the interface between γ and γ’ in ni-al alloys by hrem

7
The Interface between γ and γ’ in Ni-Al Alloys by HREM Hector A. Calderon 1, a , L. Calzado-Lopez 2,b and T. Mori 3,c 1 Depto. de Ciencia de Materiales, ESFM-IPN, Ed. 9 UPALM Zacatenco D.F., Mexico 2 Universidad de la Ciudad de Mexico. Mexico DF Mexico 3 Materials Science Centre, University of Manchester, Manchester, M1 7HS, UK a [email protected], b [email protected]. Key words: NI-Based Superalloys, Interfaces, Electron Microscopy, Micromechanics. Abstract. Ni Base superalloys owe their excellent mechanical properties to the presence of particles of γ’ phase (Ni 3 Al with an L1 2 structure) in a γ matrix (Ni–Al solid solution with an fcc structure). Besides Al, other elements are used to impart either a higher strengthening or improved corrosion properties at high temperature. The interface between γ and γ’ becomes of absolute importance for the resulting mechanical properties and technological application. Especially by considering the consequences that diffusion driven coarsening brings about to the particle distribution either with or without the influence of an applied stress or strain. In this work the interface between γ and γphases is characterized by means of measurements on phase images obtained from high resolution transmission electron microscopy images (HRTEM). Phase images represent the sample structure much more accurately than typical HREM experimental images and allow correction of spherical aberration and other residual aberrations. The investigation is performed by using a binary Ni-Al alloy as well as technical Ni base superalloys (MC2 and MCNG). While a sharp interface is developed during stress free coarsening in Ni-Al alloys, a wider volume needs to be considered when alloying elements are introduced. Measurements of lattice spacings on phase images and chemical composition from energy dispersive spectroscopy are used to show the interface characteristics in the alloys under consideration. The interface in the binary Ni-Al alloy can be described by micromechanics as a typical misfitting inclusion. In the technical alloys, the presence of concentration gradients changes the expected lattice strains in a given volume around the particles. Introduction Ni-based superalloys have a microstructure consisting of coherent particles (γ’ phase with a L1 2 structure) embedded in a solid solution matrix (γ phase with a fcc structure). Coarsening of γparticles is an important mechanism since it influences the high-temperature mechanical properties of these materials. As a consequence it has been extensively investigated both experimentally and through computer simulations. Coarsening of second phase particles involves a reduction of the total free energy of the system. In fluids, the tendency for larger particles to grow at the expense of smaller ones is driven by the reduction of surface energy. In the case of solids, the elastic strain energy becomes part of the driving force [1]. Reduction of elastic energy promotes particle alignment along elastically soft directions and changes of morphology as the precipitate volume increases. In addition according to several authors, reduction of elastic energy produces the splitting of large particles that reach a critical size [2,3]. Actually symmetric arrays of two, four, and eight γparticles have been observed during microstructural evolution of Ni alloys with low volume fraction and interpreted in terms of splitting. According to models based on numerical simulation, there is a finite distance between particles, called equilibrium distance, which minimizes the elastic strain energy [2,4]. Such an equilibrium distance arises from the interaction elastic energy which is attractive for long distances and repulsive for short distances [5,6]. An alternative mechanism for the formation of particle arrays is based on particle migration and selective coalescence [7-9]. According to this mechanism, particle arrays are formed by attractive interaction. Once the particle Solid State Phenomena Vols. 172-174 (2011) pp 254-259 Online available since 2011/Jun/30 at www.scientific.net © (2011) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.172-174.254 All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (ID: 129.186.1.55, Iowa State University, Ames, United States of America-06/10/13,07:43:00)

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Page 1: The Interface between γ and γ’ in Ni-Al Alloys by HREM

The Interface between γγγγ and γγγγ’ in Ni-Al Alloys by HREM

Hector A. Calderon1, a, L. Calzado-Lopez2,b and T. Mori3,c 1Depto. de Ciencia de Materiales, ESFM-IPN, Ed. 9 UPALM Zacatenco D.F., Mexico

2Universidad de la Ciudad de Mexico. Mexico DF Mexico

3Materials Science Centre, University of Manchester, Manchester, M1 7HS, UK

[email protected],

[email protected].

Key words: NI-Based Superalloys, Interfaces, Electron Microscopy, Micromechanics.

Abstract. Ni Base superalloys owe their excellent mechanical properties to the presence of particles

of γ’ phase (Ni3Al with an L12 structure) in a γ matrix (Ni–Al solid solution with an fcc structure).

Besides Al, other elements are used to impart either a higher strengthening or improved corrosion

properties at high temperature. The interface between γ and γ’ becomes of absolute importance for

the resulting mechanical properties and technological application. Especially by considering the

consequences that diffusion driven coarsening brings about to the particle distribution either with or

without the influence of an applied stress or strain. In this work the interface between γ and γ’

phases is characterized by means of measurements on phase images obtained from high resolution

transmission electron microscopy images (HRTEM). Phase images represent the sample structure

much more accurately than typical HREM experimental images and allow correction of spherical

aberration and other residual aberrations. The investigation is performed by using a binary Ni-Al

alloy as well as technical Ni base superalloys (MC2 and MCNG). While a sharp interface is

developed during stress free coarsening in Ni-Al alloys, a wider volume needs to be considered

when alloying elements are introduced. Measurements of lattice spacings on phase images and

chemical composition from energy dispersive spectroscopy are used to show the interface

characteristics in the alloys under consideration. The interface in the binary Ni-Al alloy can be

described by micromechanics as a typical misfitting inclusion. In the technical alloys, the presence

of concentration gradients changes the expected lattice strains in a given volume around the

particles.

Introduction

Ni-based superalloys have a microstructure consisting of coherent particles (γ’ phase with a L12

structure) embedded in a solid solution matrix (γ phase with a fcc structure). Coarsening of γ’

particles is an important mechanism since it influences the high-temperature mechanical properties

of these materials. As a consequence it has been extensively investigated both experimentally and

through computer simulations. Coarsening of second phase particles involves a reduction of the

total free energy of the system. In fluids, the tendency for larger particles to grow at the expense of

smaller ones is driven by the reduction of surface energy. In the case of solids, the elastic strain

energy becomes part of the driving force [1]. Reduction of elastic energy promotes particle

alignment along elastically soft directions and changes of morphology as the precipitate volume

increases. In addition according to several authors, reduction of elastic energy produces the splitting

of large particles that reach a critical size [2,3]. Actually symmetric arrays of two, four, and eight γ’

particles have been observed during microstructural evolution of Ni alloys with low volume fraction

and interpreted in terms of splitting. According to models based on numerical simulation, there is a

finite distance between particles, called equilibrium distance, which minimizes the elastic strain

energy [2,4]. Such an equilibrium distance arises from the interaction elastic energy which is

attractive for long distances and repulsive for short distances [5,6]. An alternative mechanism for

the formation of particle arrays is based on particle migration and selective coalescence [7-9].

According to this mechanism, particle arrays are formed by attractive interaction. Once the particle

Solid State Phenomena Vols. 172-174 (2011) pp 254-259Online available since 2011/Jun/30 at www.scientific.net© (2011) Trans Tech Publications, Switzerlanddoi:10.4028/www.scientific.net/SSP.172-174.254

All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP,www.ttp.net. (ID: 129.186.1.55, Iowa State University, Ames, United States of America-06/10/13,07:43:00)

Page 2: The Interface between γ and γ’ in Ni-Al Alloys by HREM

separation is sufficiently small, coalescence can take place if both particles have an identical

translation order domain (TOD). Thus, particle groups can be formed either by splitting of larger

particles or by migration of individual particles. Interestingly splitting of large particles is not

supported by experimental evidence as the TODs in many particle pairs show that most of them do

not match. This suggests that particles migrate or coalesce and form arrays of several particles [8].

Despite such evidence simulations based on the phase field method have predicted particle splitting

[10] but also coalescence [11]. In this method, the interface is assumed to have a width and becomes

a tunable parameter. In this investigation, the nature of interfaces is investigated both structurally

and compositionally in binary Ni-Al alloys and the technical superalloys MC2 and MCNG. This is

done by using HRTEM techniques and measuring lattice spacings directly from aberration corrected

phase images as well as determining chemical compositions by direct techniques (energy dispersive

spectroscopy, EDS) as a function of position at the interface region.

Experimental Procedure

Binary Ni-12 at. % Al, MC2 and MCNG alloys are used for this investigation. Binary alloys have

been aged for 5 h at 1113 K and 1700 h at 923 K. MC2 and MCNG alloys have several different

alloying additions that are given in [11]. They are investigated in the standard condition (for details

of heat treatment see [11]). Sufficiently thin samples can be prepared by manually polishing the

samples with diamond impregnated paper and subsequent electrolytic polishing. HRTEM has been

performed in different microscopes including a CM300 and an image Cs corrected Titan (FEI®).

Approximately 20 experimental images (each one with a different defocus setting) are acquired to

determine the phase and the amplitude images by using the software TrueImage (FEI ®). The results

are further corrected to compensate for residual aberrations (focus, astigmatismus and coma).

Results and Discussion

Figure 1 shows representative images of γ’ particles after aging in the alloys under investigation.

Figs. 1a,b show typical arrays of particles that are formed by migration of particles and coalescence.

The particles in the array have different TODs and further coalescence is unlikely. Fig. 1b shows a

representative STEM image from the particle distribution in the alloy MC2 after aging without

stress (Standard Condition), the volume fraction is considerably high which most likely reduces

migration leaving coalescence to the randomly distributed TODs. Fig. 1c shows a STEM image of

the γ’ particles after coarsening under stress, the rafts develop as a function of creep time. The

interfaces that are shown below are taken from such particle distributions.

Figure 2 shows a phase image corresponding to the binary Ni-Al alloy, the image plane is

perpendicular to the crystallographic direction [001]. The contrast is relatively low but the γ’ region

can be recognized by the ordered pattern with a stronger intensity every two lattice planes. This

image can be quantitatively evaluated giving rise to the results in Figs. 2b and 2c. The measurement

of lattice spacings is performed parallel (direction a) and perpendicularly (direction c) to the

interface as a function of unit cell. The tendency is clearly shown in the above figures but averaging

the results for each unit cell (image columns) in the selected area produces clearer results as shown

in Fig. 2d. There is a continuous value for the averaged lattice parameter a, but a distinctive

discontinuity in the case of the parameter c due to the conditions imposed by strain continuity

around an inclusion. The lattice parameter mismatch (δ = ap-ap/am, m=matrix, p=particle) normally

used for characterization of a Ni alloy can be readily derived from the data in Fig. 2d as shown

elsewhere [13]. In this case there is no appreciable difference in the contrast at either side of the

interface suggesting that the involved volume tends to zero i.e., the interface tends to be a surface.

This is supported by the high long range order parameter of the Ni3Al compound and the fact that

order is kept up to the melting point. Without ruling out a given Al or Ni composition profile

outside the particle, the clear behavior of the γ’ particles as an inclusion, shows that the interface is

rather thin and judging only from experimental or processed HREM images, the interface is rather

sharp.

Solid State Phenomena Vols. 172-174 255

Page 3: The Interface between γ and γ’ in Ni-Al Alloys by HREM

Fig. 1. γ’ particles in alloys

under investigation. (a,b) Ni-12

at.% at. Al after 5 h at 1113 K

and 1700 h at 923 K. (c)

MCNG alloy in the standard

condition. (d) MC2 alloy after

120 h at 1150 oC and under 150

MPa. Arrows indicate <001>,

the electron beam is traveling

along [ ]100 . Line in (d)

indicates direction for strain

measurement.

Figure 2. (a) Phase image after exit wave

reconstruction procedure. (b) Measurement of

lattice spacings a parallel to interface. (c)

Measurement of lattice spacings c perpendicularly

to the interface. (d) Averaged c spacings as a

function of position. Area for measurements is

sketched. Units = nm. The arrows indicate cube

directions.

0.2 µm

1 µm

(a)

(b)

(c)

(d)

(a)

(b)

(c)

(d)

a Spacings

c Spacings

256 Solid-Solid Phase Transformations in Inorganic Materials

Page 4: The Interface between γ and γ’ in Ni-Al Alloys by HREM

Figure 3. Alloy MC2 in the standard condition. (a) Phase

image after exit wave reconstruction. (b) Area for

measurement of lattice spacings. (c) Representative

measurement of a spacings (in pixels) as a function of

location. (d) Relative variation of c and (e) a spacings

after averaging along a given column or cell in the phase

image.

The interface between γ and γ’ in the technical alloys under consideration is affected by the

alloying additions. While in the case of the binary alloy micromechanics can be directly applied to

describe the local strains i.e., via consideration of a misfitting inclusion, the situation is apparently

changed once alloying elements are added. Figure 3a shows the phase image obtained after exit

wave reconstruction for a representative interface in alloy MC2 in the standard condition. The area

for quantitative analysis is only a section and given in Fig. 3b. The corresponding measurement is

shown in Fig. 3c. In this case measurements perpendicular or parallel to the interface show similar

patterns with an approximately homogeneous lattice spacing and frequent peaked variations at

specific locations. Figs. 3d,e show averaged (along each column in the image) a and c parameters as

a function of position for the image in Fig. 3b, the values have been normalized to the average. This

is taken as a measure of strain as a function of position. Apparently the effect of alloying addition in

these alloys with a variety of alloying elements is to modify the strain field and the measurements

reflect the distribution of alloying elements around the interface. Such profiles for the different

alloying elements are likely to depend on aging condition up to certain degree which reflects on the

need to determine the conditions for stability represented by the standard condition.

(a)

(b)

(c)

(e)

(d)

Solid State Phenomena Vols. 172-174 257

Page 5: The Interface between γ and γ’ in Ni-Al Alloys by HREM

Figure 4. Alloy MCNG in the standard condition. (a)

Phase image containing the γ−γ’ interface as it is shown

by the (b) intensity profiles from selected columns. (c)

Measurement of lattice parameter as a function of

position, peak high (in pixels) is equivalent to lattice

parameter. (d,e) strain (measured with c and a spacings,

respectively).

Figure 4 shows results regarding alloy MCNG. A representative phase image is given in Fig. 4a

where clear intensity maxima are seen for the lattice positions but the typical contrast differences in

the γ’ particle (due to ordering) are absent. This is an effect produced by the extremely thin sample

in use. However there are intensity differences that make possible to locate the interface region as is

shown by the intensity profiles in Fig 4b. The profiles are taken from rows as indicated by a

number. Apparently both the ordering in the L12 structure and the chemical element distribution

affect the blob intensity in the image. There is a characteristic intensity variation at the γ’-γ interface

but also the intensity varies inside the γ or the γ’ phase indicating a variation of the chemical

composition of the atomic columns since a considerable thickness variation is rather unlikely in the

reduced area under observation. Fig. 4 shows that the interface position is not unique and thus it is

not atomically flat. However ordering ends abruptly at each column in the image. Additionally the

distribution of the alloying elements produces a transition zone or volume characteristic of the

interface. On the other hand, measurements give similar results to those shown for the MC2 alloy,

but the lattice parameter differences are smaller giving rise to lower localized strains as can be

qualitatively seen in Fig. 4c and quantitatively in Figs 4d ,e. Interestingly the strains shown in Figs

3d,e and Fig. 4d,e are considerably different. While in the case of alloy MC2, relatively high local

strain fields can be found reaching a maximum of around 4-5%, those in alloy MCNG are much

lower and around 1-1.5%. This impacts the lattice resistance to plastic deformation and dislocation

motion but it will be discussed elsewhere.

(c)

(d)

(e)

(a)

(b)

258 Solid-Solid Phase Transformations in Inorganic Materials

Page 6: The Interface between γ and γ’ in Ni-Al Alloys by HREM

The use of HREM allows characterization of the γ−γ’ interface up to a very fine scale. In the past,

HREM has been limited by the aberrations in the microscope lenses. Image delocalization and aber-

rations reduce precision in the measurement of lattice spacings. The use of aberration corrected mi-

croscopes and determination of the corresponding exit wave from reconstruction procedures allow

quantitative determination of strain fields at γ−γ’ interfaces. In the case of the binary Ni-Al alloy, the

strain field predicted from micromechanicas can be reproduced. As for the technical MC2 and

MCNG alloys, the interface region becomes a volume due to the distribution of alloying elements.

Such a distribution can be readily measured by spectroscopic techniques as shown in Fig. 5 where

the elemental composition is determined across an interface in the MC2 alloy (standard condition).

The profile has been measured from the region in the inset (see line). It clearly shows that the

chemical distribution of most elements extends over a distance few nanometers wide as observed by

HREM.

References

[1] A. J. Ardell, R. B. Nicholson, J.D. Eshelby: Acta Metall. Vol. 14(1966) p. 1295.

[2] M. M. Doi, T. Miyazaki and t. Wakatsuki: Mater. Sci. Eng. Vol. 67 (1984), p. 247.

[3] Y. Wang, L-Q. Chen, A. G. Khachaturyan: Acta Metall. Mater. Vol. 41, p. 279.

[4] A.G. Khachaturyan, S.V. Semenoskaya and J.W. Morris: Acta Metall. Vol. 36 (1988), p. 1563.

[5] C. S. Su and P. W. Voorhees: Acta Mater. Vol. 44 (1996), p.2016.

[6] W. C. Johnson and J. K. Lee: Metall. Trans. Vol. 10A (1979), p. 1141.

[7] H. A. Calderon, J. G. Cabanas-Moreno, T. Mori: Phil. Mag. Lett. Vol. 80 (2000), p. 669.

[8] Y. Wang, D. Banerjee, C. Su, A. G. Khachaturyan: Acta Mater. Vol. 46 (1998), p. 2983.

[9] D. Banerjee, R. Banerjee and Y. Wang: Scripta Mater. Vol. 41 (1999), p. 1023.

[10] T. Miyazaki, T. Koyama, T. Kozakai: Mat. Sci. Eng. A312 (2001) p. 38.

[11] Y.H Wen, J.P. Simmons, C. Woodward: Model. Simul. Mat. Sci. Eng. Vol. 18 (2010) 055002.

[12] R.Glas, M.Jouiad, P.Caron, N. Clement, H.O.K. Kirchner: Acta. Mater. Vol 44(1996), p. 4917.

[13] T. Mori to be published.

[14] Grant 58133 (CONACYT), CEMES-CNRS (France) and NCEM (USA) are gratefully

acknowledged.

Fig. 5. Concentration profiles of W, Ta, Ni, Co, Cr, Ti and Al in the MC2 alloy (standard

condition). The inset shows the region for the line scan.

100

200

300

0 50 100 150 Position (nm)

Inte

nsi

ty (

cou

nts

)

Solid State Phenomena Vols. 172-174 259

Page 7: The Interface between γ and γ’ in Ni-Al Alloys by HREM

Solid-Solid Phase Transformations in Inorganic Materials 10.4028/www.scientific.net/SSP.172-174 The Interface between γ and γ’ in Ni-Al Alloys by HREM 10.4028/www.scientific.net/SSP.172-174.254