- Email: [email protected]
Ordering, deformation and microstructure in L10 type FePt
PII:
Acta mater. Vol. 46, No. 18, pp. 6485±6495, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00311-5 1359-6454/98 $19.00 + 0.00
ORDERING, DEFORMATION AND MICROSTRUCTURE IN L10 TYPE FePt S. H. WHANG{, Q. FENG and Y.-Q. GAO Polytechnic University, Six Metrotech Center, Brooklyn, NY 11201, U.S.A. (Received 27 March 1997; accepted 21 August 1998) AbstractÐThe long-range order (LRO) parameters for L10 FePt have been determined by the X-ray diraction method from powder specimens as functions of time and temperature from 773 K to near 1573 K. By and large, the ordering takes place rapidly below Tc and reaches as high as 0.85 even at 773 K in the ®rst 30 min. The LRO value is about 0.81 near Tc (1573 K) before it drops abruptly to zero at 1573 K. As a result, the order±disorder transformation in FePt is concluded to be a ®rst-order phase transformation. Deformation behavior in an L10 type FePt alloy was investigated through both compressive and tensile deformation from room temperature (RT) to 1073 K. The negative temperature dependence of yield stress in this alloy contrasts with the positive dependence in L10 type TiAl. The elongation increases exponentially with temperature and reaches 06% at 873 K. The strain rate sensitivity parameter against temperature is similar to those found in silver and copper, where the non-zero minimum is centered in a broad basin. This indicates that the temperature-dependent deformation in the range of RT to 1073 K is analogous to that of some face-centered cubic metals, but signi®cantly dierent from that of L10 TiAl. The deformation structure investigated by TEM shows that slip and twinning are the two major deformation mechanisms. The identi®ed slip systems include 1=2110f111g; h101f111g and 1=2h112f111g. The 112h112f111g slip system, however, is only active at very low temperatures, e.g. 77 K. The twin system was identi®ed as f111gh112 type. No pseudo-twinning was found in this alloy. The deformation below RT is mainly carried out by both superdislocations and ordinary dislocations, while above 673 K, it is carried out mainly by ordinary dislocations. The morphology of these dislocations in the entire temperature range indicates that the dislocations do not experience a high Peierls stress contrary to that observed in TiAl. No self-dissociation of superdislocations or APB cross-slip onto cube planes was observed under weak beam conditions. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
The majority of L10 type compound alloys are composed of combinations of transition metals, but there are a few exceptions such as TiAl and TiGa. The atomic size dierence between the two component atoms is less than 15% in L10 compounds. This indicates that the size factor is important for the stability of L10 compounds. Second, the majority of these compounds have a c/a ratio less than unity. The lattice parameter contraction occurs along the c-axis which is perpendicular to the superstructure layers. Hence, the atomic distance between unlike atoms is an important factor for the formation of the L10 structure. The exceptions for the c/a ratio include TiAl and TiHg where the c/a ratio is larger than unity. In TiAl, a directional bond between Ti and Ti was found [1] and may be responsible for the unusual c/a ratio. Stoichiometric FePt, as shown in the Fe±Pt phase diagram (Fig. 1) [2, 3], solidi®es into a disordered face-centered cubic (f.c.c.) phase from the melt and transforms into an ordered face-centered tetragonal (f.c.t.) phase (L10) near 1573 K. The characteristics of this transformation include (i) the order±disorder {To whom all correspondence should be addressed.
transformation takes place at a very high temperature, and (ii) the ordering does not require a longrange diusion process. The transformation rapidly takes place in the absence of long-range diusion. In ordered FePt, Fe and Pt atoms are disposed in alternative layers of {200} planes. The L10 FePt phase region in the FePt phase diagram is surrounded by narrow two-phase regions (shown in Fig. 1), which indicates that the formation of the FePt superstructure may be characterized by a ®rstorder phase transition [4]. Cabri et al. found that the ordered phase of FePt was an f.c.t. phase (L10 type) con®rmed by the X-ray powder diraction method [5]. Vlasova and Sapozhkova studied the type and kinetics of ordering in Fe±32 at.% Pt alloy and found that the ordering in FePt took place rapidly as a result of annealing at 1023 K [6]. In order to determine the long-range order (LRO) parameter (S) of FePt, the present study is made with regard to the time and temperature dependence of ordering by X-ray diraction measurements. The deformation behavior of L10 FePt has not been investigated in the past. Previous studies of L10 TiAl show that anomalous hardening occurs at high temperatures [7±9]. Therefore, fundamental questions have been raised whether the anomaly is
6485
6486
WHANG et al.: L10 TYPE FePt
Fig. 1. Binary phase diagram of Fe±Pt.
an intrinsic property of the L10 structure or is associated with a particular bond, or is limited to L10 TiAl. Further, it is important to understand the relationship among deformation behavior, resulting dislocation structure, other planar defect structures, and chemical and physical properties in L10 type compounds. FePt may be considered a model L10 compound representing a group of L10 transition± transition metal compounds. In addition, FePt has a high order±disorder transition temperature of 1573 K as shown Fig. 1 [2], thus allowing an investigation of deformation behavior at high temperatures. Therefore, a study of L10 FePt will answer some of these questions, which include (i) the temperature-dependent deformation behavior of FePt; (ii) ®ne dislocation structures resulting from the deformation and their relationship with work hardening behavior. The deformation behavior of FePt from 77 to 873 K and the corresponding microstructures are reported in this paper. 2. EXPERIMENTAL PROCEDURE
FePt alloy buttons were made in an arc furnace from powders of pure iron (99.99%) and pure platinum (99.9%). The arc melted buttons, which were placed in an argon-pressurized quartz tube, were homogenized at 1643 K for 24 h, and subsequently quenched immediately into ice±water. The powder specimens were prepared by ®ling o the quenched FePt buttons. The powders were found to have a disordered f.c.c. phase (g) identi®ed by X-ray diffraction. The powder specimens were sealed in very
narrow, evacuated quartz tubes and annealed at various temperatures until the ordering had been achieved at a saturated level, followed by quenching into a 10% NaCl ice±water solution. The annealing for ordering was performed at various temperatures from 773 K to near 1573 K; and from 0.5 to 256 h. The X-ray diraction patterns and data were obtained from powder specimens using a Philips XRG Cu ÿ Ka1 X-ray diractometer at 40 kV, 20 mA. For tensile specimens, the arc melted buttons were re-melted by induction power in an argon atmosphere and cast into a ceramic tube. The cast rods were annealed at 1173 K for homogenization. The L10 structure was con®rmed by X-ray diraction and TEM. The alloy was cast in a copper mold followed by annealing at 1173 K for 115 h. The long rods were machined into non-standard tensile specimens with a gauge length of 22 mm and diameter of 3 mm. The machined tensile specimens were polished with ®ne polishing paper, and electropolished. The initial strain rate of 3 10ÿ4/s was used for tensile testing in argon atmosphere. Compression specimens were machined from the rods into cylindrical specimens of 4.3 mm in diameter and 9 mm in length. The compression tests were carried out with strain rates of 3 10ÿ3 and 3 10ÿ4/s at room temperature (RT), 473, 673 and 873 K, respectively. TEM work was performed with a 200 keV TEM at RT. Also, two-beam and weak-beam techniques were performed with g±3g conditions using a 200 keV TEM.
WHANG et al.: L10 TYPE FePt 3. RESULTS
6487
Fundamental: Ff 2 fFe fPt ; h k even, hkl unmixed
3.1. Ordering transformation Typical X-ray patterns obtained from the powder specimens of FePt are shown in Fig. 2. In Fig. 2(a), the pattern from the homogenized specimen at 1643 K exhibits the diraction lines of a disordered g-single phase. In Fig. 2(b), however, the diraction patterns from a specimen annealed at 1473 K for 1 h show many superlattice diraction lines of the ordered g2 phase (L10 type). The LRO parameter, S, for the L10 FePt structure can be determined from the integrated intensity ratio of the superlattice diraction peaks and the fundamental diraction peaks. The structure factor F takes two forms:
Superlattice: Fs 2S fFe ÿ fPt ; h k even, hkl mixed: The integrated intensity, I, per unit length of powder-pattern diraction line is given by I km LP F 2 A yT yE y
1
where k is a constant for all re¯ections; m, the multiplicity of the re¯ecting planes; LP, the Lorenz polarization factor; A(y), the absorption correction factor; T(y), the Debye temperature factor; and E(y), the extinction factor. The LRO parameter S
Fig. 2. X-ray diraction pattern obtained with Cu ÿ Ka1 radiation of Fe±50Pt alloy: (a) the disordered g phase; (b) the ordered g1 phase.
6488
WHANG et al.: L10 TYPE FePt
Fig. 3. Long-range order parameter S for FePt as a function of annealing time at 773 K.
Fig. 5. Plot of strain and stress curves using alternating strain rates in FePt.
may be given by As m LP fFe fPt 2 A yT yE yf S2 : Af m LP fFe fPt 2 A yT yE ys
intensity of the superlattice diraction, Is,T, and the fundamental diraction of the (111) plane, I(111),T at a given temperature, T
2
where As and Af are the peak areas of superlattice diraction and fundamental re¯ections, respectively. When the LRO in the alloy is perfect, S is equal to unity; when the atomic arrangement is completely random, S is equal to zero. In principle, the value of S can be determined by calculating the bracketed term in equation (2). Nevertheless, in practice, the bracketed term in equation (2) cannot be accurately calculated due to (i) lack of understanding of the preferred orientation in the powder specimens; (ii) the diculty in estimating the Honl L-electron correction for the Pt scattering factor; and (iii) the uncertainty in estimating the extinction factor and degree of powder size uniformity. Alternatively, if one uses the X-ray patterns corresponding to nearly perfect ordering as a reference, S values for various ordered states may be determined by comparing the superlattice intensity data of an investigating state with that of the reference state. In other words, S values can be determined simply from the relative ratio of the integrated
Fig. 4. Long-range order parameter S for FePt as a function of temperature.
Is,T =I 111,T S: Is,773 K =I 111,773 K
3
The latter method was adopted in the current experiments. For FePt, the reference patterns were obtained from the specimen annealed at 0.5Tc. The 0.5Tc for FePt is located 0773 K. Four superlattice diraction patterns were selected to calculate the LRO parameter S. They were re¯ections from (001), (110), (021), (112) planes. The value of S at each temperature was optimized by the method of least squares. Also, St,773 K was obtained by annealing the specimens at 773 K for various durations from 0.5 to 256 h. The results shown in Fig. 3 indicate that the S value asymptotically approaches a ®nite value, i.e. S = 1. Therefore, the S value for 773 K, 256 h, where the
Fig. 6. Temperature dependence of compressive yield stress.
WHANG et al.: L10 TYPE FePt
6489
Fig. 7. Strain rate sensitivity as a function of temperature.
Fig. 10. Plot of tensile yield stress as well as UTS vs temperature in FePt.
Fig. 8. Fractured tensile specimen of FePt at RT.
S value approaches a saturated value, was conveniently taken as a reference point. Further, the ratio Is,773 K =I 111,773 K may be de®ned to establish the relationship between S and the fundamental diffraction intensity in this alloy. Various S values at dierent temperatures were determined, which are shown in Fig. 4. The S curve in Fig. 4 shows that the LRO parameter decreases very slowly with increasing temperature to a value of about 0.81 at 1567 K, where it drops sharply to zero at the transition temperature, 1573 K.
was investigated from 77 to 1073 K as shown in Fig. 6. The sharp rise in yield stress at 77 K was noted. In contrast, the decline in yield stress with increasing temperature is moderate from RT to 1073 K. The absence of anomalous hardening in FePt is also characteristic, in contrast to the case of TiAl where anomalous hardening occurs from RT to 1073 K [9]. The strain rate sensitivity parameter, z 1=T @ ln s=@ ln e_ from the compression test adopts a ``U'' shape centered near 600 K (Fig. 7). The minimum value of S occurs between 400 and 700 K, followed by a sharp increase above 900 K. Despite a lower value of z, the z value never approaches zero in contrast to the zero minimum near 600 K in L10 TiAl.
3.2. Compression test The compression tests were performed at the initial strain rate of 3 10ÿ4/s (``a'' in Fig. 5) which was subsequently changed to an alternate rate, 3 10ÿ3/s (``b'' in Fig. 5), and thereafter the strain rate oscillated between the two strain rates. The temperature dependence of yield stress (0.2% oset)
Fig. 9. Stress±strain curves under compression for RT to 1073 K.
Fig. 11. Elongation vs temperature in FePt.
6490
WHANG et al.: L10 TYPE FePt
3.3. Tensile test The non-standard tensile specimens of an FePt alloy tested in an argon atmosphere from RT to 1073 K did not show any necking resulting from RT testing (Fig. 8). The stress±strain curves at various temperatures are shown in Fig. 9. The stress± strain curves do not show a sharp yield point which is normally observed in carbon steels. Nevertheless, the trend in temperature dependence is consistent with that observed in the compression test (Fig. 5). The tensile yield stress from RT to 1073 K (Fig. 10) tends to be slightly higher than the counterpart
compressive yield stress. The measured grain sizes in the tensile and compression specimens were 019 and 039 mm in diameter, respectively. The UTS/YS ratio varies from 1.15 to 1.47 and the highest value occurred near 673 K. The UTS curves follow the tensile yield stress curve as shown in Fig. 10. The elongation suddenly increases above 673 K (Fig. 11) which is coincident with the rapid decline of the UTS values above this temperature (Fig. 10). The elongation shown in Fig. 11 is as large as approximately 6% at 873 K where the fracture surface showed increased ductile characteristics.
Fig. 12. (a) Fracture surface pro®le of FePt alloy from tensile deformation at 77 K. (b) Fracture surface pro®le of FePt alloy from tensile deformation at RT. (c) Fracture surface pro®le by SEM from tensile deformation at 673 K. (d) Fracture surface pro®le of SEM from tensile deformation of FePt at 1073 K.
WHANG et al.: L10 TYPE FePt
6491
Fig. 13. [101] dislocations (arrow sign) in the tensile deformation of specimens at 77 K: (a) bright ®eld, B 011, g 200; (b) weak beam, B 011, g 200.
3.4. Fractography The fractured surfaces of tensile specimens deformed at 77 K were investigated by SEM. Figure 12(a) clearly shows predominant intergranular fracture (indicated with arrow sign) accompanied by some transgranular fracture. At RT, the fracture morphology essentially remains in an intergranular mode as shown in Fig. 12(b). When the deformation temperature was increased to 673 K, the fracture morphology [Fig. 12(c)] changed from a predominant intergranular mode to a mixed-mode of intergranular (left-middle, arrow sign) and transgranular cracking (top-center, arrow sign). The elongation increased to 3% at this temperature from 1% at RT. In the middle part of Fig. 12(c), surface dimple patterns started to appear. The frequency of such dimples signi®cantly increased on the fracture surface of the alloy deformed at 1073 K [Fig. 12(d)], where the elongation was found to be as large as 9.5%. 3.5. Dislocation structure
Fig. 14. Twins ( 111112) found in FePt alloy deformed under tensile stress at 77 K: (a) bright ®eld; (b) diraction pattern; (c) sketch of twin and matrix diraction.
The deformation structure of FePt reveals that slip and twinning are the main deformation mechanisms. Predominant superdislocations with Burgers' vector of [101] [arrow sign in Fig. 13(a)] were observed in the specimens compressively deformed at 77 K. In addition, a signi®cant amount of 1=2110 ordinary dislocations and a small amount of 1=2112 superdislocations were also observed in the same specimen. The weak-beam image of [101] superdislocations in Fig. 13(b) shows no sign of self-dissociation in contrast to the dissociation of [101] superdislocations in L10 type TiAl. The twins found in the deformed alloy at 77 K (Fig. 14) were identi®ed to be of the 111112 system. In other words, pseudo-twin systems of f111g121 or f111g211 were not observed in this alloy. The alloy deformed at RT exhibits an increased amount of ordinary dislocations and a decreased amount of superdislocations (b 101) [Figs 15(a) and (b)] compared to the distribution in the 77 K alloy [Figs 13(a) and (b)], while the mor-
6492
WHANG et al.: L10 TYPE FePt
Fig. 15. Dislocations in FePt deformed in tension at RT: (a) bright ®eld, b 1=2110, B 001, g 200; (b) bright ®eld, b 101, B 001, g 200.
Fig. 16. Dislocations in FePt deformed in tension at 873 K: (a) bright ®eld, b 1=2110, B 001, g 200; (b) bright ®eld, b 101, B 001, g 200.
Fig. 17. Dislocations in FePt deformed in tension at 1073 K: (a) bright ®eld, b 1=2110, B 211, g 022; (b) bright ®eld, b 101, B 112, g 110.
WHANG et al.: L10 TYPE FePt
phology of these dislocations remains unchanged regardless of deformation temperature. The density of ordinary dislocations (b 1=2110) was found to be further increased in the alloy deformed at 873 K [Fig. 16(a)]. Also, superdislocations with Burgers' vectors of [101] [Fig. 16(b)] were still found in this alloy, though the density is much less than that of ordinary dislocations. The morphology of these dislocations is similar to that of dislocations found in the deformed alloy at RT as shown in Fig. 16(b). When the alloy was deformed at 1073 K, the density of 1=2110 dislocations became further dominant over that of [101] superdislocations, while the morphology remained unchanged [Fig. 17(a)] except that the dislocations were elongated. In other words, the morphology and characteristics of the ordinary dislocations do not change with increasing temperature, while the fraction of ordinary dislocations increases with temperature. Those [101] type dislocations observed in the alloy deformed at 1073 K [in Fig. 17(b)] exhibit morphology similar to those found in the specimens deformed at lower temperatures. Neither the decomposition of superdislocations into perfect dislocations nor the self-dissociation into partial dislocations was observed throughout the entire temperature range investigated, which is in contrast to those cases in L10 TiAl. 4. DISCUSSION
The LRO parameter S for FePt was determined as a function of time (t) and temperature. The plot of S vs t based on experimental data at 773 K is shown in Fig. 3. The S value for 773 K, 256 h was assumed to be a reference point, i.e. S 773 K,256 h 1. The regression curve of the data based on a power equation was obtained, i.e. S atn , where constants ``a'' and ``n'' were found to be 0.86 and 0.025, respectively. A series of S values can be calculated by this formula at dierent annealing times at 773 K, for example, S 130 h 0:960 and S 256 h 0:988. Therefore, it is reasonable to regard the S value for 773 K, 256 h as a reference point within experimental error, i.e. S 773 K,256 h 1. From the curve in Fig. 3, it is seen that the LRO parameter S reaches 0.85 after annealing only for 0.5 h at 773 K. This means that the major portion of ordering takes place very rapidly in the early stage of annealing at this temperature. Earlier, Vlasova and Sapozhkova [6] arrived at the same conclusion that rapid ordering at an intermediate temperature occurs in an Fe± 32 at.% Pt alloy. The reason for this is that the long-range diusion is not required in the process of ordering in FePt. The transformation from f.c.c. to f.c.t. (L10) requires exchange among the nearest neighbor atoms. The LRO S in L10 FePt approaches a perfect value (S = 1) after 256 h annealing at this temperature (Fig. 3).
6493
The S value for FePt as a function of temperature shows that the S value decreases very slowly with temperature from nearly unity at 773 K to 0.81 at 1567 K (Fig. 4). The ®rst derivation of the S curve yields a discontinuity at 1573 K, and therefore it can be concluded that the order±disorder transition from g to g2 in FePt is of a ®rst-order type. This is in agreement with the prediction made by Kussman and Rittbeg [4]. This curve shown in Fig. 4 is very similar to those of other L10 binary alloys such as CuAu [10] and CoPt [11]. Roberts [10] found that the order±disorder transformation for CuAu I was ®rst-order as con®rmed by the X-ray diraction method, though initially Young's modulus measurements for L10 type CoPt indicated that the order±disorder transformation was a secondorder transformation [12]. In Fig. 18, the plot of the LRO parameter S for FePt against T/Tc is compared to those for CuAu and CoPt, which were also determined by the X-ray diraction method. All the curves show that the degree of ordering decreases slowly with increasing temperature at ®rst, remains high as the temperature approaches a critical point, Tc, and then suddenly drops to zero at Tc. The yield strengths of FePt and TiAI at RT are similar in magnitude. At high temperatures, however, the yield strength of FePt monotonously decreases, showing negative temperature dependence, whereas TiAl exhibits strong positive dependence at high temperatures [7±9]. In general, the yield stress may be a sum of contributions from various intrinsic and extrinsic terms. The yield stress may be expressed in a linear form for a small deformation, i.e. sy sp sd si sj sq
4
where sp, sd, si, sj, sq, are stress contributions from Peierls force, production of all dislocations, dislocation intersection, production of jogs, and production of defects other than dislocations, re-
Fig. 18. Long-range order parameter S expressed as a function of the critical temperature Tc for CuAu, CoPt and FePt.
6494
WHANG et al.: L10 TYPE FePt
spectively. For a very small plastic deformation (strain), however, si, sj, sq may be negligible. The dislocation±dislocation interaction and dislocation± twin interaction are minimal in a small strain. Hence, equation (4) may be reduced to sy sp sd for small deformation. The large yield stress below RT appears to be associated with high Peierls stress, and the diculty of nucleating ordinary dislocations. The relatively low density of ordinary dislocations in FePt at cryogenic temperatures indicates that the Peierls stress increase is more pronounced in ordinary dislocations than in superdislocations at these temperatures. The observation of dislocation structure in the deformed alloys at various temperatures shows that superdislocation slip (101f111g) is dominant at 77 K and signi®cant at RT, while at high temperatures (e.g. 1073 K), ordinary dislocation slip (1=2110f111g) is dominant. Therefore, the contribution of dislocations in equation (4) may be divided into two terms, i.e. a term representing the superdislocation contribution and a term representing the ordinary dislocation contribution. Therefore, sp and sd may be again broken into two components sp f0 s0p fs ssp where f0 fs 1 sd f0 s0d fs ssd where f0, fs are fractions of ordinary dislocations and superdislocations; s0p, ssp are Peierls stress components from ordinary dislocations and superdislocations; and s0d and ssd are glide stress components from ordinary dislocations and superdislocations. Therefore, equation (4) may be written as sy f0 s0p fs ssp f0 s0d fs ssd :
the tensile yield stress must be larger than the compressive yield stress due to the Hall±Petch eect as well as surface eects. The strain rate sensitivity parameter (z) near 0.4Tm, where Tm is melting temperature, does not vanish in FePt, similar to that of copper and silver where the z values do not vanish, but are as low as 10±20 10ÿ6/K at this temperature (Fig. 19). The strain rate sensitivity values for polycrystalline FePt, polycrystalline Ti±55Al±10V and single crystal Cu are compared in Fig. 19. The dierence between the two curves is clear. The curve for FePt is nearly constant from 273 to 673 K similar to those of Cu and Ag (f.c.c.), while that of TiAl approaches a near-zero value at 0.4Tm. The discrepancy between the deformation mode of FePt and that of TiAl may be explained from the dislocation characteristics. First, deformation in FePt appears to be governed by ordinary dislocations and twins since the superdislocations were rarely observed at RT and 673 K. Second, the ordinary dislocations at RT and 673 K were curved without any other special character, whereas in TiAl the ordinary dislocation lines at RT are straight implying the existence of a high Peierls stress, but they become curved at high temperatures. In FePt, the superdislocations (b 101) show neither self-dissociation nor cross-slip onto cube planes at RT and 673 K, though it is possible that the dissociation or cross-slip width might be too small (<1 nm) to be identi®ed by weak-beam imaging using a 200 keV TEM. As a result, no dislocation pinning structure, which might be responsible for anomalous strain hardening, was found in this investigation to be consistent with the defor-
5
At very low temperatures, equation (5) may be written as sy ssp ssd :
50
At very high temperatures, equation (5) may be written as sy s0p s0d :
50
The yield stress determined from the compression test is systematically higher than that by the tensile test. The reason for such inconsistency appears to originate from the specimen geometry and grain size. The compression specimens have a cylindrical shape of 4.3 mm in diameter compared to that of 3 mm in diameter for the tensile specimens. Therefore, the surface to volume ratio is higher in the tensile specimens than in the compression specimens. Second, the grain sizes in the compression and tensile specimens are 39 and 19 mm in diameter, respectively. Hence, it is anticipated that
Fig. 19. Comparison among strain rate sensitivity parameters for FePt, Cu and TiAl.
WHANG et al.: L10 TYPE FePt
mation behavior of FePt. In other words, neither does the deformation lead to any anomalous hardening at high temperatures nor does the strain rate sensitivity parameter ever become vanishingly small in L10 FePt. 5. CONCLUSIONS
1. The isothermal annealing of FePt at 773 K shows that the LRO parameter S reaches 0.85 in the ®rst 30 min of annealing and thereafter, it slowly approaches unity. 2. The LRO parameter S for FePt decreases slowly with increasing temperature and vanishes discontinuously from the value of 0.81 near the critical temperature (1573 K). As a result, the order±disorder transition of FePt is concluded to be a ®rst-order phase transformation. 3. The yield stress of polycrystalline FePt decreases monotonically with increasing temperature from RT to 1073 K, showing negative temperature dependence. The strain sensitivity curve as a function of temperature forms a ``U'' shaped basin with its minimum value between 250 and 670 K. The values in the ``U'' basin remain in the range of 10±20 10ÿ6/K, similar to those of silver and copper. The value, however, rapidly increases above 673 K (0.4Tm). 4. Deformation in FePt is carried out by a combination of slip and twinning. The slip involves three dierent types of slip systems: 1=2h110f111g, h101f111g and 1=2h112f111g. The morphology of ordinary dislocations shows no special features, similar to those found in f.c.c. metals. At cryogenic temperatures (e.g. 77 K), h101f111g slip becomes dominant whereas at high temperatures, 1=2h110f111g slip is dominant. 5. The dominant fracture mode from 77 K to RT was found to be intergranular, but above 673 K, the mode becomes transgranular. Dimple mor-
6495
phology becomes predominant in the fracture at 1073 K, indicating the ductile fracture operating at this temperature. 6. Neither were 1=2h101 superdislocations found to be self-dissociated nor did they cross-slip onto cube planes. No potential locking mechanism was observed by TEM in FePt, which is consistent with the deformation behavior of this alloy. 7. It is demonstrated that the signi®cant disparity between the temperature dependent deformation of FePt and that of TiAl, i.e. positive or negative, originates from the dierence in the atomic bond characteristics.
AcknowledgementsÐThe authors acknowledge the support of the Division of Basic Energy Science, Department of Energy (Contract No. DE-FG02-93ER45499) for this research.
REFERENCES 1. Fu, C. L. and Yoo, M. H., Phil. Mag. Lett., 1990, 62, 159. 2. Massalski, T. B., Murray, J. L., Bennet, L. H. and Baker, H. (ed.), Binary Phase Diagrams. ASM International, Materials Park, Ohio, 1986, p. 1096. 3. Hausen, M. and Anderko, K., Constitution of Binary Alloys. McGraw-Hill, New York, 1958. 4. Kussman, A. and Rittberg, G., Z. Metallk., 1953, 10, 243. 5. Cabri, L. J., Owens, D. R. and La¯amme, J. H. G., Can. Mineral., 1973, 12, 21. 6. Vlasova, Ye. N. and Sapozhkova, T. P., Fiz. phys. Met. Metallogr., 1970, 5, 85. 7. Kawabata, T., Kanai, T. and Izumi, O., Acta metall., 1985, 33, 1355. 8. Li, Z. C. and Whang, S. H., Mater. Sci. Engng, 1992, A152, 182. 9. Hahn, Y. D. and Whang, S. H., Metall. Mater. Sci. Trans. A, 1995, 26A, 113. 10. Roberts, B. W., Acta metall., 1954, 2, 597. 11. Rudman, P. S. and Averbach, L. A., Acta metall., 1957, 5, 67. 12. Gebhart, E. and Koster, W., Z. Metallk., 1940, 32, 253.
Acta mater. Vol. 46, No. 18, pp. 6485±6495, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00311-5 1359-6454/98 $19.00 + 0.00
ORDERING, DEFORMATION AND MICROSTRUCTURE IN L10 TYPE FePt S. H. WHANG{, Q. FENG and Y.-Q. GAO Polytechnic University, Six Metrotech Center, Brooklyn, NY 11201, U.S.A. (Received 27 March 1997; accepted 21 August 1998) AbstractÐThe long-range order (LRO) parameters for L10 FePt have been determined by the X-ray diraction method from powder specimens as functions of time and temperature from 773 K to near 1573 K. By and large, the ordering takes place rapidly below Tc and reaches as high as 0.85 even at 773 K in the ®rst 30 min. The LRO value is about 0.81 near Tc (1573 K) before it drops abruptly to zero at 1573 K. As a result, the order±disorder transformation in FePt is concluded to be a ®rst-order phase transformation. Deformation behavior in an L10 type FePt alloy was investigated through both compressive and tensile deformation from room temperature (RT) to 1073 K. The negative temperature dependence of yield stress in this alloy contrasts with the positive dependence in L10 type TiAl. The elongation increases exponentially with temperature and reaches 06% at 873 K. The strain rate sensitivity parameter against temperature is similar to those found in silver and copper, where the non-zero minimum is centered in a broad basin. This indicates that the temperature-dependent deformation in the range of RT to 1073 K is analogous to that of some face-centered cubic metals, but signi®cantly dierent from that of L10 TiAl. The deformation structure investigated by TEM shows that slip and twinning are the two major deformation mechanisms. The identi®ed slip systems include 1=2110f111g; h101f111g and 1=2h112f111g. The 112h112f111g slip system, however, is only active at very low temperatures, e.g. 77 K. The twin system was identi®ed as f111gh112 type. No pseudo-twinning was found in this alloy. The deformation below RT is mainly carried out by both superdislocations and ordinary dislocations, while above 673 K, it is carried out mainly by ordinary dislocations. The morphology of these dislocations in the entire temperature range indicates that the dislocations do not experience a high Peierls stress contrary to that observed in TiAl. No self-dissociation of superdislocations or APB cross-slip onto cube planes was observed under weak beam conditions. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
The majority of L10 type compound alloys are composed of combinations of transition metals, but there are a few exceptions such as TiAl and TiGa. The atomic size dierence between the two component atoms is less than 15% in L10 compounds. This indicates that the size factor is important for the stability of L10 compounds. Second, the majority of these compounds have a c/a ratio less than unity. The lattice parameter contraction occurs along the c-axis which is perpendicular to the superstructure layers. Hence, the atomic distance between unlike atoms is an important factor for the formation of the L10 structure. The exceptions for the c/a ratio include TiAl and TiHg where the c/a ratio is larger than unity. In TiAl, a directional bond between Ti and Ti was found [1] and may be responsible for the unusual c/a ratio. Stoichiometric FePt, as shown in the Fe±Pt phase diagram (Fig. 1) [2, 3], solidi®es into a disordered face-centered cubic (f.c.c.) phase from the melt and transforms into an ordered face-centered tetragonal (f.c.t.) phase (L10) near 1573 K. The characteristics of this transformation include (i) the order±disorder {To whom all correspondence should be addressed.
transformation takes place at a very high temperature, and (ii) the ordering does not require a longrange diusion process. The transformation rapidly takes place in the absence of long-range diusion. In ordered FePt, Fe and Pt atoms are disposed in alternative layers of {200} planes. The L10 FePt phase region in the FePt phase diagram is surrounded by narrow two-phase regions (shown in Fig. 1), which indicates that the formation of the FePt superstructure may be characterized by a ®rstorder phase transition [4]. Cabri et al. found that the ordered phase of FePt was an f.c.t. phase (L10 type) con®rmed by the X-ray powder diraction method [5]. Vlasova and Sapozhkova studied the type and kinetics of ordering in Fe±32 at.% Pt alloy and found that the ordering in FePt took place rapidly as a result of annealing at 1023 K [6]. In order to determine the long-range order (LRO) parameter (S) of FePt, the present study is made with regard to the time and temperature dependence of ordering by X-ray diraction measurements. The deformation behavior of L10 FePt has not been investigated in the past. Previous studies of L10 TiAl show that anomalous hardening occurs at high temperatures [7±9]. Therefore, fundamental questions have been raised whether the anomaly is
6485
6486
WHANG et al.: L10 TYPE FePt
Fig. 1. Binary phase diagram of Fe±Pt.
an intrinsic property of the L10 structure or is associated with a particular bond, or is limited to L10 TiAl. Further, it is important to understand the relationship among deformation behavior, resulting dislocation structure, other planar defect structures, and chemical and physical properties in L10 type compounds. FePt may be considered a model L10 compound representing a group of L10 transition± transition metal compounds. In addition, FePt has a high order±disorder transition temperature of 1573 K as shown Fig. 1 [2], thus allowing an investigation of deformation behavior at high temperatures. Therefore, a study of L10 FePt will answer some of these questions, which include (i) the temperature-dependent deformation behavior of FePt; (ii) ®ne dislocation structures resulting from the deformation and their relationship with work hardening behavior. The deformation behavior of FePt from 77 to 873 K and the corresponding microstructures are reported in this paper. 2. EXPERIMENTAL PROCEDURE
FePt alloy buttons were made in an arc furnace from powders of pure iron (99.99%) and pure platinum (99.9%). The arc melted buttons, which were placed in an argon-pressurized quartz tube, were homogenized at 1643 K for 24 h, and subsequently quenched immediately into ice±water. The powder specimens were prepared by ®ling o the quenched FePt buttons. The powders were found to have a disordered f.c.c. phase (g) identi®ed by X-ray diffraction. The powder specimens were sealed in very
narrow, evacuated quartz tubes and annealed at various temperatures until the ordering had been achieved at a saturated level, followed by quenching into a 10% NaCl ice±water solution. The annealing for ordering was performed at various temperatures from 773 K to near 1573 K; and from 0.5 to 256 h. The X-ray diraction patterns and data were obtained from powder specimens using a Philips XRG Cu ÿ Ka1 X-ray diractometer at 40 kV, 20 mA. For tensile specimens, the arc melted buttons were re-melted by induction power in an argon atmosphere and cast into a ceramic tube. The cast rods were annealed at 1173 K for homogenization. The L10 structure was con®rmed by X-ray diraction and TEM. The alloy was cast in a copper mold followed by annealing at 1173 K for 115 h. The long rods were machined into non-standard tensile specimens with a gauge length of 22 mm and diameter of 3 mm. The machined tensile specimens were polished with ®ne polishing paper, and electropolished. The initial strain rate of 3 10ÿ4/s was used for tensile testing in argon atmosphere. Compression specimens were machined from the rods into cylindrical specimens of 4.3 mm in diameter and 9 mm in length. The compression tests were carried out with strain rates of 3 10ÿ3 and 3 10ÿ4/s at room temperature (RT), 473, 673 and 873 K, respectively. TEM work was performed with a 200 keV TEM at RT. Also, two-beam and weak-beam techniques were performed with g±3g conditions using a 200 keV TEM.
WHANG et al.: L10 TYPE FePt 3. RESULTS
6487
Fundamental: Ff 2 fFe fPt ; h k even, hkl unmixed
3.1. Ordering transformation Typical X-ray patterns obtained from the powder specimens of FePt are shown in Fig. 2. In Fig. 2(a), the pattern from the homogenized specimen at 1643 K exhibits the diraction lines of a disordered g-single phase. In Fig. 2(b), however, the diraction patterns from a specimen annealed at 1473 K for 1 h show many superlattice diraction lines of the ordered g2 phase (L10 type). The LRO parameter, S, for the L10 FePt structure can be determined from the integrated intensity ratio of the superlattice diraction peaks and the fundamental diraction peaks. The structure factor F takes two forms:
Superlattice: Fs 2S fFe ÿ fPt ; h k even, hkl mixed: The integrated intensity, I, per unit length of powder-pattern diraction line is given by I km LP F 2 A yT yE y
1
where k is a constant for all re¯ections; m, the multiplicity of the re¯ecting planes; LP, the Lorenz polarization factor; A(y), the absorption correction factor; T(y), the Debye temperature factor; and E(y), the extinction factor. The LRO parameter S
Fig. 2. X-ray diraction pattern obtained with Cu ÿ Ka1 radiation of Fe±50Pt alloy: (a) the disordered g phase; (b) the ordered g1 phase.
6488
WHANG et al.: L10 TYPE FePt
Fig. 3. Long-range order parameter S for FePt as a function of annealing time at 773 K.
Fig. 5. Plot of strain and stress curves using alternating strain rates in FePt.
may be given by As m LP fFe fPt 2 A yT yE yf S2 : Af m LP fFe fPt 2 A yT yE ys
intensity of the superlattice diraction, Is,T, and the fundamental diraction of the (111) plane, I(111),T at a given temperature, T
2
where As and Af are the peak areas of superlattice diraction and fundamental re¯ections, respectively. When the LRO in the alloy is perfect, S is equal to unity; when the atomic arrangement is completely random, S is equal to zero. In principle, the value of S can be determined by calculating the bracketed term in equation (2). Nevertheless, in practice, the bracketed term in equation (2) cannot be accurately calculated due to (i) lack of understanding of the preferred orientation in the powder specimens; (ii) the diculty in estimating the Honl L-electron correction for the Pt scattering factor; and (iii) the uncertainty in estimating the extinction factor and degree of powder size uniformity. Alternatively, if one uses the X-ray patterns corresponding to nearly perfect ordering as a reference, S values for various ordered states may be determined by comparing the superlattice intensity data of an investigating state with that of the reference state. In other words, S values can be determined simply from the relative ratio of the integrated
Fig. 4. Long-range order parameter S for FePt as a function of temperature.
Is,T =I 111,T S: Is,773 K =I 111,773 K
3
The latter method was adopted in the current experiments. For FePt, the reference patterns were obtained from the specimen annealed at 0.5Tc. The 0.5Tc for FePt is located 0773 K. Four superlattice diraction patterns were selected to calculate the LRO parameter S. They were re¯ections from (001), (110), (021), (112) planes. The value of S at each temperature was optimized by the method of least squares. Also, St,773 K was obtained by annealing the specimens at 773 K for various durations from 0.5 to 256 h. The results shown in Fig. 3 indicate that the S value asymptotically approaches a ®nite value, i.e. S = 1. Therefore, the S value for 773 K, 256 h, where the
Fig. 6. Temperature dependence of compressive yield stress.
WHANG et al.: L10 TYPE FePt
6489
Fig. 7. Strain rate sensitivity as a function of temperature.
Fig. 10. Plot of tensile yield stress as well as UTS vs temperature in FePt.
Fig. 8. Fractured tensile specimen of FePt at RT.
S value approaches a saturated value, was conveniently taken as a reference point. Further, the ratio Is,773 K =I 111,773 K may be de®ned to establish the relationship between S and the fundamental diffraction intensity in this alloy. Various S values at dierent temperatures were determined, which are shown in Fig. 4. The S curve in Fig. 4 shows that the LRO parameter decreases very slowly with increasing temperature to a value of about 0.81 at 1567 K, where it drops sharply to zero at the transition temperature, 1573 K.
was investigated from 77 to 1073 K as shown in Fig. 6. The sharp rise in yield stress at 77 K was noted. In contrast, the decline in yield stress with increasing temperature is moderate from RT to 1073 K. The absence of anomalous hardening in FePt is also characteristic, in contrast to the case of TiAl where anomalous hardening occurs from RT to 1073 K [9]. The strain rate sensitivity parameter, z 1=T @ ln s=@ ln e_ from the compression test adopts a ``U'' shape centered near 600 K (Fig. 7). The minimum value of S occurs between 400 and 700 K, followed by a sharp increase above 900 K. Despite a lower value of z, the z value never approaches zero in contrast to the zero minimum near 600 K in L10 TiAl.
3.2. Compression test The compression tests were performed at the initial strain rate of 3 10ÿ4/s (``a'' in Fig. 5) which was subsequently changed to an alternate rate, 3 10ÿ3/s (``b'' in Fig. 5), and thereafter the strain rate oscillated between the two strain rates. The temperature dependence of yield stress (0.2% oset)
Fig. 9. Stress±strain curves under compression for RT to 1073 K.
Fig. 11. Elongation vs temperature in FePt.
6490
WHANG et al.: L10 TYPE FePt
3.3. Tensile test The non-standard tensile specimens of an FePt alloy tested in an argon atmosphere from RT to 1073 K did not show any necking resulting from RT testing (Fig. 8). The stress±strain curves at various temperatures are shown in Fig. 9. The stress± strain curves do not show a sharp yield point which is normally observed in carbon steels. Nevertheless, the trend in temperature dependence is consistent with that observed in the compression test (Fig. 5). The tensile yield stress from RT to 1073 K (Fig. 10) tends to be slightly higher than the counterpart
compressive yield stress. The measured grain sizes in the tensile and compression specimens were 019 and 039 mm in diameter, respectively. The UTS/YS ratio varies from 1.15 to 1.47 and the highest value occurred near 673 K. The UTS curves follow the tensile yield stress curve as shown in Fig. 10. The elongation suddenly increases above 673 K (Fig. 11) which is coincident with the rapid decline of the UTS values above this temperature (Fig. 10). The elongation shown in Fig. 11 is as large as approximately 6% at 873 K where the fracture surface showed increased ductile characteristics.
Fig. 12. (a) Fracture surface pro®le of FePt alloy from tensile deformation at 77 K. (b) Fracture surface pro®le of FePt alloy from tensile deformation at RT. (c) Fracture surface pro®le by SEM from tensile deformation at 673 K. (d) Fracture surface pro®le of SEM from tensile deformation of FePt at 1073 K.
WHANG et al.: L10 TYPE FePt
6491
Fig. 13. [101] dislocations (arrow sign) in the tensile deformation of specimens at 77 K: (a) bright ®eld, B 011, g 200; (b) weak beam, B 011, g 200.
3.4. Fractography The fractured surfaces of tensile specimens deformed at 77 K were investigated by SEM. Figure 12(a) clearly shows predominant intergranular fracture (indicated with arrow sign) accompanied by some transgranular fracture. At RT, the fracture morphology essentially remains in an intergranular mode as shown in Fig. 12(b). When the deformation temperature was increased to 673 K, the fracture morphology [Fig. 12(c)] changed from a predominant intergranular mode to a mixed-mode of intergranular (left-middle, arrow sign) and transgranular cracking (top-center, arrow sign). The elongation increased to 3% at this temperature from 1% at RT. In the middle part of Fig. 12(c), surface dimple patterns started to appear. The frequency of such dimples signi®cantly increased on the fracture surface of the alloy deformed at 1073 K [Fig. 12(d)], where the elongation was found to be as large as 9.5%. 3.5. Dislocation structure
Fig. 14. Twins ( 111112) found in FePt alloy deformed under tensile stress at 77 K: (a) bright ®eld; (b) diraction pattern; (c) sketch of twin and matrix diraction.
The deformation structure of FePt reveals that slip and twinning are the main deformation mechanisms. Predominant superdislocations with Burgers' vector of [101] [arrow sign in Fig. 13(a)] were observed in the specimens compressively deformed at 77 K. In addition, a signi®cant amount of 1=2110 ordinary dislocations and a small amount of 1=2112 superdislocations were also observed in the same specimen. The weak-beam image of [101] superdislocations in Fig. 13(b) shows no sign of self-dissociation in contrast to the dissociation of [101] superdislocations in L10 type TiAl. The twins found in the deformed alloy at 77 K (Fig. 14) were identi®ed to be of the 111112 system. In other words, pseudo-twin systems of f111g121 or f111g211 were not observed in this alloy. The alloy deformed at RT exhibits an increased amount of ordinary dislocations and a decreased amount of superdislocations (b 101) [Figs 15(a) and (b)] compared to the distribution in the 77 K alloy [Figs 13(a) and (b)], while the mor-
6492
WHANG et al.: L10 TYPE FePt
Fig. 15. Dislocations in FePt deformed in tension at RT: (a) bright ®eld, b 1=2110, B 001, g 200; (b) bright ®eld, b 101, B 001, g 200.
Fig. 16. Dislocations in FePt deformed in tension at 873 K: (a) bright ®eld, b 1=2110, B 001, g 200; (b) bright ®eld, b 101, B 001, g 200.
Fig. 17. Dislocations in FePt deformed in tension at 1073 K: (a) bright ®eld, b 1=2110, B 211, g 022; (b) bright ®eld, b 101, B 112, g 110.
WHANG et al.: L10 TYPE FePt
phology of these dislocations remains unchanged regardless of deformation temperature. The density of ordinary dislocations (b 1=2110) was found to be further increased in the alloy deformed at 873 K [Fig. 16(a)]. Also, superdislocations with Burgers' vectors of [101] [Fig. 16(b)] were still found in this alloy, though the density is much less than that of ordinary dislocations. The morphology of these dislocations is similar to that of dislocations found in the deformed alloy at RT as shown in Fig. 16(b). When the alloy was deformed at 1073 K, the density of 1=2110 dislocations became further dominant over that of [101] superdislocations, while the morphology remained unchanged [Fig. 17(a)] except that the dislocations were elongated. In other words, the morphology and characteristics of the ordinary dislocations do not change with increasing temperature, while the fraction of ordinary dislocations increases with temperature. Those [101] type dislocations observed in the alloy deformed at 1073 K [in Fig. 17(b)] exhibit morphology similar to those found in the specimens deformed at lower temperatures. Neither the decomposition of superdislocations into perfect dislocations nor the self-dissociation into partial dislocations was observed throughout the entire temperature range investigated, which is in contrast to those cases in L10 TiAl. 4. DISCUSSION
The LRO parameter S for FePt was determined as a function of time (t) and temperature. The plot of S vs t based on experimental data at 773 K is shown in Fig. 3. The S value for 773 K, 256 h was assumed to be a reference point, i.e. S 773 K,256 h 1. The regression curve of the data based on a power equation was obtained, i.e. S atn , where constants ``a'' and ``n'' were found to be 0.86 and 0.025, respectively. A series of S values can be calculated by this formula at dierent annealing times at 773 K, for example, S 130 h 0:960 and S 256 h 0:988. Therefore, it is reasonable to regard the S value for 773 K, 256 h as a reference point within experimental error, i.e. S 773 K,256 h 1. From the curve in Fig. 3, it is seen that the LRO parameter S reaches 0.85 after annealing only for 0.5 h at 773 K. This means that the major portion of ordering takes place very rapidly in the early stage of annealing at this temperature. Earlier, Vlasova and Sapozhkova [6] arrived at the same conclusion that rapid ordering at an intermediate temperature occurs in an Fe± 32 at.% Pt alloy. The reason for this is that the long-range diusion is not required in the process of ordering in FePt. The transformation from f.c.c. to f.c.t. (L10) requires exchange among the nearest neighbor atoms. The LRO S in L10 FePt approaches a perfect value (S = 1) after 256 h annealing at this temperature (Fig. 3).
6493
The S value for FePt as a function of temperature shows that the S value decreases very slowly with temperature from nearly unity at 773 K to 0.81 at 1567 K (Fig. 4). The ®rst derivation of the S curve yields a discontinuity at 1573 K, and therefore it can be concluded that the order±disorder transition from g to g2 in FePt is of a ®rst-order type. This is in agreement with the prediction made by Kussman and Rittbeg [4]. This curve shown in Fig. 4 is very similar to those of other L10 binary alloys such as CuAu [10] and CoPt [11]. Roberts [10] found that the order±disorder transformation for CuAu I was ®rst-order as con®rmed by the X-ray diraction method, though initially Young's modulus measurements for L10 type CoPt indicated that the order±disorder transformation was a secondorder transformation [12]. In Fig. 18, the plot of the LRO parameter S for FePt against T/Tc is compared to those for CuAu and CoPt, which were also determined by the X-ray diraction method. All the curves show that the degree of ordering decreases slowly with increasing temperature at ®rst, remains high as the temperature approaches a critical point, Tc, and then suddenly drops to zero at Tc. The yield strengths of FePt and TiAI at RT are similar in magnitude. At high temperatures, however, the yield strength of FePt monotonously decreases, showing negative temperature dependence, whereas TiAl exhibits strong positive dependence at high temperatures [7±9]. In general, the yield stress may be a sum of contributions from various intrinsic and extrinsic terms. The yield stress may be expressed in a linear form for a small deformation, i.e. sy sp sd si sj sq
4
where sp, sd, si, sj, sq, are stress contributions from Peierls force, production of all dislocations, dislocation intersection, production of jogs, and production of defects other than dislocations, re-
Fig. 18. Long-range order parameter S expressed as a function of the critical temperature Tc for CuAu, CoPt and FePt.
6494
WHANG et al.: L10 TYPE FePt
spectively. For a very small plastic deformation (strain), however, si, sj, sq may be negligible. The dislocation±dislocation interaction and dislocation± twin interaction are minimal in a small strain. Hence, equation (4) may be reduced to sy sp sd for small deformation. The large yield stress below RT appears to be associated with high Peierls stress, and the diculty of nucleating ordinary dislocations. The relatively low density of ordinary dislocations in FePt at cryogenic temperatures indicates that the Peierls stress increase is more pronounced in ordinary dislocations than in superdislocations at these temperatures. The observation of dislocation structure in the deformed alloys at various temperatures shows that superdislocation slip (101f111g) is dominant at 77 K and signi®cant at RT, while at high temperatures (e.g. 1073 K), ordinary dislocation slip (1=2110f111g) is dominant. Therefore, the contribution of dislocations in equation (4) may be divided into two terms, i.e. a term representing the superdislocation contribution and a term representing the ordinary dislocation contribution. Therefore, sp and sd may be again broken into two components sp f0 s0p fs ssp where f0 fs 1 sd f0 s0d fs ssd where f0, fs are fractions of ordinary dislocations and superdislocations; s0p, ssp are Peierls stress components from ordinary dislocations and superdislocations; and s0d and ssd are glide stress components from ordinary dislocations and superdislocations. Therefore, equation (4) may be written as sy f0 s0p fs ssp f0 s0d fs ssd :
the tensile yield stress must be larger than the compressive yield stress due to the Hall±Petch eect as well as surface eects. The strain rate sensitivity parameter (z) near 0.4Tm, where Tm is melting temperature, does not vanish in FePt, similar to that of copper and silver where the z values do not vanish, but are as low as 10±20 10ÿ6/K at this temperature (Fig. 19). The strain rate sensitivity values for polycrystalline FePt, polycrystalline Ti±55Al±10V and single crystal Cu are compared in Fig. 19. The dierence between the two curves is clear. The curve for FePt is nearly constant from 273 to 673 K similar to those of Cu and Ag (f.c.c.), while that of TiAl approaches a near-zero value at 0.4Tm. The discrepancy between the deformation mode of FePt and that of TiAl may be explained from the dislocation characteristics. First, deformation in FePt appears to be governed by ordinary dislocations and twins since the superdislocations were rarely observed at RT and 673 K. Second, the ordinary dislocations at RT and 673 K were curved without any other special character, whereas in TiAl the ordinary dislocation lines at RT are straight implying the existence of a high Peierls stress, but they become curved at high temperatures. In FePt, the superdislocations (b 101) show neither self-dissociation nor cross-slip onto cube planes at RT and 673 K, though it is possible that the dissociation or cross-slip width might be too small (<1 nm) to be identi®ed by weak-beam imaging using a 200 keV TEM. As a result, no dislocation pinning structure, which might be responsible for anomalous strain hardening, was found in this investigation to be consistent with the defor-
5
At very low temperatures, equation (5) may be written as sy ssp ssd :
50
At very high temperatures, equation (5) may be written as sy s0p s0d :
50
The yield stress determined from the compression test is systematically higher than that by the tensile test. The reason for such inconsistency appears to originate from the specimen geometry and grain size. The compression specimens have a cylindrical shape of 4.3 mm in diameter compared to that of 3 mm in diameter for the tensile specimens. Therefore, the surface to volume ratio is higher in the tensile specimens than in the compression specimens. Second, the grain sizes in the compression and tensile specimens are 39 and 19 mm in diameter, respectively. Hence, it is anticipated that
Fig. 19. Comparison among strain rate sensitivity parameters for FePt, Cu and TiAl.
WHANG et al.: L10 TYPE FePt
mation behavior of FePt. In other words, neither does the deformation lead to any anomalous hardening at high temperatures nor does the strain rate sensitivity parameter ever become vanishingly small in L10 FePt. 5. CONCLUSIONS
1. The isothermal annealing of FePt at 773 K shows that the LRO parameter S reaches 0.85 in the ®rst 30 min of annealing and thereafter, it slowly approaches unity. 2. The LRO parameter S for FePt decreases slowly with increasing temperature and vanishes discontinuously from the value of 0.81 near the critical temperature (1573 K). As a result, the order±disorder transition of FePt is concluded to be a ®rst-order phase transformation. 3. The yield stress of polycrystalline FePt decreases monotonically with increasing temperature from RT to 1073 K, showing negative temperature dependence. The strain sensitivity curve as a function of temperature forms a ``U'' shaped basin with its minimum value between 250 and 670 K. The values in the ``U'' basin remain in the range of 10±20 10ÿ6/K, similar to those of silver and copper. The value, however, rapidly increases above 673 K (0.4Tm). 4. Deformation in FePt is carried out by a combination of slip and twinning. The slip involves three dierent types of slip systems: 1=2h110f111g, h101f111g and 1=2h112f111g. The morphology of ordinary dislocations shows no special features, similar to those found in f.c.c. metals. At cryogenic temperatures (e.g. 77 K), h101f111g slip becomes dominant whereas at high temperatures, 1=2h110f111g slip is dominant. 5. The dominant fracture mode from 77 K to RT was found to be intergranular, but above 673 K, the mode becomes transgranular. Dimple mor-
6495
phology becomes predominant in the fracture at 1073 K, indicating the ductile fracture operating at this temperature. 6. Neither were 1=2h101 superdislocations found to be self-dissociated nor did they cross-slip onto cube planes. No potential locking mechanism was observed by TEM in FePt, which is consistent with the deformation behavior of this alloy. 7. It is demonstrated that the signi®cant disparity between the temperature dependent deformation of FePt and that of TiAl, i.e. positive or negative, originates from the dierence in the atomic bond characteristics.
AcknowledgementsÐThe authors acknowledge the support of the Division of Basic Energy Science, Department of Energy (Contract No. DE-FG02-93ER45499) for this research.
REFERENCES 1. Fu, C. L. and Yoo, M. H., Phil. Mag. Lett., 1990, 62, 159. 2. Massalski, T. B., Murray, J. L., Bennet, L. H. and Baker, H. (ed.), Binary Phase Diagrams. ASM International, Materials Park, Ohio, 1986, p. 1096. 3. Hausen, M. and Anderko, K., Constitution of Binary Alloys. McGraw-Hill, New York, 1958. 4. Kussman, A. and Rittberg, G., Z. Metallk., 1953, 10, 243. 5. Cabri, L. J., Owens, D. R. and La¯amme, J. H. G., Can. Mineral., 1973, 12, 21. 6. Vlasova, Ye. N. and Sapozhkova, T. P., Fiz. phys. Met. Metallogr., 1970, 5, 85. 7. Kawabata, T., Kanai, T. and Izumi, O., Acta metall., 1985, 33, 1355. 8. Li, Z. C. and Whang, S. H., Mater. Sci. Engng, 1992, A152, 182. 9. Hahn, Y. D. and Whang, S. H., Metall. Mater. Sci. Trans. A, 1995, 26A, 113. 10. Roberts, B. W., Acta metall., 1954, 2, 597. 11. Rudman, P. S. and Averbach, L. A., Acta metall., 1957, 5, 67. 12. Gebhart, E. and Koster, W., Z. Metallk., 1940, 32, 253.