Thermal stability of dispersoids in ferritic oxide-dispersion-strengthened alloys

Materials Science and Engineering, A 177 (1994) 199-208

199

Thermal stability of dispersoids in ferritic oxide-dispersionstrengthened alloys P. Krautwasser, A. Czyrska-Filemonowicz*, M. Widera** and F. Carsughi*** Forschungszentrum (KFA) Jiilich, D-52425 Jiilich (Germany) (Received March 24, 1993; in revised form July 27, 1993)

Abstract To study the thermal stability of dispersoids in ferritic oxide-dispersion-strengthened (ODS) alloys, specimens of two alloys, F10 (PM2000) and F11 (PM2010), with different yttria contents were isothermally annealed at 1623 K or isochronally heat treated for 110 h in the temperature range 1273-1623 K. The investigations were performed by small angle neutron scattering and transmission electron microscopy in conjunction with quantitative image analysis. Two overlapping size distributions of dispersoids in the approximate size ranges from 5 to 40 nm and from 40 to 100 nm were found. In addition, a third size distribution of larger particles with a mean value above 250 nm was present. The size distributions determined with each technique are in good agreement for the dispersoids in the size range 5-100 nm, which influence the creep resistance of ODS alloys at high temperatures. Since the aluminium content of the dispersoids increases with increasing heat treatment time and phase transformations take place, the coarsening behaviour of the dispersoids cannot be described by Ostwald ripening. A variation in yttria content in the alloys of between 0.5% and 1% had no significant effect on the number density of dispersoids smaller than 100 nm.

I. Introduction

Owing to their excellent hot corrosion resistance and good mechanical strength at temperatures up to 1500 K, oxide-dispersion-strengthened (ODS) alloys are promising materials for high temperature components [ 1] operating in aggressive environments. Oxide dispersoids in an austenitic or ferritic matrix act as obstacles to dislocations and therefore improve the creep behaviour at high temperatures. The pinning of dislocations by disperoids is mainly determined by the number density of and the mean free path between dispersoid particles. For a given volume fraction of dispersoids, coarsening of the particles leads to a decrease in number density, an increase in mean particle distance and thus to lower creep resistance. As an example of dislocation-dispersoid interaction, Fig. l(a) shows the microstructure of Incoloy alloy MA956 after tensile deformation at 1173 K [2]. Extraction of the dispersoids as replicas (Fig. l(b)) allows accurate measurements of dispersoid size by means of quantita*On leave from Academy of Mining and Metallurgy, Cracow, Poland. **Present address: RWE Energie AG, Essen, Germany. ***Present address: Universith di Ancona, Instituto di Science Fisiche, Ancona, Italy.

0921-5093/94/$7.00 SSDI 0921-5093(93)09396-Z

tive image analysis (Fig. l(c)). The tensile properties of alloy MA956 were investigated in an earlier study [3]. Although dispersoids are more stable than strengthening precipitates (e.g. carbides or intermetallic phases) in conventional superalloys, coarsening of yttriaalumina particles at high temperatures has been observed. Several experimental studies of the thermal growth of dispersoids have been performed on austenitic ODS alloys [4-9], but only a few papers have been published about particle coarsening in ferritic ODS materials [10-12]. As part of a study to compare two techniques for microstructural analysis (small angle neutron scattering (SANS) and transmission electron microscopy (TEM)), the thermal stability of dispersoids in ferritic ODS alloys was investigated by these methods. In an earlier study TEM and small angle X-ray scattering (SAXS) were used to determine the size distribution of dispersoids in the austenitic alloy MA754 [13].

2. Experimental details

The chemical compositions of the ferritic ODS alloys investigated are listed in Table 1. Alloy F 10 is a precursor alloy of the commercial alloy PM2000. © 1994 - Elsevier Sequoia. All rights reserved

200

P. Krautwasser et al.

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Dispersoids in ferritic ODS alloys

TABLE 1. Nominal chemical compositions (wt.%)of the ferritic ODS alloys investigated

t

UU

3o

Tensile test

s

I i i,m

i'-'-1

b,Q

L', | ..................................

25"0 L 20Q. ¢/) °~

"o

15-

.~

10-

E ¢" Q) n,"

5-

C

ilm

Mo,~:11..~ .-1,8 i

i i ilU i il

:: ~

:: C:: :: ~ i i

Alloy

Fe

Cr

AI

TGi

C

Y203

F 10 (PM2000) F11 (PM2010)

Bal. Bal.

20 20

5.5 5.5

0.5 0.5

< 0.04 < 0.04

0.5 1.0

The specimens for structural investigations were machined from recrystallized bars. Details of the alloys and their manufacture are described in ref. 14. The specimens were isothermally annealed in air at 1623 K for up to 300 h or isochronally heat treated for 110 h in the range 1273-1623 K. The techniques used for microstructural investigations were optical microscopy, scanning electron microscopy (SEM), TEM and SANS. The transmission electron microscopy studies were performed using a Jeol JEM 200CX microscope equipped with an energy-dispersive X-ray (EDX) system of Tracor Northern. Thin foils for T E M investigations were prepared by conventional double-jet electropolishing in a solution of 10% perchloric acid in glacial acetic acid (temperature less than 288 K, voltage about 50 V) or by ion beam milling. Extraction double-replicas were prepared by evaporation of carbon on to both sides of thin specimens (thickness about 0.05 mm) followed by dissolution of the metallic matrix in a 10% solution of bromine in ethanol [15]. Replicas were used for E D X analysis of the dispersoid composition and measurements of dispersoid size distributions. Quantitative microstructural analysis (QMA) of T E M images was carried out using an interactive image analysis system (IBAS) of Kontron Co. The SANS measurements were performed with the KWS1 instrument in the neutron guide laboratory ELLA at the FRJ2 reactor of the Forschungszentrum (KFA) Jiilich and in part with the SANS-2 facility at the FRG1 reactor of the Forschungszentrum (GKSS) Geesthacht using neutrons with 2 = 0.7 nm moderated by a cold source. The sample-to-detector distance range was 2-20 m, corresponding to scattering vectors of 0.02~
iiil iiiiiiil I

I

i

2

4-

6

I

i

I

i

l

10

2

4

6

i

100

Diameter I'nm] Fig. 1. Microstructure of Incoloy alloy MA956 after tensile deformation at 1173 K with a strain rate of 10 s s-1 [2]: (a) dislocation-dispersoid interaction; (b) fine AI-Y oxides in extraction double-replica; (c) size distribution of extracted dispersoids determined by TEM-QMA.

P. Krautwasseret al. / Dispersoidsin ferritic ODS alloys The measured SANS intensities were corrected for background and detector efficiency and then normalized to absolute units of the differential cross-section by using a lupolene standard precalibrated with vanadium. The macroscopic differential cross-section dZ/df2 for the two-phase system can be written as oo

dZ

d~(Q)=(Ap)2

f

N(R)Ve(R)IF(Q,R)] 2dR

0

where Ap represents the difference in scattering length density between the matrix and the dispersoids, N(R) is the size distribution (number of dispersoids per unit volume with dimensions between R and R + dR), V(R) is the volume of a single precipitate and F(Q,R) is the form factor of the scattering dispersoids. Since the dispersoids are mostly spherical, the corresponding form factor was used for the calculations. As shown in ref. 16, the unknown N(R) function can be written as a linear combination of cubic spline functions ~k, i.e.

N(R)= E C~bk(Z) k=t

where z = In R and the N~+ 3 knots are equispaced in z. The SANS cross-section can be written as - - ( Q ) = ( A p ) e Z Ck~k(O) dQ k I where oo

~k(Q) = f V2(R)IF(Q, R)12q~k(lnR) dR 0

A constrained least-squares fit of the experimental data was used to optimize the Ck coefficients and then to obtain N(R) in absolute units once the Ap factor was known. It is also possible to estimate the error of N(R) by randomly changing each experimental value inside a Gaussian distribution whose standard deviation corresponds to its experimental error. By repeating this procedure (typically 50-100 times), it is possible to obtain different N(R) functions and hence an estimation of the error on N(R). The best fit for each specimen, i.e. the size distribution with the lowest Z 2, is used in the following diagrams.

3. Results and discussion

The microstructure of the alloys investigated consists of dispersoids (mixed Y-AI oxides), larger particles of pure AI203 and titanium carbonitrides

201

Ti(C,N) in a ferritic matrix. The Y-A1 oxides which may be present in ODS alloys are yttrium-aluminium monoclinic, YAM (Y4A1209), YAH (hexagonal, YAIO3), YAP (perovskite, YA103) and YAG (garnet, Y3AIsO12). Since the various dispersoids have different compositions and structures, their neutron-scattering contrast varies by up to 30%. In the case of more than two phases the inverse solution to calculate the size distribution cannot be determined. The calculations were therefore performed with a single value of the scattering contrast for the different dispersoids. This means that the heights of the size distributions of each type of dispersoids may vary by up to 30% from low to high A1 content. Systematic Q M A performed using TEM and IBAS showed that four distinct size distributions of particles are present in all specimens. As an example, Fig. 2 represents the extracted particles of PM2010 after annealing for 110 h at 1623 K. The micrographs show three size ranges of mixed Y-A1 oxides from about 5 t o 500 nm and large Ti(C,N) particles a few microns in l diameter. In Fig. 3 the volume fractions as a function of particle diameter of the three size ranges of dispersoids in the as-received material (Fig. 3(a)) and after 110 h at 1623 K (Fig. 3(b)) are represented by the histograms obtained by Q M A of TEM images of replicas. For comparison, the corresponding volume distributions derived from SANS measurements are shown as solid curves. Since the number density of the particles was not determined quantitatively by TEM, the maximum of each histogram was fitted to the number density of the SANS result. The findings of SANS and TEM investigations are in good agreement with the dominant small dispersoids of size ranges 1 and 2, whereas the difference for size range 3 indicates the upper detection limit of the SANS measurements for dispersoid sizes of about 200-500 nm [17]. A comparison of the diagrams in Figs. 3(a) and 3(b) shows a decrease in the number of particles in size range 1 and increase in the number density for size range 2. The size distributions of dispersoids determined by TEM during isothermal annealing (Fig. 4) and isochronal heat treatment (Fig. 5) show coarsening of the dispersoids in size ranges 1 and 2 at high temperatures. Since the number density of dispersoids of size range 2 is smaller by a factor of about 100 than that of size range 1 and the two size distributions are partially overlapping, particles of size range 2 can only be distinguished from those of size range i in the volume distribution plots in Figs. 3(a) and 3(b) but not in the size distribution histograms of Figs. 4 and 5. In contrast to this, the second size range shows up in the size distributions derived from SANS

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Dispersoids in ferritic ODS alloys

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Fig. 3. Volume distribution of dispersoids in alloy PM2010 (Fll): (a) as-received condition; (b) after heat treatment for 110 h at 1623 K; comparison of TEM-QMA and SANS results.

Fig. 2. Multimodal size distribution of particles from alloy PM2010 (Fll) heat treated at 1623K for ll0h: (a) dispersoids--three sizes of mixed Y-A1 oxides; (b) titanium carbonitride.

measurements (Fig. 6). The corresponding volume distributions are plotted in Fig. 7. 3.1. I s o c h r o n a l heat t r e a t m e n t

The variation in mean dispersoid diameter for size ranges 1 and 2 as a function of isochronal annealing for 110 h is shown in Fig. 8. A straight line was fitted to the data points of the heat-treated specimens. An extrapolation of this line to the value of 15 nm, being the mean diameter of the as-received material, indicates a substantial coarsening of the dispersoids starting at temperatures above about 1400 K. In general the observed coarsening of particles can be compared with the Ostwald ripening mechanism according to the Lifshitz-Slyozov-Wagner (LSW) theory [18, 19] to discriminate between interface and volume diffusion-controlled coarsening. In our case this is not possible, because the dispersoids become

richer in A1 content with increasing heat treatment time and phase transformations take place, shifting the main composition towards YAG. Since our measurements do not allow determination of the volume fraction variations of the different phases with heat treatment time, correction of the mean dispersoid diameter for the effect of structural volume changes is not possible. In the case of changing composition of the dispersoids, SANS measurements cannot determine whether the total volume fractions of dispersoids in size ranges 1 and 2 are decreasing with heat treatment time or whether the measured decrease in their volume fraction is an artefact due to reductions in neutron-scattering length density caused by phase transformations as described above. If all or part of the decrease in the volume fraction is real and caused by dissolution of small dispersoids of size range 1, this would not only lead to coarsening of the overlapping size range 2 but would also contribute to the growth of dispersoids of size range 3. In the late case the requirement for Ostwald ripening, i.e. a constant volume fraction for the size distribution in question, would not be fulfilled.

P. Krautwasser et al. r-~

Dispersoids in ferritic ODS alloys

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[nm]

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"0

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ii ¸ i

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The variation in size of particles in size range 3 does not follow the LSW theory, since a decomposition of large particles (with high aluminium content) was observed in TEM investigations and a strong decrease in volume fraction of about 50% was measured by SANS.

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Dicmeter

3.2. Isothermal heat treatment

In Fig. 9 the variation in particle size is displayed as a function of isothermal annealing at 1623 K. The

1oo

[nm]

Fig. 5. Size distribution of dispersoids in alloy PM2010 (F 11 ) isochronally heat treated for 110 h at 1473, 1573 and 1623 K: T E M - Q M A results.

P. Krautwasser et al.

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..+.. ~

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Dispersoids in ferritic ODS alloys

I_ ,,_~,o,~

i iii

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~50.

iiii

................

o.

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;....i...i..;

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Fig. 6. Size distribution of dispersoids in alloy PM2010 (FI 1) isothermally heated treated at 1623 K: SANS results.

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: ......

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2

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Dispersoid d i a m e t e r [ n m ]

Fig. 7. Volume distribution of dispersoids in alloy PM2010 (F 11 ) isothermally heated treated at 1623 K: SANS results.

35

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i

l ......... ! ......... ~.......... r ......... ; ......... i ....

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2

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.... ._

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800

1000

1200 Temperoture [ K ]

[~( o~/1'~ ~ l+d2/j

....... i ....

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2

d i a g r a m c o m p a r e s the m e a n d i a m e t e r s of the c o m b i n e d size ranges 1 and 2 as d e t e r m i n e d by TEM-QMA with those m e a s u r e d by S A N S (lower part). T h e u p p e r part of the plot shows the variation in size range 3 derived f r o m T E M - Q M A investigations. A s discussed earlier, in addition to an attractive d i s p e r s o i d - d i s l o c a t i o n interaction [20], the pinning of dislocations (Fig. l(a)) d e p e n d s strongly on the n u m b e r of density, size and local distribution of dispersoids. T h e m e a n interparticle spacings it of particles of size ranges 1 and 2 on o n e h a n d and of size range 3 on the o t h e r h a n d w e r e calculated i n d e p e n d e n t l y as a function of annealing time a c c o r d i n g to the f o r m u l a [21] 2=1.18d

O.

2 IO0 '10 Annealing time +1 [h]

'

Fig. 9. Coarsening of dispersoids with isothermal annealing at 1523 K: top, T E M - Q M A results for size range 3; bottom, comparison of T E M - Q M A and SANS results for combined size ranges 1 and 2.

.2_ 0 P 20-

._w

4.'

!

1400

1600

Fig. 8. C o a r s e n i n g of d i s p e r s o i d s of c o m b i n e d size ranges 1 and 2 with isochronal annealing for 110 h: T E M - Q M A results.

for the d i s p e r s o i d v o l u m e fractions f d e t e r m i n e d by S A N S a n d the m e a n particle d i a m e t e r s d of the distributions derived f r o m T E M - Q M A studies (solid lines), disregarding the s t a n d a r d deviation a of the distributions. In addition, the width of the distributions was t a k e n into a c c o u n t (dashed lines) and plotted in

P. Krautwasser et al.

~

........ i ........ i .... i ; ; •

/

Dispersoids in ferritic ODS alloys

..... ; ........ : ~ ; i i ~ = / ,

t

........ ........ ........ i ........ ~ :: ::::i r~:o~ . . ; . . . i . - - ~

...... t .... i........ ]

........ ........

....... ii

',~ ........ ~........ i iii:~::

205

ill

i

..... '..... : ~~~ ......... ~........

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p,

....... ~........ :,.... !

200.

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.

.

i~ ........ ? ........ -

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.

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i

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: : k , ~

Fig. 11. TEM image of extracted dispersoids of alloy PM2010 ( F l l ) heat treated for 3 0 0 h at 1623 K showing the fragmentation of large particles with high aluminium content.

.

150-

0

E o

.... i i i ..........

i

i ii

4-

10

4

6

100

2

~ .'-+-- Siz, ~o&g~:2

T~

PM20,10 i

t,..J •" "

4.

KI

HTr: 300 h, 1623

I

50!

3-

4

;

Anneding time +1 [h] Fig. 10. Variation in mean dispersoid distance with isothermal heat treatment at 1623 K. The calculations were performed for all size ranges with the dispersoid volume fraction f determined by SANS for the mean particle parameter (solid lines) as well as for the mean particle diameter and standard deviation o (dashed lines). Also shown are the results calculated with the mean diameter and dispersoid volume fraction both measured by SANS.

Fig. 10. For comparison, the mean distances for the combined size ranges 1 and 2 calculated from the mean diameters d measured by SANS are shown. It is evident that particles of size range 3 do not contribute significantly to the overall mean distance. Figure 11 shows the dispersoids extracted from a specimen heat treated for 300 h at 1523 K, where some of the larger particles are disintegrating into a large number of small fragments. This is reflected in a high n u m b e r density of small particles ( 1 0 - 4 0 nm) (Fig. 12) measured by SANS in the same specimen. In contrast, these fragments were neglected during Q M A measurements, leading to a discrepancy between the SANS and T E M results in this size range. T h e decomposition of the large particles of size range 3 may be the cause of the decreasing volume fraction of these particles with increasing heat treatment time as shown in Fig. 13. In addition, this plot demonstrates the coarsening of the dispersoids: the volume fraction of size range 1 particles decreases,

o

~.

Q

E

-

m o 0

si~,,~o~g;,, .~-L. ,",

!

i i i

! iilS: ilil' :

i

:

:

i

i

:

i

:

i

.

,,........

i

1

10

i

i

:

:

i

i

100

Diamefer

:

E'" ::3 ~ ~> @

i

1000

[nm]

Fig. 12. Volume distribution of dispersoids in alloy PM2010 (F11) after heat treatment for 300 h at 1623 K: comparison of T E M - Q M A and SANS results.

5,~_ ~.: i ! PM2010 . . . . . ~ ~-K i :: IN: 1623 K 4-. ~ ~ ~.-~..U..z~.•~--~- 1 •- .-.x ...... : ~ .................. i .................. \ :

~r,,g,,÷2

i

-""'--.*-~-"-,--~---.,.

~_ 2 .................... ]................... i ............. : ::*'"""

S~l(~) 0

0.1

i

10 lime [h]

...........

~"

100

1000

Fig. 13. Variation in volume fractions of different dispersoid size ranges with increasing heat treatment time at 1623 K.

206

P. Krautwasser et al.

/

Dispersoids in ferritic ODS alloys

whereas that of size range 2 particles increases. The slight overall decrease in both sizes with increasing annealing time may be real or an artefact due to changes in neutron-scattering contrast caused by the increase in aluminium content with increasing annealing time as discussed above. In Fig. 14 X-ray diffraction patterns of extracted dispersoids of specimens in the as-received condition and after heat treatment for 2000 h at 1123 K are displayed [22]. This figure shows that the number density of the aluminium-rich dispersoids (YAG, Y3AIsO~2 ) is increasing at the expense of the particles with the AIYO3 (YAP) phase. At the same time the number of AI20 ~ particles decreases. This is in agreement with our findings from T E M - E D X analysis, which showed for the dispersoids an increase in aluminium content as the particles grew during prolonged exposure at higher temperatures. The largest particles were pure AI20 3. The dispersoid growth can be interpreted as an uptake of aluminium with the tendency for transformation to YAG as well as the dissolution of small particles.

A comparison of the ODS alloys F 10 (PM2000) and F 11 (PM2010) with different yttria contents is given in Figs. 15 and 16. Both plots demonstrate that in this case a higher yttria content does not lead to a significantly higher number density of the dispersoids in size ranges 1 and 2. Therefore an improvement in the creep strength at high temperatures with increasing Y203 content is not to be expected. Since yttria contents above 0.5 wt.% lead to inferior corrosion resistance [23] but not to an improvement in creep strength, a lower Y203 content as in alloy P M 2 0 0 0 is favourable. 4. S u m m a r y

To study the thermal stability of dispersoids in ferritic ODS alloys, specimens of two alloys, F10

•Es:

~

!

--Fll--

1 h

~FIO--

1 h

---Fll--

1Oh

--FIO--

1Oh

t.t/i

0

,'-.6,

[.l\i

,.o :,,,-,, Y-oxi,, .......

a

i ...........

! ......

...........

,

o~

a b c d e

YAIO3 AIYO3 TiC cub :, AIsY3012 AI203

E ~0 50

1O0

150

e

200

250

300

350

Dlometer rnm]

Fig. 15. Comparison of dispersoid volume distributions of alloys PM2000 (F 10) and PM2010 (F 11 ) measured by SANS. Q_ 0

# C-

b

.4o C n

~

2000 h / 1123 K b a

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Fig. 16. Changes in volume fractions of different dispersoid size ranges of alloys PM2000 (F10) and PM2010 (F 11 ) with increasing heat treatment time at 1623 K.

P. Krautwasser et al.

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Dispersoids in.lerritic ODS alloys

(PM2000) and F11 (PM2010) of Metallwerk Plansee Co., with different yttria contents were isothermally annealed at 1623 K or isochronally heat treated for 110 h in the temperature range 1 2 7 3 - 1 6 2 3 K. The investigations were performed by small angle neutron scattering and transmission electron microscopy in conjunction with quantitative image analysis. It could be shown that three different size distributions of dispersoids ranging from about 5 to 1000 nm are present: dominant small dispersoids in the size range from 5 to about 40 nm (size range 1), medium size dispersoids from about 40 to 100 nm (size range 2) and large particles with diameters above 120 nm (size range 3). The number densities of the different size ranges decrease by a factor of about 100 for each size interval from size range 1 to 2 and from size range 2to3. A comparison of the size distributions determined with both investigative techniques showed good agreement for the dispersoids in the size range 5 - 1 0 0 nm, which influence the creep resistance at high temperatures. In contrast, the sensitivity of SANS measurements decreases strongly for particle sizes larger than 300 nm. This leads to discrepancies in this size range between the two methods. The temperature at which substantial coarsening of the dispersoids during heat treatment occurred was found to be a b o v e 1400 K.

The dispersoid growth can be interpreted as an uptake of aluminium with the tendency for transformation to YAG as well as the dissolution of small particles. Owing to phase transformation and particle disintegration with increasing heat treament time, the volume fractions for each dispersoid size range varied with time. The LSW theory cannot therefore be applied to describe the dispersoid growth. The calculation of the mean interparticle spacing of the dispersoids showed that the group of large particles ( D > 120 rim) does not contribute significantly to the mean particle distance. The comparison of the size distributions of the two alloys with different yttria contents (0.5 and 1.0 wt.%) exhibited hardly any difference for the dispersoids smaller than 50 nm, but showed an increase in the number density of particles larger than 120 nm with higher yttria content.

Acknowledgments The authors wish to thank H. Schuster for initiating the investigations and helpful discussions. Ms. R. Harscheidt for providing the IBAS analysis as well as D. Schwahn, H. Eckerlebe and R. Kampmann for

207

making the SANS facilities at Jiilich and Geesthacht available.

References I E. Arzt and L. Schultz (eds.), New Materials by Mechanical Alloying Techniques, DGM-Informationsgesellschaft, Oberursel, 1989. 2 A. Czyrska-Filemonowicz, E J. Ennis and M. Wrobel, in S. Gorczyca et al. (eds.), 8th Conf. on Electron Microscopy of Solid State, Wroclaw, 1993, p. 271. 3 E J. Ennis~ M. E. Abd EI-Azim and H. Schuster, in E. Bachelet (ed.), High Temperature Materials for Power Engineering, Lifge, September 1990, Part 1, Kluwer, Dordrecht, 199(L p. 143. 4 J. S. Benjamin, T. E. Volin and J. H. Weber, ttigh Temp. -High Press., 6 ( 1974) 443. 5 R. C. Benn and S. K. Kang, l'roc. ('ol!]i Superalloys 84. (7mmpion, PA. 1984, Metallurgical Society of AIME, Warrendale, PA, 1984, p. 319. 6 G. B. Schaffer, M. H. Loretto, R. E. Smallman and J. W. Brooks, Acre Metall., 37(1989) 255 I. 7 R. F. Singer, R. C. Benn and S. K. Kang~ in J. S. Benjamin and R. C. Benn (eds.), Frontiers of ttigh Temperature Materials H,

Proc. 2nd Int. Conf. on ODS Superalloys by Mechanical Alloying, London, May 1983, IncoMAP, INCO Alloys International, New York, 1983, p. 336. 8 R. C. Benn and P. K. Marchandani, in E. Arzt and L. Schultz (eds.), New Materials" by Mechanical Alloying Techniques, DGM-Informationsgesellschaft, Oberursel, 1989, p. 19. 9 U. Gennari, E. Kny and T. Garner, in H. Bildstein and M. Ortner (eds.), Proc. Int. Phmsee Seminar, Reutte, May 1989, Metallwerk Plansee, Reutte, 1989, Vol. 3, p. 567. 10 J. Nutting, S. Ubhi and T. A. Hughes, in J. S. Benjamin (ed.), l"rontiers ~f ltigh Temperature Materials" 1, Proc. 1st Int. Con[i on ODS Superalloys by Mechanical Alloying, New York, 1981, lncoMAP, INCO Alloys International, New York, 1982, p. 33. 11 R. Petkovic-Luton, D. J. Srolovitz and M. J. Luton, in J. S. Benjamin and R. C. Benn (eds.), f)'ontiers of High Temperature Materials 11, Proc. 2nd Int. Conf. on ODS Superalloys by Mechanical Alloying, London, May 1983, lncoMAP, INCO Alloys International, New York, 1983, p. 73. 12 J. K. Tien, in J. S. Benjamin and R. C. Benn (eds.), krontiers of ttigh Temperature Materials" 1L I¥oc. 2nd Int. CotTfi on ODS Superalloys by Mechanical Alloying, London, May 1983, l n c o M A E INCO Alloys International, New York, 1983, p. 114. 13 J.J. Stephens and S. Spooner, Acre Metall., 34 (1986) 303. 14 A. Czyrska-Filemenowicz and E J. Ennis, unpublished results. 15 A. Czyrska-Filemonowicz, K. Spiradek and S. Gorczyca, Prakt. Metall., 22 ( 1991 ) 217. 16 M. Magnani, E Puliti and M. Stefanou, NucL lnstrum. Methods Phys. Res. A, 271 (1988) 611. 17 A. Czyrska-Filemonowitz, P. Krautwasser, M. Widewra and H. Schuster, Hauptversammlung 1991 der Deutschen Gesellschaft ji)r Materialkunde, Graz, May 1991, unpublished results. 18 I. M. Lifshitz and V. V. Slyozov, J. Phys. ('hem. Solids, 19 (1961)35. 19 C. Wagner, Z. Electrochem., 65 ( 1961 ) 581.

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20 J. R6sler and E. Arzt, Acta Metall. Mater., 38 ( 1988) 671. 21 M. E Ashby and R. Ebeling, Trans. Metall. Soc. AIME, 236 (1966) 1396. 22 A. Weisbrodt, H. J. Penkalla, F. Schubert and H. Nickel, Beitrag zur Phasenkinetik in neuen Superlegierungen, Dissertation, RWTH Aachen, 1990.

23 H. Nickel and W. J. Quadakkers, in K. Natesan and D. J. Tillack (eds.), Proc. First Int. Conf. on Heat Resistant Materials, Fontana, W1, September 1991, American Society for Metals, Metals Park, OH, 1991, p. 87.