SISF).
g0 L12 Precipitate Leading Partial Dislocation (LPD) Stacking Fault (SISF) Co Ni Cr Al W Ti Ta
40.1 ± 0.5 34.7 ± 0.4 3.2 ± 0.2 9.7 ± 0.3 6.5 ± 0.2 3.6 ± 0.2 0.6 ± 0.1
44.9 ± 0.8 30.5 ± 0.7 7.3 ± 0.3 7.3 ± 0.4 5.1 ± 0.3 3.1 ± 0.2 0.5 ± 0.1
41.7 ± 0.5 32.8 ± 0.4 3.4 ± 0.3 8.5 ± 0.3 7.6 ± 0.3 3.5 ± 0.2 0.7 ± 0.1
partition inbetween the SISF, the LPD and the adjacent bulk region, tending towards local chemical equilibrium. Second, the compositional gradients of solutes along the fault plane reflects sluggish diffusion of some of the solutes such as W. The solute pile up ahead of the LPD provide evidence of solute diffusion and its direction. Both aspects indicate a transient state during exposure to a high temperature of ~850 C and creep loading such as in the current experiments, as depicted in Fig. 6. The change in the local composition ahead of the partial can have a dramatic effect on its local stability and ease of displacement of the partial dislocation. The obtained structural information of the planar defect shows that it corresponds to a SISF. Recently, using detailed conventional and HR-(S)TEM experiments, Eggeler et al. [40] revealed shearing of g0 in a similar CoNi-based superalloys and proposed a ½{111} shear mechanism [53,54]. Assuming that the same shearing mechanism is associated with the present alloy system, we elucidate the local solute diffusion process occurring during g0 shearing. Fig. 6(aec) outlines the genesis of the planar imperfection and some of the associated kinetic consequences during high temperature creep deformation. It shows a set of parallel ð111Þ planes containing a coherent interface shared between the g matrix and the g0 L12 precipitate (Fig. 6(a)). Under the influence of the applied stress, first the matrix dislocations interact at the g/g0 interface and dissociate into a LPD and a TPD such that the LPD moves inside the g0 precipitate (Fig. 6(b)) and the trailing one remains at the interface. The overall reaction at the g/g0 interface involves a known
dissociation where 1 ½101 2
reaction
1 ½101 2
þ 12 ½011/13 ½112 þ SISF þ 16 ½112
1 ½112 and 1 ½112 are leading and trailing partials while 3 6 and 12 ½011 are g fcc matrix dislocations [53,54]. As
mentioned in section 3.3, we were able to capture the SISF and its associated LPD, but not the TPD. Hence, we reduce here the discussion to the formation of the SISF and to the solute diffusion directions across the LPD. Further effects associated with these mechanisms such as the TPD entering the g0 precipitate and the associated conversion of the SISF into an antiphase boundary (APB) are discussed in Ref. [40]. The SISF forms, as indicated in Fig. 6(c), a region with a DO19 ordered structure on an a hexagonal-close-packed (hcp) lattice within the surrounding fcc-derived L12 ordered lattice structure [40,41]. A Cottrell atmosphere enriched with Co and Cr with respect to the surrounding L12 ordered lattice structure develops in the vicinity of the LPD. Cr and Co are stabilizers of the g matrix and may have been carried by the LPD as it crossed the g/g0 interface. This change in the atomic arrangement will lead to a high-energy structural configuration that drives the phase dependent compositional partitioning of the solutes. The composition profiles along the fault plane and the solute pile up ahead of the LPD (Fig. 5(d)) demonstrate in-plane diffusion of solute W and Co from the region ahead of the leading partial towards the newly formed DO19 ordered SISF behind, Fig. 6(c). In contrast, Al/Ni diffuse in opposite direction from the D019 ordered SISF into the region ahead of the LPD. This selective diffusive solute migration is driven by the thermodynamic trend to form, at the SISF behind the LPD, a DO19 ordered structure where the stoichiometry Co3W is more stable, thereby leading to a reduction in the stacking fault energy (SFE) [41,45]. Since, W and Al occupy the same lattice site in the ordered structure, W solute atoms prefer to replace only Al atoms at their sites in the DO19 structure. Several recent reports show evidence of chemical fluctuations in stacking faults [40,41,45,55] and at the bounding partial dislocation [56] through energy dispersive X-ray spectroscopy (EDS) in aberration-corrected TEMs in STEM mode. Such observations were enabled by the development of the Super-X EDS system (ChemiSTEM) which is equipped with four silicon drift
Fig. 6. Deformation mechanism: (a) Perfect g fcc dislocation approaching the g/g0 interface. (b) Dissociation of the dislocation into two partials when crossing the g/g0 interface with the leading partial becoming enriched with Cr and Co with respect to the L12 ordered surrounding lattice. (c) Shearing of the ordered precipitate with the movement of the leading partial dislocation (LPD). It leaves behind a SISF with DO19 ordered structure. Subsequent to this process in-plane diffusion of W/Co atoms towards the SISF and Al/Ni diffusion in opposite direction (towards adjoining g0 L12 lattice) take place. This results in a solute pile up ahead of the LPD (dark green region). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
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detectors (SDD) and thus provides high signal collection efficiency [57]. The experiments were carried out while observing the stacking faults edge on (mainly along a zone axis) with chemical measurements performed across the fault plane. Since the thickness of the faults is in the range of 2e3 nm, EDS probing in a TEM may encounter certain limits in identifying chemical fluctuations along the fault plane, a feature type that can indeed be resolved by the current correlative approach combining TEM and APT analysis. Diffusion of solutes to the faults, until now, were discussed in terms of three mechanisms, namely, 1) diffusion occurring perpendicular to the fault plane, i.e. from the bulk of the surrounding lattice [40,41]; 2) diffusion occurring along the fault, i.e. to and from the g channel; and 3) diffusion along the dislocation core, i.e. pipe diffusion [58]. However, as mentioned earlier, these were never proven due to the composition detection limits discussed above. As examples, in a commercial Ni-Al based superalloy (CMSX-4), Viswanathan et al. [42] showed solute segregation of Co/Cr/W and depletion of Ni/Al at SISFs. The authors proposed, out of these three mechanisms, mass transport occurs through pipe diffusion via the extended shearing partial dislocations. Following this, Smith et al. [55,56] investigated the microstructures of two crept Ni-Al based superalloys (ME3 and ME501) and proposed that the stacking faults in ME3 show segregation of g stabilizers, namely Co/Cr/Mo, and in ME501, the fault enriches with Nb/Ta/Ti promoting the formation of new h phase with an ordered DO24 crystal structure. Additionally in ME3 alloy, they report Co/Cr enrichment (as g fcc atmosphere) near to the leading partial and suggested that this effect may avoid formation of high energy complex stacking faults (CSFs). The driving force for the observed solute segregations was attributed to lowering of the stacking fault energy (SFE), however, the actual solute diffusion mechanisms were not revealed. Similarly, in CoNi-based single crystal superalloys, Titus et al. [41] and Eggeler et al. [40] showed Co/Cr/W enrichment and deficiency of Ni/Al at SISF. Both the groups suggested that the diffusion of these solute species occurs perpendicular to the fault plane (mechanism (1)) inbound from the surrounding bulk g0 lattice. Here, we show evidence of solute pile up ahead of the LPD, which implies that solutes diffuse along the fault plane, thereby constituting an alternative diffusional mechanism. More specifically, the APT results indeed reveal in-plane solute diffusion inbetween the SISF and the L12 lattice ahead of the LPD. This solute migration reduces the SFE and hence facilitates further movement of the leading partial during creep deformation. Simultaneously, the stacking fault area will increase, behind the moving partial, with an hcp-based D019 ordering and hence drives further in-plane solute diffusion of W and Co towards the fault while Al and Ni are repelled and diffuse in opposite direction. Hence the rate of g0 shearing depends on the rate of diffusion of segregating solutes to the defect. For an estimate of the diffusion length required for this segregation mechanism, the diffusivity of Al, W and Cr were considered. From the experimental activation energy (Q) and frequency pre-factor (D0) reported for Al and W in Ref. [59], the interdiffusion coefficients at 850 C were calculated to be 3 1017 m2 =s pffiffiffiffiffiffi and 2 1018 m2 =s respectively. Using the relationship x ¼ 2 Dt Al diffuses in 1 s over a distance of ~5.5 nm while W travels only ~1.4 nm. As expected, W is the slowest diffusing solute. However, the movement of the bounding partial dislocation, being decorated by a Cr atmosphere, also depends on the diffusivity of the Cottrell solute through the lattice during shearing. This was postulated by Titus et al. [41] who did not show the evidence of a Cottrell cloud of solute at the shearing LPD. The velocity of the shearing partial dislocation, dragged by the Cottrell cloud, was estimated using the
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equation [41,60]
wPD ¼
ðtbÞDkB T 2:1C0 b2
where, wPD is the partial dislocation velocity, t is the relative shear stress on the Cottrell solute atmosphere, b is the magnitude of the Burgers vector, D is the diffusivity of the solute, kB is the Boltzmann constant, T is the temperature (K), C0 is solute composition in the Cottrell atmosphere and b is a factor related to the interaction energy between the dislocation and the solutes. The inter-diffusion coefficient for Cr solutes at 850 C was taken as ~12.83 1018 m2/s (DCr) [59]. By taking the values: t ¼ 188.4 MPa, b ¼ 0.293 nm, T ¼ 1123 K (850 C), C0 ¼ 5.8 1027 m3 (for 7.3 at.% Cr in Cottrell cloud) and b ¼ 1.6 1029 Nm2, the partial dislocation velocity is estimated to be ~3.3 nm/s. This is approximately ~2.5 times higher than the diffusing distance estimated for the slowest solute W. Hence, the slowest and thus rate limiting process is the diffusion of W at 850 C. This is in contrast to the results reported by Smith et al. [44] in the ME3 Ni-Al based superalloy. Indeed, they have shown that the Cottrell cloud of Co/Cr slow down the shearing partial dislocation and act as the principle late limiting factor for the formation of stacking fault during intermediate temperature creep (760 C). However, they observed the same solutes (Co/Cr) enrichment even in the stacking faults and found their diffusion towards the stacking fault is faster than the Cottrell cloud. In the present case, we observe segregation of W to the stacking faults which is found to be the slowest diffusing solute at 850 C in comparison to even the velocity of shearing LPD with Co/Cr Cottrell cloud. In summary, from the microstructure (Fig. 1(a)), it is readily visible that shearing occurs simultaneously throughout the alloy. The associated solute segregation to the faults reduces the SFE, further facilitating the movement of the leading partial, which in turns lowers the external stress required for g0 shearing and hence larger plastic deformation can be accommodated in the microstructure during high temperature creep without failure. 5. Conclusions We revealed the near-atomic structure and composition of crystalline imperfections by TEM and APT in a creep-deformed CoNi-based alloy. We proposed a new solute diffusion mechanism occurring along the fault plane such that the W/Co migrates towards the SISF, while Al/Ni are pushed ahead of the LPD. Segregation of W and Co to the fault lowers the SFE and facilitates further movement of the LPD. Additionally, Cr/Co enrichment was found to the LPD that induce a solute drag effect during its movement. The diffusivity of W is found to be the slowest and thus represents the rate limiting process in the g0 shearing process. Acknowledgments The authors are grateful to U. Tezins and A. Sturm for their technical support of the atom probe tomography and focused ion beam facilities at the Max-Planck-Institut für Eisenforschung. ML and ES thank Yolita Eggeler and Julian Müller for support on the HR STEM investigations and fruitful discussions. SKM and CHZ acknowledge financial support from the Alexander von Humboldt Foundation. The authors acknowledge financial support from the DFG SFB TR 103 through projects A1, A4, A7 and B3. SKM, BG and DR are grateful to the MPG for the funding of the Laplace Advanced Atom Probe Tomography project.
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