- Email: [email protected]
A high resolution electron microscope study of the omega transformation in ZrNb alloys
A HIGH RESOLUTION ELECTRON MICROSCOPE STUDY OF THE OMEGA TRANSFORMATION IN Zr-Nb ALLOYS A. L. J. WANG
and S. L. SASS
Department of Materials Science and Engineering and Materials Science Center. Cornell University, Ithaca, NY 13553. U.S.A. and FV. KRAKOW Department of Applied and Engineering Physics, Cornell University, now with the Xerox Corp.. Rochester, NY. U.S.A.
Abstract-High resolution direct lattice imaging and dark field electron microscopy were used to examine the omega phase transformation in Zr-Nb alloys. Direct lattice imaging demonstrated the existence of subvariants within an omega variant. The existence of an ordered sequence of subvariants. which is the basic premise of recent diffuse intensity calculations which seek to explain diffuse diffraction observations in high Nb content alloys, could not be checked because of the small size of the omega regions. fn the low Nb content alloys dark field electron microscopy was used to show that the w phase consists of large domains (ICC-XO A dia.) the interior of which contains features that are 3-6A dia. As the Nb content is increased the omega domains decrease in size until only 3-5 A images are observed in alloys containing 15 wt.0,6 Nb or more. The isolated images are present over the range of composition from 8 to 30 wt. ‘?!,Nb. Time sequence dark field micrographs show that these small images change with time. The diffuse w reflections are believed due in part to the existence of a (111) linear detect. consisting of groups of (111) rows of atoms which are displaced from b.c.c. to (I) positions for short periods of time.
R&urn&-On a utiiist la visualisation directe des plans reticufaires en microscopic efectronique B haute resolution et la microscopic Plectronique en champ sombre pour etudier la transformation de phase omega dans les affiages Zr-Nb. La visualisation directe des plans reticulaires montre f’existence de sous-variantes dans une variante de la phase omega. Par suite de la petite taiffe des regions omega. on n’a pas pu verifier f‘existence dune sequence ordonnee des sous-variantes. qui est a la base de cafcufs recents de f’fntensiti diffuse caicuis cherchant a expfiquer fes intensitis diffuses observks dans ies diagrammes de diffraction des aiiiages a forte teneur en niobium. La microscopic Cfectronique en champ sombre a permis de montrer que. dans fes alfiages a faibie teneur en niobium. la phase omega se presente sous la forme de grands domaines (100-100A de diametre) a I’intirieur desquels on observe des details de 3 a 6 A de diambtre. Lorsqu’ on augmente la teneur en niobium, la taille des domaines omega diminue. si bien que f’on n’observe plus que des images de 3 a 5 A dans les afiiages contenant 15% de niobium (en poids) ou plus. On observe ces images isoiees dans tout fe domaine de com~sition aiiant de 8 a 30% de niobium (en poids). Une serie de champs sombres montre que ces petits details de structure evofuent dans ie temps. On pense que les rtffexions diffuses de type omega sont dues en partie a i’existence dun dtsordre lintaire (f 1i}, consistant en des groupes de rangees d’atomes (111) diplacees de la position c.c.~ la position omega pour un temps assez bref.
Z~mmenfa~~~Die Omega-Tr~sfo~ation wurde mitteis direkter G~tterabbifdung und Dunkefrefd-~iektronenmikroskop~e in Zr-Nb Legierungen untersucht. Die direkte Gitterabbifdung fffhrte auf die Existenz von Subvarianten innerhalb der Omega-Variante. Ob geordnete Sequenzen der Subvarianten existieren. konnte wegen der Kfeinheit der Omega-Gebiete nicht geprtift werden. Eine solche Sequenz ist die Grundannahme bei kiirzlichen Berechnungen der diffusen Streuintensitiit, welche Beobachtungen in Legierungen mit groDen Nb-Anteil zu erkllren versuchen. Bei Legierungen mit geringem Nb-Gehalt wurde Dunkelfefdabbiidung angewendet. Hier besteht die Omega-Phase aus grof3en Bereichen (I#-200 A Durchmesserf, die im Innem Einzefheiten mit 3-6 A Durcfunesser enthalten. Mit ansteigendem Nb-Gehaft verringert sich die GrijBe der o_Bereiche; in Legierungen mit 15 Gew.-?; Yb I oder mehr wurden Kontraste von nur noch 3-.5A Durcfunesser beobachtet. Nach Zeitfofgen von Dunkelfeld-Bildem lndem sich diese kfeinen Kontraste mit der Zeit. Die diffusen Omega-Reflexe werden teilweise der Existenz einer linearen ( 111)-Entordnung zugeschrieben, welche aus Gruppen von (11 l)-Reihen besteht. in denen die Atome kurzzeitig von k.r.z. zu w-lagen verschoben sind. 29
30
CHANG er al: THE OMEGA TRANSFORMATION I&-Zr-Nb ALLOYS
LNTRODUCTION When certain Ti and Zr base alloys (e.g. Zr-Nb, Ti-V) containing a limited amount of solute are quenched, the decomposition of the high temperature fl (or b.c.c.1 phase to the low temperature equilibrium phase(s) is repressed and instead the fl phase transforms partially to a metastable structure called the omega phase. Over the past few years a substantial amount of experimental and theoretical work has been carried out in order to understand the various aspects of the athermdl o transformation, e.g. microstructure, atomic mechanism. diffuse scattering. In this paper we shall consider several of the questions that still remain concerning the transformation and use lattice resolution and dark field electron microscopy in an attempt to answer these questions. The structure of the o phase has been shown by Silcock et al. [l] and Bagaryatskii et al. [2] to be based on the unit ceil in Fig. 1 which is viewed looking along a [ill] direction in the b.c.c. parent phase. The b.c.c. structure, described in terms of this unit ceil, has three atoms located at (O,O,O);(213, l/3, l/3); (l/3, 2/3, 2/3). The omega phase has the same unit cell with atoms B and C shifted along the [ 111) direction to the new positions along the C axis, l/3 + u and 2!3- tl, respectively, where u is equal to or less than l/6, depending upon alloy content [3]. Since there are four equivalent ( 111) directions in the b.c.c. parent phase, there can be four orientations (or variants) of the omega phase. A useful framework for discussing the p-w transfo~ation is the w, curve which represents the variation of the start temperature of the transformation with solute content. Such a curve for the Zr-Nb system, determined by Cometto, Houze and Hehemann [4]. is given in Fig. 2. Here we see that the o, temperature decreases with increasing Nb content, with the 20 wt.% Nb alloy having an w, in the vicinity of room temperature. The electron diffraction patterns at 25% from alloys containing both the beta and omega phases display two kinds of reflections [I-7] (a) sharp reflections common to beta and omega, and (b) reflections due only to omega phase. For low solute content alloys the omega reflections are sharp and with increasing solute content become more diffuse and are shifted along the C*-axis [OOOl] of each variant away from the expected omega positions. The direction of the shift is either in or out, depending on the particular reflection. The ma~itude of the peak shift increases with increasing Nb content, or apparently with decreasing undercooling relative to the camcurve. The diffuse w peaks have also been observed outside of the w, curve, i.e. in the region which should be a b.c.c. solid solution. Moss et al. [83 reported diffuse 0 reflections from a Zr-20 wtTO Nb alloy at 1000°C, or 975” above the tlla tem~rature. Dawson and Sass [63 observed diffuse o reflections in 22 and 3Owt.% Nb alloys at 25C. An important question
that is as yet unanswered concerns the origin of these diffuse peaks from alloys outside of the I’J, curve and their relationship to the /3-cg transformation. Since diffuse w reflections are observed from alloys that are far outside of the 10, curve one can question whether this curve has any significance. Yet measurements of the variation of hardness with Nb content [I] and elastic constants with temperaturetg] are in agreement with the general shape of the curve. In an attempt to explain the diffuse diffraction observations a structural model was proposed recently [lo, 111 which assumes that there are three possible ways to form each variant of omega phase in the b.c.c. parent phase, since the unshifted atom can be chosen to be any one of the three atoms in unit cell, i.e. there are three subvariants per variant. This model is illustrated in Fig. 3. where it is shown that the three subvariants of the omega phase can be formed by (1) keeping the A-plane hxed, letting the B-plane shift up and the C-plane shift down by the same amount, u, (2) keeping the B-plane lixed, letting the C-plane shift up and the A-plane shift down and (3) keeping the C-plane fived letting the A-plane shift up and the B-plane shift down. The diffuse intensity calculations based on such a model show that a preferred order of the subvariants along the [OOOl] direction (e.g. wi followed by ebb, followed by (0~1 will give rise to the observed diffuse peak shifts. The physical basis for this model which is the existence of subvariants, has not yet been demonstrated. The existence of the ordered sequence of subvariants also has not been confirmed. Using dark field electron microscopy Dawson and Sass[6] showed that the omega phase exists as 10-15 A dia. particles in rows along (11 I> directions. In low solute content alloys the small particles are arranged in large omega domains and with increasing solute content these domains decrease in size. Dark field microscopy using the diffuse scattering from alloys which are just inside or are outside of the o, curve was not successful possibly because of insufficient resolution and/or contrast. Using a new technique involving ;Cfossbauer scattering to analyze the energy spectrum of the diffuse intensity from a Zr-20 wt.% Nb allo:; Batterman et RI. Cl23 were able to show that a substantial fraction of the diffuse scattering has undergone an energy shift. This indicates that a substantial fraction of the omega regions which are giving rise to the diffuse scattering are dynamic in nature, i.e. varying with time. The life time of the omega regions was estimated to be of the order of lo-‘see [13]. With this X-ray technique any dynamic process with a longer lifetime would appear static. Dark field microscopy, where the image is formed from the diffuse scattering, should provide information both on its origin, and on the dynamic nature of the omega regions. We should be able to image the omega regions if the dynamic process has a sufficiently wide spectrum of lifetimes extending to the
CHASG er af:
THE OMEGA TRANSFORMATION
-
[OI ilp
1 [TZio],
\,
Eloilp
Fig. 1. The b.c.c. atomic arrangement viewed along a [ 11l] direction. The numbers withk the atoms are the coordinates along the C axis (out of the plane of the paper). The w phase is produced by shifting atoms B and C by an amount u along the C direction toward the ) plane. order of seconds. Dark field observations on the microstructure inside and outside of the o, curve should provide information on the changes that occur on passing through the W$ temperature. Direct lattice imaging should allow the confi~ation of the existence of the subvariants.
EXPERIMENTAL
PROCEDURE
In order to demonstrate the existence of the three subvariants it is necessary to show that within a particular omega variant there is present the same w unit cell, but with its origin translated. The translation vector, t, between subvariants, is written in terms of the basis vectors of the unit cell in Fig. 1 as & @AA,+ )a2 + _sCt.The direct lattice image formed by admitting the 000 and 0001 w reflections in to the objective aperture will exhibit the (0001) or C spacing. If different subvariants are present they should manifest themselves by giving rise, when a boundary is 550.
I
I
I
I
I
/
I
,
<
I
/
,
I
I
500450
crossed, to shifts of the httice fringes equal to the component of t along the 0001 reciprocal lattice vector, or 2 j1 Cl. If it is desired to check the magnitude of the remainder of t, then it would be necessary to form a direct lattice image using a reciprocal lattice vector inclined to [OOOl). The 0001 reflection is an allowed reflection with the largest planar spacing in the w unit cell and. therefore, using this reflection to form the iattice image would minim& the experimental difficulties associated with this technique. A dark field image formed with the 0001 o reflection will serve to define the location of a particular (L,variant and then a lattice image wili be formed of that variant. Electron microscope samples of Zr-8, 12, 15, 20 and 30 wt.7, Nb alloys were prepared by chemical polishing [6]. A Siemens 101 electron microscope operated at 100 kV was used with a special objective lens with spherical aberration coefficient C, = i.35 mm and focaI Length of 2.35 mm. Lattice images were formed by admitting the 000 and 0001 w reflections into a 3Oy objective aperture (0.011 rad aperture haff angle). Dark field images were formed by using either Oool or 0002 reflections with a 20~ objective aperture (0.~73 rad aperture half angle) which limited resolution in this mode of imaging to _ 4 A for best focus. Tilted illumination was used in both cases and images were formed at electron optical magnifications of from 300,000 to 425,000 x . Lattice fringes were observed directly on the final viewing screen by adjusting the objective lens defocus and astigmatism, and beam tilt to optimize contrast. The divergence of the illumination is 1.2 x 10T3 rad with a 400~ condenser aperture. An image reversal technique was used to enhance the contrast of the lattice image during enlargement [ 14,15]. This technique c 8 ,’
A C 8
/
..
I
*y ‘ \
#’
j
L.--..+ -
31
IN Zr-Nb ALLOYS
‘\
400 0 a_. 350
C 5 A
~~I
(,,1\\,,,, /
,(,),
0
4
8 NIOBIUM
I2
I6 CONTENT
20
24
28
(wt.%)
Fig. 2. The camcurve for the Zr-Nb system [4]. This curve represents the variation v&h Nb content of the start temperature of the athermal fi--+w transformation. The dashed line is an extrapolation of the experimentally determined w, curve.
C
-.-
6-----d
:
A-----J :
-4
p---?y
M 1
C-
‘..
:
\
:
A --?
8 A
‘.
/
,---\
,
:
w,
\
I
_I_’
F I
Fig. 3. The ( 11l} stacking sequence of the b.c.c. structure showing the formation of three subvariants of o phase within a single variant [lo]. The subscripts refer to the three subvariants.
32
CHANG rr at: THE OMEGA TRANSFOR,~A~O~
has been found very effective in averaging out long range contrast without disturbing the short range contrast (the fringe pattern). EXPERIMENTAL
RESULTS
Figure 4(a) is a dark field image showing omega domains in a Zr-8 wt.% Nb alloy and Fig. 4(b) is an (0001) fattice image from the same area. The imaging conditions used are given in Figs. 4(c) and Id). In most cases we see that as expected the fringes end at the edge of an omega domain. The dark regions in Fig. 4(a) are either other omega variants or untransformed fl phase. From the dark held image we see that the omega domains contain a fine internal structure with features that are 3-6A dia. Furthermore, a high density of 3-9A images is present in the background where no large omega domains are observed. In the dark field micrograph, what appear to be boundaries between omega domains are indicated by a’, 6’. c’. These regions are also marked in the lattice image and it is evident that the fringe pattern shifts when crossing such a boundary. In order to measure the shift a further enlargement shown in Fig. 4(e) was made of the region in the vicinity of a’. From micrographs such as this the fringe shift was measured to be 0.35 k b05 of the fringe spacing. A similar set of micrographs obtained for a Zr-12 wt.% Nb alloy are presented in Figs. 5(a) and (b). The omega domains are smaller in this alloy than in the 8 wt.% Nb alloy and they still appear to contain an internal structure with extremely small features. Upon comparison of Figs. S(a) and (b) we note that fringe shifts occur when crossing boundaries, and the shift was measured to be 0.35 + 00.5 of the fringe spacing. Since the fringe spacing is /Cl and the fringe shift predicted from Fig 3 is 1/3/C/, our observations in Figs. 4 and 5 provide a direct confirmation of the existence of subvariants within an omega variant. In the rest of the paper the term ‘omega domain’ will be used interchangeably with ‘omega subvariant.’ Both the 8 and 12wt.% Nb alloys do not show a peak shift [6] and in the higher Nb content alloys which do show a diffuse peak shift, the omega particles are sufficiently small (as will be shown) so as to make lattice imaging impossible. Thus it is not possibIe to directly demonstrate that ordered sequences of subvariants occur. We do note in the Zr-8 and 12 wt.% Nb alloys that the subvariants are not arrayed along the [OOOl], but rather along a direction normal to [OOOlJ. At present, however, we cannot confirm the possibility of ordered sequences of subvariants in high Nb content alloys. Dark field images from Zr-15, 20 and 30 wt.% Nb alloys are presented in Figs. 6(a-c). The imaging-conditions are shown in Figs. 6(d-f). We note that images are present with widths of 3-5A and by comparing Figs. 4(a), 5(a), 6(a-c), it is apparent that a high density of these images is present in all the alloys. The large omega domains in the 8 wt.% Nb alloy decrease
IN Zr-Nb ALLOYS
in size in the 12 wt.?; Nb alloy and are no longer present in alloys that contain 15 wt,“.; ti or more. Examination of the small images on the viewing screen indicated that they were changing rapidly with time. This effect is illustrated in the time sequence shown in Figs. 7(a-c) taken from a Zr-3Owt.x Nb alloy. Many of the small images. which are 3-5 A dia., are seen to have a lifetime of the order of IOsec. Time sequence dark field micrographs of the Zr-8 wt.?; Nb ahoy also showed that the tine structure within the omega domains as well as the isolated images were changing rapidly with time.
DISCUSSION The dark field results demonstrate that in alloys well beneath the o, curve large omega domains are present. As the Nb content increases (and the undercooling relative to the w$ curve decreases), the domains decrease in size. While a ( 111) row structure had been observed previously by Dawson and Sass [6], the higher resolution mi~o~aphs taken in this study show a fess obvious (111) d~ectionality; see for example, Figs. 4(a) and S(a). As well, in both the 8 and 12 wt.?,; Nb alloys, the boundaries between subvariants have a (111) trace. Comparison of the microstructure of the Zr-15 and 30 wt.% Nb alloys shows that there is very little difference between the two alloys. This is quite surprising since, as is evident upon comparison of Figs. 6(d) and (f), the o reflections in the 15 wt.Yi Nb alloy are much sharper than in the 30 wt.Pa Nb alloy. We will attempt to understand this observation after the discussion of other results. The presence of smafl images with diameters of 3-5A was noted in all the alloys. The o unit cell has dimensions of approx. 3 by 5 A, and since a finite size objective aperture is being used, the image of the unit cell should be larger than 3 by 5A. It is possible, therefore, that these images are due to deviations from the perfect crystal periodicity on the scale of either single or small groups of atoms [16]. The sequence in Fig. 7 demonstrates that these small images change with time. It is possible that these observations are related either directly or indirectly to the dynamic behavior fl2] of the omega regions which scatter to the diffuse reflections. The observations in Fig. 7 could be directly related to the results of Batterman et al. [ITJ, if there was a wide spectrum of lifetimes extending to seconds associated with the dynamic process. The diffuse reflections are in the form of discs and it is expected that the regions which scatter to them are in the form of rods. Since the images in Fig 7 are equiaxed they cannot he the result of solely imaging the diffuse reflections. It is conceivable that associated with a Iocalized imperfection which gives rise to the small images are groups of (11 I> rows of atoms which are displaced from b.c.c. to o positions for a period of lo-'sec. This
33
CH.WG
et al:
THE OMEGA TRASSFORMGTION
IS Zr .Xb ALLOYS
Fig. S. Observations on a Zr-12 wt.?; Nb alloy. (a) Dark field image formed with the objective aperture located as in Fig. 4(c). The gaps between w domains are indicated by a’, b’, c’. d’. (b) (WI) lattice image of the same area in (a) formed with the objective aperture located as in Fig. 4(d).
second posssibility can be thought of as a linear de. feet or structurai fluctuation associated with a point defect such as a vacancy [17-J. If this linear defect exists with a very short lifetime then a dark field electron microscope image will not reveal its presence since the photographic plate will only record re@ons having lifetimes on the order of seconds. Assuming that short lifetime dynamic processes also occur in the Zr-15 and 3Owt.74 Nb alioys (measurements are in progress [18]), we can then suggest an explanation of why the microstructures of the Zr-15 and 30 wt.?; Nb alloys in Fig. 6 appear similar, while their diffraction patterns are quite different. We believe that the relatively sharp w reflections from the Zr-I5 wt:;, Nb alloy are due to short lifetime structural fluctuations having a spatial extent of several unit cells. In the diffraction pattern the fluctuation will appear frozen in, while in the dark -field image the fluctuation will not be visible. In the Zr-30wt.Y; Nb alloy presumably the short lifetime fluctuations have a much smaller spatiai extent then in the 15 wt.% Nb alloy, thereby
diffuse reflections.
giving rise to more
The origin of the fine structure within the omega domains is not known. The dynamic nature ass& ciated with this fine structure is also not understood. It should be pointed out that the possibility exists that all the dynamic observations of this study may be induced or stimulated by the electron beam, as suggested by NeIson [19] for the migration of point defects. CONCLUSIONS 1. Using direct lattice imaging the existence of three subvariants in an omega variant has been comfirmed. It was not possible to check on the existena: of ordered sequences of subvariants which has recently been proposed to explain diffuse intensity observations. 2. Large omega domains present in the low solute content alloys contain a fine structure MA dia. The omega domains decrease in size with increasing solute content and are no longer observed in alloys containing 15 wt.“; Nb or more. All of the alloys exhibit 3-5A dia. isolated images.
Zr-15
w/~Nb
Zr-20
w/~ Nb
Zr-30
w/o Nb
Fig. 6. Dark field images from ta) Zr-IS wt.?, Nb alloy formed with the objective aperture placed as shown in (d\. ib) Zr-I!Ov,r.“, Nb alloy formed with thz objective aperture placed as show in (2). (c) Zr-30wt.“, Nb alloy uirh the objective aperture placed as shown in ifI. Id). (e). (t) Diffraction pattsrns corrrjponding to dark field images in (al. (bl. IC/. rcspectivsly.
t= I5
t=o
t=30
set Zr-30
w/o Nb
Fig. 7. [s-c) Time sequence dark field microgzaphs taken from a Zr-30 ~i.“~ ir;h alloy using an ima@ng condition similar to that shown in Fig. 610. Landmarks that are retativ&y unchanged \*ith time ;tr~ indicated b) thz ietters ci.!?.The time of sach micrograph is indicated beneath each figure. The lengh of each exposure ~~‘9s6 sec.
36
CHXNG er ut:
THE OMEGA T~~SFOR~fATIO~
3. The internaf structure of the large omega domains and the isolated images are observed to change markedly with time at room temperature over a period of the order of ten seconds. 4. High resolution dark field micrographs from alloys with compositions between 15 and 30 wt.?; Nb exhibit .3-S.+%dia images, while the diffraction patterns show relatively sharp (L)peaks, which with increasing solute content, become diffuse. This apparent disagreement between the appearance of the image and the duration pattern is resolved by noting that the omega regions in these alloys have a dynamic nature. it is suggested that the w diffraction peaks are due to short lifetime structural fluctuations which cannot be imaged. rlcknoH,(edgetnmts-This research was supported by the Metallurgy Branch of the Ofice of Naval Research under Contract N00014-67-A-0077-012, NR031-74. Additional support was received from the National Science Foundation through the use of the technical facilities of the Materials Science Center at Cornell Universitv. We are mateful to Professor B. Siegel for allowing the use of his cemens IO1 electron microscope.
REFERENCES 1. Silcock J. 111..Davies hf. H. and Hardy H. K., Sytnp.
Iz; Zr-Nb ALLOYS
2. Bagaryatskii Iu. A., Nosova G. I. and Tagunova T. V.. D&i. .-L&d. &‘alB SSSR. 105. 1225 i 1955’1. 3. Sass S. L. and Boric B., J. .qpl. Cr.w.‘>,-536 (1972). 4. Cometto D. J., Houze G. L.. Jr. and Hehemann R. F.. Trims. rlI.\fE 233. 30 rl965). 5. Sass S. L., J. less common Merals 28, 157 (1972). 6. Dawson C. W. and Sass S. L., Met. Trans. 1, 2225 (1970). 7. McCabe K. K. and Sass S. L., PM .V\ic1<3.23, 957 (1971). 8. Moss S. C.. Keating D. T. and Axe J. D., Phcrse Tmnsirivns 1973. Pennsvfvania State Universirv. 23-25 Mav. ’ Pergamon Press.‘Oxford (I973f. 9. Goasdoue C., Ho P. S. and Sass S. L., &fir hfet. 20, 725 (19721. 10. Borie B., Sass S. L. and Andreassen A.. Acrrr Crp. IA) 29, 585 (1973). 11. Botie B.. Sass S.’ L. and Andreassen A.. rlcrcr Crq’st. 1.4) 29, 593 (1973). 12. Batterman B. W.. Maracci G.. Merlini A. and Pace S.. Ph_ss. Rev. Left. 31, 277 (1973). 13. Batterman B. W. Private communication. 14. Ottensmeyer F. P., Schmidt E. E., Jack T. and Powell J., J. Ulrrttsrrtrcr. Res. u), 546 (1972). t5. Gonzales F.. i. Ceil Biol. II 1% (196% 16. Krakow &‘.. Chang A. L. J. and Sass S. L., To be published. 17. Kuan T. S. and Sass S. L. To be publbhzd. 18. Lin W., Spait H. and Batterman B. \v.. To be published. 19. Nelson R. S.. Phil. Mug. 10, 723 (196-h.
and S. L. SASS
Department of Materials Science and Engineering and Materials Science Center. Cornell University, Ithaca, NY 13553. U.S.A. and FV. KRAKOW Department of Applied and Engineering Physics, Cornell University, now with the Xerox Corp.. Rochester, NY. U.S.A.
Abstract-High resolution direct lattice imaging and dark field electron microscopy were used to examine the omega phase transformation in Zr-Nb alloys. Direct lattice imaging demonstrated the existence of subvariants within an omega variant. The existence of an ordered sequence of subvariants. which is the basic premise of recent diffuse intensity calculations which seek to explain diffuse diffraction observations in high Nb content alloys, could not be checked because of the small size of the omega regions. fn the low Nb content alloys dark field electron microscopy was used to show that the w phase consists of large domains (ICC-XO A dia.) the interior of which contains features that are 3-6A dia. As the Nb content is increased the omega domains decrease in size until only 3-5 A images are observed in alloys containing 15 wt.0,6 Nb or more. The isolated images are present over the range of composition from 8 to 30 wt. ‘?!,Nb. Time sequence dark field micrographs show that these small images change with time. The diffuse w reflections are believed due in part to the existence of a (111) linear detect. consisting of groups of (111) rows of atoms which are displaced from b.c.c. to (I) positions for short periods of time.
R&urn&-On a utiiist la visualisation directe des plans reticufaires en microscopic efectronique B haute resolution et la microscopic Plectronique en champ sombre pour etudier la transformation de phase omega dans les affiages Zr-Nb. La visualisation directe des plans reticulaires montre f’existence de sous-variantes dans une variante de la phase omega. Par suite de la petite taiffe des regions omega. on n’a pas pu verifier f‘existence dune sequence ordonnee des sous-variantes. qui est a la base de cafcufs recents de f’fntensiti diffuse caicuis cherchant a expfiquer fes intensitis diffuses observks dans ies diagrammes de diffraction des aiiiages a forte teneur en niobium. La microscopic Cfectronique en champ sombre a permis de montrer que. dans fes alfiages a faibie teneur en niobium. la phase omega se presente sous la forme de grands domaines (100-100A de diametre) a I’intirieur desquels on observe des details de 3 a 6 A de diambtre. Lorsqu’ on augmente la teneur en niobium, la taille des domaines omega diminue. si bien que f’on n’observe plus que des images de 3 a 5 A dans les afiiages contenant 15% de niobium (en poids) ou plus. On observe ces images isoiees dans tout fe domaine de com~sition aiiant de 8 a 30% de niobium (en poids). Une serie de champs sombres montre que ces petits details de structure evofuent dans ie temps. On pense que les rtffexions diffuses de type omega sont dues en partie a i’existence dun dtsordre lintaire (f 1i}, consistant en des groupes de rangees d’atomes (111) diplacees de la position c.c.~ la position omega pour un temps assez bref.
Z~mmenfa~~~Die Omega-Tr~sfo~ation wurde mitteis direkter G~tterabbifdung und Dunkefrefd-~iektronenmikroskop~e in Zr-Nb Legierungen untersucht. Die direkte Gitterabbifdung fffhrte auf die Existenz von Subvarianten innerhalb der Omega-Variante. Ob geordnete Sequenzen der Subvarianten existieren. konnte wegen der Kfeinheit der Omega-Gebiete nicht geprtift werden. Eine solche Sequenz ist die Grundannahme bei kiirzlichen Berechnungen der diffusen Streuintensitiit, welche Beobachtungen in Legierungen mit groDen Nb-Anteil zu erkllren versuchen. Bei Legierungen mit geringem Nb-Gehalt wurde Dunkelfefdabbiidung angewendet. Hier besteht die Omega-Phase aus grof3en Bereichen (I#-200 A Durchmesserf, die im Innem Einzefheiten mit 3-6 A Durcfunesser enthalten. Mit ansteigendem Nb-Gehaft verringert sich die GrijBe der o_Bereiche; in Legierungen mit 15 Gew.-?; Yb I oder mehr wurden Kontraste von nur noch 3-.5A Durcfunesser beobachtet. Nach Zeitfofgen von Dunkelfeld-Bildem lndem sich diese kfeinen Kontraste mit der Zeit. Die diffusen Omega-Reflexe werden teilweise der Existenz einer linearen ( 111)-Entordnung zugeschrieben, welche aus Gruppen von (11 l)-Reihen besteht. in denen die Atome kurzzeitig von k.r.z. zu w-lagen verschoben sind. 29
30
CHANG er al: THE OMEGA TRANSFORMATION I&-Zr-Nb ALLOYS
LNTRODUCTION When certain Ti and Zr base alloys (e.g. Zr-Nb, Ti-V) containing a limited amount of solute are quenched, the decomposition of the high temperature fl (or b.c.c.1 phase to the low temperature equilibrium phase(s) is repressed and instead the fl phase transforms partially to a metastable structure called the omega phase. Over the past few years a substantial amount of experimental and theoretical work has been carried out in order to understand the various aspects of the athermdl o transformation, e.g. microstructure, atomic mechanism. diffuse scattering. In this paper we shall consider several of the questions that still remain concerning the transformation and use lattice resolution and dark field electron microscopy in an attempt to answer these questions. The structure of the o phase has been shown by Silcock et al. [l] and Bagaryatskii et al. [2] to be based on the unit ceil in Fig. 1 which is viewed looking along a [ill] direction in the b.c.c. parent phase. The b.c.c. structure, described in terms of this unit ceil, has three atoms located at (O,O,O);(213, l/3, l/3); (l/3, 2/3, 2/3). The omega phase has the same unit cell with atoms B and C shifted along the [ 111) direction to the new positions along the C axis, l/3 + u and 2!3- tl, respectively, where u is equal to or less than l/6, depending upon alloy content [3]. Since there are four equivalent ( 111) directions in the b.c.c. parent phase, there can be four orientations (or variants) of the omega phase. A useful framework for discussing the p-w transfo~ation is the w, curve which represents the variation of the start temperature of the transformation with solute content. Such a curve for the Zr-Nb system, determined by Cometto, Houze and Hehemann [4]. is given in Fig. 2. Here we see that the o, temperature decreases with increasing Nb content, with the 20 wt.% Nb alloy having an w, in the vicinity of room temperature. The electron diffraction patterns at 25% from alloys containing both the beta and omega phases display two kinds of reflections [I-7] (a) sharp reflections common to beta and omega, and (b) reflections due only to omega phase. For low solute content alloys the omega reflections are sharp and with increasing solute content become more diffuse and are shifted along the C*-axis [OOOl] of each variant away from the expected omega positions. The direction of the shift is either in or out, depending on the particular reflection. The ma~itude of the peak shift increases with increasing Nb content, or apparently with decreasing undercooling relative to the camcurve. The diffuse w peaks have also been observed outside of the w, curve, i.e. in the region which should be a b.c.c. solid solution. Moss et al. [83 reported diffuse 0 reflections from a Zr-20 wtTO Nb alloy at 1000°C, or 975” above the tlla tem~rature. Dawson and Sass [63 observed diffuse o reflections in 22 and 3Owt.% Nb alloys at 25C. An important question
that is as yet unanswered concerns the origin of these diffuse peaks from alloys outside of the I’J, curve and their relationship to the /3-cg transformation. Since diffuse w reflections are observed from alloys that are far outside of the 10, curve one can question whether this curve has any significance. Yet measurements of the variation of hardness with Nb content [I] and elastic constants with temperaturetg] are in agreement with the general shape of the curve. In an attempt to explain the diffuse diffraction observations a structural model was proposed recently [lo, 111 which assumes that there are three possible ways to form each variant of omega phase in the b.c.c. parent phase, since the unshifted atom can be chosen to be any one of the three atoms in unit cell, i.e. there are three subvariants per variant. This model is illustrated in Fig. 3. where it is shown that the three subvariants of the omega phase can be formed by (1) keeping the A-plane hxed, letting the B-plane shift up and the C-plane shift down by the same amount, u, (2) keeping the B-plane lixed, letting the C-plane shift up and the A-plane shift down and (3) keeping the C-plane fived letting the A-plane shift up and the B-plane shift down. The diffuse intensity calculations based on such a model show that a preferred order of the subvariants along the [OOOl] direction (e.g. wi followed by ebb, followed by (0~1 will give rise to the observed diffuse peak shifts. The physical basis for this model which is the existence of subvariants, has not yet been demonstrated. The existence of the ordered sequence of subvariants also has not been confirmed. Using dark field electron microscopy Dawson and Sass[6] showed that the omega phase exists as 10-15 A dia. particles in rows along (11 I> directions. In low solute content alloys the small particles are arranged in large omega domains and with increasing solute content these domains decrease in size. Dark field microscopy using the diffuse scattering from alloys which are just inside or are outside of the o, curve was not successful possibly because of insufficient resolution and/or contrast. Using a new technique involving ;Cfossbauer scattering to analyze the energy spectrum of the diffuse intensity from a Zr-20 wt.% Nb allo:; Batterman et RI. Cl23 were able to show that a substantial fraction of the diffuse scattering has undergone an energy shift. This indicates that a substantial fraction of the omega regions which are giving rise to the diffuse scattering are dynamic in nature, i.e. varying with time. The life time of the omega regions was estimated to be of the order of lo-‘see [13]. With this X-ray technique any dynamic process with a longer lifetime would appear static. Dark field microscopy, where the image is formed from the diffuse scattering, should provide information both on its origin, and on the dynamic nature of the omega regions. We should be able to image the omega regions if the dynamic process has a sufficiently wide spectrum of lifetimes extending to the
CHASG er af:
THE OMEGA TRANSFORMATION
-
[OI ilp
1 [TZio],
\,
Eloilp
Fig. 1. The b.c.c. atomic arrangement viewed along a [ 11l] direction. The numbers withk the atoms are the coordinates along the C axis (out of the plane of the paper). The w phase is produced by shifting atoms B and C by an amount u along the C direction toward the ) plane. order of seconds. Dark field observations on the microstructure inside and outside of the o, curve should provide information on the changes that occur on passing through the W$ temperature. Direct lattice imaging should allow the confi~ation of the existence of the subvariants.
EXPERIMENTAL
PROCEDURE
In order to demonstrate the existence of the three subvariants it is necessary to show that within a particular omega variant there is present the same w unit cell, but with its origin translated. The translation vector, t, between subvariants, is written in terms of the basis vectors of the unit cell in Fig. 1 as & @AA,+ )a2 + _sCt.The direct lattice image formed by admitting the 000 and 0001 w reflections in to the objective aperture will exhibit the (0001) or C spacing. If different subvariants are present they should manifest themselves by giving rise, when a boundary is 550.
I
I
I
I
I
/
I
,
<
I
/
,
I
I
500450
crossed, to shifts of the httice fringes equal to the component of t along the 0001 reciprocal lattice vector, or 2 j1 Cl. If it is desired to check the magnitude of the remainder of t, then it would be necessary to form a direct lattice image using a reciprocal lattice vector inclined to [OOOl). The 0001 reflection is an allowed reflection with the largest planar spacing in the w unit cell and. therefore, using this reflection to form the iattice image would minim& the experimental difficulties associated with this technique. A dark field image formed with the 0001 o reflection will serve to define the location of a particular (L,variant and then a lattice image wili be formed of that variant. Electron microscope samples of Zr-8, 12, 15, 20 and 30 wt.7, Nb alloys were prepared by chemical polishing [6]. A Siemens 101 electron microscope operated at 100 kV was used with a special objective lens with spherical aberration coefficient C, = i.35 mm and focaI Length of 2.35 mm. Lattice images were formed by admitting the 000 and 0001 w reflections into a 3Oy objective aperture (0.011 rad aperture haff angle). Dark field images were formed by using either Oool or 0002 reflections with a 20~ objective aperture (0.~73 rad aperture half angle) which limited resolution in this mode of imaging to _ 4 A for best focus. Tilted illumination was used in both cases and images were formed at electron optical magnifications of from 300,000 to 425,000 x . Lattice fringes were observed directly on the final viewing screen by adjusting the objective lens defocus and astigmatism, and beam tilt to optimize contrast. The divergence of the illumination is 1.2 x 10T3 rad with a 400~ condenser aperture. An image reversal technique was used to enhance the contrast of the lattice image during enlargement [ 14,15]. This technique c 8 ,’
A C 8
/
..
I
*y ‘ \
#’
j
L.--..+ -
31
IN Zr-Nb ALLOYS
‘\
400 0 a_. 350
C 5 A
~~I
(,,1\\,,,, /
,(,),
0
4
8 NIOBIUM
I2
I6 CONTENT
20
24
28
(wt.%)
Fig. 2. The camcurve for the Zr-Nb system [4]. This curve represents the variation v&h Nb content of the start temperature of the athermal fi--+w transformation. The dashed line is an extrapolation of the experimentally determined w, curve.
C
-.-
6-----d
:
A-----J :
-4
p---?y
M 1
C-
‘..
:
\
:
A --?
8 A
‘.
/
,---\
,
:
w,
\
I
_I_’
F I
Fig. 3. The ( 11l} stacking sequence of the b.c.c. structure showing the formation of three subvariants of o phase within a single variant [lo]. The subscripts refer to the three subvariants.
32
CHANG rr at: THE OMEGA TRANSFOR,~A~O~
has been found very effective in averaging out long range contrast without disturbing the short range contrast (the fringe pattern). EXPERIMENTAL
RESULTS
Figure 4(a) is a dark field image showing omega domains in a Zr-8 wt.% Nb alloy and Fig. 4(b) is an (0001) fattice image from the same area. The imaging conditions used are given in Figs. 4(c) and Id). In most cases we see that as expected the fringes end at the edge of an omega domain. The dark regions in Fig. 4(a) are either other omega variants or untransformed fl phase. From the dark held image we see that the omega domains contain a fine internal structure with features that are 3-6A dia. Furthermore, a high density of 3-9A images is present in the background where no large omega domains are observed. In the dark field micrograph, what appear to be boundaries between omega domains are indicated by a’, 6’. c’. These regions are also marked in the lattice image and it is evident that the fringe pattern shifts when crossing such a boundary. In order to measure the shift a further enlargement shown in Fig. 4(e) was made of the region in the vicinity of a’. From micrographs such as this the fringe shift was measured to be 0.35 k b05 of the fringe spacing. A similar set of micrographs obtained for a Zr-12 wt.% Nb alloy are presented in Figs. 5(a) and (b). The omega domains are smaller in this alloy than in the 8 wt.% Nb alloy and they still appear to contain an internal structure with extremely small features. Upon comparison of Figs. S(a) and (b) we note that fringe shifts occur when crossing boundaries, and the shift was measured to be 0.35 + 00.5 of the fringe spacing. Since the fringe spacing is /Cl and the fringe shift predicted from Fig 3 is 1/3/C/, our observations in Figs. 4 and 5 provide a direct confirmation of the existence of subvariants within an omega variant. In the rest of the paper the term ‘omega domain’ will be used interchangeably with ‘omega subvariant.’ Both the 8 and 12wt.% Nb alloys do not show a peak shift [6] and in the higher Nb content alloys which do show a diffuse peak shift, the omega particles are sufficiently small (as will be shown) so as to make lattice imaging impossible. Thus it is not possibIe to directly demonstrate that ordered sequences of subvariants occur. We do note in the Zr-8 and 12 wt.% Nb alloys that the subvariants are not arrayed along the [OOOl], but rather along a direction normal to [OOOlJ. At present, however, we cannot confirm the possibility of ordered sequences of subvariants in high Nb content alloys. Dark field images from Zr-15, 20 and 30 wt.% Nb alloys are presented in Figs. 6(a-c). The imaging-conditions are shown in Figs. 6(d-f). We note that images are present with widths of 3-5A and by comparing Figs. 4(a), 5(a), 6(a-c), it is apparent that a high density of these images is present in all the alloys. The large omega domains in the 8 wt.% Nb alloy decrease
IN Zr-Nb ALLOYS
in size in the 12 wt.?; Nb alloy and are no longer present in alloys that contain 15 wt,“.; ti or more. Examination of the small images on the viewing screen indicated that they were changing rapidly with time. This effect is illustrated in the time sequence shown in Figs. 7(a-c) taken from a Zr-3Owt.x Nb alloy. Many of the small images. which are 3-5 A dia., are seen to have a lifetime of the order of IOsec. Time sequence dark field micrographs of the Zr-8 wt.?; Nb ahoy also showed that the tine structure within the omega domains as well as the isolated images were changing rapidly with time.
DISCUSSION The dark field results demonstrate that in alloys well beneath the o, curve large omega domains are present. As the Nb content increases (and the undercooling relative to the w$ curve decreases), the domains decrease in size. While a ( 111) row structure had been observed previously by Dawson and Sass [6], the higher resolution mi~o~aphs taken in this study show a fess obvious (111) d~ectionality; see for example, Figs. 4(a) and S(a). As well, in both the 8 and 12 wt.?,; Nb alloys, the boundaries between subvariants have a (111) trace. Comparison of the microstructure of the Zr-15 and 30 wt.% Nb alloys shows that there is very little difference between the two alloys. This is quite surprising since, as is evident upon comparison of Figs. 6(d) and (f), the o reflections in the 15 wt.Yi Nb alloy are much sharper than in the 30 wt.Pa Nb alloy. We will attempt to understand this observation after the discussion of other results. The presence of smafl images with diameters of 3-5A was noted in all the alloys. The o unit cell has dimensions of approx. 3 by 5 A, and since a finite size objective aperture is being used, the image of the unit cell should be larger than 3 by 5A. It is possible, therefore, that these images are due to deviations from the perfect crystal periodicity on the scale of either single or small groups of atoms [16]. The sequence in Fig. 7 demonstrates that these small images change with time. It is possible that these observations are related either directly or indirectly to the dynamic behavior fl2] of the omega regions which scatter to the diffuse reflections. The observations in Fig. 7 could be directly related to the results of Batterman et al. [ITJ, if there was a wide spectrum of lifetimes extending to seconds associated with the dynamic process. The diffuse reflections are in the form of discs and it is expected that the regions which scatter to them are in the form of rods. Since the images in Fig 7 are equiaxed they cannot he the result of solely imaging the diffuse reflections. It is conceivable that associated with a Iocalized imperfection which gives rise to the small images are groups of (11 I> rows of atoms which are displaced from b.c.c. to o positions for a period of lo-'sec. This
33
CH.WG
et al:
THE OMEGA TRASSFORMGTION
IS Zr .Xb ALLOYS
Fig. S. Observations on a Zr-12 wt.?; Nb alloy. (a) Dark field image formed with the objective aperture located as in Fig. 4(c). The gaps between w domains are indicated by a’, b’, c’. d’. (b) (WI) lattice image of the same area in (a) formed with the objective aperture located as in Fig. 4(d).
second posssibility can be thought of as a linear de. feet or structurai fluctuation associated with a point defect such as a vacancy [17-J. If this linear defect exists with a very short lifetime then a dark field electron microscope image will not reveal its presence since the photographic plate will only record re@ons having lifetimes on the order of seconds. Assuming that short lifetime dynamic processes also occur in the Zr-15 and 3Owt.74 Nb alioys (measurements are in progress [18]), we can then suggest an explanation of why the microstructures of the Zr-15 and 30 wt.?; Nb alloys in Fig. 6 appear similar, while their diffraction patterns are quite different. We believe that the relatively sharp w reflections from the Zr-I5 wt:;, Nb alloy are due to short lifetime structural fluctuations having a spatial extent of several unit cells. In the diffraction pattern the fluctuation will appear frozen in, while in the dark -field image the fluctuation will not be visible. In the Zr-30wt.Y; Nb alloy presumably the short lifetime fluctuations have a much smaller spatiai extent then in the 15 wt.% Nb alloy, thereby
diffuse reflections.
giving rise to more
The origin of the fine structure within the omega domains is not known. The dynamic nature ass& ciated with this fine structure is also not understood. It should be pointed out that the possibility exists that all the dynamic observations of this study may be induced or stimulated by the electron beam, as suggested by NeIson [19] for the migration of point defects. CONCLUSIONS 1. Using direct lattice imaging the existence of three subvariants in an omega variant has been comfirmed. It was not possible to check on the existena: of ordered sequences of subvariants which has recently been proposed to explain diffuse intensity observations. 2. Large omega domains present in the low solute content alloys contain a fine structure MA dia. The omega domains decrease in size with increasing solute content and are no longer observed in alloys containing 15 wt.“; Nb or more. All of the alloys exhibit 3-5A dia. isolated images.
Zr-15
w/~Nb
Zr-20
w/~ Nb
Zr-30
w/o Nb
Fig. 6. Dark field images from ta) Zr-IS wt.?, Nb alloy formed with the objective aperture placed as shown in (d\. ib) Zr-I!Ov,r.“, Nb alloy formed with thz objective aperture placed as show in (2). (c) Zr-30wt.“, Nb alloy uirh the objective aperture placed as shown in ifI. Id). (e). (t) Diffraction pattsrns corrrjponding to dark field images in (al. (bl. IC/. rcspectivsly.
t= I5
t=o
t=30
set Zr-30
w/o Nb
Fig. 7. [s-c) Time sequence dark field microgzaphs taken from a Zr-30 ~i.“~ ir;h alloy using an ima@ng condition similar to that shown in Fig. 610. Landmarks that are retativ&y unchanged \*ith time ;tr~ indicated b) thz ietters ci.!?.The time of sach micrograph is indicated beneath each figure. The lengh of each exposure ~~‘9s6 sec.
36
CHXNG er ut:
THE OMEGA T~~SFOR~fATIO~
3. The internaf structure of the large omega domains and the isolated images are observed to change markedly with time at room temperature over a period of the order of ten seconds. 4. High resolution dark field micrographs from alloys with compositions between 15 and 30 wt.?; Nb exhibit .3-S.+%dia images, while the diffraction patterns show relatively sharp (L)peaks, which with increasing solute content, become diffuse. This apparent disagreement between the appearance of the image and the duration pattern is resolved by noting that the omega regions in these alloys have a dynamic nature. it is suggested that the w diffraction peaks are due to short lifetime structural fluctuations which cannot be imaged. rlcknoH,(edgetnmts-This research was supported by the Metallurgy Branch of the Ofice of Naval Research under Contract N00014-67-A-0077-012, NR031-74. Additional support was received from the National Science Foundation through the use of the technical facilities of the Materials Science Center at Cornell Universitv. We are mateful to Professor B. Siegel for allowing the use of his cemens IO1 electron microscope.
REFERENCES 1. Silcock J. 111..Davies hf. H. and Hardy H. K., Sytnp.
Iz; Zr-Nb ALLOYS
2. Bagaryatskii Iu. A., Nosova G. I. and Tagunova T. V.. D&i. .-L&d. &‘alB SSSR. 105. 1225 i 1955’1. 3. Sass S. L. and Boric B., J. .qpl. Cr.w.‘>,-536 (1972). 4. Cometto D. J., Houze G. L.. Jr. and Hehemann R. F.. Trims. rlI.\fE 233. 30 rl965). 5. Sass S. L., J. less common Merals 28, 157 (1972). 6. Dawson C. W. and Sass S. L., Met. Trans. 1, 2225 (1970). 7. McCabe K. K. and Sass S. L., PM .V\ic1<3.23, 957 (1971). 8. Moss S. C.. Keating D. T. and Axe J. D., Phcrse Tmnsirivns 1973. Pennsvfvania State Universirv. 23-25 Mav. ’ Pergamon Press.‘Oxford (I973f. 9. Goasdoue C., Ho P. S. and Sass S. L., &fir hfet. 20, 725 (19721. 10. Borie B., Sass S. L. and Andreassen A.. Acrrr Crp. IA) 29, 585 (1973). 11. Botie B.. Sass S.’ L. and Andreassen A.. rlcrcr Crq’st. 1.4) 29, 593 (1973). 12. Batterman B. W.. Maracci G.. Merlini A. and Pace S.. Ph_ss. Rev. Left. 31, 277 (1973). 13. Batterman B. W. Private communication. 14. Ottensmeyer F. P., Schmidt E. E., Jack T. and Powell J., J. Ulrrttsrrtrcr. Res. u), 546 (1972). t5. Gonzales F.. i. Ceil Biol. II 1% (196% 16. Krakow &‘.. Chang A. L. J. and Sass S. L., To be published. 17. Kuan T. S. and Sass S. L. To be publbhzd. 18. Lin W., Spait H. and Batterman B. \v.. To be published. 19. Nelson R. S.. Phil. Mug. 10, 723 (196-h.