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Microstructural evolution and mechanical properties of rapidly solidified AlZrV alloys at high temperatures
Acta metall, mater. Vol. 38, No. 5, pp. 771-780, 1990 Printed in Great Britain. All rights reserved
0956-7151/90 $3.00 + 0.00 Copyright @~1990 Pergamon Press plc
MICROSTRUCTURAL EVOLUTION A N D MECHANICAL PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V ALLOYS AT HIGH TEMPERATURES Y. C. CHEN~', M. E. FINE and J. R. WEERTMAN Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208, U.S.A.
(Received 6 October 1988; in revised form 27 September 1989) Abstract--Aging studies of three melt spun, A I - Z ~ V alloys showed that precipitation of the metastable LI 2 phase occurs prior to that of the equilibrium phases, AI~0V and A13Zr(D023). The precipitation mechanism of the LI2 phase can be either continuous or discontinuous. The former is favored by decreasing the value of the Zr/V ratio in the L12 phase precipitate or by increasing the aging temperature. Decreasing the Zr/V ratio also makes it easier to obtain a supersaturated solid solution during melt spinning. Coarsening of the spherical L l:-structured m13(Zr0.25V0.75) particles was studied at 425,450 and 500°C. The activation energy calculated from the particle coarsening kinetics is close to that for diffusivity of Zr in the AI matrix, suggesting that volume diffusion of Zr is the rate controlling mechanism for the Ostwald ripening of the L12-structured A13(Zr025V07~). The LI 2 phase precipitate in extruded samples of this alloy grew slowly at 425°C, close to the rate measured in the melt spun ribbon. Dislocation cross-slip occurred inside the LI2 phase particles from a (l l l) plane to a {100} plane during creep. R6sum~---Des 6tudes sur le vieillissement de trois alliages A I ~ r - V 61abor6s par trempe rapide sur roue, montrent que la pr6cipitation de la phase m6tastable LI 2 se produit avant celle des phases d'6quilibre, AI~0V et AI3Zr(D023 ). Le m6canisme de pr6cipitation de la phase L12 peut 6tre continu ou discontinu. Le premier est favoris6 par la diminution du rapport Zr/V dans la phase pr6cipit6 LI 2 ou par l'augmentation de la temp6rature de vieillissement. La diminution du rapport Zr/V rend aussi plus facile l'obtention d'une solution solide sursatur6e pendant la trempe rapide. On &udie le grossissement des particules sph6riques de Al~(Zr0,2~V0.7~) de structure LI 2 ~i 425,450 et 500°C. L'6nergie d'activation calcul~e fi partir des cin6tiques de grossissement des particules est voisine de celle de la diffusion du Zr dans la matrice d'A1, ce qui laisse supposer que la diffusion en volume du Zr est le m6canisme qui contr61e la vitesse pour le mfirissement d'Ostwald du AI 3(Zr0.25Vo,75) de structure L 12. Le pr6cipit6 de phase L 1~dans les 6chantillons extrud6s de cet alliage croit lentement fi 425' C, ~, une vitesse voisine de celle qu'on mesure sur le ruban tremp6. Le glissement d6vi6 des dislocations se produit a l'int6rieur des particules de phase L12 d'un plan (111) vers un plan {100} pendant le fluage. Zusammenfassung--Auslagerungsexperimente an drei schmelzgesponnenen Al~r-V-Legierungen zeigten, dab sich die metastabile L12-Phase vor den Gleichgewichtsphasen AI~0V und AI3Zr(D023) ausscheidet. Der Ausscheidungsmechanismus f/Jr die Ll2-Phase kann entweder kontinuierlich oder diskontinuierlich sein. Der kontinuierliche Mechanismus wird begfinstigt, wenn das Zr/V-Verh/iltnis in der LI,Ausscheidungsphase verringert oder die Auslagerungstemperatur erh6ht wird. Verringerung des Zr/VVerh/iltnisses erleichtert es auch, w~ihrend des Schmelzspinnens eine iiberstfittigte feste L6sung zu erhalten. Die Vergr6berung der kugelf6rmigen A13(Zr0.25V0.ys)-Teilchen (L12-Struktur) wurde bei 425, 450 und 500"C untersucht. Die aus der Vergr6berungskinetik der Teilchen berechnete Aktivierungsenergie liegt nahe der ffir die Diffusivitfit von Zr in der A1-Matrix; dieser Befund legt nahe, dab hier die Volumdiffusion des Zr der ratenvestimmende Mechanismus der Ostwaldreifung dieser Teilchen ist. Die L12-Ausscheidungsphase in stranggeprel3ten Proben dieser Legierung wuchsen bei 425:C langsam mit einer Rate nahe der, die in dem schmelzgesponnenen Band gemessen wurde. W~ihrend des Kriechens trat Quergleitung xon Versetzung innerhalb der Ll2-Phasenteilchen von einer (111)- in eine {100}-Ebene auf.
INTRODUCTION
Co o f its rate-controlling element is small. A m o n g the transition metals, Z r has the smallest diffusion flux DCo in AI [2]. The A13Zr phase in binary AIZr alloys can form either in a semicoherent tetragonal equilibrium phase or a coherent metastable LI 2 form [3-6]. Both forms are expected to form a low energy interface with the matrix, with that of the latter being especially low. Partial substitution of V for Z r in the precipitate increases the thermal stability of the metastable L12 phase in very dilute alloys a n d also slows d o w n the particle coarsening rate at 425°C [7].
T o improve the strength of A1 alloys at high temperatures a microstructure c o n t a i n i n g a thermally stable a n d coarsening resistant dispersoid is required [1, 2]. According to theory, a particle will be resistant to coarsening if the energy o f its interface with the matrix is low, a n d if the diffusivity D a n d solubility tPresent address: Materials Research Laboratories, 195-5 Chung-Hsing Road. Chutung, Hsinchu, Taiwan 31015, R.O.China. 771
772
CHEN et al.:
PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES
In the present study, three more concentrated A I - Z r - V alloys with different Zr/V ratios were prepared by melt spinning. These alloys were used to study the influence of the Zr/V ratio on the melt spun structure and the thermal stability of the metastable L12 phase with respect to phase change and Ostwald ripening. Coarsening and mechanical properties of extrusions made from the melt spun ribbons were evaluated at 425°C. A very slow growth rate of LI2structured precipitate at 425°C was reported in melt spun ribbons of A1-5 vol.% A13(Zr0.25V0.75) [8]. The temperatures have now been extended to 500°C. EXPERIMENTAL
Three A I - Z r - V alloys were prepared by melt spinning. The compositions of the alloys were chosen with the goal of producing material with 5 vol.% Ll2-structured Al3ZrxV~_x. The solid solubilities of both Zr and V in AI are 0.1wt% at 425°C. Assuming the solubilities to be the same in the ternary A1-Zr-V system, alloys were designed with values of x of 0.75, 0.5 and 0.25, designated alloys 2, 3 and 4, respectively. The nominal compositions are given in Table 1. Extruded bars having dimensions 6.4 × 14mm were made from the ribbons through vacuum hot degassing, hot pressing and hot extrusion. The processing details are described elsewhere [9]. The specimens for optical microscopy were prepared by electropolishing and etching in Keller's reagent. TEM foils were prepared in a double jet polishing apparatus with a solution of 25 vol. % nitric acid in methanol. The temperature of the electrolyte was kept between - 4 0 to - 6 0 ° C and a current of 100mA was applied. The melt spun ribbons were mechanically polished with 600-grit abrasive paper before the electropolishing. Some specimens were selectively polished on one surface with the other protected by lacquer to examine the near surface microstructure. After electropolishing the lacquer was dissolved in acetone. Sometimes ion milling was applied at 2 mA and 4 kV for half an hour at a grazing angle of 15° to remove surface contamination. Extracted second phase particles were obtained by electrolytic dissolution of the A1 matrix using an electrolyte which contains 60 g benzoic acid, 15g 8-hydroxyquinoline, 60ml chloroform and 160 ml methanol [10]. The sizes of the L12 precipitates were measured from TEM negatives taken by the central dark field (DF) technique using superlattice reflection spots. More than 400 particles were measured for each Table 1. Nominal compositions of three AI-Zr V alloys Alloy
Composition (wt%)
2 3 4
AI-3.12 Zr4).66 V AI-2.14 Zr-1.24 V Al-l.13 Zr 1.82V
Fig. 1. Typical TEM microstructures of (a) alloy 2, (b) alloy 3 and (c) alloy 4 ribbons in the as-melt spun condition. determination of average particle radius. The precipitate free zone (PFZ) widths were measured with the beam direction lying in the grain boundary plane. Creep tests were conducted at 4 2 5 _ 2°C on a dead load creep machine. Strain was monitored with a calibrated linear variable differential transducer. Tensile tests were carried out on an MTS machine at 425°C. Both kinds of mechanical tests were done in an Ar atmosphere. The specimens had gage dimensions of 20 by 6,6 by 0.9 mm (S1) or 20 by 4 by 2 mm ($2).
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED A1-Z~V AT HIGH TEMPERATURES RESULTS
Microstructures and phase transformations The melt spun ribbons were approximately 20-60/tm thick and 2 3 mm wide. Occasionally small globules of previously molten metal were stuck to them. The specimens investigated in this study were chosen to avoid such regions. The microstructure of alloy 4 varied slightly from place to place, depending upon the contact of the ribbon with the chilled Cu wheel. Columnar grams are observed where the chill surface is flat. Howcver, in some areas where the chill surface is unexen, a dark-etching region is present near the free surface of the ribbon. A lower cooling rate is associated with such an uneven contact surface. TEM selccted area diffraction (SAD) showed that the second phase particles present near the free surface of alloy 4 have the structure of All0V [11]. No such particles were found in the as-received ribbons of alloys 2 or 3. Typical microstructures from the areas close to the middle of ribbons of alloys 2, 3 and 4 are shown in Fig. 1. In alloys 2 and 3, which have higher Zr/V ratios, cellular fan shaped precipitates
773
form inside the grains [Fig. l(a,b)]. The SAD patterns from these cellular precipitates are consistent with the LI 2 structure, such as reported for metastable AI3Zr in AIZr alloys [4, 5]. The grain boundaries are rather irregular and island grains surrounded by other grains are often observed in these two alloys. In alloy 4, no precipitates were noted inside the grains and the grain boundaries are more regular, as seen in Fig. l(c). Some small second phase particles, approximately 0.1/~m in size, are present along the grain boundaries. Selected area diffraction from the larger of these particles proved that they have the Al~0V structure. This precipitate is likely the source of the etch pits seen at grain boundaries. In alloys 2 and 3 spherical precipitates were interspersed with the cellular precipitates. In TEM dark field micrographs, Fig. 2(a) and (b), this phase is clearly distinguishable. A TEM diffraction pattern from this area and a schematic drawing of it are shown in Fig. 2(c) and (d). The electron beam direction is close to the ( 100 ) of the AI matrix. Since all spherical phase particles in one grain always light up at the same time with one operating reflection, they have the same crystallographic orientation.
•
•
o
•
•
•
•
o
o
110Lie o
•
0 000
o o
•
o
• 111
(d)
•
Fig. 2. TEM dark field micrographs from as-spun alloy 2 ribbon and its associated SAD patterns where (a) was taken with [110] spot of the LI 2 phase, (b) was taken with [111] spot of the unidentified phase, (c) is SAD and (d) is a schematic drawing of SAD. The open circle represents reflections from the cellular LI, phase and the AI matrix and the closed circle represents reflections from the unidentified phase.
774
CHEN
et al.:
PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
However, as seen in Fig. 2(d), a cube/cube orientation relationship with the A1 matrix does not exist for these particles. Figure 3 illustrates two other selected area electron diffraction patterns from these particles. The electron diffraction patterns shown in Fig. 2(c) and Fig. 3(a) and (b) can be indexed as (112), (100) and (110), respectively, based on a simple cubic
structure with the lattice constant equal to 4.06/~. As seen in Fig. 3(a), where the specimen was tilted to near the [100] zone axis, the intensities of {002} spots are stronger than those for {001} spots. These may be Ll2-structured particles but they do not have the normal cube-cube orientation with the matrix. On aging at 500°C for 1 h, All0V and D023structured A13Zr particles formed on the grain boundaries [11]. After aging alloy 4 at 350°C for 3 h, some cellular L12 phase had formed from grain boundaries, as shown in Fig. 4. The cellular LI 2 phase
Fig. 5. High resolution electron micrograph of an AI~0V particle.
Fig. 3. Selected area electron diffraction patterns from the unidentified phase where (a) can be indexed as (100) and (b) as (110).
Fig. 4. TEM of alloy 4 ribbon aged at 350°C for 3 h.
Fig. 6. High resolution electron micrograph of an L12 particle in A1 matrix from alloy 4 ribbon aged at 600°C for 1.5h.
CHEN et aL:
PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES
has been suggested by Nes and Billdal [12] to form by discontinuous precipitation resulting from grain boundary movement. The presence of the original traces of the grain boundaries, indicated by arrows in Fig. 4, left behind by the cellular decomposition confirms this. In other areas SAD shows the matrix in alloy 4 to be still free of any precipitate. During further aging to 30 h at the same temperature [11], small spherical L12 phase particles form by continuous
14.01 j
6.c~-
~lt-
4.0
2 O~
precipitation inside all the previously undecomposed grains. The fraction of the cellular LI2 precipitate decreases with increase of the aging temperature [7, 11]. On aging at 600°C, cellular precipitates are not found at all. During aging of alloy 4 ribbon, All0V and some D023-structured AI3Zr, precipitate on the grain boundaries leading to the formation of LI 2 precipitate free zones (PFZs). Both grain boundary
0 T : 425* C
12.0~I lo.o~
'o
775
A T : 450"• C u T:500 C
.~ f /
.
0 T
~ •
~ i
f l
~ I
~ I
I
I
0 50 100 150 200 250 300 350 400 450 Annealing time (hrs.)
Fig. 7. Growth kinetics of LI 2 A%(Zr0.25V075)at various temperatures in alloy 4 ribbon. A pre-aging of 2.5 h at 500°C was done for the specimens aged at 425 and 450°C.
Fig. 8. Llz-structured A13(Zr025V07s) particles in alloy 4 ribbon after aging at 500°C for (a) 20 and (b) 206 h.
Fig. 9. Typical microstructures (TEM) of (a) alloy 2, (b) alloy 3,.and (c) alloy 4 after extrusion. Extrusion direction is normal to the micrograph in each case.
776
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
precipitate size and P F Z width increase with aging time [8, 13]. A high resolution TEM (HRTEM) micrograph of AI~0V, Fig. 5, reveals microtwins and a stacking fault. The existence of twins is also indicated by the twin spots shown in the electron diffraction pattern (inset in Fig. 5). Using the energy dispersive X-ray spectrometry (EDS) attachment to the STEM, the average composition of extracted "Al~0V'' particles analyzed to approximately 11 at.% V and 1 at.% Zr. A standardless analysis technique and thin film approximation were used. For simplicity, the nominal designation AI,0V is used for this ternary grain boundary precipitate throughout this paper. An electron diffraction pattern and the HRTEM micrograph of L12-A13(Zr, V) after aging at 600°C for 1.5 h are shown in Fig. 6 where each white spot corresponds to the image from the projection of AI atoms stacked along the [110] direction. The presence of order in the LI2 structure is clearly shown. A perfect lattice match between the LI2 particle and the AI matrix and a stepped interface are also seen in this figure. The growth kinetics of the spherical L12 particles in alloy 4 ribbons were determined at 450 and 500°C. Detailed coarsening results at 425°C have been reported previously [8]. Figure 7 shows the coarsening kinetics of the L12 particles at these three temperatures, where the cube of the average particle radius, f, is plotted against aging time. The linear correlation coefficient, R, and the slopes for the best fitting lines are given in Table 2. Typical L12 particles after aging at 500°C for 20 and 206 h are shown in Fig. 8(a) and (b), respectively. A planar fault exists inside many of the LI2 phase particles seen in Fig. 8(b). These planar faults are always parallel to the {100} planes of the LI2 phase. TEM study indicated them to be antiphase boundaries (APBs) which are associated with a displacement of a / 2 ( l 1 0 ) on a {100} plane, where 20.0
Fellure
T : 425"C
18.0
C~t : 17 MPcl
16.0 14.0 --
I loy 2
12.0 c 10.0
Fig. 11. Dark field TEM micrograph from alloy 4 after creep at 425°C. Shearing of L12 phase particles has resulted in formation of faults.
a is the lattice constant of the LI z phase [11]. In an L12 precipitate containing an APB, preferential growth along the APB is noted. This causes some L12 precipitates to appear as rounded parallelepipeds rather than rounded cuboids. Microstructures and mechanical properties o f extruded alloys The microstructures of the three alloys after consolidation and extrusion look similar to one another in the optical microscope. Coarse second phase particles of 1 #m in size are observed in some areas. Presumably, these regions came from the previously mentioned globules of molten metal which splashed onto already solidified ribbons. More of these coarse second phase particles were found in extruded alloy 2 than in the other two alloys. TEM observation reveals that the grain sizes, or subgrain sizes, in all three alloys were reduced to around 0.5 #m after extrusion, as shown in Fig. 9. In extruded alloy 4, all the L12 phase particles are spherical. No cellular precipitation is observed. While the size of the L12 particles is quite small, the number of particles in the extruded material appears to be somewhat less than in aged melt spun ribbon. In some grains, stable tetragonal AI3Zr particles are also found. The growth of the spherical L12 particles at 425°C in extruded alloy is nearly the same as that in ribbons, 1.1 × 10-28m3/h compared to 1.0 x 10 -28 m3/h.
8.0
6.0
FOil ure /
4.0
Ter mi notecl J
2.0 o lO "1
~ I 1 creep
l lO time
o
y
4 lO z
lO 3
(hrs.)
Fig. 10. Creep curves of three extruded AI-Zr-V alloys tested at 425°C with an initial stress of 17 MPa.
Table 2. Measured volumetric coarsening rate constants K and coefficientsof linearity R for LI2AI3(Zr0.25Vo75) precipitatesin alloy 4 aged at 425, 450 and 500°C Temperature K (°C) (m3/h) R 425a 1.03 x 10-2s 0.997 450m 6.13 x l0 28 0.997 500 1.28 x 10-26 0.996 ~Specimenswere preagedat 500°C for 2.5 h to prevent cellularprecipitation.
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES Tensile properties at 425°C for extruded alloys 2 and 4 are given in Table 3. Different strain rates as indicated were employed to acquire yield stress and tensile strength in accordance with A S T M E 151-64. Each data set was obtained from the average of two tests. At 425°C alloy 2 has a lower yield strength than alloy 4, 13.8 MPa compared to 42.7 MPa. The elongations obtained at 425'~C for both alloys are less than those for room temperature which were around 30% [91. Creep tests were done at 425°C with an initial stress of 17 MPa. Typical creep curves for the three extruded alloys, using the SI specimen geometry, are shown in Fig. 10. As seen. alloy 2 creeps much faster than alloys 3 and 4 correlating with the lower yield stress. The fracture surfaces are also different for alloy 2. A dimple structure and a large reduction in area are seen indicating a ductile fracture. In alloys 3 and 4 the failure began in an intergranular fracture mode but changed to the ductile dimple mode with a large reduction in area. In extruded alloy 4, spherical L12 phase particles were sheared along {100} planes during creep as seen in Fig. 1 1. Since the slip planes in the AI matrix are { 111 }, it implies that dislocations cross-slipped from {111} planes to {100~ planes inside the L12 particles during creep. Occasionally, a pile-up of dislocations is observed within a PFZ, suggesting that localization of plastic deformation may occur in this soft region. The steady state creep rate, ~s, at 425°C was determined as a function of creep stress for extruded alloy 4 using the thicker $2 specimen. A partial correction tbr the change in stress caused by the decrease in cross-sectional area during the creep test was made using cr = cri/(l -- E)
(1)
where ~r~ is the initial stress, a is the corrected stress and E is the strain at the beginning of steady state creep. The total change in area during the test was only 1 or 2°'0. The normalized steady state creep rate vs normalized stress (7i is plotted in Fig. 12. The normalized creep rate is defined as ~skT/DGb, where Table 3. Tensileproperties of extruded alloys 2 and 4 tested at 425°C in Ar Yield stress~ Ultimatetensileb (0.2'~0 offset) strength Elongation Alloy (MPa) (MPa) (%) 2 13.8 22.0 16.5 4 42.7 66.8 21.8 ~Measured with ~ - 1 x 10 4/s. bMeasured with ( = 1 x 10 */s.
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1 0 .7
.~
lo .8
777
///
10-9
1 0 -1° I 7X1(~ 4
I
I 10.3
2 X 1 ~ :)
Fig. 12. Log-log plot of normalized creep rate ~skT/DGb vs normalized creep stress tUG for alloy 4. D is the lattice diffusion coefficient of A1 and b is the Burgers vector, and the creep stress is normalized by dividing it by the shear modulus G. The stress exponent n ( ~ s ~ a " ) is found to be close to 14; however, the normal value for A1 may be obtained by introducing a threshold stress, as discussed later. DISCUSSION
Microstructure of the rapidly solidified A I - Z r - V alloys The as-melt spun microstructures show precipitation has occurred in alloys 2 and 3 but not in alloy 4. Dissolved zirconium increases the lattice constant of the AI matrix. However, vanadium behaves oppositely. The change in lattice constant of the A1 solid solution on dissolution of 1 wt% of Zr or V are +0.0005 nm and -0.00033 nm, respectively [14-16]. Based on these values and Vegard's law, the changes of the lattice constant, Aa, for the three alloys in the supersaturated solid solution state were calculated and are listed in Table 4. Alloys 2 and 3 have larger lattice expansions over unalloyed AI than alloy 4. For this reason, supersaturated alloys 2 and 3 are expected to be more unstable to decomposition by precipitation than supersaturated alloy 4. The lower liquidus temperature of alloy 4 due to the lower Zr content also plays an important role in suppressing precipitation. The metastable Ll2-structured A13(Zr, V) phase may form in cellular or spherical morphology. It is
Table 4. Calculatedchange of lattice constant. Aa, of super-saturated solid solutions of AI-Zr-V alloys from that of unalloyed AI Aa (A)a Aa (,~)b Aa (A) Alloy (from Zr) (from V) (total) 2. AI-3.50wt% Zr4).54 wt% V +0.0175 -0.0016 +0.0159 3. AI-2.15wt% Z~1.22 wt% V +0.0108 - - 0 . 0 0 3 7 +0.0071 4. AI 1.16wt% Zr-l.76 wt% V +0.0058 -0.0053 +0.0005 ~Aa/wt% Zr = + 0.005 ,~/wt%. ~Aa/wt% V = -0.0033 ,~/wt%.
778
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
expected that discontinuous precipitation is kinetically faster than continuous precipitation at low aging temperatures since the diffusion path is shorter. Consequently, L1 z phase in cellular precipitation is dominant in alloy 4 ribbon when aged at 350°C. With increased aging temperature the matrix diffusion rate increases leading to the result that the spherical precipitation mode is dominant in alloy 4 ribbon after aging at 500°C or higher. As a result, the cellular L12 phase may be suppressed by an initial treatment at a higher aging temperature as found by Zedalis in a much more dilute A I - Z r - V alloy [7], Alloy composition may also play a role in the precipitation kinetics of the L12 phase. In alloys 2 and 3, which have lower V/Zr ratios, the cellular L12 phase forms on cooling after solidification. As shown by Zedalis [7], the L 12-structured AI 3(Zrx V 1 _ x ) phase with the highest V content and smallest x has the best lattice matching with the AI matrix. Thus, the effect of the V/Zr ratio on the continuous precipitation of the L12 phase may be explained by a lower energy barrier to homogeneous nucleation. In addition to the cellular L12 phase, an unidentified simple cubic precipitate is also present in alloys 2 and 3. It does not show the normal cube/cube orientation relationship with the AI matrix observed with the L12 phase. This has not been reported previously in binary AI-Zr alloys. Due to the similarity of lattice structures between the L12-structured A13(Zr, V) and this unidentified phase, it is probably also L12-structured A13(Zr, V ). If this is true, the different orientation relationship with the AI matrix might occur by sequential precipitation, as suggested by the reviewer. After solidification the LI 2 phase particles formed firstly by continuous precipitation. Later on discontinuous precipitation took place associated with the movement of the grain boundaries, which resulted in formation of the cellular L12 phase leaving the initially precipitated LI2 phase particles embedded in the cellular regions. The original orientation relation with the matrix changed as the grain boundary swept over them. However, most of the unidentified phase particles formed in a linear array instead of being statistically homogeneous, which seems at variance with this explanation.
Coarsening kinetics of the spherical Lle phase in the melt spun ribbon of alloy 4 The growth behavior of the L12 particles at 425, 450 and 500°C shows good agreement with the ~3 vs t relationship, f 3 _ 73° = Kt, which suggests that an LSW-type volume diffusion controlling process I2] is the coarsening mechanism. Here f is the average particle radius after aging for time t and ~o is an integration constant. If the effect of the precipitate volume fraction on Ostwald ripening is taken into account, the coarsening rate constant K for an alloy containing 5 vol.% of precipitates is given by [17] K =
12yDV2 CO 9RT
(2)
-SB
-6C
- Q / R : -35.38 x 103 Q=294 kJ//mole I
-62
{:
-64
-66
-6e
i
1.20
I
J
I
1.30
i
1.40
1.50
~/T(x~O-~ K-~)
Fig. 13. Plot of ln(KT) vs lIT for the coarsening of the L12-structured A13(Zro.25V075). where y is the interfacial energy, Vm is the molar volume of the precipitate, and R is the gas constant. This equation has been used by others to determine the activation energy of diffusivity of solute in the matrix from the experimental coarsening results [18-20]. After rearrangement, equation (2) becomes
ln(KT/Co) = In A - Q/RT
(3)
where A is given by A = 12~DoVL
9R
(4)
and Q is the activation energy of diffusivity for the rate-controlling element in the matrix. The value of A is a slowly varying function of temperature and usually may be regarded as a constant over a small temperature range. Therefore, the value of Q can be experimentally determined from a plot of ln(KT/Co) vs the reciprocal of temperature. In the present alloy, neither the concentration C0 of V or of Zr in the solid solution AI matrix in equilibrium with the metastable LIz phase is accurately known. For a first order approximation, Co is assumed to be constant in the temperature range from 425 to 500°C. From the plot of ln(KT) vs I/T shown in Fig. 13, Q is calculated to be 294 kJ/mol. This value is closer to the activation energy for diffusion of Zr in AI, 242 kJ/mol, than to the value for V, 82kJ/mol [3,21]. Consequently, the volume diffusion of Zr is likely the rate controlling process for the Ostwald ripening of the Llz-structured A13(Zr0.25V0.75), as expected originally. In binary A1-Zr alloys, the solvus line of the stable phase A%Zr(D023), Co, can be expressed by the following equation, in wt% [22]
Co = 1180 exp(-64.96 kJ/RT).
(5)
If it is assumed that the Zr concentration of the AI matrix in equilibrium with the metastable L12 phase,
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES C~, does not change much from Co, then a better calculation for Q can be obtained by considering the temperature dependence of C'o from Equation (5). The final result for Q is 229 kJ/mol. This value agrees even better with the activation energy of diffusivity for Zr in A1 measured in a binary AI-Zr alloy [21]. It was also found that the growth kinetics of the PFZ at 425 ° in alloy 4 ribbon is controlled by the volume diffusion of Zr in the AI matrix [13].
779
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Mechanical properties elevated temperature
of
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at
i~ 0 -1611
The poor creep resistance of extruded alloy 2 (Fig. 10) is related to its low yield stress, 14 MPa at 425°C, which is lower than the creep stress, 17 MPa. The plastic tearing seen on the fracture surface after failure by creep confirms this. It is of interest to know why this alloy possesses a low strength compared to the other two extruded AI-Zr-V alloys since they all contain around 5 vol.% of precipitates. TEM investigation disclosed that in alloy 2 somewhat larger amounts of the equilibrium phases compared to the other two extruded alloys were present after annealing of 1 h at 425°C. This difference was not seen in the melt spun ribbons of the three alloys. Thus it indicates that after hot extrusion the thermal stability of the L12-structured A13(Zr075V,.25), with a high Zr/V ratio is lower. Another factor contributing to the weakness of alloy 2 at 425°C is its inhomogeneity. Regions with coarse second phase particles were often observed in alloy 2. Intergranular fracture in creep is expected to be more prominent in particle hardened alloys. This is confirmed by the present results for alloys 3 and 4. Localization of plastic deformation along the PFZs, which are softer than the precipitation hardened matrix, and sliding of the grain boundaries or subgrain boundaries are assumed to be the mechanisms for the intergranular fracture in these two alloys. As the crack deepens the creep stress increases because of the reduction in cross-section, becoming 34 MPa when the cross-section is one-half of its original value which is comparable to the 0.2% offset yield stress. 43 MPa, of alloy 4 at 425°C. Accordingly, a transition is expected from an intergranular to a transgranular fracture exhibiting a dimple structure as observed. Creep theories for particle strengthened alloys [24-27] suggest that processes occurring in the matrix are rate controlling. Thus, the creep rate equation for particle strengthened alloys is expected to be in the same form as that for the matrix material substituting an effective creep stress 0"ef t equal to a --ao for the actual stress a, where a o is the threshold stress. The creep rate for extruded alloy 4 may be expected to fit the equation DGB a ~
gs=AD k T
-
-
a o 4.4
(6)
1612
16~3 1(]TM 10
I 12
I
I
I
14
16
18
I
I
I
20 22 24 O" ( M P a )
I
I
26
28
I
30 32
Fig. 14. Comparison of creep strain rate vs stress for alloy 4 (data points) with prediction from equation (6) with a o equal to 14.2 MPa (curve). where the exponent 4.4 is that obtained for AI in the power law regime and AD is an empirical constant, 3.4 x 106 for AI [28]. By linear regression, the value of ao is calculated to be 14.2 MPa. A plot of creep rate vs creep stress for the data obtained from alloy 4 is compared with the curves predicted by equation (6) in Fig. 14. The actual strain rates are off by one order of magnitude from the predicted curve but the curves are essentially parallel. A higher creep rate compared with that from equation (6) might be caused by localization of deformation within the PFZs. SUMMARY AND CONCLUSIONS 1. In rapidly solidified AI-5 vol.% A13(ZrxV~ x) alloys, the metastable L12 phase precipitates initially instead of the equilibrium phases, Alj0V and A13Zr(D023). The precipitation mechanism of the L12 phase depends on alloy composition and aging temperature. In the alloys with a Zr/V ratio equal to 3 or 1, a cellular L12 phase forms by discontinuous precipitation immediately after solidification. However, a supersaturated solid solution is obtained in an alloy with a Zr/V ratio equal to 1/3. In the latter, the cellular mode is dominant when aging is carried out at low temperatures, but spherical LI 2 particles form when aging is done at higher temperatures. 2. A good fit linear relationship between the cube of the average particle radius and aging time is obtained for the L12-structured A13(Zr0.25V0.7s) precipitate aged at 425, 450 or 500°C. Volume diffusion of Zr in the AI matrix appears to be the rate controlling mechanism for coarsening. The coarsening rate of the L12 phase precipitate at 425°C is 1.1 x 10-28m3/h. 3. In extruded samples of the AI 5v01.% A13(Zro.25Vo.75), alloy 4, L12 phase precipitates were
780
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
sheared by dislocations on { 100} planes during creep, suggesting that cross slip of dislocations from a { 111 } plane to a {100} plane occurred inside the LI 2 phase precipitate. With the introduction of a threshold stress of 14.2 MPa, the data may be fit to a creep equation with the stress exponent obtained for pure AI, 4.4. However, the experimental creep rates are faster than those predicted. Acknowledgements--This research was supported by the Air Force Office of Scientific Research Grant No. AFOSR85-0337 under the direction of Dr Alan H. Rosenstein. The authors are particularly grateful to R. E. Lewis and D. D. Crooks, Lockheed Missiles and Space Company, for preparing the materials. The use of the facilities of Northwestern University's Materials Research Center sponsored under the NSF-MRL Grant No. DMR9520280 is greatly appreciated. The authors would like to thank Professor L, D. Marks and Dr J. P. Zhang of Northwestern University for assistance with the high resolution electron microscope.
REFERENCES
1, L. Angers, Y. Chen, M. E. Fine, J. R. Weertman and M. S. Zedalis, Rational design of high temperature aluminum alloys, in Aluminum Alloys--Physical and Mechanical Properties (edited by E. A. Starke Jr and T. H. Sanders Jr), Vol. I, pp. 321-337. EMAS, U.K. (1986). 2. C. M. Adam, Structure/property relationships and applications of rapidly solidified aluminum alloys, in Rapidly Solidified Amorphous and Crystalline Alloy (edited by B. H. Kear, B. C. Giessen and M. Cohen), pp. 411-422. Elsevier, Amsterdam (1982). 3. N. Ryum, Acta metall. 17, 269 (1969). 4. O. Izumi and D. Oelschlagel, Scripta metall. 3, 619 (1969). 5. E. Ness, Acta metall. 20, 499 (1972). 6. K. S. Vecchio and D. B. Williams, Acta metalL 35, 2959 (1987). 7. M. S. Zedalis, Ph.D. thesis, Northwestern University, Evanston, I11. (1985). 8. Y. C. Chen, M. E. Fine, J. R. Weertman and R. E. Lewis, Scripta metall. 21, 1003 (1987).
9. R. E. Lewis, D. D. Crooks, Y. C. Chen, M. E. Fine and J. R. Weertman, High temperature A1-Zr-V alloys using rapid solidification processing, in Proc. o f 3rd Int. Conf. on Creep and Fracture of Engineering Materials and Structures (edited by B. Wilshire and R. W. Evans), pp. 331-346. Inst. of Metals, London (1987). 10. F. H. Cocks, M. L. Shepard and H. G. Chilton, J. Mater, Sci. 12, 494 (1977). I1. Y. C. Chen, Ph.D. thesis, Northwestern University, Evanston, I11. (1988). 12. E. Nes and H. Billdal, Acta metall. 25, 1039 (1977). 13, Y. C, Chen, M. E. Fine and J. R. Weertman, Grain boundary precipitate free zone in melt spun A1-Zr-V alloy, M R S Int. Symp. Advanced Materials, Tokyo, Japan (1988). 14. E, D. Zakharov and G. L. Shneyder, Russian Metall., pp. 141-145 (1978). 15. N. I. Varich, L. M. Burov, K. Y. Kolesnichenko and A. P. Maksimenko, Fizika Metalloved. 15, 292 (1963). 16. L. M. Burov and A. A, Yakunin, Russian J. Phys. Chem. 42, 540 (1968). 17. P. W. Voorhees and M. E. Glicksman, Metall. Trans. 15A, 1081 (1984). 18. P. K. Footner and B. P. Richards, J. Mater. Sci. 17, 214l (1982). 19. A. J. Ardell and R. B. Nicholson, J. Phys. Chem. Solids 27, 1793 (1966). 20. D. B. Williams and J. W. Edington, J. Metal Sci. 9, 529 (1975). 21. T. Marumo, S. Fujikawa and K. Hirano, J. Japan Inst. Light Metals 23, 17 (1973). 22. W. L. Fink and L. A. Willey, Trans. Am. Inst. Min. Engrs 133, 69 (1939). 23. M. F. Ashby, C. Gandhi and D. M. R. Taplin, Acta metall. 27, 699 (1979). 24. B. Reppich, H. Buglar, R. Leistner and M. Schutze, Application of the microstructural concept of creep and rupture lifetime to a precipitating Ni-base alloy--I. Yielding and Creep Behavior. Proc. 2nd Int. Conf. on Creep and Fracture of Engineering Materials and Structures, pp. 279-297, Swansea, England (1984). 25. P. J. Henderson and M. McLean, Acta metall. 31, 1203 (1983). 26. R. W. Lund and W. D. Nix, Acta metall. 24, 469 (1976). 27. M. McLean, Acta metall. 33, 545 (1985). 28. A. K. Mukherjee, J. E. Bird and J. E. Dorn, Am. Soc. Metall. Trans. 62, 155 (1969).
0956-7151/90 $3.00 + 0.00 Copyright @~1990 Pergamon Press plc
MICROSTRUCTURAL EVOLUTION A N D MECHANICAL PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V ALLOYS AT HIGH TEMPERATURES Y. C. CHEN~', M. E. FINE and J. R. WEERTMAN Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208, U.S.A.
(Received 6 October 1988; in revised form 27 September 1989) Abstract--Aging studies of three melt spun, A I - Z ~ V alloys showed that precipitation of the metastable LI 2 phase occurs prior to that of the equilibrium phases, AI~0V and A13Zr(D023). The precipitation mechanism of the LI2 phase can be either continuous or discontinuous. The former is favored by decreasing the value of the Zr/V ratio in the L12 phase precipitate or by increasing the aging temperature. Decreasing the Zr/V ratio also makes it easier to obtain a supersaturated solid solution during melt spinning. Coarsening of the spherical L l:-structured m13(Zr0.25V0.75) particles was studied at 425,450 and 500°C. The activation energy calculated from the particle coarsening kinetics is close to that for diffusivity of Zr in the AI matrix, suggesting that volume diffusion of Zr is the rate controlling mechanism for the Ostwald ripening of the L12-structured A13(Zr025V07~). The LI 2 phase precipitate in extruded samples of this alloy grew slowly at 425°C, close to the rate measured in the melt spun ribbon. Dislocation cross-slip occurred inside the LI2 phase particles from a (l l l) plane to a {100} plane during creep. R6sum~---Des 6tudes sur le vieillissement de trois alliages A I ~ r - V 61abor6s par trempe rapide sur roue, montrent que la pr6cipitation de la phase m6tastable LI 2 se produit avant celle des phases d'6quilibre, AI~0V et AI3Zr(D023 ). Le m6canisme de pr6cipitation de la phase L12 peut 6tre continu ou discontinu. Le premier est favoris6 par la diminution du rapport Zr/V dans la phase pr6cipit6 LI 2 ou par l'augmentation de la temp6rature de vieillissement. La diminution du rapport Zr/V rend aussi plus facile l'obtention d'une solution solide sursatur6e pendant la trempe rapide. On &udie le grossissement des particules sph6riques de Al~(Zr0,2~V0.7~) de structure LI 2 ~i 425,450 et 500°C. L'6nergie d'activation calcul~e fi partir des cin6tiques de grossissement des particules est voisine de celle de la diffusion du Zr dans la matrice d'A1, ce qui laisse supposer que la diffusion en volume du Zr est le m6canisme qui contr61e la vitesse pour le mfirissement d'Ostwald du AI 3(Zr0.25Vo,75) de structure L 12. Le pr6cipit6 de phase L 1~dans les 6chantillons extrud6s de cet alliage croit lentement fi 425' C, ~, une vitesse voisine de celle qu'on mesure sur le ruban tremp6. Le glissement d6vi6 des dislocations se produit a l'int6rieur des particules de phase L12 d'un plan (111) vers un plan {100} pendant le fluage. Zusammenfassung--Auslagerungsexperimente an drei schmelzgesponnenen Al~r-V-Legierungen zeigten, dab sich die metastabile L12-Phase vor den Gleichgewichtsphasen AI~0V und AI3Zr(D023) ausscheidet. Der Ausscheidungsmechanismus f/Jr die Ll2-Phase kann entweder kontinuierlich oder diskontinuierlich sein. Der kontinuierliche Mechanismus wird begfinstigt, wenn das Zr/V-Verh/iltnis in der LI,Ausscheidungsphase verringert oder die Auslagerungstemperatur erh6ht wird. Verringerung des Zr/VVerh/iltnisses erleichtert es auch, w~ihrend des Schmelzspinnens eine iiberstfittigte feste L6sung zu erhalten. Die Vergr6berung der kugelf6rmigen A13(Zr0.25V0.ys)-Teilchen (L12-Struktur) wurde bei 425, 450 und 500"C untersucht. Die aus der Vergr6berungskinetik der Teilchen berechnete Aktivierungsenergie liegt nahe der ffir die Diffusivitfit von Zr in der A1-Matrix; dieser Befund legt nahe, dab hier die Volumdiffusion des Zr der ratenvestimmende Mechanismus der Ostwaldreifung dieser Teilchen ist. Die L12-Ausscheidungsphase in stranggeprel3ten Proben dieser Legierung wuchsen bei 425:C langsam mit einer Rate nahe der, die in dem schmelzgesponnenen Band gemessen wurde. W~ihrend des Kriechens trat Quergleitung xon Versetzung innerhalb der Ll2-Phasenteilchen von einer (111)- in eine {100}-Ebene auf.
INTRODUCTION
Co o f its rate-controlling element is small. A m o n g the transition metals, Z r has the smallest diffusion flux DCo in AI [2]. The A13Zr phase in binary AIZr alloys can form either in a semicoherent tetragonal equilibrium phase or a coherent metastable LI 2 form [3-6]. Both forms are expected to form a low energy interface with the matrix, with that of the latter being especially low. Partial substitution of V for Z r in the precipitate increases the thermal stability of the metastable L12 phase in very dilute alloys a n d also slows d o w n the particle coarsening rate at 425°C [7].
T o improve the strength of A1 alloys at high temperatures a microstructure c o n t a i n i n g a thermally stable a n d coarsening resistant dispersoid is required [1, 2]. According to theory, a particle will be resistant to coarsening if the energy o f its interface with the matrix is low, a n d if the diffusivity D a n d solubility tPresent address: Materials Research Laboratories, 195-5 Chung-Hsing Road. Chutung, Hsinchu, Taiwan 31015, R.O.China. 771
772
CHEN et al.:
PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES
In the present study, three more concentrated A I - Z r - V alloys with different Zr/V ratios were prepared by melt spinning. These alloys were used to study the influence of the Zr/V ratio on the melt spun structure and the thermal stability of the metastable L12 phase with respect to phase change and Ostwald ripening. Coarsening and mechanical properties of extrusions made from the melt spun ribbons were evaluated at 425°C. A very slow growth rate of LI2structured precipitate at 425°C was reported in melt spun ribbons of A1-5 vol.% A13(Zr0.25V0.75) [8]. The temperatures have now been extended to 500°C. EXPERIMENTAL
Three A I - Z r - V alloys were prepared by melt spinning. The compositions of the alloys were chosen with the goal of producing material with 5 vol.% Ll2-structured Al3ZrxV~_x. The solid solubilities of both Zr and V in AI are 0.1wt% at 425°C. Assuming the solubilities to be the same in the ternary A1-Zr-V system, alloys were designed with values of x of 0.75, 0.5 and 0.25, designated alloys 2, 3 and 4, respectively. The nominal compositions are given in Table 1. Extruded bars having dimensions 6.4 × 14mm were made from the ribbons through vacuum hot degassing, hot pressing and hot extrusion. The processing details are described elsewhere [9]. The specimens for optical microscopy were prepared by electropolishing and etching in Keller's reagent. TEM foils were prepared in a double jet polishing apparatus with a solution of 25 vol. % nitric acid in methanol. The temperature of the electrolyte was kept between - 4 0 to - 6 0 ° C and a current of 100mA was applied. The melt spun ribbons were mechanically polished with 600-grit abrasive paper before the electropolishing. Some specimens were selectively polished on one surface with the other protected by lacquer to examine the near surface microstructure. After electropolishing the lacquer was dissolved in acetone. Sometimes ion milling was applied at 2 mA and 4 kV for half an hour at a grazing angle of 15° to remove surface contamination. Extracted second phase particles were obtained by electrolytic dissolution of the A1 matrix using an electrolyte which contains 60 g benzoic acid, 15g 8-hydroxyquinoline, 60ml chloroform and 160 ml methanol [10]. The sizes of the L12 precipitates were measured from TEM negatives taken by the central dark field (DF) technique using superlattice reflection spots. More than 400 particles were measured for each Table 1. Nominal compositions of three AI-Zr V alloys Alloy
Composition (wt%)
2 3 4
AI-3.12 Zr4).66 V AI-2.14 Zr-1.24 V Al-l.13 Zr 1.82V
Fig. 1. Typical TEM microstructures of (a) alloy 2, (b) alloy 3 and (c) alloy 4 ribbons in the as-melt spun condition. determination of average particle radius. The precipitate free zone (PFZ) widths were measured with the beam direction lying in the grain boundary plane. Creep tests were conducted at 4 2 5 _ 2°C on a dead load creep machine. Strain was monitored with a calibrated linear variable differential transducer. Tensile tests were carried out on an MTS machine at 425°C. Both kinds of mechanical tests were done in an Ar atmosphere. The specimens had gage dimensions of 20 by 6,6 by 0.9 mm (S1) or 20 by 4 by 2 mm ($2).
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED A1-Z~V AT HIGH TEMPERATURES RESULTS
Microstructures and phase transformations The melt spun ribbons were approximately 20-60/tm thick and 2 3 mm wide. Occasionally small globules of previously molten metal were stuck to them. The specimens investigated in this study were chosen to avoid such regions. The microstructure of alloy 4 varied slightly from place to place, depending upon the contact of the ribbon with the chilled Cu wheel. Columnar grams are observed where the chill surface is flat. Howcver, in some areas where the chill surface is unexen, a dark-etching region is present near the free surface of the ribbon. A lower cooling rate is associated with such an uneven contact surface. TEM selccted area diffraction (SAD) showed that the second phase particles present near the free surface of alloy 4 have the structure of All0V [11]. No such particles were found in the as-received ribbons of alloys 2 or 3. Typical microstructures from the areas close to the middle of ribbons of alloys 2, 3 and 4 are shown in Fig. 1. In alloys 2 and 3, which have higher Zr/V ratios, cellular fan shaped precipitates
773
form inside the grains [Fig. l(a,b)]. The SAD patterns from these cellular precipitates are consistent with the LI 2 structure, such as reported for metastable AI3Zr in AIZr alloys [4, 5]. The grain boundaries are rather irregular and island grains surrounded by other grains are often observed in these two alloys. In alloy 4, no precipitates were noted inside the grains and the grain boundaries are more regular, as seen in Fig. l(c). Some small second phase particles, approximately 0.1/~m in size, are present along the grain boundaries. Selected area diffraction from the larger of these particles proved that they have the Al~0V structure. This precipitate is likely the source of the etch pits seen at grain boundaries. In alloys 2 and 3 spherical precipitates were interspersed with the cellular precipitates. In TEM dark field micrographs, Fig. 2(a) and (b), this phase is clearly distinguishable. A TEM diffraction pattern from this area and a schematic drawing of it are shown in Fig. 2(c) and (d). The electron beam direction is close to the ( 100 ) of the AI matrix. Since all spherical phase particles in one grain always light up at the same time with one operating reflection, they have the same crystallographic orientation.
•
•
o
•
•
•
•
o
o
110Lie o
•
0 000
o o
•
o
• 111
(d)
•
Fig. 2. TEM dark field micrographs from as-spun alloy 2 ribbon and its associated SAD patterns where (a) was taken with [110] spot of the LI 2 phase, (b) was taken with [111] spot of the unidentified phase, (c) is SAD and (d) is a schematic drawing of SAD. The open circle represents reflections from the cellular LI, phase and the AI matrix and the closed circle represents reflections from the unidentified phase.
774
CHEN
et al.:
PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
However, as seen in Fig. 2(d), a cube/cube orientation relationship with the A1 matrix does not exist for these particles. Figure 3 illustrates two other selected area electron diffraction patterns from these particles. The electron diffraction patterns shown in Fig. 2(c) and Fig. 3(a) and (b) can be indexed as (112), (100) and (110), respectively, based on a simple cubic
structure with the lattice constant equal to 4.06/~. As seen in Fig. 3(a), where the specimen was tilted to near the [100] zone axis, the intensities of {002} spots are stronger than those for {001} spots. These may be Ll2-structured particles but they do not have the normal cube-cube orientation with the matrix. On aging at 500°C for 1 h, All0V and D023structured A13Zr particles formed on the grain boundaries [11]. After aging alloy 4 at 350°C for 3 h, some cellular L12 phase had formed from grain boundaries, as shown in Fig. 4. The cellular LI 2 phase
Fig. 5. High resolution electron micrograph of an AI~0V particle.
Fig. 3. Selected area electron diffraction patterns from the unidentified phase where (a) can be indexed as (100) and (b) as (110).
Fig. 4. TEM of alloy 4 ribbon aged at 350°C for 3 h.
Fig. 6. High resolution electron micrograph of an L12 particle in A1 matrix from alloy 4 ribbon aged at 600°C for 1.5h.
CHEN et aL:
PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES
has been suggested by Nes and Billdal [12] to form by discontinuous precipitation resulting from grain boundary movement. The presence of the original traces of the grain boundaries, indicated by arrows in Fig. 4, left behind by the cellular decomposition confirms this. In other areas SAD shows the matrix in alloy 4 to be still free of any precipitate. During further aging to 30 h at the same temperature [11], small spherical L12 phase particles form by continuous
14.01 j
6.c~-
~lt-
4.0
2 O~
precipitation inside all the previously undecomposed grains. The fraction of the cellular LI2 precipitate decreases with increase of the aging temperature [7, 11]. On aging at 600°C, cellular precipitates are not found at all. During aging of alloy 4 ribbon, All0V and some D023-structured AI3Zr, precipitate on the grain boundaries leading to the formation of LI 2 precipitate free zones (PFZs). Both grain boundary
0 T : 425* C
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.~ f /
.
0 T
~ •
~ i
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~ I
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Fig. 7. Growth kinetics of LI 2 A%(Zr0.25V075)at various temperatures in alloy 4 ribbon. A pre-aging of 2.5 h at 500°C was done for the specimens aged at 425 and 450°C.
Fig. 8. Llz-structured A13(Zr025V07s) particles in alloy 4 ribbon after aging at 500°C for (a) 20 and (b) 206 h.
Fig. 9. Typical microstructures (TEM) of (a) alloy 2, (b) alloy 3,.and (c) alloy 4 after extrusion. Extrusion direction is normal to the micrograph in each case.
776
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
precipitate size and P F Z width increase with aging time [8, 13]. A high resolution TEM (HRTEM) micrograph of AI~0V, Fig. 5, reveals microtwins and a stacking fault. The existence of twins is also indicated by the twin spots shown in the electron diffraction pattern (inset in Fig. 5). Using the energy dispersive X-ray spectrometry (EDS) attachment to the STEM, the average composition of extracted "Al~0V'' particles analyzed to approximately 11 at.% V and 1 at.% Zr. A standardless analysis technique and thin film approximation were used. For simplicity, the nominal designation AI,0V is used for this ternary grain boundary precipitate throughout this paper. An electron diffraction pattern and the HRTEM micrograph of L12-A13(Zr, V) after aging at 600°C for 1.5 h are shown in Fig. 6 where each white spot corresponds to the image from the projection of AI atoms stacked along the [110] direction. The presence of order in the LI2 structure is clearly shown. A perfect lattice match between the LI2 particle and the AI matrix and a stepped interface are also seen in this figure. The growth kinetics of the spherical L12 particles in alloy 4 ribbons were determined at 450 and 500°C. Detailed coarsening results at 425°C have been reported previously [8]. Figure 7 shows the coarsening kinetics of the L12 particles at these three temperatures, where the cube of the average particle radius, f, is plotted against aging time. The linear correlation coefficient, R, and the slopes for the best fitting lines are given in Table 2. Typical L12 particles after aging at 500°C for 20 and 206 h are shown in Fig. 8(a) and (b), respectively. A planar fault exists inside many of the LI2 phase particles seen in Fig. 8(b). These planar faults are always parallel to the {100} planes of the LI2 phase. TEM study indicated them to be antiphase boundaries (APBs) which are associated with a displacement of a / 2 ( l 1 0 ) on a {100} plane, where 20.0
Fellure
T : 425"C
18.0
C~t : 17 MPcl
16.0 14.0 --
I loy 2
12.0 c 10.0
Fig. 11. Dark field TEM micrograph from alloy 4 after creep at 425°C. Shearing of L12 phase particles has resulted in formation of faults.
a is the lattice constant of the LI z phase [11]. In an L12 precipitate containing an APB, preferential growth along the APB is noted. This causes some L12 precipitates to appear as rounded parallelepipeds rather than rounded cuboids. Microstructures and mechanical properties o f extruded alloys The microstructures of the three alloys after consolidation and extrusion look similar to one another in the optical microscope. Coarse second phase particles of 1 #m in size are observed in some areas. Presumably, these regions came from the previously mentioned globules of molten metal which splashed onto already solidified ribbons. More of these coarse second phase particles were found in extruded alloy 2 than in the other two alloys. TEM observation reveals that the grain sizes, or subgrain sizes, in all three alloys were reduced to around 0.5 #m after extrusion, as shown in Fig. 9. In extruded alloy 4, all the L12 phase particles are spherical. No cellular precipitation is observed. While the size of the L12 particles is quite small, the number of particles in the extruded material appears to be somewhat less than in aged melt spun ribbon. In some grains, stable tetragonal AI3Zr particles are also found. The growth of the spherical L12 particles at 425°C in extruded alloy is nearly the same as that in ribbons, 1.1 × 10-28m3/h compared to 1.0 x 10 -28 m3/h.
8.0
6.0
FOil ure /
4.0
Ter mi notecl J
2.0 o lO "1
~ I 1 creep
l lO time
o
y
4 lO z
lO 3
(hrs.)
Fig. 10. Creep curves of three extruded AI-Zr-V alloys tested at 425°C with an initial stress of 17 MPa.
Table 2. Measured volumetric coarsening rate constants K and coefficientsof linearity R for LI2AI3(Zr0.25Vo75) precipitatesin alloy 4 aged at 425, 450 and 500°C Temperature K (°C) (m3/h) R 425a 1.03 x 10-2s 0.997 450m 6.13 x l0 28 0.997 500 1.28 x 10-26 0.996 ~Specimenswere preagedat 500°C for 2.5 h to prevent cellularprecipitation.
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED A1-Zr-V AT HIGH TEMPERATURES Tensile properties at 425°C for extruded alloys 2 and 4 are given in Table 3. Different strain rates as indicated were employed to acquire yield stress and tensile strength in accordance with A S T M E 151-64. Each data set was obtained from the average of two tests. At 425°C alloy 2 has a lower yield strength than alloy 4, 13.8 MPa compared to 42.7 MPa. The elongations obtained at 425'~C for both alloys are less than those for room temperature which were around 30% [91. Creep tests were done at 425°C with an initial stress of 17 MPa. Typical creep curves for the three extruded alloys, using the SI specimen geometry, are shown in Fig. 10. As seen. alloy 2 creeps much faster than alloys 3 and 4 correlating with the lower yield stress. The fracture surfaces are also different for alloy 2. A dimple structure and a large reduction in area are seen indicating a ductile fracture. In alloys 3 and 4 the failure began in an intergranular fracture mode but changed to the ductile dimple mode with a large reduction in area. In extruded alloy 4, spherical L12 phase particles were sheared along {100} planes during creep as seen in Fig. 1 1. Since the slip planes in the AI matrix are { 111 }, it implies that dislocations cross-slipped from {111} planes to {100~ planes inside the L12 particles during creep. Occasionally, a pile-up of dislocations is observed within a PFZ, suggesting that localization of plastic deformation may occur in this soft region. The steady state creep rate, ~s, at 425°C was determined as a function of creep stress for extruded alloy 4 using the thicker $2 specimen. A partial correction tbr the change in stress caused by the decrease in cross-sectional area during the creep test was made using cr = cri/(l -- E)
(1)
where ~r~ is the initial stress, a is the corrected stress and E is the strain at the beginning of steady state creep. The total change in area during the test was only 1 or 2°'0. The normalized steady state creep rate vs normalized stress (7i is plotted in Fig. 12. The normalized creep rate is defined as ~skT/DGb, where Table 3. Tensileproperties of extruded alloys 2 and 4 tested at 425°C in Ar Yield stress~ Ultimatetensileb (0.2'~0 offset) strength Elongation Alloy (MPa) (MPa) (%) 2 13.8 22.0 16.5 4 42.7 66.8 21.8 ~Measured with ~ - 1 x 10 4/s. bMeasured with ( = 1 x 10 */s.
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I
I 10.3
2 X 1 ~ :)
Fig. 12. Log-log plot of normalized creep rate ~skT/DGb vs normalized creep stress tUG for alloy 4. D is the lattice diffusion coefficient of A1 and b is the Burgers vector, and the creep stress is normalized by dividing it by the shear modulus G. The stress exponent n ( ~ s ~ a " ) is found to be close to 14; however, the normal value for A1 may be obtained by introducing a threshold stress, as discussed later. DISCUSSION
Microstructure of the rapidly solidified A I - Z r - V alloys The as-melt spun microstructures show precipitation has occurred in alloys 2 and 3 but not in alloy 4. Dissolved zirconium increases the lattice constant of the AI matrix. However, vanadium behaves oppositely. The change in lattice constant of the A1 solid solution on dissolution of 1 wt% of Zr or V are +0.0005 nm and -0.00033 nm, respectively [14-16]. Based on these values and Vegard's law, the changes of the lattice constant, Aa, for the three alloys in the supersaturated solid solution state were calculated and are listed in Table 4. Alloys 2 and 3 have larger lattice expansions over unalloyed AI than alloy 4. For this reason, supersaturated alloys 2 and 3 are expected to be more unstable to decomposition by precipitation than supersaturated alloy 4. The lower liquidus temperature of alloy 4 due to the lower Zr content also plays an important role in suppressing precipitation. The metastable Ll2-structured A13(Zr, V) phase may form in cellular or spherical morphology. It is
Table 4. Calculatedchange of lattice constant. Aa, of super-saturated solid solutions of AI-Zr-V alloys from that of unalloyed AI Aa (A)a Aa (,~)b Aa (A) Alloy (from Zr) (from V) (total) 2. AI-3.50wt% Zr4).54 wt% V +0.0175 -0.0016 +0.0159 3. AI-2.15wt% Z~1.22 wt% V +0.0108 - - 0 . 0 0 3 7 +0.0071 4. AI 1.16wt% Zr-l.76 wt% V +0.0058 -0.0053 +0.0005 ~Aa/wt% Zr = + 0.005 ,~/wt%. ~Aa/wt% V = -0.0033 ,~/wt%.
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CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
expected that discontinuous precipitation is kinetically faster than continuous precipitation at low aging temperatures since the diffusion path is shorter. Consequently, L1 z phase in cellular precipitation is dominant in alloy 4 ribbon when aged at 350°C. With increased aging temperature the matrix diffusion rate increases leading to the result that the spherical precipitation mode is dominant in alloy 4 ribbon after aging at 500°C or higher. As a result, the cellular L12 phase may be suppressed by an initial treatment at a higher aging temperature as found by Zedalis in a much more dilute A I - Z r - V alloy [7], Alloy composition may also play a role in the precipitation kinetics of the L12 phase. In alloys 2 and 3, which have lower V/Zr ratios, the cellular L12 phase forms on cooling after solidification. As shown by Zedalis [7], the L 12-structured AI 3(Zrx V 1 _ x ) phase with the highest V content and smallest x has the best lattice matching with the AI matrix. Thus, the effect of the V/Zr ratio on the continuous precipitation of the L12 phase may be explained by a lower energy barrier to homogeneous nucleation. In addition to the cellular L12 phase, an unidentified simple cubic precipitate is also present in alloys 2 and 3. It does not show the normal cube/cube orientation relationship with the AI matrix observed with the L12 phase. This has not been reported previously in binary AI-Zr alloys. Due to the similarity of lattice structures between the L12-structured A13(Zr, V) and this unidentified phase, it is probably also L12-structured A13(Zr, V ). If this is true, the different orientation relationship with the AI matrix might occur by sequential precipitation, as suggested by the reviewer. After solidification the LI 2 phase particles formed firstly by continuous precipitation. Later on discontinuous precipitation took place associated with the movement of the grain boundaries, which resulted in formation of the cellular L12 phase leaving the initially precipitated LI2 phase particles embedded in the cellular regions. The original orientation relation with the matrix changed as the grain boundary swept over them. However, most of the unidentified phase particles formed in a linear array instead of being statistically homogeneous, which seems at variance with this explanation.
Coarsening kinetics of the spherical Lle phase in the melt spun ribbon of alloy 4 The growth behavior of the L12 particles at 425, 450 and 500°C shows good agreement with the ~3 vs t relationship, f 3 _ 73° = Kt, which suggests that an LSW-type volume diffusion controlling process I2] is the coarsening mechanism. Here f is the average particle radius after aging for time t and ~o is an integration constant. If the effect of the precipitate volume fraction on Ostwald ripening is taken into account, the coarsening rate constant K for an alloy containing 5 vol.% of precipitates is given by [17] K =
12yDV2 CO 9RT
(2)
-SB
-6C
- Q / R : -35.38 x 103 Q=294 kJ//mole I
-62
{:
-64
-66
-6e
i
1.20
I
J
I
1.30
i
1.40
1.50
~/T(x~O-~ K-~)
Fig. 13. Plot of ln(KT) vs lIT for the coarsening of the L12-structured A13(Zro.25V075). where y is the interfacial energy, Vm is the molar volume of the precipitate, and R is the gas constant. This equation has been used by others to determine the activation energy of diffusivity of solute in the matrix from the experimental coarsening results [18-20]. After rearrangement, equation (2) becomes
ln(KT/Co) = In A - Q/RT
(3)
where A is given by A = 12~DoVL
9R
(4)
and Q is the activation energy of diffusivity for the rate-controlling element in the matrix. The value of A is a slowly varying function of temperature and usually may be regarded as a constant over a small temperature range. Therefore, the value of Q can be experimentally determined from a plot of ln(KT/Co) vs the reciprocal of temperature. In the present alloy, neither the concentration C0 of V or of Zr in the solid solution AI matrix in equilibrium with the metastable LIz phase is accurately known. For a first order approximation, Co is assumed to be constant in the temperature range from 425 to 500°C. From the plot of ln(KT) vs I/T shown in Fig. 13, Q is calculated to be 294 kJ/mol. This value is closer to the activation energy for diffusion of Zr in AI, 242 kJ/mol, than to the value for V, 82kJ/mol [3,21]. Consequently, the volume diffusion of Zr is likely the rate controlling process for the Ostwald ripening of the Llz-structured A13(Zr0.25V0.75), as expected originally. In binary A1-Zr alloys, the solvus line of the stable phase A%Zr(D023), Co, can be expressed by the following equation, in wt% [22]
Co = 1180 exp(-64.96 kJ/RT).
(5)
If it is assumed that the Zr concentration of the AI matrix in equilibrium with the metastable L12 phase,
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES C~, does not change much from Co, then a better calculation for Q can be obtained by considering the temperature dependence of C'o from Equation (5). The final result for Q is 229 kJ/mol. This value agrees even better with the activation energy of diffusivity for Zr in A1 measured in a binary AI-Zr alloy [21]. It was also found that the growth kinetics of the PFZ at 425 ° in alloy 4 ribbon is controlled by the volume diffusion of Zr in the AI matrix [13].
779
18a Id 2
16E T
o
o
o
1(~7 --
I(59 -
Mechanical properties elevated temperature
of
extruded
materials
at
i~ 0 -1611
The poor creep resistance of extruded alloy 2 (Fig. 10) is related to its low yield stress, 14 MPa at 425°C, which is lower than the creep stress, 17 MPa. The plastic tearing seen on the fracture surface after failure by creep confirms this. It is of interest to know why this alloy possesses a low strength compared to the other two extruded AI-Zr-V alloys since they all contain around 5 vol.% of precipitates. TEM investigation disclosed that in alloy 2 somewhat larger amounts of the equilibrium phases compared to the other two extruded alloys were present after annealing of 1 h at 425°C. This difference was not seen in the melt spun ribbons of the three alloys. Thus it indicates that after hot extrusion the thermal stability of the L12-structured A13(Zr075V,.25), with a high Zr/V ratio is lower. Another factor contributing to the weakness of alloy 2 at 425°C is its inhomogeneity. Regions with coarse second phase particles were often observed in alloy 2. Intergranular fracture in creep is expected to be more prominent in particle hardened alloys. This is confirmed by the present results for alloys 3 and 4. Localization of plastic deformation along the PFZs, which are softer than the precipitation hardened matrix, and sliding of the grain boundaries or subgrain boundaries are assumed to be the mechanisms for the intergranular fracture in these two alloys. As the crack deepens the creep stress increases because of the reduction in cross-section, becoming 34 MPa when the cross-section is one-half of its original value which is comparable to the 0.2% offset yield stress. 43 MPa, of alloy 4 at 425°C. Accordingly, a transition is expected from an intergranular to a transgranular fracture exhibiting a dimple structure as observed. Creep theories for particle strengthened alloys [24-27] suggest that processes occurring in the matrix are rate controlling. Thus, the creep rate equation for particle strengthened alloys is expected to be in the same form as that for the matrix material substituting an effective creep stress 0"ef t equal to a --ao for the actual stress a, where a o is the threshold stress. The creep rate for extruded alloy 4 may be expected to fit the equation DGB a ~
gs=AD k T
-
-
a o 4.4
(6)
1612
16~3 1(]TM 10
I 12
I
I
I
14
16
18
I
I
I
20 22 24 O" ( M P a )
I
I
26
28
I
30 32
Fig. 14. Comparison of creep strain rate vs stress for alloy 4 (data points) with prediction from equation (6) with a o equal to 14.2 MPa (curve). where the exponent 4.4 is that obtained for AI in the power law regime and AD is an empirical constant, 3.4 x 106 for AI [28]. By linear regression, the value of ao is calculated to be 14.2 MPa. A plot of creep rate vs creep stress for the data obtained from alloy 4 is compared with the curves predicted by equation (6) in Fig. 14. The actual strain rates are off by one order of magnitude from the predicted curve but the curves are essentially parallel. A higher creep rate compared with that from equation (6) might be caused by localization of deformation within the PFZs. SUMMARY AND CONCLUSIONS 1. In rapidly solidified AI-5 vol.% A13(ZrxV~ x) alloys, the metastable L12 phase precipitates initially instead of the equilibrium phases, Alj0V and A13Zr(D023). The precipitation mechanism of the L12 phase depends on alloy composition and aging temperature. In the alloys with a Zr/V ratio equal to 3 or 1, a cellular L12 phase forms by discontinuous precipitation immediately after solidification. However, a supersaturated solid solution is obtained in an alloy with a Zr/V ratio equal to 1/3. In the latter, the cellular mode is dominant when aging is carried out at low temperatures, but spherical LI 2 particles form when aging is done at higher temperatures. 2. A good fit linear relationship between the cube of the average particle radius and aging time is obtained for the L12-structured A13(Zr0.25V0.7s) precipitate aged at 425, 450 or 500°C. Volume diffusion of Zr in the AI matrix appears to be the rate controlling mechanism for coarsening. The coarsening rate of the L12 phase precipitate at 425°C is 1.1 x 10-28m3/h. 3. In extruded samples of the AI 5v01.% A13(Zro.25Vo.75), alloy 4, L12 phase precipitates were
780
CHEN et al.: PROPERTIES OF RAPIDLY SOLIDIFIED AI-Zr-V AT HIGH TEMPERATURES
sheared by dislocations on { 100} planes during creep, suggesting that cross slip of dislocations from a { 111 } plane to a {100} plane occurred inside the LI 2 phase precipitate. With the introduction of a threshold stress of 14.2 MPa, the data may be fit to a creep equation with the stress exponent obtained for pure AI, 4.4. However, the experimental creep rates are faster than those predicted. Acknowledgements--This research was supported by the Air Force Office of Scientific Research Grant No. AFOSR85-0337 under the direction of Dr Alan H. Rosenstein. The authors are particularly grateful to R. E. Lewis and D. D. Crooks, Lockheed Missiles and Space Company, for preparing the materials. The use of the facilities of Northwestern University's Materials Research Center sponsored under the NSF-MRL Grant No. DMR9520280 is greatly appreciated. The authors would like to thank Professor L, D. Marks and Dr J. P. Zhang of Northwestern University for assistance with the high resolution electron microscope.
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