modulated structure and magnetic properties of age-hardenable fe-mn-al-c alloys

Modulated Properties

Structure and Magnetic of Age-Hardenable Fe-Mn-AI-C Alloys

KAZUNORI SATO, KAZUHIRO TAGAWA, and YASUNOBU INOUE

Microstructural changes in the age-hardenable Fe-(30 to 34) wt pct Mn-(8 to 11) wt pct Al-(0.9 to 1.0) wt pct C alloys during aging in the temperature range between 773 and 823 K have been investigated by means of transmission electron microscopy (TEM) and X-ray diffraction. The wavelength of the modulated structure was found to be nearly constant for short aging times and then to increase on further aging, whereas the compositional modulation amplitude was found to increase rapidly from the beginning of aging. The growth of a spinodally modulated structure along the orthogonal (100) directions results in a periodic arrangement of the K-carbide precipitates, (Fe, Mn)3A1Cx, in the austenite matrix. The increases in hardness and residual and saturation inductions in the early stage of aging were in accord with the increase in the amplitude.

I. INTRODUCTION

THE austenitic Fe-Mn-A1-C alloys, whose compositions are about 30 pct Mn, 8 to 11 pct A1, and about 1 pct C, show age hardening when aged at 723 to 923 K. H'2,31 The structure of the alloys after aging revealed a precipitation of coherent metastable particles of the K-carbide, (Fe, Mn)3A1Cx, in the austenite matrix. I21 Recently, Han et al. t41 observed a modulated structure along the orthogonal (100) directions with superlattice reflections by TEM for an Fe-30Mn-8AI-1.3C* alloy aged

*All in weight percent unless specified otherwise.

at 823 K for 300 minutes. They indicated that the superlattice reflections result from the L' 12 ordering which oc- curred in the carbon-rich zones. Oshima and Wayman I51 also observed the same type of carbide, Fe3A1Cx, which is finely precipitated in the retained austenite of a quenched Fe-7A1-2C alloy. In both cases, it was suggested that the precipitation of the K-carbide results from spinodal decomposition. More recently, we have shown that a coherency strain accompanied by spinodal decomposition can be responsible for the age hardening of an Fe-30Mn- 9A1-0.9C alloy, t61 However, the evidence which shows spinodal decomposition of the Fe-Mn-AI-C system seems rather limited.

Thus, we have made a detailed examination of changes in microstructures during aging of Fe-Mn-AI-C alloys. Further, changes in hardness and magnetic properties have also been investigated as functions of the structural variations: modulation amplitude and wavelength.

KAZUNORI SATO, Research Associate, and YASUNOBU INOUE, Professor, are with the Analysis Center, Nagaoka University of Tech- nology, Nagaoka, Niigata 940-21, Japan. KAZUHIRO TAGAWA, formerly Graduate Student of Mechanical Engineering, Nagaoka University of Technology, is with Tochigi Research Laboratory, Kao Corporation, Tochigi 321-34, Japan.

Manuscript submitted July 6, 1988.

II. EXPERIMENTAL PROCEDURES

A. Specimen Preparation and Heat Treatments

Fe-Mn-A1-C ingots were initially prepared by air induction melting. High-purity iron, electrolytic man- ganese, high-purity aluminum, and electrode-grade graphite were cast into approximately 10-kg ingots. They were homogenized at 1473 K for 60 minutes and sub- sequently forged to 25-mm-thick slabs. Specimens 15 mm in diameter and about 5 mm in thickness were machined from the slabs. They were heated in air at 1273 K for 60 minutes and then water quenched. Oxidation scales were eliminated by mechanical polishing prior to aging. The compositions of the alloys are given in Table I.

In the as-quenched condition, the alloys were com- pletely austenitic (y), except alloy A, in which a slight amount of the ferrite phase ( - 1 vol pct) existed along grain boundaries of the austenite phase. Specimens were isothermally aged in a salt bath (50 pct KNO3 + 50 pct NaNO2) at 773, 793, and 823 K for various times. Aging treatments were terminated by ice-water quenching. The specimens were chemically polished in a hydrogen per- oxide and hydrofluoric acid (3 pct) solution to remove surface scales before the subsequent measurements.

B. X-Ray Diffraction Measurements

The structures of the alloys were characterized by an X-ray diffractometer with Cu-K~ radiation operated at 40 kV and 30 mA. All diffraction patterns of aged specimens were measured at the scanning speed of 1/8 deg min -~. The wavelength of the modulated structure was determined from the spacing between the sideband and the main (200) Bragg peak using the Daniel-Lipson equation:~7]

hao tan 0 ,~ = [ll AO(h 2 + k 2 + 12)

where A = the average modulation wavelength; a0 = the lattice parameter of a homogeneous alloy; 0 = the Bragg angle;

METALLURGICAL TRANSACTIONS A VOLUME 21A, JANUARY 1990--5

Table I. Chemical Compositions of Fe-Mn-AI-C Alloys

(At. Pct)

Alloy (Fe, Mn) AI C Mn A1 C

(Wt Pct)

Si P Phases* A 75.0 20.9 3.93 32.0 11.8 0.986 0.114 0.0034 0.0026 y + a B 76.4 19.6 3.93 34.3 11.0 0.98 0.054 0.009 0.0040 3' C 79.4 16.8 3.85 29.5 9.2 0.94 0.027 0.016 0.0010 y D 81.9 14.4 3.65 30.8 7.8 0.88 0.053 0.015 0.005 y E 84.5 13.4 2.11 33.7 7.1 0.50 0.018 0.009 0.027 y F 86.5 9.3 4.17 28.5 4.9 0.98 0.05 0.0026 0.0034 y

*Identified after heat treatment at 1273 K for 60 min followed by water quenching.

A0 = the angular spacing between the sideband and the main (200) Bragg peak; and

h, k, l = the Miller indices of the Bragg peak (h = 2 a n d k , / = 0).

C. Electron Microscopy

Thin foil specimens were prepared by an electropolishing in a solution of chromic (50 g) and phosphoric (400 ml) acids at 40 ~ to 70 ~ The electropolishing was conducted at 10 V and 0.5 to 1.0 A. The specimens were examined in a transmission electron microscope operated at 200 kV. Wavelengths of the modulated structure were also measured directly along the (100) and (110) directions on enlarged micrographs.

D. Optical Microscopy and Electron Microprobe Analysis

Optical microscopic observations were made for the polished specimens which had been etched in 3 pct nital solution. An electron microprobe analysis was conducted for the quantification of the precipitated phase along grain boundaries of the austenite phase.

E. Hardness and Magnetic Measurements

The changes in hardness for specimens aged at 823 K with aging time were measured by a Rockwell hardness tester using the A scale. Magnetic properties were examined by measuring magnetic hysteresis curves. The curves were measured with a vibrating sample magne- tometer (VSM) at a maximum magnetic field, 5 kOe. The residual induction, Br, and saturation induction, Bs, were obtained from the hysteresis curves.

I I I . R ES ULTS

Figure 1 shows hardness-time curves for specimens aged at 823 K. Those alloys which have both aluminum and carbon contents greater than about 8 and 0.9 pct, respectively, showed distinct age hardening. These age- hardenable alloys showed a rapid increase in hardness in the early stage of aging irrespective of their compositions. In addition to the first-stage age hardening, alloy C showed the second-stage increase in hardness for longer aging times after passing a nearly stationary hardness level (Figure l(b)), while alloys E and F, in which either aluminum or carbon content is less than about 8 and 1 pct, respectively, showed no age hardening. This result dem- onstrates that the concentrations of aluminum and carbon

atoms are indispensable for the age hardening of austenitic Fe-Mn-A1-C alloys.

Figure 2 shows variations of an X-ray diffraction profile for alloy C aged at 823 K for different aging times. After long aging times, reflections from a new phase which appears on the slightly lower angle side of the austenite matrix peak, an a-ferrite, and a fl-Mn phase were identified with superlattice (lO0) and (1 lO) reflections. These superlattice reflections were first observed

70

n " v

6O c-

10 lb. l:l

" I-

B 0

5o -= i m

I I 10 2 10 3 10 4

Aging time (rain)

(a)

n- 70

in

t--

-l- 60

C J

I I I AO. 10 2 10 3 10 4

Aging t ime (rain)

(b)

Fig. l --Age-hardening curves of (a) alloys B, D, E, F, and (b) alloy C aged at 823 K.

6 - - V O L U M E 21A, J A N U A R Y 1990 M E T A L L U R G I C A L T R A N S A C T I O N S A

A o 0

U u

(d) ~,

E 3 ,4

~ (c)

U u

>,

- / (b)

(a)

I 24 34

E

,7011) ,r(2oo)

1 I I I

46 50 I

42

2 O (deg)

Fig. 2 - -Var i a t i ons of X-ray diffraction profile for alloy C aged at 823 K for (a) 6 0 m i n , (b) 4.1 x l03min , (c) 1.3 • 104min, and (d) 3.4 x l04 min.

after 120 minutes of aging. The formation of these phases from the austenite matrix occurs slowly {1,2} compared with the rapid increase in hardness during the early stage of aging, as shown in Figure 1. The new phase has an fcc structure with lattice parameter a = 0.374 nm after 3.4 • 104 minutes of aging, whereas the austenite matrix shows a = 0.370 nm in the as-quenched condition. Thus, the lattice parameter of the newly formed phase is close to that of the ordered g phase, (Fe, Mn)3AICx, with a = 0.372 to 0.378 nm. [~,2,5}

Alloy C also represented sidebands around the main (200) Bragg line for short aging times, as shown in Figure 3. In the as-quenched condition, a sharp fcc (200) reflection was observed. The sideband peaks began to appear after about 90 minutes. On aging, the sidebands moved close to the Bragg line, with a concurrent increase in their intensity. On further aging, the sidebands disappeared, with a broadening of the Bragg line re- suiting in the appearance of the K-carbide (200) reflection. During the early stage of aging, X-ray sideband peaks were observed for alloy C at 723 to 823 K. The other age-hardenable alloys also showed sideband peaks in the same temperature range; however, these sideband peaks were not as distinct as alloy C.

Electron microscopic observations of austenite grains for alloys B, C, and D aged at 823 K for 30 minutes, 1.3 • I 0 4 minutes, and corresponding electron diffraction patterns are shown in Figures 4 through 6. Fine modulated structures along the orthogonal (100) directions were identified for these alloys. Aging for 1.3 •

t--

m i n

~k127 min 103 min

I ~ i I 5=2 I I I I I 1,8 50 48 50 52

2 8 (deg)

Fig. 3 - - X - r a y sideband profiles around the (200)v reflection of alloy C aged at 823 K for various times. The scale of X-ray intensity for the times ->1.9 • 103 min is twice as large as for those <-960 min.

10 4 minutes produced regular arrangements of fine precipitates in the austenite matrix. Electron diffraction patterns of these specimens revealed the two-phase periodic structure and the existence of an ordered phase (K-carbide), as evidenced by the {100} and {110} superlattice spots. In contrast to this, the as-quenched specimens did not show any evidence of ordering. The intensity of the ordered spots increased with aging time, suggesting progress of the ordering reaction. The wavelengths or the spacings of the periodic structure along the (100) direc- tion were about 35, 35, and 10 to 20 nm for alloys B, C, and D, respectively, after aging for 1.3 • 10 4 minutes.

Fig. 4 - -Br igh t - f i e ld TEM micrographs of alloy B aged at 823 K for (a) 30 min and (b) 1.3 • l04 min and (c) electron diffraction patterns from (b) showing [110] zone.


E e-

v

> 0

3C-

20-

10-

--e- T E M

-o- X-ray di f f ract ion , s

823 K j " ""

.," 793 K

102 103 104 Aging time (min)

Fig. 7--Isothermal wavelength growth curves for alloy C aged at 773, 793, and 823 K.

(c) Fig. 5--Bright-field TEM micrographs of alloy C aged at 823 K for (a) 30 min and (b) 1.3 x 104 min and (c) electron diffraction patterns from (b) showing [110l zone.

Figure 7 shows the wavelength growths of the modulated structure calculated from the X-ray sideband spacings for alloy C aged at 773,793, and 823 K. A nearly constant level of initial wavelength was observed at 823 K. For long aging times, linear increases in the logarithms of the wavelength and the aging time were observed.

However, the wavelength determined by TEM was longer than that expected from an extrapolation of the logarith- mic wavelength-time plot determined by X-ray diffraction.

Precipitation of a small amount of an a and/or a /3-Mn phase along grain boundaries of the austenite matrices were observed for alloys B, C, and D, as shown in Figure 8. The precipitates were found to be a /3-Mn phase for alloy B and (a + fl-Mn) phases for alloys C and D by the X-ray diffraction and electron microprobe analysis. The two-phase structure of (a + /3-Mn) consists of a lamellar-type a-ferrite and a fl-Mn matrix. It is to be noted that grain-boundary precipitation of the a and the /3-Mn phases in these alloys was first observed after about 1 x 104 minutes of aging at 823 K. Furthermore, it was found that with an increase in the fraction of the a and/3-Mn phases in the austenite, the matrix hardness of specimens further increased, probably due to the formation of the brittle/3-Mn phase, as shown in Figure l(b).

As-quenched specimens did not have ferromagnetism. However, the age-hardenable alloys showed ferromagnetism even for short aging times at the temperatures

(c)

Fig. 6--Bright-field TEM micrographs of alloy D aged at 823 K for (a) 30 min and (b) 1.3 x 104 min and (c) electron diffraction patterns from (b) showing [211] zone.

Fig. 8--Optical micrographs of(a) alloy B, (b) alloy C, and (c) alloy D aged at 823 K for 1.3 • 104 min.

8--VOLUME 21A, JANUARY 1990 METALLURGICAL TRANSACTIONS A

investigated. The degree of ferromagnetism increased with aging time. Figure 9 shows variations of the residual and saturation inductions with aging time for alloy C, aged at 823 K. Rapid increases of Br and B= were observed in the early stage of aging, which fairly resembles the hardness increase, as shown in Figure 1.

IV. DISCUSSION

A. Formation o f the Modulated Structure

The present X-ray analysis and electron microscopic observations for the age-hardenable Fe-Mn-A1-C alloys have revealed that the modulated structure with superlattice reflections is produced by spinodal decomposition during the early stage of aging at 773 to 823 K. The modulated structure is considered to be formed by composition fluctuation in the austenite matrix along the orthogonal /100) directions. The appearance of the sidebands was dependent on alloy composition, aging temperature, and aging time. Alloy C exhibited the most distinct X-ray sidebands around the (200)r line, as seen in Figure 3. After long aging times, with the disappearance of the X-ray sidebands, this alloy also showed a formation of the K-carbide which arranged regularly in the austenite matrix (Figure 5(b)).

In order to elucidate the effect of alloy composition on the formation of the periodically modulated structure, a pseudoternary isothermal section for the (Fe, Mn)-A1- C system at approximately 823 K was established, as shown in Figure 10. Concentrations of constituent com- ponents are expressed in atomic percent for quick com- prehension of comparing the atomic ratio of each component to the composition of the K-carbide, (Fe, Mn)3A1C~. Both Fe and Mn atoms are considered to be situated at the equivalent face-centered position in the perovskite-type ordered structure of the n-carbide, and the composition of (FexMn~_~) being ranged from x = 0.61 to 0.68 corresponds to a single austenitic phase in the binary Fe-Mn system at around 823 K. t81 Hence, the

x 10 3

01

2~ 0 1 (.9

m

0.5

&Q. 10 z 10 3

Aging t ime (rain)

Fig. 9- -Var ia t ions of saturation induction, B=, and residual induction, Br, for alloy C aged at 823 K for various times.

x 10 -z

- 2

- 1

3

g /

9 o I , - - '

r 10 20 Fe3A [ 30

A[ at. %

Fig. 10-- Schematic isothermal section of the pseudoternary (Fe, Mn)- AI-C system at the aging temperatures around 823 K. Hatched area indicates the region in which the spinodally modulated structure can be formed. The phase boundaries of (7 + K)/(/3-Mn + 7)/(3' + a), depending on the composition of (Fe~Mn~ x), were not shown here. The symbols o, I , and A refer to Refs. 4, 2, and 9, respectively.

Fe and Mn atoms can be safely treated as one component in this schematic isothermal section. The hatched section in Figure 10 indicates the region in which the modulated structure is formed by spinodal decomposition. Open and closed squares were obtained from the results by Han et al. I4] and Storchak and Drachinskya, {2~ respectively. Since the phase boundary between y and y + n was estimated from X-ray analysis of the nonequilibrium rapidly solidified (Fe065Mno.3~)077A10.15C0.08 alloy t91 shown in an open triangle, the equilibrium phase boundary will be located in the slightly lower carbon concentration level. In this region, alloys are a single austenitic phase at 1273 K, and their equilibrium state at lower temperatures consists of the two phases of similar structures (7 and K) and/or the a and/3-Mn phases. Thus, the modulated structure in the (Fe, Mn)-AI-C system is formed as an intermediate stage of the phase decomposition from 3' to 7 + n .

B. Kinetics o f Decomposi t ion

From the analysis of the wavelength-time curves, the kinetics of the coarsening process in the modulated structure was examined. A plot of log A vs log t for alloy C (Figure 7) showed a linear relation at 773 to 823 K. The kinetics of the wavelength growth at constant temperature obeys the relation A m oc kt , where k is a constant; the m values varied between 3.4 and 4.8 within the experimental error. According to the kinetic models proposed by Lifshitz and Slyozov [1~ and Wagner [1~} based on the coalescence of spherical particles by a volume diffusion-controlled process through the matrix phase, the m value can be 3. The m values obtained for alloy C are larger than that predicted from the Lifshitz/Slyozov and Wagner models. Butler and Thomas [~21 obtained m = 3 for Cu0.52(Nio.vFe0.3)04s alloy aged at 898 to 1048 K, whereas Hillert et al. 113] found the larger values, m = 4 to 5, for the same alloy system, Cux(Ni0.vFe0.3)l_x alloys (x = 0.15 to 0.75), aged at 895 K. Carpenter [~41 also measured the larger values of m changing from 3.2 to 9.3 with a variation of the alloy composition in the Au-Pt system aged at 773 to 875 K.


These m values larger than 3 show that the growth rate of the modulated structure of these alloys is slower than that of precipitate coarsening in spherical particles.

When alloy C was aged at 823 K for 1.3 • 10 4

minutes, the measured wavelength of the periodic structure by TEM was longer than that expected from the wavelength-time relation, as shown in Figure 7. There was no direct evidence, such as interfacial dislocations, which indicates the loss of coherency during aging; however, it seems probable that the growth rate of the wavelength in the late stages of aging increases due to the lower strain energy caused by the relief of the lattice mismatch.

The apparent activation energy for the isothermal growth of the modulated structure for alloy C can be evaluated as follows. The extent of the isothermal growth of the modulated structure at any times, r, can be represented by the equatio@ 31

r = At exp ( - Q / R T ) [2]

where A is a temperature-independent constant, t the aging time, Q the activation energy, R the gas constant, and T the absolute temperature. The extent of the growth, r, can be assumed to be the wavelength of the modulated structure. We determined the value of t as the time re- quired for the wavelength to reach 20 nm (A = 20 nm). A plot of log t vs 1 /T for the same r value is shown in Figure 11. The activation energy calculated from the slope of the plot was about 180 kJ/mol. This value can be evaluated as follows. Since there are no diffusion data for the Fe-Mn-A1-C system, activation energies for interdiffusion in the y Fe-C system and diffusion of aluminum in the austenitic iron were chosen for comparison. The activation energy for the interdiffusion in the Fe-C

"c 4.0 E

.r

t2~ O

J

3.5

3.C t I 1.3 1.2

1/T (K)-s Fig. 11--Isothermal reaction time vs reciprocal absolute temperature for alloy C.

system, depending on the composition (C = 1 to 6 at. pet), in the temperature range between 1023 and 1573 K was determined to be between 120 and 150 kJ/mol. ~ The activation energy for the diffusion of aluminum atoms in y-iron was determined to be 235 kJ/mol, t~6~ The measured activation energy for alloy C at 773 to 823 K lies between these two values. Thus, the modulated structure in the present Fe-Mn-AI-C alloys is assumed to be formed through the concomitant movement of the solute aluminum and carbon atoms.

C. Relationship between Modulation Amplitude and Age Hardening and Magnetic Properties

The modulation amplitude, which represents the extent of the spinodal decomposition completed, can be evaluated from the analysis of the X-ray sideband intensity. The model of the composition fluctuation during the earliest aging times will be given by a cosinusoidal composition modulation of three {100} waves each hav- ing amplitude, A, [171

C - Co = a(cos/3xl + cos fix2 + cos/3x3) [3]

Here, C and Co are the local atomic concentrations of the solute in the spinodal and homogeneous structures, respectively, and /3 = 2~-/A. Ditchek and Schwartz I'8~ have proposed a waveform which varies from a sine wave toward the square wave with the progress of spinodal decomposition. According to their treatment, the amplitude of one-dimensional fluctuations, A, can be obtained from the following equation:

~r + ~ - / I o - 2J,(gAes) [4] Jo(gZes)

where I0, I+, and I_ are the integrated intensities of the Bragg peak and the high- and low-angle sideband peaks, respectively, J0 and J~ are the zeroth and first-order Bessel functions of the first kind, respectively, g is a factor tak- ing into consideration the dynamic behavior of the waveform and varying from 1 to 4/7r toward the square wave, e is the nondimensional strain amplitude (e = Arl(C,j + 2Ct2)/Cll) where 7/ = (Oa/Oc)/a and Cll and C12 are the elastic constants, and s is the magnitude of the X-ray scattering vector (s = 2 sin 0/AX.ray ). Ditchek and Schwartz [~8j calculated the values of g as a function of A/Am(0), in which Am(0) represents the initial wavelength at t = 0. In alloy C, Am(0) was 13.5 nm when aged at 823 K (Figure 7). Thus, the value of g for the maximum wavelength (A = 19.8 nm) determined by X-ray analysis in alloy C was found to be 1.09. In order to obtain intensities of I0, I+, and I_, a graphical peak separation method was employed, assuming the Gaussian profile for each peak. Integrated intensities were corrected for the dependence of the atomic scattering factor, multi- plicity factor, and Lorentz-polarization factor on scattering angle. In the present alloy, solute species are considered to be both aluminum and carbon atoms. Charles etal. ~ measured the compositional depen- dency of lattice parameter in austenitic Fe-(20 to 40)Mn- (0 to 5)Al-(0 to 1)C alloys. According to their data, r/c~on and ~aluminum for alloy C were estimated to be 0.18 and 0.09, respectively. Thus, compositional fluctuation of

1 0 - - V O L U M E 21A, J A N U A R Y 1990 M E T A L L U R G I C A L TRANSACTIONS A

A

4--J O

< 10

o ~

Q. E ~ 5

o o ~

0

o ~ 0 I I

10 2 10 3

Aging time (rain) Fig. 12--Compositional modulation amplitude for alloy C aged at 823 K for different times.

carbon atoms will primarily contribute to strain amplitude, so the value of 7/c~rbo . was employed here. Other values used for the calculation of A are C~1 = 14.4 • 101~ and C~2 = 8.7 • 10 I~ N / m 2 for a single crystal of fcc Fe-30Ni, [2~ since the elastic constants of a single crystal of Fe-30Mn-9A1-0.9C are not available and s = 5.41 nm -l .

The calculated results for alloy C aged at 823 K for different aging times are shown in Figure 12. The increase of the compositional modulation amplitude, 3A, which corresponds to the concentration of carbon atoms, is rapid from the early stage of the spinodal decomposition, and it becomes almost saturated in the late stage: > 103 minutes. Comparing these amplitude data with the changes in hardness (Figure 1) and Bs and Br (Figure 9), one can see that they are in good agreement with an increase in the modulation amplitude. It is thought that the precipitation of the ferromagnetic K-carbide occurs and is followed by the spinodal decomposition. However, the ordering reaction evidenced by superlattice reflections was observed from the early stage of aging. Thus, the increase in ferromagnetism is directly related to the increase in the composition amplitude. In addition, it is to be noted that maximum amplitude was about 9 at. pct. This corresponds to the atomic concentration of the (Fe, Mn)3A1Cx=0.a-type carbide.

V. CONCLUSIONS

Age hardening was observed for the austenitic Fe-(30 to 34)Mn-(8 to 11)AI-(0.9 to 1.0)C alloys solutionized at 1273 K. These age-hardenable Fe-Mn-A1-C alloys exhibited spinodally modulated structures with a concurrent ordering reaction formed along the orthogonal (100) directions during aging. These alloys also showed increases in magnetic residual and saturation inductions with aging. The growth of the modulated structure results in

a precipitation of the K-carbide, (Fe, Mn)3A1Cx, arranged periodically in the austenite matrix. In the late stages of aging, phase decomposition from the austenite phase to the /3-Mn and the a-ferrite phases occurs at the grain boundaries. This phase decomposition contributes to a slow increase in hardness, depending on the fraction of the fl-Mn phase produced.

The rapid increase in hardness and residual and saturation inductions for short aging times is in accord with the growth of the compositional amplitude, which was obtained from the X-ray sideband analysis of Fe-30Mn- 9A1-0.9C. The kinetic data revealed that the value of the exponent m in the equation A m oc kt varied between 3.4 and 4.8 in the range of temperatures examined. The activation energy for the growth of the modulated structure was about 180 kJ/mol, corresponding to that of volume diffusion and indicating both aluminum and carbon as solute atoms. Maximum composition amplitude of carbon was estimated to be about 9 at. pct, which corresponds to the (Fe, Mn)3A1Cx=0.4 carbide.

A C K N O W L E D G M E N T S

The authors wish to thank Dr. M. Yamanka and Mr. M. Tendoh, Research and Developmental Laboratory of Nippon Steel Corporation, for preparing the mate- rials. This work was supported by Grant-in-Aid for Scientific Research from the Ministry of Education, Science and Culture, Japan.

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modulated structure and magnetic properties of age-hardenable fe-mn-al-c alloys

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