Strain induced directional coarsening in nickel based superalloys: Investigation on kinetics using the small angle neutron scattering (SANS) technique

Pergamon PII:

S1359-6454(97)00035-9

A<,/u m(lr<‘,~ Vol. 45, No. X. pp. 3177-i:!X2, 1997 !’ 1997 Act3 Metallurgica Inc. Published by Elsemer Science Ltd. All rights reserved Printed in Great Britain 1359-6454 97 s17 00 + 0.00

STRAIN INDUCED DIRECTIONAL COARSENING IN NICKEL BASED SUPERALLOYS: INVESTIGATION ON KINETICS USING THE SMALL ANGLE NEUTRON SCATTERING (SANS) TECHNIQUE M. Vl?RON’

and P. BASTIE’

‘L.T.P.C.M.. UMR CNRS no. 5614 BP 75. 38402 Saint Martin d’Htres cedex and ‘Laboratoire de Spectrometrie Physique, UMR CNRS no. 5588. BP 87. 38402 Saint Martin d‘Htres cedex. France

Abstract-Using the small angle neutron scattering technique, we have observed rafting in nickel based single crystal superalloys. Kinetics of morphological evolution of the precipitates have been studied in sitar. Therefore we used a special furnace designed for the ageing of prestrained specimens under a neutron beam. The evolution of both the precipitate aspect ratio and the distance between precipitates confirms the importance of strain in the directional coarsening process, Results are presented and discussed regarding kinetics and microstructural aspects. In such conditions, rafts seem to be different from those obtained after a creep test. tJ: 1997 Acta Metallurgica Iw. RBsumGNous avons observk la coalescence orient&e dans les superalliages monocristallins a base de nickel en utilisant la technique de diffusion aux petits angles des neutrons. La cinetique de l’evolution morphologique des precipites a Cte ttudiee “in situ”. Pour cela. nous avons utilist un four developpe sous le faisceau, d‘un echantillon specialement. pour effectuer des traitements de vieillissement, predeforme. L’evolution des facteurs de forme des prtcipites et de la distance moyenne entre ces precipiter confirme l’importance de la predeformation dans le processus de coalescence orientbe. La presentation et la discussion des rtsultats concernent a la fois les aspects cinetiques et microstructuraux. Les radeaux obtenus dans ces conditions semblent differents de ceux resultant d’un essai de fluage.

INTRODUCTION

Owing to their good mechanical properties at high temperature, nickel based superalloys are commonly used for the manufacture of single crystal turbine blades. These mechanical properties are related to the precipitation of an ordered 7 phase inside they matrix. The morphology of the precipitates depends strongly on the thermomechanical history of the sample. A well-known phenomenon is the directional coarsening of the precipitates which occurs at high temperature (typically 1323 K) under stress and results in a strong anisotropy of the microstructure [ 1,2]. After standard heat treatment, the microstructure of those alloys is composed of a matrix containing about 70% of well-aligned coherent cuboidal precipitates, with an average size of 0.45 pm [3]. Precipitates (ordered ;” phase) and matrix (7 phase) present a misfit defined as 6 = 2(q’ - q)/(q’ + a~) where a;” and q are the lattice parameters of the y’ and ; phases, respectively. Directional coarsening occurs under stress and, for example, during a tensile creep test on the ,commercial AM1 superalloy, the microstructure evolves to form “rafts” of precipitate phase perpendicular to the stress axis [4]. Other orientations of rafts with respect to the stress axis are 3277

possible: when they are parallel to the stress axis directional coarsening is called type P. and when they are perpendicular to the stress axis. type N [5]. The coarsening process has been extensively studied both experimentally and theoretically. The direction of rafts has been related to the sign of the stress (the developed microstructure is not I he same after a tensile and a compressive creep test), to the misfit sign, and to the elastic constant of both phases [6, 71. Whereas the microstructures after directional coarsening are well characterised and pertinent parameters for rafting have been put forward, the driving force is still a matter of controversy. Most of the theoretical models deal with elastic energy calculation without taking into account elastic or plastic strain. The effects of elastic strain in a regime where there is no dislocation have been shown to lead to directional coarsening [8]. However. for the temperature and stresses involved, creep does take place and therefore dislocations are induced in the material and cannot be neglected. In a plastic regime. it has been shown that inhomogeneous misfit relaxation by plastically induced dislocations provides the main driving force for the phenomenon [9], and moreover raft direction can be predicted owing to the position of dislocations and their action on

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and BASTIE: DIRECTIONAL COARSENING

Table 1. Chemical composition (in wt% and in at.%) of the AM1 superalloy Ni AM1 (wt%) AM1 (at%)

64 65.6

Ta

Cr

Co

W

Al

MO

Ti

8 1.5 2.7 8.7

6.5 6.6

5.5 5.3 1.8 11.8

2.0 1.3

1.2 1.5

coherency stresses [IO]. However, the detailed kinetics of the rafting phenomenon has not been extensively studied up to now. Recent experiments have proved that applied stresses are not necessary to obtain directional coarsening and that a simple prestrain performed at rather low temperature (typically 1123 K), could induce the phenomenon which will occur rapidly during an annealing at high temperature (typically 1323 K) [ll]. This result has been used to measure the coarsening kinetics by performing the high temperature small angle neutron scattering (SANS) technique. It has been possible to follow in situ the change in precipitate shape during the ageing of prestrain specimens. MATERIAL AND EXPERIMENTAL PROCEDURE The experiment was conducted on AM1 single crystal superalloy (composition given in Table 1). At room temperature, the volume fraction of y’ phase precipitates is about 70% [12]. For this alloy, the averaged misfit between the two phases has been shown to be slightly negative at room temperature (ay’ < ay) and to decrease when temperature increases. Its value at 1323 K is -3 x 1O-3 [13]. An AM1 single crystal was oriented, then tensile specimens with (001) main directions were cut (Fig. 1). The faces transverse to the principal tensile axis showed a deviation of 6” from ( 100). The specimen dimensions have been precisely measured in order to quantify the amount of prestrain. They were then dynamically strained in tension at 1123 K. At this temperature, coarsening is very slow and the specimens were kept for less than one hour at this temperature, so that the precipitate shape was not significantly modified. This was followed by an air quench and the specimens were measured again in order to quantify the strain. Two samples with two different strain levels were prepared. The strain was estimated to be about 0.2% for sample 1 and 0.6%

b

60 -I

for sample 2. For the latter, the strain was not uniform because one alumina piece of the holder broke during the deformation experiment. Then the specimen surfaces were mechanically polished. SANS experiments were performed on the Dll instrument at the Institut Laue Langevin (ILL) [14] using the following configuration: Sample-detector distance: Collimation length:

35 m 40 m

With the selected wavelength 1 = 12 A (full width at half maximum = 9%) the accessible scattering vector with range is 6 x 10-4&’ < q < 6 x lo-‘A-‘, q = 4~ sin 9/n and 29 the scattering angle. From the chemical composition of each phase [15], the average scattering lengths by = 0.75 x 10-‘2cm, by’ = 0.83 x 10-‘2cm were calculated. Due to this noticeable difference, the superalloys produce intense neutron scattering at small angles. Thus, with the high neutron flux of the ILL instrument, measurements could be performed in a few minutes. This time, which is short enough compared to the coalescence rate of the y’ precipitates in the studied temperature range, permits in situ measurements. These have the further advantage of using the same sample without modification of orientation with respect to the neutron beam. For the purpose of in situ measurements, a furnace was built and adapted to the facility. Its principle is similar to standard ILL furnaces [ 141 but special care was taken to prevent spurious diffuse scattering by the sample environment. Only one sapphire window is added in the neutron path and a boron nitride sample holder is used which served also as a diaphragm, thus making the parasitic scattering from the furnace negligible. The operating temperature ranges from room temperature to 1573 K. The heating of the sample can be fast and, in our case, the change from the prestraining temperature of 1173 K to the annealing temperature was less than 150 s for sample 1 (1373 K) and 100 s for sample 2 (1323 K). So the microstructure evolution of the material was negligible during the heating of the samples. The central part of the samples was cut to fit into the furnace. The SANS results are obtained as isointensity contours in reciprocal space which provide information on the precipitate size, shape and spatial correlation [16, 171. We could acquire a spectrum every 6 min. Then the specimens were etched using a solution of sulphuric acid (45%), nitric acid (42%) and phosphoric acid (13%) sod were observed on a Jeol6010 scanning electron microscope under 20 kV. RESULTS AND DISCUSSION

Fig. 1. Schematic representation of the tensile specimens (thickness 0.6 mm). Hatched part shows the region studied during the SANS experiment.

Figure 2 shows three intensity contours obtained on sample 2 at 1323 K for different annealing times. jThe first one (a) corresponds to the initial state at t = 0. It is quite similar to that obtained before

VBRON and BASTIE: DIRECTIONAL COARSENING

0a

04

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Fig. 2. SANS intensity contours obtained on sample 2 at 1323 K for different annealing times: (a) t = 0, (b) t = 4 h, (c) t = 20 h.

at room temperature. In particular, the position of correlation peaks is the same, showing that no significant change in the microstructure superlattice occurred during the heating of the specimen. However, it is worth noting that the asymptotic behaviour at large values of q is different. This point will be discussed elsewhere [18]. Intensity contours (b) and (c) correspond to annealing times of 4 h and 20 h, respectively. Before making a quantitative analysis of all the recorded patterns, several qualitative features have to be pointed out: heating

A fast change in the symmetry patterns is observed in a few hours. Such an evolution does not occur when the samples are not prestrained [16, 191. The role of the plastic deformation is primordial. ??A similar evolution is observed on sample 1. Even a plastic deformation as small as 0.2% is sufficient to induce the transformation. ??The evolution from a nearly fourfold symmetry to a twofold one corresponds to the transformation from cuboids to platelets (rafts) as confirmed by the SEM observations (Fig. 3). ??Four or two correlation peaks are observed for all patterns. The rather periodic arrangement of the precipitates is kept during the transformation. ??

It is worth noting that the fourfold symmetry of the patterns (a) at t = 0 is not perfect. This could be due to a misorientation of the crystallographic axis of the specimen with respect to the direction of the neutron beam [16]. However, although the orientation of the sample was difficult to adjust inside the furnace, several tests indicate that at least a part of the asymmetry is related to the sample and has been induced by the prestrain. Other authors [20], using small angle X-ray scattering to characterize the microstructure, observe a similar effect. It is attributed to the beginning of the loss of coherency

at the matrix/precipitate interfaces, owing to the prestrain. The microstructures before and after ageing are shown in Fig. 3 (a) and (b), respectively. Before ageing and after the prestrain, the microstructure (at the head of the specimen) is very similar to the initial one. After ageing, without any applied stress, the microstructure shows a well-developed type N coarsening. So, as for a prestrain in compression [ 111, a prestrain in tension is able to induce directional coarsening during ageing and the SANS technique is well suited to observe this transformation as the change in shape of the isointensity contours is sign of the directional coarsening process. This result validates a model, based on the role of dislocations at the matrix/precipitate interfaces, for directional coarsening. During the prestrain “relaxing” dislocations are trapped by the yly’ horizontal interfaces. creating a gradient in elastic energy between vertical and horizontal interfaces. At higher temperature, when diffusion processes become effective, the morphology evolves. Other experiments have been conducted in which all the matrix/precipitate interfaces were relaxed [20], combining prestrain in compression and in tension. In that case, they do not clearly indicate any directional coarsening. Now let us examine more quantitatively the detailed kinetics of the directional coarsening given by the in situ experiment. At t = 0 the SANS pattern shows four peaks observed along the (100) and (001) axes of the reciprocal lattice. These correlation peaks indicate that in direct space the precipitates are rather regularly arranged on the nodes of a simple rectangular “superlattice” of parameters L, parallel to the stress axis and Lz perpendicular to the stress axis. Their values, deduced from the position of the correlation peaks are L, = 0.52 pm, L) = 0.59 pm for sample 1 and Ll = 0.50 pm, L2 = 0.58 pm for sample 2, in good agreement with the SEM

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VBRON and BASTIE: DIRECTIONAL COARSENING

observations. The superlattice is not a cubic one as observed when the samples are only heat treated [16]. The prestrain has slightly modified the arrangement of the precipitates (notice that the values of L, and L2 obtained by SANS are not the size of the precipitates but the period of the superlattice). A careful examination of the SEM pictures suggests that this deformation of the superlattice is due to a change of the width of the y channel rather than to a modification of the y’ precipil tate

Fig. 3. SEM observations

During the annealing, the correlation peaks move but, due to the oriented cuboid-raft transformation, this displacement depends strongly on the directions studied. Along the (001) direction, only a small displacement is observed; it is related to a slight increase in the superlattice period perpendicular to the rafts, from 0.52 pm to 0.59 pm for sample 1 and from 0.51 pm to 0.55 pm for sample 2. This difference in the period change between the two samples seems related to the different values of the prestrain. Indeed, for the sample with the largest prestrain, SEM

of the precipitate microstructure of sample 1: (a) after prestrain ageing, (b) after ageing for 10 h at 1323 K.

and before

VERON

lO’q(A Fig.

4. Evolution

with

and BASTIE:

DIRECTIONAL

‘)

annealing

time

of the

intensity

profiles along the (100) direction for sample 2: (@) t = 0, (&r=4h,(m)t=17h.

observations show that the lateral extension of the raft is the smallest; however, this result has to be confirmed by a study at the same annealing temperature. Along the (100) direction, this evolution is much stronger and leads to the disappearance of the correlation peaks behind the beam stop as shown in Fig. 4 for sample 2 after 17 h. The time evolution of the superlattice periods has been deduced from the position of the correlation peaks and is reported in Fig. 5 where the period changes are plotted vs time for the two recorded directions. A linear increase in all the superlattice parameters with time is observed during about 8 h for sample 1 and 15 h for sample 2. Then the parameter perpendicular to the rafts stops its increase and keeps a constant value while the parameter parallel to the rafts continues to increase linearly until the correlation peaks disappear behind the beam stop, making the measurement impossible. The duration necessary to reach this moment seems strongly correlated with the end of the evolution of the parameter perpendicular to the rafts. The duration difference between the two samples is probably related to the different values of the annealing temperature (1373 K for sample 1, 1323 K for sample 2). Although more systematic studies are still necessary for a quantitative understanding of the precipitate transformation, the following qualitative explanation can be proposed. During the first stage of the annealing. the gradient in internal strain energy, created by the prestrain, provides a driving force for the diffusional mass transport between the two types of channel. The y phase is removed from the channels which were parallel to the stress axis and transported into the channels perpendicular to it, while the y’ phase moves in the opposite direction. So the observed increase of the super-lattice parameters can be attributed:

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3281

?? in one direction, to the oriented coarsening of the precipitates and describes the evolution of the lateral expansion of the 1~’precipitates when they evolve from cuboids to rafts. ?? in the other direction, to an increase in the distance between rafts owing to a broadening of the y channels. In this direction, the thickness of the y’ precipitates does not increase and even, probably, decreases slightly during the transformation. Indeed, in these materials, the coarsening generally follows a law where the volume change in the precipitates is time proportional. The observed evolution of the precipitates, which is the same for the two directions parallel to the surface of the rafts, leads to a t’ dependence. So it is probable that the thickness of the rafts decreases proportionally to t. The existence of this decrease is consistent with a simple estimation based on the saturation value of the superlattice parameter perpendicular to the rafts and on the width of the 1’channel for cuboids and rafts at a given volume fraction of 7’ phase. At the end of the

(4 0.20 -

0

2

4

6

10

8

Time (hours) (b) 0.20 .

Time (hours) Fig. 5. Variation of the precipitate superlattice parameters with annealing time: (m) along the (100) direction, (0) along the (001) direction. (a) Sample 1, annealing temperature 1373 K, (b) sample 2, annealing temperature 1323 K.

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VERON and BASTIE: DIRECTIONAL COARSENING

transformation, even with a volume fraction of 70% which is overestimated because it is the value at room temperature, the calculated thickness of the precipitates is about 0.40 pm, a value significantly smaller than the 0.50 pm calculated for the cuboids.

CONCLUSION These SANS experiments confirm the influence of a plastic prestrain on the directional coarsening of y’ precipitates in superalloys. They corroborate a mechanism based on the role of dislocations at the matrix/precipitate interfaces. Furthermore they provide information on the time evolution of the directional coarsening, a measurement difficult to perform by other methods. However, in this case, we only deal with one diffusion process during directional coarsening; under stress, during creep for example, there is another driving force to add to this one, owing to the response of the alloy to the applied stress. This effect is purely plastic, and can have the same influence on directional coarsening as the previous one or the opposite one, as previously discussed [21]. The next experimental step will be to perform an “on line” creep experiment on the SANS apparatus in order to observe possible difference in the directional coarsening and in its kinetic. Acknowledgements-Thanks are due to SNECMA for providing the material and Institut Laue Langevin for beam time allocation. The authors would particularly like to acknowledge D. Jobart and J. J. Blandin from CPM2 for prestraining the samples, L. Vuillard and A. Polsak from ILL, for their help during the neutron measurements and P. Andant and P. Martin, also from ILL, for providing the furnace. We would also like to thank J. Lajzerowicz from LSP and Y. Brechet and F. Louchet from LTPCM for fruitful discussions and C. Alonso for help in analysing the results.

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